Patent Application
20200131614
1. Filed of the Invention
The present invention relates to a manufacturing process of a Ni based superalloy, and relates more specifically to a manufacturing process of a Ni based superalloy and a member of the Ni based superalloy, a Ni based superalloy, a member of a Ni based superalloy, a forged billet of a Ni based superalloy, a component of a Ni based superalloy, a structure of a Ni based superalloy, a boiler tube, a combustor liner, a gas turbine blade, and a gas turbine disk achieving both of excellent workability in a manufacturing step of the Ni based superalloy and excellent high temperature strength of the Ni based superalloy.
Aiming to improve the efficiency of a gas turbine by raising the combustion temperature, improvement of the heat resistance temperature of gas turbine components has been required. Therefore, with respect to the gas turbine components, as a material excellent in high temperature strength, a Ni based superalloy has been used widely for a turbine disk and blade as well as combustor. The Ni based superalloy achieves excellent high temperature strength by solid solution strengthening effected by adding solid solution strengthening elements such as W, Mo, and Co and precipitation strengthening effected by adding precipitation strengthening elements such as Al, Ti, Nb, and Ta. In the Ni based superalloy of the precipitation strengthening type, the lattice of a γ′ (gamma prime) phase (L12 structure) which is a precipitation strengthening phase precipitates having continuity with a lattice of a γ (gamma) phase (FCC structure, matrix), forms a coherent interface, and thereby contributes to strengthening. Therefore, although the amount of the gamma prime phase just has to be increased in order to improve the high temperature strength, the workability deteriorates as the amount of the gamma prime phase is larger. Accordingly, there are problems that manufacturing of a large forged product becomes harder as the strength of the material becomes higher, and forging cannot be performed due to increase of the defect occurrence rate in forging, and so on.
As a technology for achieving both of the high temperature strength and the hot forgeability of the Ni based superalloy, there is one described in Patent Document (JP-A-2011-052308). In Patent Document 1, a Ni based superalloy is disclosed which contains, in terms of mass, C: 0.001 to 0.1%, Cr: 12 to 23%, Co: 15 to 25%, Al: 3.5 to 5.0%, Mo: 4 to 12%, W: 0.1 to 7.0%, the total of the content of Ti, Ta and Nb is 0.5% or less in terms of mass, and a parameter Ps expressed by an expression (1) (Ps=−7×(C amount)−0.1×(Mo amount)+0.5×(Al amount)) is 0.6 to 1.6.
Hot forging of a high strength Ni based superalloy whose solvus of the gamma prime phase is 1,050° C. or above is normally performed in the temperature range of 1,000 to 1,250° C. The reason of doing so is to reduce the precipitation amount of the gamma prime phase that is a strengthening factor and to reduce the deformation resistance by raising the working temperature to a temperature around the solvus of the gamma prime phase or thereabove. However, when forging is performed at a temperature around the solvus or thereabove, because the forging temperature comes close to the melting point of a workpiece, working crack is liable to be generated by partial melting and the like. Moreover, when the material whose solvus of the gamma prime phase is high as described above is hot-forged at the solves or above, the gamma prime phase that suppresses grain boundary migration and contributes to refinement of the crystal grain disappears, therefore the grain size of the gamma phase is coarsened, and the tensile strength and the fatigue strength in using the product deteriorate.
In of the circumstances described above, the of the present invention is to provide a manufacturing process of a Ni based superalloy member which achieves both of excellent workability in a manufacturing step of the Ni based superalloy of the precipitation strengthening type which contains much amount of the gamma prime phase and excellent high temperature strength of the Ni based superalloy.
The manufacturing process of a Ni based superalloy in relation with an aspect of the present invention includes a step for softening the Ni based superalloy and improving the workability, in which the step for softening the Ni based superalloy and improving the workability is a step for precipitating the gamma prime phase that is incoherent with a gamma phase that is a matrix by 20 vol % or more.
Further, the manufacturing process of a member of a Ni based superalloy in relation with an aspect of the present invention also includes a working step for working a Ni based superalloy obtained by the manufacturing process of a Ni based superalloy described above into a desired shape, and a solution-aging (heat) treatment step for obtaining a Ni based superalloy by performing a solution treatment for solid-dissolving a gamma prime incoherent phase and an aging treatment for re-precipitating a gamma prime coherent phase after the working step.
According to an aspect of the present invention, a Ni based superalloy and a member of a Ni based superalloy can be provided which are capable of significantly improving the workability by containing the gamma prime incoherent phase by 20 vol % or more after the softening treatment step in a high strength Ni based superalloy, and capable of achieving excellent high temperature strength equal to or better than that of a material of a related art in using a product.
Also, by using a Ni based superalloy manufactured using the manufacturing process of a Ni based superalloy or a member of a Ni based superalloy manufactured using the manufacturing process of a member of a Ni based superalloy in relation with an aspect of the present invention, a member of a Ni based superalloy, a component of a Ni based superalloy, and a structure of a Ni based superalloy having various shapes can be manufactured easily.
Hereinafter, an embodiment in relation with the present invention will be explained in detail. However, the present invention is not limited to the embodiment taken up here, and appropriate combinations and modifications are possible within a range not changing the gist.
The present inventors made intensive studies on the manufacturing process of the Ni based superalloy and the member of a Ni based superalloy capable of achieving the object described above. As a result, it was watched that the gamma prime phase precipitated incoherently with the gamma phase that was the matrix (hereinafter referred to as the gamma prime incoherent phase) did not contribute to strengthening, and it was found out that the workability in forging could be significantly improved by reducing the precipitation amount of the gamma prime phase precipitated coherently with the gamma phase (hereinafter referred to as the gamma prime coherent phase) by increasing the amount of the gamma prime incoherent phase in forging and by achieving the fine duplex phase mainly formed of the gamma phase and the gamma prime incoherent phase simultaneously. Also, it was found out that excellent high temperature strength in using a product could be achieved by performing the solution-aging treatment after working into a desired shape in this state, thereby reducing the gamma prime incoherent phase, and precipitating the gamma prime coherent phase again. The present invention is based on this knowledge.
Hereinafter, the basic thought of the present invention will be explained in more detail.
As shown in (I) of
As described above, the present invention is to improve the workability not by working in a state the gamma prime phase is reduced or eliminated, but by disabling the strengthening effect of the gamma prime phase. According to the manufacturing steps described above, the Ni based superalloy and the member of a Ni based superalloy can be obtained which can obtain a Ni based superalloy that can soften the material and can significantly improve the workability in working and have the high temperature strength equal to or greater than that of a related art in using (at the time of completion of the product).
Also, “gamma prime coherent phase” and “gamma prime incoherent phase” in the present invention will be explained.
Next, the manufacturing step of a member of a Ni based superalloy in relation with an aspect of the present invention will be explained.
Also, in the present invention, one obtained by performing the raw material preparation step (S1) is called “Ni based superalloy raw material”, one obtained by performing the softening treatment step (S2) is called “Ni based superalloy softening material”, and one obtained by performing the solution-aging treatment step (S5) is called “member of a Ni based superalloy”. Further, one obtained by performing the solution-aging treatment step (S5) after joining the Ni based superalloy using friction stir welding and the like is called “structure of a Ni based superalloy (joining structure of a Ni based superalloy)”. Furthermore, in the present invention, “Ni based superalloy” is to include “Ni based superalloy raw material” and “Ni based superalloy softening material” described above, and is to include one obtained by performing the working step (S4) by once or multiple times with respect to “Ni based superalloy softening material”.
Hereinafter, the steps S1 to S5 described above will be explained in detail.
With respect to the raw material preparation method of the Ni based superalloy, there is no limitation in particular, and a method of a related art can be used. More specifically, using a ready-made alloy after casting and a ready-made alloy after forging, steps of the softening treatment step described below and onward are performed. Also, as the composition of the Ni based superalloy raw material, one whose solvus of the gamma prime phase is 1,050° C. or above is preferably used. The reason of doing so will be described below in detail.
(S2: Softening Treatment Step)
The manufacturing process of the Ni based superalloy softening material of an aspect of the present invention which improves the workability at the time of the working step includes the first softening treatment step (S21) for hot forging at a temperature equal to or below the solvus of the gamma prime phase, and the second softening treatment step (S22) for slowly cooling the Ni based superalloy after the first softening treatment step from a temperature equal to or below the solvus of the gamma prime phase and equal to or above the hot forging finishing temperature described above and increasing the gamma prime incoherent phase.
(S21: First Softening Treatment Step)
As described above, the strengthening mechanism of the Ni based superalloy of the precipitation strengthening type contributes to strengthening by that the gamma phase and the gamma prime phase form the coherent interface (reference sign 9 of
As described above, the forging temperature T1 in the first softening treatment step should be equal to or above a temperature at which recrystallization of the gamma phase proceeds quickly. To be more specific, 1,000° C. or above is preferable and 1,050° C. or above is more preferable. When T1 is below 950° C., the gamma prime incoherent phase cannot be precipitated, and the effect of the present invention cannot be secured. Also, the upper limit temperature of T1 is equal to or below the solvus of the gamma prime phase as described above.
In the second softening treatment step, by raising the temperature to a temperature (T3) equal to or below the solvus of the gamma prime phase and equal to or above the hot forging finishing temperature in the first softening treatment step described above and solid-dissolving the gamma prime coherent phase precipitated into the gamma phase, a duplex phase structure mainly formed of the gamma phase and the gamma prime incoherent phase is achieved (
With respect to the slow cooling starting temperature T3 of the second softening treatment step, in order to achieve the duplex phase structure mainly formed of the gamma phase and the gamma prime incoherent phase, it is preferable to start slow cooling at a temperature equal to or below the solvus of the gamma prime phase and equal to or above the hot forging finishing temperature in the first softening treatment step described above. The reason is that the gamma prime coherent phase remains within the gamma phase particles when the slow cooling starting temperature T3 is lower than the forging temperature T1 of the first softening treatment step, and the gamma prime incoherent phase disappears when the slow cooling starting temperature T3 is more than the solvus of the gamma prime phase. However, even when the slow cooling starting temperature T3 is lower by 100° C. than the hot forging finishing temperature in the first softening treatment step described above, the effect of the present invention can be secured.
In the second softening treatment step described above, because the workability can be improved as the gamma prime incoherent phase is increased as described above, the amount of the gamma prime incoherent phase is preferably 20 vol % or more, and is more preferably 30 vol % or more. Here, the rate (vol %) of the content of the gamma prime incoherent phase is the rate (absolute amount) with respect to the entire alloy including the matrix and other precipitates. The amount of the gamma prime incoherent phase for securing the effect of the present invention is to be determined by such relative amount that up to which extent the rate of the gamma prime incoherent phase can be increased relative to the total amount of the gamma prime phase that can be precipitated, and is preferably 50 vol % or more of the total gamma prime phase amount, and is more preferably 60 vol % or more of the total gamma prime phase amount. Also, the temperature (T2) of the slow cooling finishing time described above should be lowered to a temperature at which the gamma prime incoherent phase precipitates by the amount described above, is preferably 1,000° C. or below, and is more preferably 900° C. or below. Further, with respect to the cooling method from the slow cooling finishing temperature T2 to the room temperature, in order to suppress precipitation of the gamma prime coherent phase during cooling, the cooling rate is preferable to be as fast as possible, air cooling is preferable, and water cooling is more preferable.
In order to secure excellent workability, the Vickers hardness (Hv) at the room temperature is preferably 400 or less and more preferably 370 or less, and the 0.2% proof stress at 900° C. is preferably 300 MPa or less, more preferably 250 MPa or less, and most preferably 200 MPa or less.
By performing the second softening treatment step described above, with respect to the Ni based superalloy softening material obtained after the second softening treatment step, one with 400 or less of the Vickers hardness (Hv) at the room temperature and with 300 MPa or less of the value of the 0.2% proof stress at 900° C. can be obtained. By the softening treatment steps described above, the working temperature lower limit that becomes an issue in hot working can be lowered, and it becomes possible to work at a temperature lower than the solvus of the gamma prime phase by 100° C. or more in the working step described below.
Although cooling is performed after the first softening treatment step, and the second softening treatment step is performed in
With respect to the Ni based superalloy softening material that has become a softened state in the softening treatment step described above, working is performed. There is no limitation in particular with respect to the working method of this time, not only forging work but also other plastic working method and welding or joining method are applicable, and repetitive working can be performed by combination with the softening treatment described above. More specifically, pressing, rolling, drawing, extruding, machining, friction stir welding, and the like are applicable. Also, by combination of the softening treatment step described above and the plastic working method and the like, a member for a thermal power generation plant such as a boiler tube, combustor liner, gas turbine blade and disk using the high strength Ni based superalloy in relation with an aspect of the present invention can be provided. Concrete examples of the member of a Ni based superalloy or the structure of a Ni based superalloy which can be provided by the present invention will be described below in detail.
In the present invention, there is no limitation in particular with respect to the condition of the solution treatment and the aging treatment, and the condition generally used can be applied.
Next, the composition of the Ni based superalloy raw material in relation with an aspect of the present invention will be explained.
It is preferable that the Ni based superalloy raw material in relation with an aspect of the present invention contains, in mass %, 10% or more and 25% or less of Cr, 30% or less of Co, 3% or more and 9% or less of the total of Ti, Nb and Ta, 1% or more and 6% or less of Al, 10% or less of Fe, 10% or less of Mo, 8% or less of W, 0.03% or less of B, 0.1% or less of C, 0.08% or less of Zr, 2.0% or less of Hf, and 5.0% or less of Re, with the balance including Ni and inevitable impurities.
One of more preferable aspects is the Ni based superalloy raw material containing, in mass %, 12.5% or more and 14.5% or less of Cr, 24% or more and 26% or less of Co, 5.5% or more and 7% or less of Ti, 1.5% or more and 3% or less of Al, 3.5% or less of Mo, 2% or less of W, 0.03% or less of B, 0.1% or less of C, and 0.08% or less of Zr, with the balance including Ni and inevitable impurities.
Also, one of other more preferable aspects is the Ni based superalloy containing, in mass %, 15% or more and 17% or less of Cr, 14% or more and 16% or less of Co, 4% or more and 6% or less of Ti, 1.5% or more and 3.5% or less of Al, 0.5% or less of Fe, 4% or less of Mo, 2% or less of W, 0.03% or less of B, 0.1% or less of C, and 0.08% or less of Zr, with the balance including Ni and inevitable impurities.
Also, one of other more preferable aspects is the Ni based superalloy raw material containing, in mass %, 15% or more and 17% or less of Cr, 7.5% or more and 9.5% or less of Co, 2.5% or more and 4.5% or less of Ti, 0.5% or more and 2.5% or less of the total of Nb and Ta, 1.5% or more and 3.5% or less of Al, 3% or more and 5% or less of Fe, 4% or less of Mo, 4% or less of W, 0.03% or less of B, 0.1% or less of C, and 0.08% or less of Zr, with the balance including Ni and inevitable impurities.
Hereinafter, the reason of the amount ratio and selection of the adding element will be shown.
Cr is an element improving oxidation resistance and high temperature corrosion resistance. In order to apply Cr to a high temperature member, addition at least 10 mass % or more is indispensable. However, because excessive addition thereof promotes formation of a harmful phase, Cr is to be made 25 mass % or less.
CO is a solid solution strengthening element having an effect of strengthening the matrix by addition thereof. Further, Co also has an effect of lowering the solvus of the gamma prime phase, and improves high temperature ductility. Co is to be made 30 mass % or less because excessive addition thereof promotes formation of a harmful phase.
Al is an indispensable element forming the gamma prime phase that is a precipitation strengthening phase. Further, Al also has an effect of improving oxidation resistance. Although the adding amount is adjusted according to the aimed precipitation amount of the gamma prime phase, excessive addition thereof deteriorates the workability because the solvus of the gamma prime phase is raised. Therefore, Al is to be made 1 mass % or more and 6 mass % or less.
Ti, Nb, and Ta is an important element stabilizing the gamma prime phase similarly to Al. However, excessive addition thereof causes formation of other intermetallic compounds including a harmful phase, and incurs deterioration of the workability by raising the solvus of the gamma prime phase. Therefore, the total of Ti, Nb, and Ta is to be made 3 mass % or more and 9 mass % or less.
Fe can be substituted to an expensive element such as Co and Ni, and reduces the cost of an alloy. However, Fe is to be made 10 mass % or less because excessive addition thereof promotes formation of a harmful phase.
Mo and W are important elements solid-dissolved into the matrix and strengthening the matrix. However, because they are elements having high density, excessive addition thereof causes increase of the density. Further, because the ductility lowers, the workability also deteriorates. Therefore, Mo is to be made 10 mass % or less, and W is to be made 8 mass % or less.
C, B, and Zr are elements effective in strengthening the grain boundary and improving high temperature ductility and creep strength. However, because excessive addition thereof deteriorates the workability, C is to be made 0.1 mass % or less, B is to be made 0.03 mass % or less, and Zr is to be made 0.08 mass % or less.
Hf is an element effective in improving oxidation resistance. However, because excessive addition thereof promotes formation of a harmful phase, Hf is preferably 2.0 mass % or less.
Re is an element solid-dissolved in the matrix and strengthening the matrix. Further, Re also has an effect of improving corrosion resistance. However, excessive addition thereof promotes formation of a harmful phase. Also, because Re is an expensive element, increase of the adding amount thereof involves cost increase of an alloy. Therefore Re is preferably 5.0 mass % or less.
Embodiments of the present invention will be explained below.
The composition of specimens is shown in Table 1.
With respect to the Ni based superalloy raw material having the composition shown in Table 1, specimens were manufactured under different manufacturing conditions, and evaluation of the workability and evaluation of high temperature strength were performed with respect to each specimen. In manufacturing each specimen, 10 kg each was molten by a vacuum induction heating melting method, was subjected to homogenizing treatment, and was hot-forged thereafter at 1,150 to 1,250° C., and thereby a round bar with 15 mm diameter was manufactured and was subjected to the first softening treatment step and the second softening treatment step described above. The condition of the first softening treatment step is shown in Table 2. Also, the solvus of the gamma prime phase and presence/absence of the gamma prime phase after the first softening treatment step were evaluated. The solvus of the gamma prime phase was calculated by a simulation based on thermodynamics calculation. Also, presence/absence of the gamma prime phase was evaluated by observation of the microstructure using an electron microscope with respect to the specimens. The result is also shown in Table 2.
In Table 2, with respect to the temperature T1 (hot forging temperature) of the first softening treatment step, when large cracks were generated in hot forging in manufacturing the specimen described above, the softening treatment step of the later stage was not performed and “−” was written, when hot forging of the first softening treatment step was not performed, “not performed” was written, and when a crack was not confirmed after hot forging, the temperature in hot forging was written.
As shown in Table 2, in the comparative examples 1 and 2, large cracks were generated at the time of hot forging in manufacturing the specimen. Although the effect of the present invention can be secured because presence of the gamma prime incoherent phase could be confirmed by observation of the structure after hot forging, the solvus of the gamma prime phase is most preferably 1,250° C. or below. The comparative example 3 is of a state immediately after manufacturing the specimen in which hot forging in the first softening treatment step is not performed, however the gamma prime incoherent phase is present because the hot forging temperature at the time of manufacturing the specimen was equal to or below the solvus of the gamma prime phase. Also, in the comparative example 4, because hot forging was performed at a temperature equal to or above the solvus of the gamma prime phase, the gamma prime incoherent phase did not precipitate after completion of forging. In contrast, in the comparative example 5, although hot forging was performed at a temperature equal to or above the solvus of the gamma prime phase, the gamma prime incoherent phase precipitated due to drop of the temperature which occurred during forging. With respect to the comparative examples 6 and 8 and the examples 1 to 9, in all specimens, because hot forging was performed at a temperature equal to or below the solvus of the gamma prime phase, presence of the gamma prime incoherent phase could be confirmed on the grain boundary of the gamma phase after completion of the first softening treatment step. In the comparative example 7, although hot forging was performed at a temperature equal to or below the solvus of the gamma prime phase, because forging was performed at a temperature below a temperature at which recrystallization of the gamma phase proceeded quickly (1,000° C. or above), the gamma prime incoherent phase did not precipitate.
From the results of the above, it was shown that the forging temperature T1 in the first softening treatment step for precipitating the gamma prime incoherent phase was preferable to be equal to or below the solvus of the gamma prime phase and equal to or above a temperature at which recrystallization of the gamma phase proceeded quickly. More specifically, forging at 1,000° C. or above is preferable, and the gamma prime incoherent phase cannot be precipitated at 950° C. or below. Therefore, the solvus of the gamma prime phase should be equal to or above a temperature at which recrystallization proceeds quickly, and 1,050° C. or above is preferable.
Next, the specimen was cooled slowly from the hot forging temperature T1 of the first softening treatment step to the slow cooling finishing temperature T2 at the cooling rate TA (° C./h) of each, and was thereafter cooled to the room temperature by water quenching. The condition of the second softening treatment step is shown in Table 3. Also, the amount of the gamma prime incoherent phase and the Vickers hardness at the room temperature after cooling were evaluated. With respect to the amount of the gamma prime incoherent phase, the content rate of the gamma prime incoherent phase was determined by observing the microstructure after casting, after hot forging, or after the softening treatment. More specifically, the area ratio of the gamma prime incoherent phase was calculated from the image observed by the electron microscope, and the content rate of the gamma prime incoherent phase was calculated by converting this area ratio to the volume ratio. Also, in order to evaluate hot workability after the softening treatment, each specimen was hot-forged at 950° C., those without a problem were evaluated to be “∘”, those in which slight cracks were generated were evaluated to be “Δ”, and those in which large cracks were generated and forging was hard were evaluated to be “x”.
As shown in Table 3, with respect to the examples 1 to 9, in all of the specimens, the amount of the gamma prime incoherent phase after the softening treatment step exceeded 20 vol %, the hardness satisfied 400 Hv or less, hot forging at 950° C. could be performed without a problem, and therefore improvement of the workability could be confirmed.
In contrast, in all of the comparative examples 3 to 6 in which the amount of the gamma prime incoherent phase was less than 20 vol % and the hardness was higher than 400 Hv, the cracks were confirmed during forging or after forging. In the comparative examples 5 and 6, although the gamma prime incoherent phase was present after the softening treatment step, the amount was not sufficient for suppressing the precipitation amount of the gamma prime coherent phase in forging. In the comparative example 7, although the gamma prime incoherent phase did not precipitate, the hardness was lower than 400 Hv, and hot forging at 950° C. could be performed. However, the comparative example 7 is not the case with the high strength Ni based superalloy that becomes a target of an aspect of the present invention because the solvus of the gamma prime phase is lower than 1050° C., and the equilibrium precipitation amount of the gamma prime coherent phase at 700° C. calculated by a simulation based on thermodynamics calculation (the precipitation amount of the gamma prime coherent phase that is stable in a thermodynamic equilibrium state) is 22 vol %. Therefore, it was confirmed that 20 vol % or more of the amount of the gamma prime incoherent phase after the softening treatment step was necessary in order to sufficiently secure the effect of the present invention.
Also, when the examples 1 and 2 or the examples 3 and 4 are compared to each other, under a condition the equilibrium precipitation amount of the gamma prime coherent phase at 700° C. is of a same degree and the slow cooling temperature range in the second softening treatment step is same, as the cooling rate is slower, the amount of the gamma prime incoherent phase increases and the hardness can be lowered. It is considered that the reason of it is that, because the gamma prime incoherent phase was made to grow more, the amount of the gamma prime coherent phase that precipitated during cooling mainly from the slow cooling finishing temperature to the room temperature could be reduced. In contrast, in the comparative example 8, although the gamma prime incoherent phase was precipitated after the first softening treatment step and the second softening treatment step was performed, the cooling rate was fast, the gamma prime incoherent phase did not grow, and therefore the effect of the present invention could not be secured sufficiently.
From the results of the above, it was shown that the slow cooling rate of the second softening treatment step was preferably slower than 50° C./h, more preferably 10° C./h or less, and the effect of the present invention could not be secured when the slow cooling rate of the second softening treatment step was faster than 100° C./h.
With respect to the examples 1 to 9, the 0.2% proof stress at 900° C. was 250 MPa or less in all of them. As an example, in the example 7, the 0.2% proof stress at 900° C. was 200 MPa, and very excellent hot workability was exhibited.
Therefore, by applying an aspect of the present invention before hot forging of the Ni based superalloy, the forging temperature can be lowered than the forging temperature of a related art by 100° C. or more, and hot forging can be performed easily. Also, in view of the excellent hot forgeability described above, it is needless to mention that the working step for the Ni based superalloy having been subjected to softening treatment in relation with an aspect of the present invention is not limited to hot forging, and that excellent workability is exhibited even in pressing, rolling, drawing, extruding, machining, and the like.
With respect to the examples 1 to 9, all of them showed a microstructure as shown in
From the results of the above, it was shown that, by applying the manufacturing method of the Ni based superalloy in relation with an aspect of the present invention, hot workability of the high strength Ni based superalloy that was hard in working could be significantly improved.
The example of the Ni based superalloy manufactured using the manufacturing process of a Ni based superalloy in relation with an aspect of the present invention will be shown below.
By using the forged billet of a Ni based superalloy 11 described above, a thin sheet 12 (with 3 mm or less thickness) using the high strength Ni based superalloy as shown in
Also, in friction stir welding, because the temperature of a member during working rises to approximately 900° C., the 0.2% proof stress at the working temperature can be made 300 MPa or less by applying an aspect of the present invention, and therefore friction stir welding also becomes possible. Thus, a structure of a Ni based superalloy joined by friction stir welding as shown in
Also, by using the Ni based superalloy in relation with an aspect of the present invention which has high workability, a boiler tube 15 as shown in
Further, because bending work of the thin sheet 12 described above is easy, by combination with friction stir welding, a combustor liner 16 as shown in
Also, because die forging is easy by using the forged billet of a Ni based superalloy 11 described above, by combination with machining, a gas turbine blade 17 excellent in high temperature strength as shown in
Also, by using the forged billet of a Ni based superalloy 11 described above, a gas turbine disk 18 as shown in
As explained above, it was proved that, according to the present invention, it was possible to provide a manufacturing process of a Ni based superalloy and a member of a Ni based superalloy which achieved both of excellent workability in a manufacturing step of the Ni based superalloy of the precipitation strengthening type which contained much amount of the gamma prime phase and excellent high temperature strength of the Ni based superalloy. Also, it was proved that, by using the manufacturing process of a Ni based superalloy in relation with an aspect of the present invention, a member of a Ni based superalloy, a component of a Ni based superalloy, and a structure of a Ni based superalloy having various shapes could be easily manufactured.
Further, the embodiments described above were explained specifically in order to assist understanding of the present invention, and the present invention is not limited to those including all configurations explained. For example, a part of a configuration of an embodiment can be replaced by a configuration of another embodiment, and it is also possible to add a configuration of another embodiment to a configuration of an embodiment. Furthermore, with respect to a part of a configuration of each embodiment, it is possible to effect deletion, replacement by another configuration, and addition of another configuration.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2014-125399 | Jun 2014 | JP | national |
The present application is a continuation of U.S. application Ser. No. 14/742,475, filed on Jun. 17, 2015, which claims priority from Japanese Patent Application No. 2014-125399, filed on Jun. 18, 2014, the entire contents of both which applications are hereby expressly incorporated by reference into this application.
| Number | Date | Country | |
|---|---|---|---|
| Parent | 14742475 | Jun 2015 | US |
| Child | 16725308 | US |
