KR20160046770A - Ni-BASED ALLOY FOR FORGING, METHOD FOR MANUFACTURING THE SAME, AND TURBINE COMPONENT - Google Patents

Abstract

A Ni-group alloy of an embodiment comprises: 0.01-0.07 mass% of C; 14-26 mass% of Cr; 10-15 mass% of Co; 5-12 mass% of Mo; 0.8-3 mass% of Ti; 0.001-0.006 mass% of B; and the remaining part consisting of Ni and unavoidable impurities. The Ni-group alloy satisfies the relationship of Mo + 0.176Cr + 0.037Co is greater than or equal to 10 mass% and less than or equal to 15 mass%. In addition, the average thickness of carbide precipitated along the grain boundary is less than or equal to 250 nm.

Description

FIELD OF THE INVENTION [0001] The present invention relates to a Ni-based alloy for forging, a method of manufacturing the same and a turbine component,

The present invention relates to a Ni-based alloy for forging, a method of manufacturing the same and a turbine component.

In recent years, from the viewpoint of reducing the amount of carbon dioxide emission in the atmosphere, the efficiency of the thermal power generation plant has been increasing. Therefore, high efficiency of steam turbines and gas turbines provided in thermal power plants is required. In addition, CO ₂ turbines, which can be installed in thermal power plants, are required to have high efficiency. Here, a CO ₂ turbine drives a turbine using CO ₂ produced by combustion of oxygen and a fuel such as natural gas as a working fluid. In the CO ₂ turbine, most of the generated CO ₂ can be easily separated, and recovered. Therefore, it is attracting attention from the viewpoint of global environment protection.

In order to increase the efficiency of each turbine, it is effective to increase the temperature of the inlet of the working fluid introduced into the turbine. For example, in the case of a steam turbine, in the future, it is expected to operate the steam at a temperature of 700 ° C or higher as a working fluid. Even in a gas turbine or a CO ₂ turbine, the inlet temperature of the working fluid to be introduced tends to rise.

Therefore, it is preferable that the components constituting the high-temperature portion of each turbine be made of a Ni-based alloy which is used in a component parts of a power generation gas turbine or an engine for an aircraft, and has a track record in use at a high temperature.

Typical examples of the Ni-based alloy include Inconel 718 and Inconel 617 (manufactured by Special Metal). The strengthening mechanism of the Ni-based alloy is roughly divided into precipitation strengthening type and solid solution strengthening type.

(Gamma prime: Ni ₃ (Al, Ti)) phase or gamma prime (gamma double prime: Ni ₃ Nb) phase by adding Al, Ti, Ta and Nb to Ni in the precipitation- By precipitating a precipitated phase called phase, the mechanical strength under high temperature is improved. As a representative precipitation-strengthening type Ni-based alloy, Inconel 718 can be mentioned.

On the other hand, in the solid solution strengthening type Ni-based alloy, the mother phase itself is strengthened by adding Co, Mo or the like to Ni. A typical example of the employment hardening type Ni-based alloy is Inconel 617 described above.

As described above, application of a Ni-based alloy as a material of component parts of a turbine, which is used under a high temperature environment, has been studied. In addition, Ni-based alloys are required to have sufficient mechanical strength under a high-temperature environment, and are required to be manufactured when producing large-sized forged parts and the like.

BRIEF DESCRIPTION OF THE DRAWINGS Fig. 1 is a diagram schematically showing a metal structure of a Ni-based alloy in the embodiment; Fig.
2 is an electron micrograph of a metal structure of a Ni-based alloy in order to explain the precipitation form of carbide precipitated on a grain boundary (grain boundary) according to the conditions of the aging treatment.
Fig. 3 is an electron micrograph of a metal structure of a Ni-based alloy in order to explain the precipitation form of carbide precipitated on grain boundaries according to the conditions of the aging treatment; Fig.

In one embodiment, the forging Ni-based alloy contains 0.01 to 0.07% of C, 14 to 26% of Cr, 10 to 15% of Co, 5 to 12% of Mo, 0.8 to 3% of Ti, 0.001 to 0.006% of B, and the balance of Ni and inevitable impurities, and satisfies the relationship of 10 mass%? Mo + 0.176 Cr + 0.037 Co 15 mass% . The average thickness of the carbide precipitated along the grain boundaries is 250 nm or less.

Hereinafter, embodiments according to the present invention will be described.

Ni-based alloys, γ 'is obtained by the addition of such as Mo, W, such as solid solution strengthening reinforcement by the element employed, and Al, Ti of: by the precipitation strengthening by fine precipitation of (gamma prime Ni ₃ (Al, Ti)) , The material strength at room temperature and high temperature is improved. On the other hand, the excessive strengthening deteriorates the workability of the material at a high temperature and deteriorates the productivity.

For example, Inconel 617 having a slight amount of precipitation strengthening by? 'Phase has a good mono-composition compared with that of Udimet 520 (produced by Special Metal Co., Ltd.), which has a large precipitation strengthening amount by?' Phase. On the other hand, in the case of Inconel 738LC (manufactured by Special Metals) having a large amount of precipitation of the γ 'phase, it is generally not molded by forging but molded by casting.

Thus, the production method of the Ni-based alloy is mainly determined according to the precipitation amount of the? 'Phase. For example, in the case of a Ni-based alloy for forging, an alloy composition which does not cause excessive precipitation of? 'Phase during forging is set.

A large member such as a steam turbine or a CO ₂ turbine turbine rotor provided in a thermal power generation plant is larger than a forged member such as a gas turbine or a jet engine in which a Ni-based alloy is conventionally used. For this reason, in order to manufacture these large members, for example, a forging member exceeding 10 tons is required.

In the forging of such a large forged member, the forged product can not be obtained even with the Inconel 617, which has been considered for forging so far due to factors such as insufficient capacity of the forging press. Thus, in the Ni-based alloy for use in a large member, it is necessary to consider not only the precipitation amount of the γ 'phase but also the solid solution strengthening amount which affects the deformation resistance at a high temperature.

The solid solution strengthening is obtained by solving (solving) other solute atoms in the solvent atoms constituting the parent phase, and then the internal deformation that occurs hinders the motion of the potential. Employment reinforcement is theoretically interpreted as a model in which dislocation moves while avoiding the disadvantages of solute atoms. According to Friedel, the solubility enhancement of the dilute solid solution is proportional to 1/2 the solute atom concentration and proportional to the 3/2 power of the misfit strain according to the atomic size difference (Advances in Physics, vol. 3, Issue 12, pp.446-507). Further, according to Labusch, in the high concentration solid solution, the solubility enhancement amount is proportional to 2/3 of the solute atom concentration and proportional to the 4/3 power of the misfit strain according to the atom size difference (Physica status solidi (b). Volume 41, Issue 2, pp. 659-669).

Further, as a factor strongly influencing the characteristics of the metal material, a microstructure of the material can be mentioned. In the Ni-based alloy, the characteristics of the material depend not only on the crystal grain but also on the grain boundary structure. Particularly, it is known that M ₂₃ C ₆ type carbide precipitated in a film form on grain boundaries lowers the toughness of the material. Therefore, in order to secure the reliability of the material, it is necessary to appropriately control the metal structure by optimizing the heat treatment conditions.

In view of the above, the present inventors have quantitatively evaluated the amount of added elements and the misfit deformation of each element added to the Ni-based alloy, thereby finding parameters indicating the amount of solid solution strengthening. In addition, various material tests were conducted on the material whose chemical composition was changed, and the chemical composition having excellent mono-composition was obtained while maintaining sufficient material strength.

As a result of investigation of the grain boundary structure of the Ni-base alloy subjected to various heat treatments, the " average thickness of the carbides on the grain boundary " was found as a factor controlling the toughness of the Ni-base alloy. Then, the range of the thickness of the carbide on the grain boundary that can secure the toughness was found.

Next, the forging Ni-based alloy of the embodiment will be specifically described.

BRIEF DESCRIPTION OF DRAWINGS FIG. 1 is a diagram schematically showing a metal structure of a Ni-based alloy in the embodiment. FIG. In the following description, percentages representing compositional components are, unless otherwise specified, in mass%.

The Ni-based alloy according to the embodiment contains 0.01 to 0.07% of C, 14 to 26% of Cr, 10 to 15% of Co, 5 to 12% of Mo, 0.8 to 3% of Al, 0.8 to 3% B: 0.001 to 0.006%, and the balance of Ni and inevitable impurities, and satisfies the relation of 10 mass%? Mo + 0.176 Cr + 0.037 Co 15 mass%.

Further, in the Ni-based alloy of the embodiment, as shown in Fig. 1, the carbide 11 precipitates along the grain boundaries. The average thickness t of the carbide 11 is preferably 250 nm or less. The carbide (11) is continuously deposited along the grain boundary system (10). In the crystal grains 12, precipitates 13 precipitate on the granular phase.

The carbide (11) is a carbide containing Cr and Mo as a main component, specifically, an M ₂₃ C ₆ type carbide. The reason why the average thickness t of the carbide 11 is preferably 250 nm or less is that the toughness does not decrease, for example, and the toughness for manufacturing a turbine component can be appropriately secured.

The precipitate 13 is composed of? '(Gamma prime: Ni ₃ (Al, Ti)) phase. The diameter of? 'phase is preferably small in view of precipitation strengthening. The average diameter of the? 'phase is preferably, for example, 150 nm or less.

Here, the Ni-based alloy in the embodiment may contain 0.05 to 0.7% of Ta in addition to the chemical composition described above. Further, the Ni-based alloy in the embodiment may contain 0.1 to 0.7% of Nb in addition to the chemical composition described above. Further, the Ni-based alloy in the embodiment may contain 0.05 to 0.7% of Ta and 0.1 to 0.7% of Nb in addition to the chemical composition described above.

Examples of the inevitable impurities include Si, Mn, N, Cu, Fe, and S. It is preferable that these inevitable impurities are made as close to 0% as possible.

The Ni-based alloy of the above embodiment is favorable as a material constituting a turbine component constituted by, for example, forging of a power generation turbine or the like, which is used under a temperature of 650 DEG C or higher, for example. Examples of the turbine component include a turbine rotor, a rotor blade, a stationary blade, a screwing member, a pipe, and the like. These forged parts are all installed in an environment of high temperature and high pressure.

Here, as the screwing member, for example, a bolt or a nut for fixing various components in the turbine casing or the turbine may be exemplified. As the piping, for example, a piping or the like which is installed on a power generation turbine plant or the like and through which a high-temperature, high-pressure working fluid passes can be exemplified.

Further, all the parts of the turbine parts of the power generation turbine may be made of the Ni-based alloy described above. Further, a part of the turbine part, which is particularly high in temperature, may be made of the Ni-based alloy described above.

The forging Ni-based alloy of the above-described embodiment is superior to the conventional forging Ni-based alloy in strength characteristics and excellent in mono-composition. Therefore, turbine components such as turbine rotors, rotor blades, stator blades, screwing members, and pipes manufactured using the forging Ni-based alloys according to the embodiments have high reliability even under a high temperature environment.

Next, reasons for limiting each composition component range in the forging Ni-based alloy of the above embodiment will be explained.

(1) C (carbon)

C is useful as a constituent element of the reinforcing phase carbide. C has a function of suppressing coarsening of crystal grains under a high temperature by the pinning effect of carbide which inhibits the movement of crystal grain boundaries. When the content of C is less than 0.01, the strengthening by carbide is not sufficient. When the content of C is less than 0.01, there is a possibility that crystal grains are coarsened because a sufficient deposition amount of carbide can not be ensured. On the other hand, if the content of C exceeds 0.07%, the mono-composition decreases. Therefore, the content of C was set to 0.01 to 0.07%. The content of C is more preferably 0.03 to 0.07%.

(2) Cr (chromium)

Cr is an indispensable element for improving the oxidation resistance, corrosion resistance and high-temperature strength characteristics of the Ni-based alloy. When the content of Cr is less than 14%, oxidation resistance and corrosion resistance are deteriorated. On the other hand, when the content of Cr exceeds 26%, precipitation of the σ phase, which causes a decrease in the creep strength, becomes remarkable, and the monocomponent composition deteriorates. Therefore, the content of Cr was set to 14 to 26%. The content of Cr is more preferably from 16 to 20%.

(3) Co (cobalt)

Co is dissolved in the parent phase in the Ni-based alloy to improve creep strength and tensile strength. When the content of Co is less than 10%, sufficient mechanical strength can not be obtained. On the other hand, when the content of Co exceeds 15%, the mono-composition decreases. Therefore, the content of Co was set to 10 to 15%. The content of Co is more preferably 11 to 14%.

(4) Mo (molybdenum)

Mo is dissolved in the Ni phase to improve creep strength and tensile strength. Further, the stability of the carbide is enhanced by substituting a part of Mo into the M ₂₃ C ₆ type carbide. If the content of Mo exceeds 12%, the hot workability deteriorates. On the other hand, when the content of Mo is less than 5%, the mechanical strength can not be improved. Therefore, the content of Mo is set to 5 to 12%. The content of Mo is more preferably 7 to 10%.

(5) Al (aluminum)

Al generates γ '(Ni ₃ Al) phase together with Ni, thereby improving the mechanical strength of the Ni-based alloy by precipitation. When the Al content is less than 0.8%, the effect of precipitation of the? 'Phase is not exerted. On the other hand, if the content of Al exceeds 3%, precipitation of the σ phase is promoted and the mechanical properties are lowered. When the content of Al exceeds 3%, the hot workability remarkably decreases. Therefore, the content of Al is set to 0.8 to 3%. The content of Al is more preferably 1 to 2%.

(6) Ti (titanium)

Ti, like Al, generates γ '(Ni ₃ (Al, Ti)) phase together with Ni, thereby improving the mechanical strength of the Ni-based alloy. When the content of Ti is less than 0.8%, the effect of precipitation of the? 'Phase is not exerted. On the other hand, if the content of Ti exceeds 3%, precipitation of σ-phase and η-phase is promoted, resulting in deterioration of mechanical properties and deterioration of hot workability. Therefore, the content of Ti was set to 0.8 to 3%. The content of Ti is more preferably 1 to 2%.

(7) B (boron)

B segregates at the grain boundaries to improve high-temperature strength characteristics. When the B content is less than 0.001%, the effect of improving the high-temperature strength characteristics is not exhibited. On the other hand, when the content of B exceeds 0.006%, grain boundary embrittlement occurs. Therefore, the B content was set to 0.001 to 0.006%. The content of B is more preferably 0.002 to 0.004%.

(8) Ta (tantalum)

Ta is solidified on γ '(Ni ₃ (Al, Ti)) to stabilize the γ' phase. When the content of Ta is less than 0.05%, the above-mentioned effect is not obtained. On the other hand, if the content of Ta exceeds 0.7%, the mono-composition decreases. Therefore, the content of Ta is set to 0.05 to 0.7%. The content of Ta is more preferably 0.08 to 0.12%.

(9) Nb (niobium)

Nb, like Ta, is dissolved on γ '(Ni ₃ (Al, Ti)) to stabilize the γ' phase. When the content of Nb is less than 0.1%, the above-mentioned effect is not obtained. On the other hand, if the content of Nb is more than 0.7%, segregation will occur at the time of melting and casting, and the mono-composition will decrease. Therefore, the content of Nb is set to 0.1 to 0.7%. The content of Nb is more preferably 0.2 to 0.5%.

(10) Mo + 0.176 Cr + 0.037 Co

As described above, the solid solution strengthening amount in the high-concentration solid solution is proportional to the 2/3 power of the solute atom concentration and proportional to the 4/3 power of the misfit strain according to the atom size difference. Therefore, in the present embodiment, the numbers of atoms per 1 mass% of Mo, Cr, and Co, which are considered to contribute to solid solution strengthening, and parameters indicating solid solution strengthening from the respective atomic radii are defined. In the present embodiment, since the content ratio of C (carbon) is small, C is excluded from the parameters.

The atomic weights of Mo, Cr and Co are 95.9, 52.0 and 58.9, respectively. The ratio of the number of atoms in the case where each element is added in the same amount is 1.84 or 1.62 for Cr and Co, respectively, when Mo is 1. The 2/3 power of these ratios is 1, 1.50, and 1.38, respectively.

In addition, the misfit deformation that occurs when each element is added is determined from the difference in atomic size with Ni atoms. The atomic radius difference between the Ni atom and the Mo, Cr, and Co atoms is 0.15 Å (angstrom), 0.03 Å, and 0.01 Å, respectively. Therefore, when Mo is 1, the ratio of the amount of misfit deformation when each element is added becomes 0.200 and 0.067, respectively, for Cr and Co. The 4/3 power of these ratios is 1, 0.117, and 0.027, respectively.

Therefore, the ratio of the solid solution strengthening amount per 1 mass% of each element is 0.176 (1.50 x 0.117 = 0.176) and 0.037 (1.38 x 0.027 = 0.037) for Cr, From these results, " Mo + 0.0176Cr + 0.037Co " was set as a parameter for expressing the solid solution strengthening amount.

When the value (content ratio) of this parameter exceeds 15%, the solute strengthening amount becomes excessive, and the deformability at the time of forging deteriorates. On the other hand, when the value of the parameter is less than 10%, the solid solution strengthening amount is remarkably lowered and sufficient strength is not obtained. Therefore, the value of the parameter is set to 10% to 15%. The value of the parameter is more preferably 11 to 13.5%.

It is considered that not only the size of the atom but also the interaction with Ni or other atoms affects the misfit deformation due to element addition. However, here, for the sake of simplicity, the unfixed strain value is uniquely determined from the difference between each solute atom and the Ni atom. It is known that Mo and Cr combine with C to form carbide. However, since the content of C is low, consumption of Mo and Cr due to carbide is ignored.

(11) A method of manufacturing a semiconductor device, comprising the steps of: (i) depositing Si (silicon), Mn (manganese), N (nitrogen), Cu (copper)

Si, Mn, N, Cu, Fe and S are inevitable impurities in the forging Ni-based alloys of the embodiments. It is preferable that these inevitable impurities are made as close to 0% as possible. It is preferable that at least Si and Mn of these inevitable impurities are suppressed to 0.1% or less and N is suppressed to 0.01% or less.

Si is usually added to supplement corrosion resistance in the case of steel. However, since the Ni-based alloy has a large Cr content, sufficient corrosion resistance can be secured. Therefore, it is preferable that the residual content of Si is set to 0.1% or less and the remaining content is made as close as possible to 0%.

In the case of ordinary steels, Mn is made of S (sulfur) due to brittleness and MnS to prevent brittleness. However, the content of S in the Ni-based alloy is extremely small and it is not necessary to add Mn. Therefore, it is preferable that the residual content of Mn is set to 0.1% or less and the remaining content is made as close as possible to 0%.

N reacts with Ti in the material to form TiN to reduce Ti contributing to the formation of the gamma prime phase. As a result, the mechanical strength is lowered. Therefore, it is preferable to set the residual content of N to 0.01% or less and to make the remaining content as close as possible to 0%.

Hereinafter, a method for manufacturing a turbine component manufactured using the forging Ni-based alloy and the forging Ni-based alloy according to the embodiment will be described.

The forging Ni-based alloy of the above embodiment is manufactured, for example, as follows.

First, the composition constituting the Ni-based alloy is subjected to vacuum induction melting (VIM), and the molten metal is injected into a predetermined mold to form an ingot. Then, the ingot is subjected to soaking, hot forging, solution treatment, aging, and the like.

The turbine rotor, which is a turbine component, is manufactured, for example, as follows.

For example, as a method (double melt), a composition component constituting the forging Ni-based alloy is subjected to vacuum induction melting (VIM) and electroslag material dissolution (ESR) . Subsequently, a soaking treatment, a forging treatment, a solution treatment treatment, an aging treatment, and the like are performed to produce a turbine rotor.

As another method (double melt), the composition constituting the forging Ni-based alloy of the embodiment is subjected to vacuum induction melting (VIM), vacuum arc remelting (VAR), and then poured into a predetermined mold. Subsequently, a soaking treatment, a forging treatment, a solution treatment treatment, an aging treatment, and the like are performed to produce a turbine rotor.

As another method (triple melt), the composition constituting the forging Ni-based alloy of the embodiment is subjected to vacuum induction melting (VIM), electroslag material dissolution (ESR), vacuum arc material dissolution (VAR) It is poured into a predetermined mold. Subsequently, a soaking treatment, a forging treatment, a solution treatment treatment, an aging treatment, and the like are performed to produce a turbine rotor.

By the above-described manufacturing method of the turbine rotor, at least a predetermined portion of the turbine rotor is manufactured. As a predetermined region, for example, a portion exposed to a high temperature of 700 DEG C or more among the turbine rotors and the like can be given. In this case, the portion of the turbine rotor exposed to a temperature of, for example, about 600 ° C is manufactured by a conventional heat-resistant alloy. A turbine rotor is constituted by joining a part made of the forging Ni-based alloy of the embodiment manufactured by the above-mentioned manufacturing method and a part made of a conventional heat-resistant alloy by, for example, welding. The method of joining the component made of the Ni-based alloy for forgings of the embodiment and the component made of the conventional heat-resistant alloy is not limited to welding, and may be fastened by bolts and nuts, for example.

By dividing the parts constituting the turbine rotor in this manner, a turbine rotor which can be used in a high-temperature environment of 700 ° C or more can be manufactured even for a Ni-based alloy having a small ingot. Further, depending on the temperature condition to be used, all of the turbine rotor may be manufactured by the manufacturing method of the turbine rotor described above.

The rotor, stator, and screwing members, which are turbine parts, are manufactured, for example, as follows.

First, the composition components constituting the forging Ni-based alloy of the embodiment are subjected to vacuum induction melting (VIM) and electro-slag solute dissolution (ESR). Subsequently, the molten alloy is poured into a predetermined mold in a reduced-pressure atmosphere to prepare an ingot, and soaking treatment is carried out. Then, the ingot is placed in a shape corresponding to the shape of the turbine component, and forging, solution treatment, aging, and the like are performed to produce a rotor, a stator, and a screwing member. That is, the rotor, the stator and the screwing member are manufactured by die forging.

As another method (double melt), for example, the composition constituting the Ni-based alloy for forgings of the embodiment is subjected to vacuum induction melting (VIM) and vacuum arc material dissolution (VAR). Subsequently, the molten alloy is poured into a predetermined mold in a reduced-pressure atmosphere to produce an ingot. Then, the ingot may be subjected to a soaking treatment, and a forging treatment, a solution treatment, an aging treatment, or the like may be carried out to produce a rotor, a stator, and a screwing member.

As another method (triple melt), the composition constituting the forging Ni-based alloy of the embodiment is subjected to vacuum induction melting (VIM), electroslag material dissolution (ESR), and vacuum arc material dissolution (VAR). Subsequently, the molten alloy is poured into a predetermined mold in a reduced-pressure atmosphere to produce an ingot. Then, the ingot may be subjected to a soaking treatment, and a forging treatment, a solution treatment, an aging treatment, or the like may be carried out to produce a rotor, a stator, and a screwing member.

The piping, which is a forged part of the embodiment, is manufactured, for example, as follows.

First, an ingot is produced by dissolving an electric furnace (EF) and performing argon-oxygen decarburization (AOD) on the composition constituting the forging Ni-based alloy of the embodiment. Then, the ingot is subjected to soaking treatment. The ingot is pierced by a vertical press to produce a base tube in the cup. Then, machining by the mandrel and dies and reheating with the horizontal press are repeated to mold the primary pipe into the pipe shape. This processing method is an Ehrhardt push bench pipe manufacturing method. Then, solution treatment, aging treatment, and the like are performed to produce a pipe.

The method of manufacturing the turbine rotor, the rotor, the stator, the screwing member, and the pipe is not limited to the above-described method. Forged parts such as turbine rotors, rotor blades, stator blades, screw-in members, and pipes can be applied to power turbines such as steam turbines, gas turbines, and CO ₂ turbines.

Hereinafter, each heat treatment at the time of manufacturing the above-mentioned forging-purpose Ni-based alloy and turbine parts will be described. The temperature in each heat treatment is set within the respective ranges shown below depending on the to-be-processed Ni-based alloy to be treated, turbine parts and the like. In addition, the time for each treatment is also set appropriately in accordance with the forged Ni-based alloy or turbine parts to be processed.

In the soaking treatment, it is necessary to heat the alloy at a high temperature for a sufficient time in order to reduce segregation of chemical components by thermal diffusion. Therefore, the soaking treatment is preferably carried out in a temperature range of 1000 to 1200 占 폚.

Forging must be performed in a range from a temperature at which sufficient deformability of the material is obtained to a ductility temperature. Therefore, the forging is preferably performed in a temperature range of 950 to 1100 캜.

In the solution treatment, it is preferable to maintain the temperature in the range of 1050 to 1200 占 폚 for 1 to 24 hours. Here, the solution treatment is carried out in order to sufficiently solidify the alloy element in the mother phase to sufficiently obtain the effect of solid solution strengthening, and to enable precipitation control of the precipitate by the subsequent heat treatment. The solution treatment may also be performed for the purpose of adjusting the crystal grain diameter.

When the temperature of the solution treatment is lower than 1050 占 폚, the alloy element is not completely dissolved in the mother phase and hardening by the solid solution strengthening element is not sufficiently performed. It is also difficult to control the precipitation form of the precipitated phase by the heat treatment after the solution treatment. On the other hand, when the temperature of the solution treatment exceeds 1200 ° C, crystal grain diameter is coarsened and mechanical strength is lowered. Therefore, the temperature of the solution treatment was set to 1050 to 1200 占 폚. It is more preferable that the temperature of the solution treatment is 1050 to 1150 캜. In addition, the solution-processed Ni-based alloy or turbine component is cooled to room temperature by, for example, water cooling or forced air cooling.

Next, the aging treatment to be performed on the Ni-based alloy or turbine parts cooled to room temperature after the solution treatment is described.

In the aging treatment, the temperature is preferably maintained at 700 to 800 DEG C for 5 to 50 hours. This aging treatment may be performed in multiple stages. After the aging treatment, the Ni-based alloy or the turbine component is cooled to room temperature by, for example, water cooling or furnace cooling.

Here, the reason why the temperature and time in the aging treatment are set in the above range will be described.

The main purpose of the aging treatment is to control the precipitation form of the γ 'phase to be precipitated in the crystal grains. The aging treatment also affects the properties of the grain boundaries. Therefore, in regard to the aging treatment, it is necessary to select the temperature and time conditions in consideration of the grain structure and grain boundary structure.

Fig. 2 and Fig. 3 are electron micrographs of the metal structure of the Ni-based alloy in order to explain the precipitation form of the carbide precipitated on the grain boundaries according to the conditions of the aging treatment. The composition of the Ni-based alloy showed 0.04% of C, 18% of Cr, 12% of Co, 9% of Mo, 1.3 of Al, 1.4% of Ti, 0.003% of B, 0.1% of Ta, Nb 0.3%, and the remainder is Ni. Fig. 2 shows a metal structure in which the aging treatment is performed at a temperature of 850 캜 for 10 hours, and Fig. 3 shows a metal structure subjected to aging treatment at a temperature of 750 캜 for 10 hours. The soaking treatment and the solubilization treatment are carried out within the above-mentioned range. 2 and 3, the precipitate 13 (phase of? ') Is also shown.

In the ordinary aging treatment, as shown in Fig. 2, the filmy carbide 11 precipitates so as to cover the grain boundaries of the Ni-based alloy. The film-like carbide 11 is a weak carbide (M ₂₃ C ₆ type carbide) mainly composed of Cr and Mo. This carbide (11) promotes fracture of grain boundaries and significantly reduces the toughness of the material. Therefore, it has been considered necessary to perform the aging treatment to prevent the deposition of the carbide 11 on the film covering the grain boundaries.

However, as shown in Fig. 3, depending on the conditions of the aging treatment, the film-like carbide covering the crystal grain boundary becomes thinner. In addition, the carbide is continuously deposited along the grain boundaries. As a result of the material test, the inventors have found that when the thickness of the carbide is sufficiently thin, no deterioration of the ductility tends to occur. The above-mentioned temperature and time are defined within a range satisfying both of fine precipitation of? 'Phase and coarsening inhibition of carbide covering grain boundaries.

When the temperature of the aging treatment is lower than 700 캜, the coarsening of the carbide covering the grain boundaries can be suppressed, but the growth of the γ 'phase is remarkably slow. Therefore, improvement in mechanical strength due to precipitation of the? 'Phase can not be obtained. On the other hand, when the temperature of the aging treatment exceeds 800 캜, fine precipitation of the γ 'phase is achieved, and sufficient strength is obtained. However, the coarsening of the carbide covering the grain boundaries is remarkable, and the toughness is lowered.

In this respect, the aging treatment temperature was set to 700 to 800 占 폚. Here, for the early precipitation of? 'Phase, the aging treatment may be heat-treated in multiple stages such as two stages. Also in this case, the temperature is set in the temperature range of the above aging treatment. The heat transfer treatment time in the multi-stage is also set in the time range of the aging treatment described above. For example, a treatment at a temperature of 800 DEG C for 10 hours and then a treatment at a temperature of 750 DEG C for 20 hours. The temperature decrease from 800 ° C to 750 ° C is performed by, for example, furnace cooling.

The cooling after the aging treatment is performed by, for example, furnace cooling and air cooling. In the case where the aging treatment is carried out in multiple stages, cooling between aging treatments is carried out by, for example, furnace cooling as described above. Then, the aging treatment is continuously performed without cooling to room temperature.

Here, the intermediate heat treatment may be performed on the Ni-based alloy or the turbine parts cooled to room temperature after the solution treatment before the aging treatment. In this intermediate heat treatment, the purpose is to intermittently form a massive carbide along the grain boundaries before the aging treatment, in order to suppress precipitation and coarsening of the carbides on the film covering the grain boundaries. This carbide is also a carbide (M ₂₃ C ₆ type carbide) mainly containing Cr and Mo.

The intermediate heat treatment is preferably carried out in a temperature range of 1000 to 1050 캜. When the intermediate heat treatment temperature is lower than 1000 占 폚 or exceeds 1050 占 폚, massive carbides do not precipitate. The time for the intermediate heat treatment is set to an arbitrary value depending on the Ni-based alloy to be treated, turbine parts, and the like.

In addition, when the content of C (carbon) is sufficiently small, precipitation of carbide on the film surface on the crystal grain is not remarkable, and this intermediate heat treatment may be omitted. The case where the content of C is sufficiently small refers to the case where the content of C is, for example, 0.04% or less although it varies depending on the crystal grain size and the like. The cooling after the intermediate heat treatment is performed by, for example, cold cooling, water cooling or forced air cooling. Then, the Ni-based alloy or the turbine part is cooled to room temperature.

(Influence of chemical composition)

Hereinafter, it is explained that the forging Ni-based alloy of the embodiment is excellent in strength characteristics and mono-composition.

Table 1 shows chemical compositions of Sample 1 to Sample 21 used for evaluation of strength characteristics, mono-composition and the like. The samples 1 to 13 shown in Table 1 were Ni-based alloys in the chemical composition range of the forging Ni-based alloys of the embodiment, and the samples 14 to 21 had the composition of the forging Ni- Is a Ni-based alloy which is not in the chemical composition range of Ni, and is a comparative example.

[Table 1]

The strength characteristics were evaluated by a tensile test and the toughness by a Charpy impact test, and the composition was evaluated by visual observation. Further, the thickness of the carbide on the film covering the grain boundaries was measured by observing the metal structure.

Test pieces used for each test were made as follows.

The Ni-base alloys of Sample 1 to Sample 21 having the chemical compositions shown in Table 1 were each dissolved in a vacuum induction melting furnace to produce ingots.

Subsequently, the ingot was subjected to soaking treatment at 1050 占 폚 for 5 hours. Thereafter, it was forged with a 500 kgf hammer forging machine at a temperature range of 950 to 1100 占 폚 (reheating temperature of 1100 占 폚). After the forging, the solution treatment was carried out at a temperature of 1100 DEG C for 4 hours, and then cooled to room temperature by air cooling. After cooling, the mixture was subjected to an intermediate heat treatment at a temperature of 1025 DEG C for 10 hours, and then cooled to room temperature by furnace cooling. After cooling, aging treatment in two stages of a temperature of 800 DEG C for 10 hours and a temperature of 750 DEG C for 20 hours was continuously performed. Thereafter, it was cooled to room temperature by air cooling to obtain a forging material.

Then, a test piece of a predetermined size for tensile test and Charpy impact test was produced from this forging material.

The tensile test was carried out in accordance with JIS Z 2241, and 0.2% proof stress and tensile strength at room temperature were measured. The Charpy impact test was carried out in accordance with JIS Z 2242, and the Charpy impact value was measured.

In the evaluation of the monoaxiality, the specimen subjected to the soaking treatment was forged by a hammer forging machine having a weight of 500 kgf to prepare a solid cylindrical test piece having a diameter of 125 mm and a length of 210 mm. The forging process was carried out until the forging ratio (forging ratio based on JIS G 0701 (method of expressing the roughing forming ratio of the steel working work)) was 3. The forging treatment was carried out in the range of 950 to 1100 占 폚. Then, when the temperature of the test piece as the forging object was lowered, that is, when the forging object was cured, the forging process was repeatedly performed by heating again to the reheating temperature of 1100 占 폚. The evaluation of the mono-composition was carried out by visually observing the presence of forging cracks after cooling the test piece.

Here, the forging can be carried out so that the cross-sectional area of the forging object perpendicular to the direction in which the forging object is stretched before the forging process is performed is determined so that the cross sectional area of the forging object perpendicular to the direction in which the forging object is stretched Minus the cross-sectional area of the object.

For the measurement of the thickness of the carbide covering the crystal grain boundaries, a forged material cooled to room temperature after the aging treatment was used. An electron microscope photograph taken at a magnification of 20,000 times using a field emission scanning electron microscope was subjected to image analysis to determine the thickness of the carbide. In each of the forgings, representative five grain boundaries were selected, and the thickness of carbide of 20 points was measured for each. Then, by arithmetically averaging them, an average thickness of carbide was obtained.

Test results and observation results are shown in Table 2. In Table 2, when there is no forgings crack, it is represented by " no ", and in order to indicate that the mono-composition is excellent, evaluation of the mono-composition is indicated by " O ". On the other hand, when there is a forgery crack, it is represented by "oil" and the evaluation of the mono-composition is indicated by "x" to indicate that the mono-composition is lowered.

[Table 2]

As shown in Table 2, the samples 1 to 13 are both 0.2% proof stress and tensile strength higher than the sample 14. The reason why the values of the 0.2% proof stress and the tensile strength were high in the samples 1 to 13 are considered to be because sufficient solid solution strengthening and precipitation strengthening were achieved. The samples 1 to 13 are also excellent in monotonicity and have a carbide thickness of 250 nm or less. From the results of the Charpy impact value, the samples 1 to 13 all exhibited a value of 50 J / cm 2 or more. Therefore, it was confirmed that the samples 1 to 13 had sufficient toughness for practical use.

On the other hand, when the value of " Mo + 0.176Cr + 0.037Co " is less than 10 mass% as in the case of the sample 21, even when the respective alloy components are in the chemical composition ranges defined in the present embodiment, sufficient 0.2% And tensile strength can not be obtained. In the samples 15 to 20, the 0.2% proof stress and the tensile strength were high, but the mono-composition was inferior. This is thought to be the result of adding excess strengthening elements.

Thus, the Ni-based alloy deviating from the chemical composition range or the range of " Mo + 0.176Cr + 0.037Co " defined in the present embodiment does not provide excellent results in both strength characteristics and mono-composition.

(Influence of heat treatment)

Here, in Sample 1, the intermediate heat treatment and the aging treatment conditions were changed, and the tensile test, the Charpy impact test, the evaluation of the mono-composition, and the measurement of the thickness of the carbide covering the crystal grain boundaries were carried out. In addition, each test, evaluation of the mono-composition, and measurement of the thickness of the carbide were the same as those described above.

Using Sample 1 shown in Table 1, heat treatment was performed under the conditions of intermediate heat treatment and aging treatment shown in Table 3. The steps other than the intermediate heat treatment and the aging treatment are the same as the above-mentioned method of producing a test piece. In Table 3, for example, " 800 DEG C x 10 h " means that heat treatment was performed at a temperature of 800 DEG C for 10 hours. In the aging treatment, in the case of performing the two-stage heat treatment, the heat treatment conditions are shown in the first and second columns.

[Table 3]

The samples 1 and 22 to 31 shown in Table 3 were heat-treated under the heat treatment conditions of the present embodiment, and the other samples were heat-treated under the conditions not included in the heat treatment conditions of the present embodiment. The test results and the observation results are shown in Table 4.

[Table 4]

Sample 1 and Sample 22 to Sample 31 are both 0.2% proof stress and tensile strength higher than Sample 32 to Sample 39. The average thickness of the carbide in the film covering the grain boundaries in the samples 1 and 22 to 31 is 250 nm or less. Sample 1 and Sample 22 to Sample 31 exhibited a high Charpy impact value as compared with Sample 34 to Sample 37 and Sample 39 because the average carbide thickness was thin.

Thus, under the aging treatment conditions defined in the present embodiment, it is possible to simultaneously suppress the fine precipitation of? 'Phase in the crystal grains and the coarsening of the carbide covering the crystal grain boundaries. As a result, a high value is obtained in both the tensile strength and the Charpy impact value.

On the other hand, in samples deviating from the aging treatment conditions prescribed in the present embodiment, excellent results of both the tensile strength and the Charpy impact value are not obtained.

According to the embodiments described above, it becomes possible to have excellent strength characteristics and mono-composition.

While the foregoing embodiments have been described, these embodiments are provided by way of example only and are not intended to limit the scope of the invention. Indeed, the embodiments described herein embrace various other forms. In addition, various omissions, substitutions and modifications in the above-described embodiments can be made without departing from the gist of the present invention. The scope of the following claims and their equivalents are intended to encompass such modifications and variations as fall within the scope and spirit of the invention.

Claims (9)

Co: 10 to 15%, Mo: 5 to 12%, Al: 0.8 to 3%, Ti: 0.8 to 3%, B: 0.001 to 3% 0.006% and the remainder being Ni, 0.01% or less of N and inevitable impurities, and satisfying the relation of 11 mass%? Mo + 0.176 Cr + 0.037 Co 13.5 mass%
The average thickness of the carbide, which is a precipitate precipitated along grain boundaries (crystal grain boundaries), is 250 nm or less,
The precipitate is composed of? 'Phase, and the average diameter of?' Phase is 150 nm or less,
Further, a solution treatment is performed at a temperature of 1050 to 1200 占 폚,
After the solution treatment, the mixture is cooled to room temperature and subjected to an intermediate heat treatment in a temperature range of 1000 to 1050 캜,
And the aging treatment is carried out at a temperature of 700 to 800 ° C after the intermediate heat treatment, whereby the carbide is precipitated. The method according to claim 1,
And further contains 0.05 to 0.7% by mass of Ta. The method according to claim 1,
And further contains 0.1 to 0.7 mass% of Nb. The method according to claim 1,
0.05 to 0.7% by mass of Ta and 0.1 to 0.7% by mass of Nb. Co: 10 to 15%, Mo: 5 to 12%, Al: 0.8 to 3%, Ti: 0.8 to 3%, B: 0.001 to 3% Based alloy material containing 0.006% of Ni and 0.01% or less of N and inevitable impurities and having a relationship of 11 mass%? Mo + 0.176 Cr + 0.037 Co 13.5 mass% is melted , A structure forming step of forming a structure having a predetermined shape,
A solution treatment step of subjecting the structure to a solution treatment at a temperature of 1050 to 1200 占 폚,
An intermediate heat treatment step carried out in a temperature range of 1000 to 1050 캜 after cooling to room temperature after the solution treatment,
An aging treatment process for aging the structure subjected to the intermediate heat treatment at a temperature of 700 to 800 ° C
And,
By performing the solution treatment process, the intermediate heat treatment process, and the aging treatment process, a carbide having an average thickness of 250 nm or less is formed as a? 'Phase along the grain boundaries, and an average diameter of?' Phase is precipitated to 150 nm or less , A method for producing a Ni-based alloy for forging. A turbine component in which at least a predetermined portion is manufactured by using the forging Ni-based alloy according to claim 1. A turbine component in which at least a predetermined portion is manufactured by using the forging Ni-based alloy according to claim 2. A turbine component in which at least a predetermined portion is manufactured by using the forging Ni-based alloy according to claim 3. A turbine component, wherein at least a predetermined portion is manufactured using the forging Ni-based alloy according to claim 4.

Images