WO2017170433A1

WO2017170433A1 - Method for producing ni-based super heat-resistant alloy

Info

Publication number: WO2017170433A1
Application number: PCT/JP2017/012447
Authority: WO
Inventors: 奈翁也佐藤; 宙也青木; 敏明野々村
Original assignee: 日立金属株式会社
Priority date: 2016-03-31
Filing date: 2017-03-27
Publication date: 2017-10-05
Also published as: JPWO2017170433A1; JP6642843B2

Abstract

Provided is a method for producing an Ni-based super heat-resistant alloy, which suppresses abnormal growth of crystal grains and enables the achievement of a fine crystal grain structure having an ASTM crystal grain size number of 7 or greater. A method for producing an Ni-based super heat-resistant alloy that is characterized by containing, in mass%, 0.5-1.0% of Al, 17-21% of Cr, 17-19% of Fe, 4.5-5.5% of Nb, 0.8-1.3% of Ti, 3.0-6.0% of W, 0.001-0.03% of B, 0.001-0.015% of C and 1.0% or less of Mo, with the balance made up of Ni and unavoidable impurities, which comprises a hot forming step wherein hot forming is carried out so that the relational expression shown below is satisfied over the entire region of a material for hot forming, said material having the above-described composition. 0 ≥ -32 + S^-0.64887 × V^-0.12809 × exp{-14592/(273 + T) + 13.631} In the expression, T represents a heating temperature (°C); S represents an equivalent strain; and V represents an equivalent strain rate (/sec).

Description

Method for producing Ni-base superalloy

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

Increasing the combustion temperature is effective for increasing the efficiency of a gas turbine. Industrial gas turbine disks have been made of heat-resistant steel with excellent manufacturability, and high-efficiency machines in Europe and the United States seem to use Ni-based super heat-resistant alloys with excellent high-temperature strength (such as Alloy 718 and Alloy 706). It has become. Among them, Alloy 718 is excellent in high-temperature strength, but is limited to the manufacture of small members because macro segregation defects are likely to occur during casting. In addition, although Alloy 706 is inferior in strength to Alloy 718, it is applied to medium- and large-sized gas turbines because it is excellent in manufacturability of large-sized products.
Thus, in general, it has been difficult for a high-strength Ni-based alloy to achieve both high-temperature strength and manufacturability of a large product. For example, Japanese Unexamined Patent Application Publication No. 2014-51698 (Patent Document 1) The Ni-based alloy proposed in (1) has been reported as an alloy having high-temperature strength equivalent to Alloy 718 and large-size manufacturability equivalent to Alloy 706.

Since high fatigue strength is required for large rotating parts of gas turbines, it is necessary to refine crystal grains to a certain level or more. For this reason, normally, after producing a billet from an ingot, the pinning effect of the delta phase (hereinafter referred to as δ phase) is used to perform hot working in a temperature range of 920 to 1000 ° C. to obtain a fine recrystallized structure, Solution heat treatment and aging treatment, or direct aging treatment are performed. However, for example, when hot working is performed under low strain conditions in stamping forging or ring rolling, pinning of the delta phase is performed during hot working, cooling after hot working, or subsequent solution treatment. As a result, abnormal grain growth (abnormal grain-growth: hereinafter referred to as “AGG”) is caused.
As a proposal for preventing the above-mentioned AGG, there is, for example, a pamphlet of WO2015 / 151808 (Patent Document 2) targeting Alloy 718 alloy. In this proposal, influencing factors for preventing AGG are specified, and an Alloy 718 alloy material at 930 to 1010 ° C. is used for all regions of hot parts, [equivalent strain] ≧ 0.139 × [equivalent strain rate (/ sec)] ⁻ It is reported that AGG can be avoided by performing hot working that satisfies the relationship of ^0.30 .
Also, for example, in Japanese Patent Laid-Open No. 2001-123257 (Patent Document 3), an influencing factor that prevents AGG is specified, and an AGG can be avoided by applying a strain of 0.125 or more in the entire region of the part. is doing.

JP 2014-051698 A WO2015 / 151808 pamphlet JP 2001-123257 A

In parts that place importance on fatigue strength, it is necessary to have a uniform and very fine grain structure of ASTM grain size number 7 or more. In the invention described in Patent Document 2, an alloy 718 alloy material of 930 to 1010 ° C. is [equivalent strain] ≧ 0.139 × [equivalent strain rate (/ sec)] ⁻⁰ with respect to the entire part in the hot forging process. It is excellent in that AGG can be avoided by performing hot working that satisfies the ^.30 relationship. However, since the hot working conditions vary depending on the alloy composition, it is necessary to derive the relationship between the equivalent strain and the equivalent strain rate for each alloy. Furthermore, the alloy described in Patent Document 1 has a problem that the relationship between the equivalent strain and the equivalent strain rate is different in the temperature range of 930 to 1010 ° C. because the recrystallization behavior is sensitive to temperature.
In addition, the invention described in Patent Document 3 is excellent in that AGG can be prevented by subsequent solution treatment by applying a strain of 0.125 or more to the entire part in the hot forging process. However, in hot working, strain is imparted at various strain rates such as stamping forging and ring rolling, and when strain of about 0.125 is imparted at low strain rate conditions, it is still hot in the region where AGG is developed. In some cases, the processing is difficult, and a fine grain structure cannot be obtained. This problem is particularly a problem when manufacturing large forged products and ring rolled products used for upsetting forging, die forging, and ring rolling. In particular, for the alloy of Patent Document 1 which is excellent in manufacturability of large ingots, no study has been made to prevent AGG.
An object of the present invention is to provide a method for producing a Ni-base superalloy that suppresses AGG and obtains a fine grain structure having an ASTM grain size number of 7 or more, particularly with respect to the alloy described in Patent Document 1. It is.

The present invention has been made in view of the above-described problems.
That is, in the present invention, in mass%, Al: 0.5 to 1.0%, Cr: 17 to 21%, Fe: 17 to 19%, Nb: 4.5 to 5.5%, Ti: 0.8 ~ 1.3%, W: 3.0 ~ 6.0%, B: 0.001 ~ 0.03%, C: 0.001 ~ 0.1%, Mo: 1.0% or less, the balance being Ni And a Ni-based super heat-resistant alloy made of inevitable impurities, the Ni-based super heat-resistant alloy having a hot working process for performing hot working so as to satisfy the following relationship throughout the hot working material having the above composition: A method for producing a heat-resistant alloy.
0 ≧ −32 + S ^−0.64887 × V ^−0.12809 × exp {−14592 / (273 + T) +13.631}, where T is the heating temperature (° C.), S is the equivalent strain, and V is the equivalent strain rate ( / Sec)
The present invention also provides a step of performing a solid solution treatment in a range of 950 to 1000 ° C. for 0.5 to 10 hours as a preferable heat treatment condition for suppressing the expression of AGG after the hot working step, and 700 to 750. Holding the temperature in the range of 2 ° C. for 2 to 20 hours, and then cooling to 600 ° C. to 650 ° C., followed by the first effect of the second time in the range of 600 to 650 ° C. for 2 to 20 hours. A method for producing a Ni-base superalloy, comprising a step of performing a two-aging treatment.

According to the present invention, it is possible to suppress the AGG of the Ni-based superalloy and particularly to obtain a uniform fine crystal grain structure having an ASTM grain size number of 7 or more for the alloy described in Patent Document 1. . The reliability of fatigue characteristics of gas turbine members and the like using this can be improved.

It is a figure which shows the relationship between an equivalent distortion and heating temperature, and a metal structure. It is a metallographic photograph which shows the presence or absence of abnormal crystal grain growth of this invention and a comparative example. It is a figure which shows the relationship between an equivalent strain and an equivalent strain rate, and a metal structure. It is a metallographic photograph which shows the presence or absence of abnormal crystal grain growth of this invention and a comparative example. It is a side surface schematic diagram of a small compression test piece.

First, the alloy composition specified in the present invention will be described. All composition ranges are mass%.
<Al: 0.5 to 1.0%>
Al is an element that forms a gamma prime phase (hereinafter referred to as γ ′ phase) such as Ni ₃ Al, and is an element that bears the strength of the γ ′ phase precipitation strengthened Ni-based alloy. It also has the effect of improving oxidation resistance. When the amount is insufficient, the precipitation amount of the γ ′ phase due to aging decreases, and sufficient high-temperature strength cannot be obtained. From this viewpoint, the lower limit of Al is 0.5%. When the amount is excessive, the appearance of a hard and brittle harmful phase is promoted, and the solid solution temperature of the γ ′ phase is increased to reduce hot forgeability, so the upper limit is made 1.0%.
<Cr: 17-21%>
Cr is an element that improves the oxidation resistance and high-temperature corrosion resistance by forming a dense oxide film made of Cr ₂ O _{3 on the} surface. In order to utilize for the high temperature member made into object by this invention, it is required to contain at least 17%. However, if the content exceeds 21%, a σ phase (sigma phase) which is a harmful phase is formed to deteriorate the ductility and fracture toughness of the material, so the upper limit is made 21%.
<Fe: 17-19%>
Fe has higher ductility than Ni, and the hot workability is improved by containing Fe. Moreover, since it is cheaper than other elements, it is also effective in reducing the cost of materials. However, if it is contained in excess, the γ ′ phase, which is a precipitation strengthening phase, becomes unstable and the high-temperature strength decreases, so the Fe range is 17-19%.

<Nb: 4.5 to 5.5%>
Nb contributes to the improvement of the high-temperature strength as an element for precipitating the γ ′ phase in the same manner as Al and Ti. However, in the present invention, Nb is converted to a gamma double prime phase (Ni ₃ Nb) having a crystal structure very similar to the γ ′ phase. Is the main contribution. The gamma double prime phase (hereinafter referred to as the γ ″ phase) acts as a precipitation strengthening phase in the same manner as the γ ′ phase and improves the high temperature strength of the material. In order to exhibit this effect, the content of 4.5% or more is necessary. However, since the segregation characteristics decrease as the content increases, the Nb range is 4.5 to 5.5%.
<Ti: 0.8 to 1.3%>
Ti is dissolved in the γ ′ phase in the form of Ni ₃ (Al, Ti) and contributes to high temperature strength. Although the effect is recognized even with a slight content, it is necessary to contain at least 0.8% from the viewpoint of improving the segregation characteristics. If excessive, intermetallic compounds other than the γ 'phase will be formed, ductility and high-temperature workability will be impaired, and similarly to Al, the solid solution temperature of the γ' phase will be raised and hot forgeability will be deteriorated. 1.3% is the upper limit.
<W: 3.0-6.0%>
W strengthens the matrix by solid solution strengthening. From the viewpoint of segregation characteristics, the content tends to be improved as the content is increased. Therefore, the content needs to be at least 3.0%. However, if it exceeds 6.0%, formation of a hard and brittle intermetallic compound phase is promoted, and high temperature forgeability is deteriorated. Therefore, the range of W is set to 3.0 to 6.0%.

<B: 0.001 to 0.03%>
B has the effect of strengthening the grain boundary and improving the creep strength when contained in a small amount. However, excessive content causes precipitation of harmful phases and partial melting leading to lowering of the melting point, so the range of B is 0.001 to 0.03%.
<C: 0.001 to 0.1%>
C dissolves in the matrix and improves the tensile strength at high temperatures, and also improves the grain boundary strength by forming carbides such as MC and M ₂₃ C ₆ . Although these effects become remarkable from about 0.001%, excessive C content causes coarse eutectic carbides and causes toughness reduction, so 0.1% is made the upper limit.
<Mo: 1.0% or less (including 0%)>
Since the effect of Mo on the strength is the same as that of W, it is contained as necessary. Mo has an effect of strengthening the matrix phase by solid solution strengthening, and an improvement in strength is recognized even with a small amount, and the effect increases with the content. However, since the segregation characteristics are significantly deteriorated with the inclusion, 1.0% can be contained at the upper limit.

The balance is Ni and inevitable impurities, but one or more of the following elements can be included. The concentration range is shown below.
<Zr: 0.05% or less>
Zr segregates at the grain boundaries and has the effect of increasing the grain boundary strength, so 0.05% can be contained at the upper limit.
<V: 0.5% or less, Ta: 0.5% or less>
V and Ta can each contain up to 0.5% in order to stabilize the γ ′ phase and the γ ″ phase and improve the strength.
<Re: 0.5% or less>
Similar to W and Mo, Re is an element effective for improving the corrosion resistance while being dissolved in the matrix, and can be contained in an upper limit of 0.5%. However, Re is expensive, has a large specific gravity, and increases the specific gravity of the alloy. Therefore, it is preferably 0.1% or less.

Next, a hot working process such as hot forging, which is the greatest feature of the present invention, will be described. The greatest feature of the present invention is that it optimizes hot working conditions for various strain rates such as stamping forging and ring rolling, and further prevents abnormal crystal grain growth by optimizing subsequent cooling conditions and heat treatment conditions. There is.
<Hot working process>
In order to obtain a fine grain structure, the hot working material is heated before hot working. By this heating, the temperature of the material for hot working is set in a range of 930 to 1000 ° C., and recrystallization is promoted during hot working such as hot forging. Note that the temperature of the heated material for hot working is a heating temperature T (° C.). Recrystallization hardly occurs when the temperature of the material for hot working by heating before hot working is less than 930 ° C. On the other hand, if the temperature of the material for hot working by heating before hot working exceeds 1000 ° C., recrystallization during hot working is promoted, but the size of the recrystallized grains to be generated increases, so that fine grains are obtained. It becomes difficult. Therefore, the temperature of the raw material for hot working by heating before hot working is set to 930 to 1000 ° C. The minimum of preferable heating temperature is 950 degreeC, More preferably, it is 970 degreeC. Moreover, the upper limit of preferable heating temperature is 990 degreeC.
In the present invention, the hot working is performed so that the following relationship is satisfied in the entire area of the hot working material.
0 ≧ −32 + S ^−0.64887 × V ^−0.12809 × exp {−14592 / (273 + T) +13.631}, where T is the heating temperature (° C.), S is the equivalent strain, and V is the equivalent strain rate ( / Sec)
The above relational expression is obtained by performing multiple observations on the relationship between the equivalent strain, the equivalent strain rate, and the heating temperature at which the crystal grain size number is 7 or more by observing the structure. From the viewpoint of high temperature creep strength, the lower limit of the right side of the relational expression is preferably −20, and more preferably close to zero.

The above relational expression is applied to an equivalent strain assumed in hot working such as hot forging including upset forging, die-casting forging, hot die forging, and isothermal forging, and 0.001-5. 0 can be set. A preferable upper limit of the equivalent strain is 4, more preferably 3.5. A preferred lower limit of the equivalent strain rate is 0.001, more preferably 0.005. A preferred upper limit of the equivalent strain rate is 1. The equivalent strain and the equivalent strain rate represent the strain and strain rate when the vertical and shear 6-axis elements are converted into a single axis.
AGG is expressed when the crystal grain size before hot working is ASTM grain size number 7 or more, and the sensitivity becomes higher as the initial crystal grain is finer. As shown in Table 2, the lower the heating temperature is, the more AGG is suppressed, and the amount of strain required for the grain size number after hot working to be 7 or more becomes smaller. This is because crystal grain growth is suppressed at lower temperatures. In addition, the preferable ASTM grain size number in which AGG is suppressed after hot working is 8 or more.
Further, as the strain rate is lower, AGG is more likely to be generated, and the amount of strain required for the grain size number after hot working to be 7 or more increases. Under low strain rate conditions, for example, strain is accumulated again in the dynamic recrystallization that occurs during stamping forging, so the grain boundaries move during the solution treatment using the accumulated energy of the grain boundaries as the driving force. caused by.

In another embodiment of the present invention, as the hot working conditions, S ≧ 0 over the entire area of the hot working material that has been subjected to heat treatment at 930 to 1000 ° C., preferably 970 to 990 ° C., before hot working. 180 × V ^−0.122 (S is equivalent strain, V is equivalent strain rate (/ sec)). Application of this relational expression is equivalent to 5 or less equivalent strain assumed in hot working such as ring mill in addition to hot forging including upset forging, die forging, hot die forging, and isothermal forging, and equivalent strain rate of 0.0001. ~ 10. A preferable upper limit of the equivalent strain is 4, more preferably 3.5. A preferred lower limit of the equivalent strain rate is 0.001, more preferably 0.005. A preferred upper limit of the equivalent strain rate is 5, more preferably 1. The equivalent strain and the equivalent strain rate represent the strain and strain rate when the vertical and shear 6-axis elements are converted into a single axis.
AGG is expressed when the crystal grain size before hot working is ASTM grain size number 7 or more, and the sensitivity becomes higher as the initial crystal grain is finer. As shown in FIG. 3, the lower the strain rate, the larger the range B in which AGG is promoted. This is because, under the condition of low strain rate, for example, strain is accumulated again in dynamic recrystallization that occurred during stamping forging. Due to moving. On the other hand, the region A is a region where crystal grain refinement by recrystallization is possible and AGG is also suppressed.
Therefore, in the present invention, an appropriate strain is applied to the entire area of the hot working material so as to satisfy the following relational expression that allows hot working in the region A, and a preferable grain size number is used to prevent AGG more reliably. Adjust to 8 or more. .
S ≧ 0.180 × V ^−0.122 (S is equivalent strain, V is equivalent strain rate (/ sec))
The relational expressions indicating the regions A and B are obtained by performing multiple observations on the relationship between the equivalent strain in which AGG occurs and the equivalent strain rate based on the results of the tissue observation.

Next, preferable heat treatment conditions in the case of performing the solution treatment and the aging treatment after the hot working process described above will be described.
<Solution treatment process>
In order to maintain the fine recrystallized structure obtained in the hot working process, the heating temperature during the solution treatment is also important. If the heating temperature of the solution treatment is less than 950 ° C., the δ phase is excessively precipitated during the solution treatment, so that the amount of γ ″ phase precipitated in the subsequent aging treatment is reduced and the overall strength is lowered. . On the other hand, when the solution treatment temperature exceeds 1000 ° C., crystal grains grow with a decrease in the pinning effect of the δ phase, and the tensile strength and fatigue strength decrease. Therefore, the solution treatment temperature is 950 to 1000 ° C. The lower limit of the preferred solution treatment temperature is 960 ° C., and the upper limit of the preferred solution treatment temperature is 990 ° C. The retention time for the solution treatment is 0.5 to 10 hours. If it is less than 0.5 hour, the solid solution effect of the compound precipitated during cooling after the end of hot working is low. On the other hand, treatment exceeding 10 hours is economically inefficient and may cause growth of fine crystal grains. The lower limit of the preferable retention time of the solution treatment is 1 hour, and the upper limit of the preferable retention time of the solution treatment is 4 hours.

<Aging process>
The Ni-base superalloy subjected to solution heat treatment is held at 700 to 750 ° C. for 2 to 20 hours, then cooled to 600 to 650 ° C., and then held at 600 to 650 ° C. for 2 to 20 hours. Perform aging treatment. The purpose of the aging treatment is to obtain a high strength at a high temperature by finely precipitating the γ ′ phase and γ ″ phase of the precipitation strengthening phase. Only the second aging treatment on the low temperature side takes too much time to precipitate the precipitation strengthening phase, so as the first temporary treatment, aging treatment is performed on the high temperature side to promote precipitation of the γ 'phase and γ''phase. Let If the temperature of the first temporary effect treatment is less than 700 ° C., the effect of promoting precipitation is insufficient, so that the effect of precipitation strengthening is reduced. On the other hand, when the temperature of the first temporary treatment exceeds 750 ° C., precipitation is further promoted, but the size of the precipitated particles is increased and the effect of precipitation strengthening is reduced, and the γ ″ phase has no precipitation strengthening ability. Transformation into δ phase. Therefore, the temperature of the first temporary treatment is in the temperature range of 700 to 750 ° C. The lower limit of the temperature of the preferred first-effect treatment is 710 ° C., and the upper limit of the temperature of the preferred first-effect treatment is 730 ° C. On the other hand, if the time of the first temporary treatment is less than 2 hours, the precipitation of the γ ′ phase and the γ ″ phase becomes insufficient. On the other hand, if the time of the first temporary treatment exceeds 20 hours, the effect of precipitation of the γ ′ phase and the γ ″ phase is saturated, which is not economical. Therefore, the holding time of the first temporary treatment is set in the range of 2 to 20 hours. The lower limit of the holding time of the preferred first-effect treatment is 4 hours, and the upper limit of the holding time of the preferred first-effect treatment is 15 hours.
A second aging process is performed after the first temporary effect process described above. If the temperature of the second aging treatment is less than 600 ° C., it takes too much time to precipitate the γ ′ phase and γ ″ phase, which is not efficient. On the other hand, if the temperature of the second aging treatment exceeds 650 ° C., the temperature difference from the temperature of the first aging treatment is small, so that the driving force for precipitation is insufficient and the amount of precipitation is reduced. Therefore, the temperature of the second aging treatment is in the temperature range of 600 to 650 ° C. The minimum of the temperature of a preferable 2nd aging treatment is 610, and the upper limit of the temperature of a preferable 2nd aging treatment is 630 degreeC. The holding time of the second aging treatment is set to 2 to 20 hours for the same reason as the first temporary aging treatment described above. The lower limit of the preferred second aging treatment is 4 hours, and the preferred upper retention time of the second aging treatment is 15 hours.

About 2 tons of billet manufactured from a large ingot having a chemical composition corresponding to the Ni-base superalloy having the composition shown in Patent Document 1 shown in Table 1 is subjected to upset forging at a temperature range of 950 to 1000 ° C., and then The sample was held at 970 ° C. for 2.5 hours and then air-cooled to produce a small compression test piece shown in FIG. Using this small compression test piece as a test material, a hot working test was conducted to investigate factors affecting the generation of AGG. The crystal grain size of the test material was an average crystal grain size number 10 as measured by ASTM-E112.

Regarding the factors causing AGG, the effects of strain amount, strain rate and temperature were investigated.
As the first of the investigation, heating temperature 927 ° C., 954 ° C., 982 ° C., reduction ratio 30%, nominal strain rate calculated by compression rate with respect to test piece height before compression, 0.5 / second, 0.05 / second, The compression test was performed under the conditions of 0.005 / second and a cooling rate after compression of 540 ° C./min. Then, the solution treatment for 1 hour was performed at the same temperature as heating temperature, the structure of the longitudinal section was observed with an optical microscope, and the crystal grain size number was measured by a comparative method. The equivalent strain and the equivalent strain rate at the position where the structure was observed were calculated by inputting the heating temperature, the compression rate, the nominal strain rate, and the cooling rate after compression using a commercially available forging analysis software DEFORM. When the grain size number after the solution treatment was less than 7, it was determined that AGG was not suppressed.
Table 2 shows the determination results of AGG, and FIG. No. 7 metallographic photograph and Comparative Example No. 26 metallographic photographs are shown. From the results shown in Table 2, the relationship between the equivalent strain having a grain size number of 7 or more, the equivalent strain rate, and the heating temperature was calculated by multiple regression to obtain the following relational expression.
0 ≧ −32 + S ^−0.64887 × V ^−0.12809 × exp {−14592 / (273 + T) +13.631}, where T is the heating temperature (° C.), S is the equivalent strain, and V is the equivalent strain rate ( / Sec)
Further, FIG. 1 shows an appropriate range of the equivalent strain and the equivalent strain amount with respect to the target grain size number 7 or more after hot working. FIG. 1 also shows that AGG can be prevented by applying the manufacturing method defined in the present invention. In addition, when manufacturing an actual product, using the above-described commercially available forging analysis software DEFORM, from the results calculated by inputting the heating temperature, compression rate, nominal strain rate, and cooling rate after compression, Manufacturing conditions can be determined by obtaining hot working conditions that satisfy the relational expression over the entire area.

A compression experiment was conducted at a nominal strain rate of 0.05 / second assuming a general die forging rate with a target grain size number of 8 or more that suppresses AGG as a target. No. in the experimental data shown in Table 2. For 1 to 4, 6 to 10, 12, 13, 23, 25, 26, 29 and 30, the results of investigating the effects of strain amount, strain rate and temperature on AGG generation are also shown. The compression test was performed under the conditions of heating temperatures of 927 ° C., 954 ° C., 982 ° C., a reduction rate of 30%, and a cooling rate after compression of 540 ° C./min. Then, the solution treatment for 1 hour was performed at the same temperature as heating temperature, the structure of the longitudinal section was observed with an optical microscope, and the crystal grain size number was measured by a comparative method. The equivalent strain and the layer equivalent strain rate at the position where the structure was observed were calculated by inputting the heating temperature, the compression rate, the nominal strain rate, and the cooling rate after compression using a commercially available forging analysis software DEFORM. When the crystal grain size number after the solution treatment was less than 8, it was determined that AGG was not suppressed.
With respect to the above test piece, the following relational expression can be obtained by calculating the relationship among the equivalent strain, the equivalent strain rate, and the heating temperature with a grain size number of 8 or more by multiple regression.
0 ≦ 78.212 + 2.4612 × 10 ⁻⁵ × (T + 273) ² −8.2603 × 10 ⁻² × (T + 273) -11.914 × (1-V) ² + 13.0682 × (1-V) +4. ^{8646 (1-R) 2 -22.135} (1-R), where, T is heating temperature (° C.), S is equivalent strain, V is equivalent strain rate (/ sec)

As the second part of the investigation, the influence of strain and strain rate was investigated on the factors causing AGG.
Compression test under the conditions of a heating temperature of 982 ° C., a reduction rate of 30%, a nominal strain rate of 0.005 to 0.5 / second calculated by the compression rate with respect to the test piece height before compression, and a cooling rate of 540 ° C./minute after compression. Went. Then, the solution treatment for 1 hour was performed at 982 degreeC, the structure | tissue was observed for the vertical cross section with the optical microscope, and it image | photographed in arbitrary positions. The crystal grain size number at the photographed position was converted to the crystal grain size number after marking the crystal grain boundary and performing image analysis to obtain the equivalent circle diameter. The equivalent strain and the equivalent strain rate at the position where the structure was observed were calculated by reproducing a hot working test using a commercially available forging analysis software DEFORM. When the grain size number after the solution treatment was less than 7, it was determined that AGG was not suppressed. Table 2 shows the determination results of AGG, and FIG. No. 17 metallographic photograph and Comparative Example No. 27 metal structure photographs are shown.
From the results shown in Table 2, the relationship of the metal structure exerted by the relationship between the equivalent strain and the equivalent strain rate in FIG. 3 was derived. Region A is a region where AGG is suppressed, and region B is a region where AGG is not suppressed. As shown in FIG. 3, it can be seen that the equivalent strain range in which AGG occurs is larger as the equivalent strain rate is smaller. From these results, the grain size number at which AGG is suppressed after hot working is 7 or more, and the preferred grain size number at which AGG is more reliably suppressed is 8 or more. Therefore, the relationship between the equivalent strain at which the grain size number is 8 or more and the equivalent strain rate was calculated by multiple regression to obtain the following relational expression. The region A in FIG. 3 satisfies the following relational expression, and it has been confirmed that AGG can be suppressed when hot working is performed so as to satisfy the region A over the entire area of the processed material.
S ≧ 0.180 × V ^−0.122 (S is equivalent strain, V is equivalent strain rate (/ sec))
It can also be seen from FIG. 4 that AGG can be prevented by applying the manufacturing method defined in the present invention.
As described above, when the manufacturing method of the present invention is applied, even when hot working is performed under a low strain condition, AGG of the Ni-base superalloy is suppressed, and the ASTM grain size number is 7th. As described above, it can be seen that a fine grain structure of No. 8 or more is obtained.

Using the Ni-based superalloy having the composition shown in Table 1 above, No. Upset forging was performed under 22 conditions to obtain a hot-worked material having a diameter of 1300 mm and a thickness of 200 mm. Thereafter, a solid solution treatment is performed at 968 ° C. for 2.5 hours, and after holding at 718 ° C. for 8 hours as the first temporary effect treatment, it is cooled to 621 ° C., and then as a second aging treatment at 621 ° C. for 8 hours. An aging treatment was performed.
A test piece for measuring the crystal grain size number and a test piece for AGG confirmation were collected from the above-mentioned aging treatment material, and when the crystal grain size and the presence or absence of AGG were confirmed, the ASTM crystal grain size number was 9.5 and the occurrence of AGG was confirmed. There wasn't.

As described above, the Ni-base superalloy that has undergone the hot working step, solution treatment step, and aging treatment step specified in the present invention, even when hot working under low strain conditions, It can be seen that the AGG of the Ni-base superalloy is suppressed, and a fine grain structure having an ASTM grain size number of 7 or more is obtained. From this, the reliability of the fatigue characteristics of a jet engine, a gas turbine member, etc. can be improved.

Claims

In mass%, Al: 0.5 to 1.0%, Cr: 17 to 21%, Fe: 17 to 19%, Nb: 4.5 to 5.5%, Ti: 0.8 to 1.3%, W: 3.0 to 6.0%, B: 0.001 to 0.03%, C: 0.001 to 0.1%, Mo: 1.0% or less, the balance being Ni and inevitable impurities In a method for producing a Ni-base superalloy having a composition, after the hot working material having the above composition is heated in a temperature range of 930 to 1000 ° C., the following relationship is applied to the entire area of the hot working material having the above composition. A method for producing a Ni-base superalloy having a hot working step of performing hot working so as to satisfy the above.
0 ≧ −32 + S −0.64887 × V −0.12809 × exp {−14592 / (273 + T) +13.631}
Here, T is the heating temperature (° C.), S is the equivalent strain, and V is the equivalent strain rate (/ sec).
After the hot working step, a solution treatment for 0.5 to 10 hours in a range of 950 to 1000 ° C., a holding in a range of 700 to 750 ° C. for 2 to 20 hours, and then to 600 to 650 ° C. The Ni base according to claim 1, comprising a step of performing a first temporary effect treatment for cooling, and a step of performing a second aging treatment for 2 to 20 hours in the range of 600 to 650 ° C following the first temporary effect treatment. Manufacturing method of super heat-resistant alloy.