NI-BASE SUPERALLOY AND MANUFACTURING METHOD THEREOF

Abstract

According to the present invention, a Ni-base superalloy is comprised 2.0 to 6.0 wt. % of cobalt (Co), 8.0 to 12.0 wt. % of chromium (Cr), 5.0 to 9.0 wt. % of tungsten (W), 3.5 to 6.0 wt. % of aluminum (Al), 3.0 wt. % or less of titanium (Ti), 5.0 to 10.0 wt. % of tantalum (Ta), 0.05 to 0.15 wt. % of carbon (C), 0.02 wt. % or less of boron (B), 0.05 wt. % or less of zirconium (Zr), with the remainder being composed of nickel (Ni) and unavoidable impurities.

Description

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims priority to Korean Patent Application No. 10-2014-0139830, filed on Oct. 16, 2014 and entitled “NI-BASE SUPERALLOY AND MANUFACTURING METHOD THEREOF”. The disclosures of the foregoing applications are incorporated herein by reference in their entirety.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to a Ni-base superalloy and a manufacturing method thereof, and more particularly, to a Ni-base superalloy having a high-temperature creep properties and remarkably reduced in manufacturing cost by not adding rare elements such as Re, Ru, and Mo and a manufacturing method thereof.

Further, the present invention relates to a Ni-base superalloy and a manufacturing method of the Ni-base superalloy which can maximize the solid-solution hardening effect and improve the directional solidification castability, oxidation resistance, and creep properties while maintaining a phase stability at high temperatures by optimizing a composition ratio of alloying elements and controlling a size and volume fraction of γ′ precipitates.

2. Description of the Related Art

A Ni-base superalloy which exhibits superior creep resistance has been widely used for structural components serviced at elevated temperatures, such as turbine blades and vanes of aircraft jet engines and industrial gas turbines for power generation.

As an environmental problem such as global warming has emerged, new power generation methods for reducing or removing the emission of CO₂ have been studied and a demand for a method to improve the efficiency of conventional power generation system has been gradually increased. Consequently, turbine inlet temperature of the gas turbines has been continuously increased in order to improve the efficiency of gas turbines.

In a gas turbine, atmospheric air flows through a compressor that brings to high pressure. Then, energy is added by spraying fuel into the compressed air and igniting them, so the combustion generates high-temperature flow. This high-temperature and high-pressure gas rotates the turbine, so that an electric power is produced by driving the device such as an electric generator that is coupled to the turbine.

Therefore, a turbine blade or vane, which operates at high-temperature, has a three-dimensionally complicated aerodynamic design including a complicatedly shaped cooling passage within the component to obtain higher efficiency under given conditions. For this reason, turbine blades and the vanes are usually manufactured by a casting process.

Further, the turbine blade operated at high temperatures receives a centrifugal force caused by high-speed rotation of turbines, and thus a creep resistance of the alloy is of great importance for enduring the centrifugal force at high temperatures.

An alloy manufactured by a typical casting process has a grain boundary vulnerable to high-temperature creep. So, in order to improve a creep resistance of the alloy, a directional solidification (DS) process that removes a grain boundary perpendicular to the direction of stress and a single crystal (SC) casting process which can completely remove the grain boundary have been developed since 1970 and used to manufacture turbine blades and vanes.

With such a process development, specialized alloys adequate to the poly-crystalline casting process, directional solidification process, and single crystal casting process have been developed and used for each casting process, respectively.

Among them, Ni-base superalloys manufactured by the directional solidification process exhibit excellent mechanical properties at temperatures above 760° C. compared with those manufactured by conventional casting process. These DS Ni-base superalloys usually contain various alloying elements such as chromium (Cr), cobalt (Co), aluminum (Al), titanium (Ti), molybdenum (Mo), tungsten (W), tantalum (Ta), niobium (Nb), carbon (C), boron (B), zirconium (Zr), rhenium (Re), ruthenium (Ru), etc.

The property of such a Ni-base superalloy varies depending not only on the type and content of each alloying element but also on the combination of specific elements. Therefore, a lot of alloy design studies have been continuously carried out to develop an alloy that exhibits excellent performance.

Ni-base superalloys are strengthened by the precipitation of ordered γ′ (Ni₃(Al,Ti,Ta), L1₂ structure) phase within a disordered matrix γ phase, which is a face-centered cubic solid-solution, by the addition of various alloying elements such as W, Mo, Re, Cr, and Co, etc. A grain boundary of Ni-base superalloys is also strengthened by the precipitation of discrete fine particles by the minor additions of C, B, and Zr.

In recent years, in order to meet the necessity of an alloy having excellent temperature capability and creep resistance, suppressing the additional input of expensive alloying elements as low as possible and adjusting the amounts of other alloying elements are considered to be an effective method in alloy design of Ni-base superalloys.

In particular, in the case of a structural component used at high temperatures, a creep life to reach rupture is important, but if the component has plastic deformation, it cannot be continuously used for its original purpose or the efficiency of the component can be decreased. Therefore, a resistance to initial creep deformation is a crucial factor to be considered in alloy design.

In this regard, there have been continuous attempts to obtain an alloy having excellent tensile and creep strength at high temperatures by adjusting the amounts of alloying elements. For example, KR 2012-0105693 A describes a single crystal Ni-base superalloy improved in creep resistance by adjusting the content of aluminum or titanium.

The superalloy described in KR 2012-0105693 A having the following composition in weight percent: Co: 11.5 to 13.5%, Cr: 3.0 to 5.0%, Mo: 0.7 to 2.0%, W: 8.5 to 10.5%, Al: 3.5 to 5.5%, Ti: 2.5 50 3.5%, Ta: 6.0 to 8.0%, Re: 2.0 to 4.0%, Ru: 0.1 to 2.0%, and the rest including Ni and other unavoidable impurities. However, in KR 2012-0105693 A, high-priced Re and Ru are contained. The addition of about 3 wt. % Re and Ru accounts for about half of the total alloy cost, and thus it is difficult to reduce the price of the alloy containing Re or Ru.

SUMMARY OF THE INVENTION

The objective of the present invention is to provide a Ni-base superalloy having excellent high-temperature creep and oxidation resistance and remarkably reduced in manufacturing cost by excepting high-priced Re, Ru, and Mo from its alloy composition, and a manufacturing method thereof in order to satisfy the demand for the development of a superalloy and solve the problems of the conventional technology.

Another objective of the present invention is to provide a Ni-base superalloy improved in directional solidification castability by optimizing a composition ratio of alloying elements and a manufacturing method thereof.

Still another objective of the present invention is to provide a Ni-base superalloy improved in high-temperature oxidation resistance and a manufacturing method thereof.

In order to achieve the above objectives, an exemplary embodiment of the present invention provides a Ni-base superalloy comprised of cobalt (Co) of 2.0 to 6.0 wt. %, chromium (Cr) of 8.0 to 12.0 wt. %, tungsten (W) of 5.0 to 9.0 wt. %, aluminum (Al) of 3.5 to 6.0 wt. %, titanium (Ti) of 3.0 wt. % or less, tantalum (Ta) of 5.0 to 10.0 wt. %, carbon (C) of 0.05 to 0.15 wt. %, boron (B) of 0.02 wt. % or less, zirconium (Zr) of 0.05 wt. % or less, and the rest including nickel (Ni) and other unavoidable impurities.

The Ni-base superalloy according to the present invention is comprised of cobalt (Co) of 4.0 wt. %, chromium (Cr) of 10.0 wt. %, tungsten (W) of 7.0 wt. %, aluminum (Al) of 5.0 wt. %, titanium (Ti) of 1.0 wt. %, tantalum (Ta) of 7.5 wt. %, carbon (C) of 0.07 wt. %, boron (B) of 0.015 wt. %, zirconium (Zr) of 0.01 wt. %, and the rest including nickel (Ni) and other unavoidable impurities.

The Ni-base superalloy has M₂₃C₆- and M₆C-type precipitates at a grain boundary, and uniformly distributed γ′ precipitates of 0.3 to 0.4 μm in average size within a γ matrix.

The Ni-base superalloy has a creep life of 600 hours or more under the creep condition of 871° C./310 MPa.

The Ni-base superalloy has a creep life of 150 hours or more under the creep condition of 982° C./187 MPa.

The Ni-base superalloy has a weight change of 10 mg/cm² or less when a cyclic oxidation test, in which the Ni-base superalloy is kept at 1100° C. for 1 hour and cooled to room temperature, is repeated for 200 times.

Another exemplary embodiment of the present invention provides a manufacturing method of a Ni-base superalloy including: a material preparing step of preparing an alloy comprised of cobalt (Co) of 2.0 to 6.0 wt. %, chromium (Cr) of 8.0 to 12.0 wt. %, tungsten (W) of 5.0 to 9.0 wt. %, aluminum (Al) of 3.5 to 6.0 wt. %, titanium (Ti) of 3.0 wt. % or less, tantalum (Ta) of 5.0 to 10.0 wt. %, carbon (C) of 0.05 to 0.15 wt. %, boron (B) of 0.02 wt. % or less, zirconium (Zr) of 0.05 wt. % or less, and the rest including nickel (Ni) and other unavoidable impurities; a casting step of manufacturing a casting through a directional solidification process to the alloy; a solution treatment step of performing a homogenization heat treatment to the casting at 1280° C. for 4 hours; a first aging step of performing an aging process to the casting at 1080° C. for 4 hours; and a second aging step of heat treating the casting at 871° C. for 24 hours to complete the Ni-base superalloy.

The heat treatment steps are carried out in a vacuum or an inert gas atmosphere.

In the first and second aging steps, M₂₃C₆- and M₆C-type carbides are precipitated at a grain boundary, and a γ matrix has uniformly distributed γ′ precipitates of 0.3 to 0.4 μm in average size.

The Ni-base superalloy and the manufacturing method thereof have advantages of excellent high-temperature creep and oxidation resistance, and remarkably reduced in manufacturing cost.

Further, there is an advantage of improvement in directional solidification castability by optimizing the composition ratio of alloying elements.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a table showing the chemical compositions of a Ni-base superalloy in an exemplary embodiment and Comparative Examples according to the present invention;

FIG. 2 is a transmission electron microscopic image of a grain boundary of the Ni-base superalloy according to the present invention;

FIG. 3 is a process flowchart showing the manufacturing method of the Ni-base superalloy according to the present invention;

FIG. 4 is a graph showing the heat treatment conditions in each step of the manufacturing method of the Ni-base superalloy according to the present invention;

FIG. 5 provides the microstructures in dendritic core and interdendritic area after directional solidification of the Ni-base superalloy according to the present invention;

FIG. 6 is a table showing heat treatment conditions of the manufacturing method of the Ni-base superalloy in Example and Comparative Examples according to the present invention;

FIG. 7 and FIG. 8 are graphs showing the creep curves of the alloys heat treated under the conditions of the manufacturing method of the Ni-base superalloy in Example and Comparative Examples according to the present invention;

FIG. 9 to FIG. 12 provide micrographs in each step of the manufacturing method of the Ni-base superalloy in an exemplary embodiment of the present invention;

FIG. 13 is a graph showing the result of cyclic oxidation test carried out for the Ni-base superalloy according to the present invention;

FIG. 14 provides graphs comparing the creep properties between the Ni-base superalloy according to the present invention and Comparative Examples under the creep condition of 815° C. and 483 MPa;

FIG. 15 provides graphs comparing the creep properties between the Ni-base superalloy according to the present invention and Comparative Examples under the creep condition of 871° C. and 310 MPa;

FIG. 16 provides graphs comparing the creep properties between the Ni-base superalloy according to the present invention and Comparative Examples under the creep condition of 982° C. and 187 MPa; and

FIG. 17 is a table summarizing the results shown in FIG. 14 to FIG. 16.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Hereinafter, referring to FIG. 1 and FIG. 2, a composition and an internal structure of the Ni-base superalloy according to the present invention will be explained.

Terms or words used in this specification and claims are not limited to a general definition or a dictionary definition, and should be interpreted as a meaning or a concept conforming to the technical idea of the present invention based on a principle in which an inventor can properly define the terms or the words in order to describe his/her invention in the best way.

Therefore, Example described in the present specification and the configurations illustrated in the drawings are just one exemplary embodiment of the present invention and do not represent all the technical concepts of the present invention, and thus, it should be understood that various equivalents and modifications may be present which can replace them at the time of application of the present invention.

FIG. 1 is a table showing the chemical compositions of a Ni-base superalloy in an exemplary embodiment and Comparative Examples according to the present invention. Alloys used in Comparative Example 1 and Comparative Example 2 are the directionally solidified commercial Ni-base superalloys IN738LC and GTD-111 typically used as turbine blades in gas turbine model 501F by Siemens-Westinghouse and GE 7FA by General Electric, respectively.

A superalloy according to the present invention is developed so as to be applicable to a turbine blade of a gas turbine which is operated at elevated temperatures. As shown in FIG. 1, the superalloy is designed to improve the directional solidification castability by optimizing a composition ratio of alloying elements, and as compared with Comparative Examples, the superalloy is further designed to improve high-temperature creep and oxidation resistance, and to reduce the manufacturing cost by excepting high-priced rare elements such as molybdenum (Mo), rhenium (Re), and ruthenium (Ru) from its alloy chemistry.

That is, the superalloy is comprised of cobalt (Co) of 2.0 to 6.0 wt. %, chromium (Cr) of 8.0 to 12.0 wt. %, tungsten (W) of 5.0 to 9.0 wt. %, aluminum (Al) of 3.5 to 6.0 wt. %, titanium (Ti) of 3.0 wt. % or less, tantalum (Ta) of 5.0 to 10.0 wt. %, carbon (C) of 0.05 to 0.15 wt. %, boron (B) of 0.02 wt. % or less, zirconium (Zr) of 0.05 wt. % or less, and the rest including nickel (Ni) and other unavoidable impurities.

To be more specific, the superalloy is comprised of cobalt (Co) of 4.0 wt. %, chromium (Cr) of 10.0 wt. %, tungsten (W) of 7.0 wt. %, aluminum (Al) of 5.0 wt. %, titanium (Ti) of 1.0 wt. %, tantalum (Ta) of 7.5 wt. %, carbon (C) of 0.07 wt. %, boron (B) of 0.015 wt. %, zirconium (Zr) of 0.01 wt. %, and the rest including nickel (Ni) and other unavoidable impurities.

Furthermore, as shown in FIG. 2, a grain boundary of the superalloy is decorated with discretely precipitated M₂₃C₆- and M₆C-type fine carbides.

Due to the above-described characteristics, the superalloy can be improved in high-temperature creep and oxidation resistance.

Comparative Examples 1 and 2 shown in FIG. 1 are materials used for manufacturing a turbine blade of a conventional gas turbine. They contain certain amount of molybdenum and have higher contents of cobalt and chromium, but lower contents of tungsten and tantalum compared to the Example.

The chemical composition of the superalloy is determined from a superalloy design program (Program Registration No. C-2013-004753) by considering several factors such as the solid-solution hardening, precipitation strengthening, microstructural stability related to the formation of TCP (Topologically Close-Packed) phase, oxidation resistance, directional solidification castability, and the price of raw materials. Each element is limited for the following reasons:

Cobalt (Co): 2.0 to 6.0 wt. %

Cobalt is an essential alloying element for a Ni-base superalloy currently used and is known to increase high-temperature strength not only by solid-solution hardening of a matrix γ of the Ni-base superalloy but also by reducing the stacking fault energy of the matrix. In the case of a cobalt content of less than 2.0%, a solid-solution hardening effect of the alloy is reduced, and thus, improvement in creep resistance cannot be anticipated. In the case of cobalt addition of more than 6.0%, the formation of detrimental TCP phase is promoted, and thus a high-temperature microstructural stability and a mechanical property of the alloy can be decreased.

Chromium (Cr): 8.0 to 12.0 wt. %

The main purpose of the chromium addition in a Ni-base superalloy is to improve hot corrosion and oxidation resistance. Further, chromium is a main forming element of M₂₃C₆- and M₆C-type carbides, which are usually formed at grain boundary, thus improving a creep resistance by suppressing grain boundary sliding under a high-temperature creep condition.

In the case of the chromium addition of less than 8.0%, hot corrosion resistance of the alloy is decreased. In the case of the chromium addition of more than 12%, a solid-solution hardening effect is decreased and a TCP phase formation may be rapidly increased when the alloy is exposed at high temperatures.

Tungsten (W): 5.0 to 9.0 wt. %

As a dense refractory element, tungsten has a very low diffusivity in nickel so that this element greatly contributes to the solid-solution hardening of the Ni-base superalloy. Tungsten also increases the melting point of the alloy. Furthermore, tungsten is one of the main forming elements (Cr, Mo, and W) of M₂₃C₆- and M₆C-type carbides and thus contributing to the grain boundary strengthening by forming fine carbide phases on grain boundaries.

Despite of such advantages, tungsten has a strong tendency to form a brittle TCP phase. In addition, because tungsten severely segregates to a solid phase during directional solidification and single crystal casting process, increase of the tungsten addition increases the possibility for the formation of freckle defect.

Therefore, in order to improve creep strength through an appropriate solid-solution hardening effect, tungsten of 5.0% or more is added, and in order to suppress negative effects in terms of castability and high-temperature mechanical property caused by TCP phase formation, a content of tungsten is limited to 9.0%.

Molybdenum (Mo): Not contained

Molybdenum contributes to improvement in high-temperature mechanical property of the alloy by solid-solution hardening of the matrix, but it is one of the main elements to form a TCP phase with chromium and tungsten, and decreases a phase stability and oxidation resistance. Thus, molybdenum is not contained in Example of the present invention.

Aluminum (Al): 3.5 to 6.0 wt. %

Aluminum is a main element for forming a γ′ (Ni₃Al) phase of the Ni-base superalloy. In the Ni-base superalloy, aluminum improves a creep strength of the alloy by precipitation hardening of γ′ phase and contributes to the improvement in oxidation resistance of the alloy by forming a dense oxide layer. In the case of the alloy added with aluminum of less than 3.5%, a volume fraction of the γ′ precipitates is decreased so that less contribution of precipitation hardening on the creep resistance of the alloy is expected. In the case of aluminum addition of more than 6.0%, an excessive amount of γ′ phase is precipitated, and thus the creep strength of the alloy can be decreased. Furthermore, excessive amount of aluminum addition leads to the increase of γ′ solvus temperature so that it is difficult to carry out a solution heat treatment without incipient melting of the alloy due to the narrow solution heat treatment window (a temperature range between the γ′ solvus temperature and the incipient melting point of the alloy).

Titanium (Ti): 3.0 wt. % or less

Titanium is also a γ′ forming element and reinforces the γ′ phase by substituting for aluminum in the γ′ phase. Further, titanium is an element that improves the hot corrosion resistance of the alloy. As the titanium addition is increased, a volume fraction of the γ′ precipitates increases, which results in the increase of the creep resistance of the alloy. However, in the case of the alloy that having a high aluminum content of 5.0% or more, an increase of titanium addition can lead to the excessive increase of γ′ volume fraction and the formation of coarse eta (Ni₃Ti) phase in interdendritic area during the solidification of the alloy. Therefore, a phase stability and mechanical property of the alloy can be decreased. For this reason, titanium content is limited to 3.0% in Example of the present invention.

Tantalum (Ta): 5.0 to 10.0 wt. %

Tantalum is an element that not only contributes to solid-solution hardening of the matrix γ phase but also strengthens the γ′ phase by substituting for aluminum in the γ′ phase. Further, as a dense refractory element, tantalum has a strong segregation tendency to liquid phase during solidification, which leads to the increase of the liquid density at the interdendritic area. As a result, tantalum decreases the buoyancy force of the interdendritic liquid, which plays a significant role in the suppression of freckle defect during directional solidification or single crystal solidification.

Therefore, in the alloy having 5.0 to 9.0% tungsten content, tantalum addition of 5.0% or more is preferable. However, tantalum addition of more than 10.0% promotes a formation of a TCP phase such as a Mu phase, so that the creep resistance of the alloy can be deteriorated.

Carbon (C): 0.05 to 0.15 wt. %

Carbon contributes to the grain boundary strengthening of the Ni-base superalloy by forming M₂₃C₆- or M₆C-type fine carbides at grain boundaries in combination with Cr, Mo, and W. In the case of carbon addition of less than 0.05%, insufficient carbide is formed at the grain boundary. In the case of carbon addition in an amount of more than 0.15%, carbon combines with Ti and Ta during solidification and forms an excessive amount of coarse MC-type carbides within grain structure, and thus decreasing the strength of the alloy. Furthermore, in this case, film-like continuous M₂₃C₆- or M₆C-type carbides may form at the grain boundary, thereby decreasing the strength of the grain boundary. For this reason, the maximum content of carbon is limited to 0.15%.

Boron (B): 0.02 wt. % or less

Similar to carbon, boron reinforces a grain boundary and substitutes for carbon from M₂₃C₆- or M₆C-type carbide formed at the grain boundary. However, in the case of boron addition in an excessive amount, boron lowers an incipient melting point of the alloy and thus may cause local melting near eutectic microstructures during solution heat treatment. Therefore, a content thereof is limited to 0.02% or less.

Zirconium (Zr): 0.05 wt. % or less

Zirconium plays a significant role in the grain boundary strengthening together with carbon and boron. Further, if sulfur (S) as an impurity element that remarkably decreases the creep resistance of the Ni-base superalloy is introduced, zirconium forms sulfide phase by combining with sulfur, and thus contributing to the prevention of a decrease in creep strength caused by the grain boundary segregation of sulfur.

However, zirconium in an amount of more than 0.05% causes excessive grain boundary segregation of zirconium, and thus decreasing the high-temperature strength of the grain boundary. Therefore, the maximum content thereof is limited to 0.05%.

Hereinafter, a manufacturing method of the superalloy will be explained with reference to FIG. 3 and FIG. 4.

FIG. 3 is a process flowchart showing the manufacturing method of the Ni-base superalloy according to the present invention, and FIG. 4 is a graph showing temperature and time conditions in each step of the manufacturing method of the Ni-base superalloy according to the present invention.

As shown in FIG. 3 and FIG. 4, the superalloy is manufactured by a material preparing step of preparing an alloy having the above-described composition (S100); a casting step of manufacturing a casting through a directional solidification process to the alloy (S200); a solution treatment step of performing a homogenization heat treatment to the casting at 1280° C. for 4 hours (S300); a first aging step of performing an aging heat treatment to the casting at 1080° C. for 4 hours (S400); and a second aging step of heat treating the casting at 871° C. for 24 hours to complete the Ni-base superalloy (S500).

That is, the superalloy is completely manufactured by preparing an alloy having the above-described composition, manufacturing a material through directional solidification, and performing a solution treatment and two-step aging heat treatment.

FIG. 5 provides the microstructures in dendritic core and interdendritic area after directional solidification of the Ni-base superalloy according to the present invention. There is a difference in composition between the dendritic core and the interdendritic area due to the segregation of the alloying elements during directional solidification. In particular, the precipitation of γ′ phase is promoted in the interdendritic area where the γ′ forming elements are severely segregated, and thus higher fraction of coarse γ′ is precipitated compared with the dendritic core area. At a later stage of solidification, a γ/γ′ eutectic is formed at the interdendritic area.

Accordingly, the solution treatment step (S300) is carried out to homogenize the segregation of the alloying elements and to dissolve coarse γ′ precipitates after the casting step (S200). For the solution treatment step (S300), three different conditions are selected based on the result of differential scanning calorimetry (DSC) by considering the dissolution of γ′ phase, time, and temperature related to diffusion.

FIG. 6 is a table showing the heat treatment conditions of the manufacturing method of the Ni-base superalloy in Example and Comparative Examples according to the present invention. The solution treatment step (S300) in Example is carried out at 1280° C. for 4 hours, and the first aging step (S400) is carried out at 1080° C. for 4 hours-followed by the second aging step (S500) at 871° C. for 24 hours in a vacuum or an inert gas atmosphere.

In a manufacturing method according to Comparative Example 1, temperature increases in a step-like form during a solution treatment step, and in a manufacturing method according to Comparative Example 2, only two-step aging heat treatment is carried out without a solution treatment step.

The results thereof are shown in FIGS. 7 and 8.

FIG. 7 and FIG. 8 are graphs showing the creep curves of the alloys heat treated under the conditions of the manufacturing method of the Ni-base superalloy in Example and Comparative Examples according to the present invention. It was confirmed that the superalloy manufactured according to preferred Example has a significantly longer creep life than the alloys of Comparative Examples.

To be more specific, the Ni-base superalloy has a creep life of 600 hours or more under the conditions of 871° C./310 MPa and 150 hours or more under the conditions of 982° C./187 MPa.

Hereinafter, referring from FIG. 9 to FIG. 12, microstructural change in each step will be reviewed. FIG. 9 to FIG. 12 provide micrographs in each step of the manufacturing method of the Ni-base superalloy in an exemplary embodiment of the present invention, and each shows a dendritic core, an interdendritic area, and a grain boundary, respectively.

It can be seen in FIG. 9 and FIG. 10 that coarse and irregular-shaped γ′ precipitates observed at dendritic core and interdendritic area in as-cast state are dissolved to matrix γ and their size becomes decreased by the solution treatment step (S300). However, the microstructure of grain boundary is not affected by the solution treatment step (S300) and is found to be composed of coarse MC-type carbides and γ′ precipitates.

FIG. 11 shows the microstructure after the first aging step (S400). It can be seen in FIG. 11 that the difference in size of γ′ precipitates between dendritic core and interdendritic area becomes uniform through the first aging step (S400). The average size of γ′ precipitates in dendritic core and interdendritic area is measured as 0.3 to 0.4 μam.

Furthermore, it can be seen that the first aging step (S400) leads to the precipitation of fine M₂₃C₆- and M₆C-type carbides on the grain boundary. Therefore, after the first aging step (S400), the microstructure of grain boundary is composed of M₂₃C₆- and M₆C-type carbides as well as MC-type carbides and γ′ precipitates.

FIG. 12 shows the microstructure after the second aging step (S500). It can be seen in FIG. 11 and FIG. 12 that a heat treatment performed at 871° C. for 24 hours has little effect on the size and morphology of the γ′ precipitates in dendritic core and interdendritic area. However, it is found that the second aging step (S500) increases not only a volume fraction of the γ′ precipitates in dendritic core and interdendritic area but also a fraction of fine M₂₃C₆^- and M₆C-type carbides precipitated on the grain boundary.

It can be seen that the superalloy manufactured as described above is improved in high-temperature oxidation resistance as shown in FIG. 13.

That is, FIG. 13 is a graph showing the result of cyclic oxidation test for 200 cycles in which the Ni-base superalloy according to the present invention is kept at 1100° C. for 1 hour and cooled to room temperature. It can be said that a smaller weight decrement per unit area (mg/cm²) means a better oxidation resistance.

As shown in FIG. 13, the superalloy manufactured as described above exhibits a similar oxidation resistance to an alloy IN738LC, which is one of Comparative Example, but is greatly improved in high-temperature oxidation resistance as compared with the other Comparative Example, an alloy GTD-111.

Hereinafter, a comparison of the creep resistance between the Ni-base superalloy according to the present invention and Comparative Examples will be explained with reference from FIG. 14 to FIG. 17.

FIG. 14 to FIG. 16 provide graphs comparing a creep life and a creep strain between the Ni-base superalloy according to the present invention and Comparative Examples in different creep conditions, and FIG. 17 is a table summarizing the results shown in FIG. 14 to FIG. 16.

Further, the lower graph in each of FIG. 14 to FIG. 16 is an enlarged view of a part of the upper graph to check a time to reach 1.0% of creep strain.

As can be seen from FIG. 14 to FIG. 17, the Ni-base superalloy according to the present invention is remarkably improved in creep-rupture life and time to reach 1.0% of creep strain compared with the most widely used alloys IN738LC and GTD-111 of Comparative Examples under various temperature and stress conditions.

As described above, it can be seen that the superalloy according to the present invention is greatly improved in oxidation resistance and creep strength at elevated temperatures as compared with Comparative Examples.

The scope of the present invention is not limited to the above-described Example, and the present invention can be modified in various ways by those skilled in the art based on the present invention within the above-described technical scope.

Claims

1. A Ni-base superalloy comprising:

2.0 to 6.0 wt. % of cobalt (Co), 8.0 to 12.0 wt. % of chromium (Cr), 5.0 to 9.0 wt. % of tungsten (W), 3.5 to 6.0 wt. % of aluminum (Al), 3.0 wt. % or less of titanium (Ti), 5.0 to 10.0 wt. % of tantalum (Ta), 0.05 to 0.15 wt. % of carbon (C), 0.02 wt. % or less of boron (B), 0.05 wt. % or less of zirconium (Zr), with the remainder being composed of nickel (Ni) and unavoidable impurities.

2. A Ni-base superalloy comprising:

4.0 wt. % of cobalt (Co), 10.0 wt. % of chromium (Cr), 7.0 wt. % of tungsten (W), 5.0 wt. % of aluminum (Al), 1.0 wt. % of titanium (Ti), 7.5 wt. % of tantalum (Ta), 0.07 wt. % of carbon (C), 0.015 wt. % of boron (B), 0.01 wt. % of zirconium (Zr), with the remainder being composed of nickel (Ni) and unavoidable impurities.

3. The Ni-base superalloy of claim 1,

wherein the Ni-base superalloy has M23C6- and M6C-type precipitates at a grain boundary, and uniformly distributed γ′ precipitates of 0.3 to 0.4 μm in average size within a γ matrix.

4. The Ni-base superalloy of claim 3,

wherein the Ni-base superalloy has a creep life of 600 hours or more under the creep conditions of 871° C./310 MPa.

5. The Ni-base superalloy of claim 3,

wherein the Ni-base superalloy has a creep life of 150 hours or more under the creep conditions of 982° C./187 MPa.

6. The Ni-base superalloy of claim 3,

wherein the Ni-base superalloy has a change in weight of 10 mg/cm2 or less when an cyclic oxidation test, in which the Ni-base superalloy is kept at 1100° C. for 1 hour and cooled to room temperature, is repeated 200 times.

7. A manufacturing method of for preparing a Ni-base superalloy, comprising:

performing a material preparing step comprising preparing an alloy comprised of 2.0 to 6.0 wt. % of cobalt (Co), 8.0 to 12.0 wt. % of chromium (Cr), 5.0 to 9.0 wt. % of tungsten (W), 3.5 to 6.0 wt. % of aluminum (Al), 3.0 wt. % or less of titanium (Ti), 5.0 to 10.0 wt. % of tantalum (Ta), 0.05 to 0.15 wt. % of carbon (C), 0.02 wt. % or less of boron (B), 0.05 wt. % or less of zirconium (Zr), with the remainder being composed of nickel (Ni) and unavoidable impurities;

performing a casting step comprising manufacturing a casting through a directional solidification process with the alloy;

performing a solution treatment step comprising performing a homogenization heat treatment on the casting at 1280° C. for 4 hours;

performing a first aging step comprising performing an aging heat treatment on the casting at 1080° C. for 4 hours; and

performing a second aging step comprising heat treating the casting at 871° C. for 24 hours to complete the Ni-base superalloy.

8. The manufacturing method of claim 7, wherein the heat treatment steps are carried out in a vacuum or an inert gas atmosphere.

9. The manufacturing method of claim 8, wherein in the first aging step and the second aging step, M23C6- and M6C-type carbides are precipitated at a grain boundary, and a γ matrix has uniformly distributed γ′ precipitates having an average size of 0.3 to 0.4 μm.

10. The Ni-base superalloy of claim 2, wherein the Ni-base superalloy has M23C6- and M6C-type precipitates at a grain boundary, and uniformly distributed γ′ precipitates of 0.3 to 0.4 μam in average size within a γ matrix.