JPWO2019189576A1 - Alloy plate and manufacturing method thereof - Google Patents

Abstract

This alloy plate is% by mass, C: 0.0200% or less, Si: 0.02 to 2.00%, Mn: 0.02 to 2.00%, P: 0.050% or less, S: 0. 0.0100% or less, Cr: 12.00 to 30.00%, Ni: 35.00 to 60.00%, N: 0.0200% or less, V: 0.02 to 1.00%, Nb: 2. > 00%, 3.50% or less, Al: more than 2.60%, 4.00% or less, and optionally one of Ti, Mo, W, Cu, Co, B, Zr, Ca, Mg. It contains the above, the balance is 10.00% or more of Fe, and a chemical composition that is an impurity, the microstructure includes an austenite phase, and the austenite phase has an average crystal grain size of 20 μm or more. The average value of L1 / L2, which is the aspect ratio of the crystal grains of the phase, is 1.0 to 3.0. In the microstructure, no more than 15.0% is the area ratio of γ'-phase is 0.2% yield strength at room temperature 400MPa or more.

Description

The present invention relates to an alloy plate and a method for manufacturing the same.
The present application claims priority based on Japanese Patent Application No. 2018062547 filed in Japan on Mar. 28, 2018 and 2018-223493 filed in Japan on Nov. 29, 2018, and the contents thereof are claimed. Is used here.

  Regarding an engine exhaust system of an automobile, exhaust components such as an exhaust manifold, a turbocharger, a catalytic converter, and an exhaust pipe are connected from an exhaust gas outlet of the engine using flanges and the like. At the connection part such as the flange, a gasket or the like is inserted to ensure the airtightness so that the exhaust gas does not leak outside.

  The gasket (exhaust system gasket) used in the exhaust system as described above is a seal component used as a heat-resistant member in an engine exhaust system member of an automobile, and includes a metal thin plate having a plate thickness of 0.1 to 0.3 mm and a bead. It is manufactured by forming a so-called stepped portion. The parts are used by being sandwiched between connecting parts such as flanges in an exhaust system of an automobile, and play a role of preventing exhaust gas from leaking from the connecting parts by a repulsive force of a bead. Since the exhaust system gasket is used under severe conditions in which it is exposed to an environment where high temperature exhaust gas is present for a long time, it is required to have high high temperature strength even after being kept at a high temperature for a long time. In particular, the turbo gasket used for the connection portion with the turbocharger is used in a high temperature environment exceeding 700 ° C. Therefore, as a material of the turbo gasket, a cold-rolled thin plate of precipitation hardening Ni-based alloy such as Inconel 718 is currently used. However, in recent years, due to the increase in operating temperature, existing heat-resistant alloys such as Inconel 718 have been unable to obtain a predetermined high-temperature strength.

  For example, in Patent Document 1, C: 0.15% or less, Si: 1.0% or less, Mn: 0.3% or less, Ni: 30 to 49%, Cr: 10 to 18%, Al in mass%. : 1.6 to 3.0%, 1.5 to 8.0% of one or more elements selected from the IVa group and the Va group in total, and the balance essentially except impurities. An Fe-Ni-Cr-based superheat-resistant alloy characterized by comprising Fe is disclosed.

In Patent Document 2, C: 0.02 to 0.30%, Si: 0.02 to 3.5%, Mn: 0.02 to 2.5%, Ni: 10 to 50% by mass%. Cr: 12-25%, Ti: 1.0-5.0%, Al: 0.002-1.0%, and Nb: 0.1-3.0%, B: 0.001- 0.01, Mo: heat-resistant stainless steel containing one or more elements selected from 0.1 to 4.0% and having a total content of Ti, Al and Nb of 3.0 to 7.0%. In steel, the η phase [Ni ₃ Ti: hcp structure] that precipitates at grain boundaries and the γ ′ phase [Ni ₃ (Al, Ti, Nb)] that precipitates in the γ phase (austenite) crystal grains that are the matrix weight ratio "{eta phase _[Ni 3 Ti: hcp structure] / gamma prime phase _{[Ni 3 (Al, Ti,} Nb)]} × 100 (%) " is more than 0.01% 30.00% more than , And the is not less 800 N / mm ² or more hot tensile strength at 600 ° C., also a base γ phase (austenite) gamma prime phase precipitates in crystal grains _{[Ni 3 (Al, Ti,} Nb)] Disclosed is a heat resistant stainless steel characterized in that the spherical particles have a diameter of 1 nm or more and 20 nm or less.

  In Patent Document 3, C: 0.01 to 0.1%, Si: 2% or less, Mn: 2% or less, Cr: 12 to 25%, Nb + Ta: 0.2 to 2.0% in mass%. , Ti: less than 1.5%, Al: 0.5 to 3.0%, Ni: 25 to 45%, Cu: 0.1 to 5.0%, W: 3% or less, Mo: 3 % Or less, V: 1% or less, Co: 5% or less, B: 0.001 to 0.01%, Zr: 0.001 to 0.1%, Ca + Mg: 0.001 to 0.01%. Cold workability and over-aging characteristics including one or more kinds and satisfying Ti / Al = 0.115 to 1.0, 1 / 2W + Mo + V <3%, Ti + Al + Nb + Ta: 4.5 to 7.0%. Excellent heat resistant alloys are disclosed.

Japanese Patent Laid-Open No. 7-109539 Japanese Patent Laid-Open No. 2000-109955 Japanese Patent Laid-Open No. 10-130790

  According to the results of studies by the present inventors, in the gaskets using the alloys disclosed in these prior arts, when used for a long time at a high temperature required for exhaust system gaskets in recent years, It has been found that sufficient high temperature strength cannot be secured due to a decrease in high temperature strength.

  The present invention has been made in view of the above problems. It is an object of the present invention to provide an alloy plate and a method for producing the same, which ensures sufficient high temperature strength even when used as a gasket that is used at high temperature for a long time.

That is, the gist of the present invention is as follows.
(1) The alloy plate according to one aspect of the present invention is, in mass%, C: 0.0200% or less, Si: 0.02 to 2.00%, Mn: 0.02 to 2.00%, P: 0.050% or less, S: 0.0100% or less, Cr: 12.00-30.00%, Ni: 35.00-60.00%, N: 0.0200% or less, V: 0.02- 1.00%, Nb: more than 2.00%, 3.50% or less, Al: more than 2.60%, 4.00% or less, Ti: 0 to 0.80%, Mo: 0 to 2.00% , W: 0 to 2.00%, Cu: 0 to 1.00%, Co: 0 to 1.00%, B: 0 to 0.0100%, Zr: 0 to 0.0100%, Ca: 0 to 0.0050%, Mg: 0 to 0.0050%, the balance is 10.00% or more of Fe, and the chemical composition is an impurity, and the microstructure is austenite. Phase, the average crystal grain size of the austenite phase is 20 μm or more, and when the major axis length of the austenite phase crystal grains is L1 and the minor axis length is L2, the aspect ratio is L1. The average value of / L2 is 1.0 to 3.0, the area ratio of the γ'phase in the microstructure is 15.0% or less, and the 0.2% proof stress at room temperature is 400 MPa or more. .

(2) In the alloy plate described in (1) above, the aspect ratio may be 1.0 to 2.0.

(3) In the alloy sheet according to (1) or (2), when the aging treatment is performed at 750 ° C. for 100 hours, the area ratio of the γ ′ phase in the microstructure after the aging treatment is 15.0%. As described above, the average equivalent circle diameter of the γ ′ phase may be 30 nm or more and 70 nm or less, and the area ratio of the precipitation phase other than the γ ′ phase may be 5.0% or less.

(4) A method for manufacturing an alloy plate according to another aspect of the present invention is a method for manufacturing the alloy plate according to any one of (1) to (3), in which a rolled alloy material is formed from 1050 to 1050. It includes a first step of solution heat treatment at 1200 ° C. for 30 seconds to 600 seconds and a second step of finish cold rolling at a sheet thickness reduction rate of 5% or more and 45% or less.

(5) A method for manufacturing an alloy plate according to another aspect of the present invention is a method for manufacturing the alloy plate according to any one of (1) to (3), in which a rolled alloy material is formed from 1050 to 1050. It includes a third step of solution heat treatment at 1200 ° C. for 30 seconds to 600 seconds, and a fourth step of aging treatment at 700 to 850 ° C. for 30 seconds to 600 seconds.

  (6) In the method for manufacturing an alloy plate according to (5) above, after the third step and before the fourth step, finish cold rolling is performed at a reduction rate of a sheet thickness reduction rate of 45% or less. 5 steps may be included.

  The alloy plate according to the above aspect of the present invention is used for an exhaust system gasket or the like, and ensures sufficient high temperature strength even when exposed to high temperatures for a long time.

  The present inventors conducted a detailed investigation on the composition of the alloy and the manufacturing process, and examined an alloy plate suitable for an exhaust system gasket that can withstand the severe environment of use required in recent years.

The gasket used at high temperature for a long time is formed using a heat-resistant alloy thin plate. When the heat-resistant alloy thin plate is exposed to the high temperature environment of the exhaust gas process, a fine γ '(gamma prime) phase is precipitated in the alloy, and the γ'phase ensures high temperature strength. The γ'phase is a precipitate represented by Ni ₃ (Al, Ti, Nb), and becomes Ni ₃ (Al, Nb) when Ti is not contained in the alloy. The γ'phase has a high interface consistency with the austenite phase, which is the matrix, and is more easily finely dispersed in the matrix grains than other precipitation strengthening phases. Since it has an inverse dependency with respect to, it has a feature of high precipitation strengthening ability especially at high temperature. Since the state of the precipitation phase changes depending on the composition, it is possible to promote proper precipitation of the γ'phase in a high temperature environment by optimizing the alloy composition.

  Even after the gasket is mounted on an automobile and exposed to high temperature for a long time (hereinafter also referred to as “high temperature and long time aging”), the γ ′ phase is aged for a long time at high temperature in order to sufficiently maintain the airtightness retaining function. It must be stable afterwards, so that high temperature strength is maintained. However, as a result of the study by the present inventors, when a conventionally known heat-resistant alloy plate is applied to a gasket, particularly when it is exposed to an extremely high temperature for a long time like a turbo gasket, the high temperature strength is lowered. Became clear.

  As a result of the present inventors' investigation of the reason why the high temperature strength of conventionally known heat-resistant alloys is reduced, it was found that the conventional heat-resistant alloys have a high temperature long-term aging treatment (high-temperature long-term aging heat treatment). It was found that the high temperature strength was lowered due to the precipitation of compounds other than the γ'phase (unreinforced phase) such as the Laves phase. It is considered that since a precipitation-free zone in which the γ'phase does not exist is formed around the Laves phase, as a result, the existence ratio of the γ'phase decreases and the high temperature strengthening ability deteriorates remarkably. That is, as a result of the study by the present inventors, it was found that suppressing the precipitates (non-strengthened phase) other than the γ ′ phase such as the Laves phase during aging at high temperature for a long time is important for maintaining the high temperature strength.

  As a result of a detailed investigation of precipitation behavior of the γ ′ phase which is a strengthening phase and non-reinforced phases such as Laves phase during high temperature long time aging from the alloy plate after high temperature long time aging, the γ ′ phase was the parent phase. It was found that the Laves phase was precipitated as grain nucleation sites and dislocations as nucleation sites, while dispersed over the entire surface and finely precipitated. Therefore, it was found that by reducing these nucleation sites as much as possible, the precipitation of the non-reinforced phase such as the Laves phase can be suppressed and the γ ′ phase which is the strengthened phase can be maintained.

  In order to maintain the high temperature strength even after aging at high temperature for a long time, 1) control the components so as to promote the precipitation of the γ'phase and suppress the precipitation of non-reinforced phases such as the Laves phase, and 2) heat treatment (solution To reduce the grain boundaries, thereby reducing precipitation sites of non-strengthened phases such as the Laves phase, and 3) reducing the reduction rate of the plate thickness in the final finish cold rolling. Therefore, it is important to reduce dislocations in the crystal grains of the alloy plate and reduce precipitation sites of non-reinforced phases such as Laves phase. The present inventors have also studied these process improvements and completed the present invention.

Hereinafter, an alloy plate according to an embodiment of the present invention (alloy plate according to the present embodiment) will be described.
First, the component composition (chemical composition) of the alloy plate according to this embodiment will be described. Unless otherwise specified,% relating to the content of components means% by mass.

C: 0.0200% or less C is an element that forms a carbide by combining with Ti and Nb. When Ti and Nb form carbides, the amount of the γ'phase, which is the strengthening phase, is reduced. Therefore, a lower C content is desirable. Therefore, the C content is 0.0200% or less.
On the other hand, if the C content is less than 0.0002%, the manufacturing cost will remarkably increase. Therefore, the C content may be 0.0002% or more.

Si: 0.02% or more and 2.00% or less Si is an element effective as a deoxidizing element during refining. Further, Si is an element that improves the protection property of the film that improves the oxidation resistance of the alloy. In order to obtain these effects, the Si content is set to 0.02% or more.
On the other hand, if the Si content is too high, the hot workability is significantly deteriorated, cracks in the ears occur, and the maintenance cost increases. Therefore, the Si content is set to 2.00% or less.

Mn: 0.02% or more and 2.00% or less If the Mn content is too large, the hot workability is deteriorated and the oxidation resistance at high temperature is significantly deteriorated. Therefore, the Mn content is 2.00% or less.
On the other hand, since Mn is mixed from raw material scraps, it is necessary to reduce the use of scrap in order to greatly reduce the Mn content, resulting in an increase in cost. Therefore, the Mn content is set to 0.02% or more.

P: 0.050% or less P is an impurity element contained in ferrochrome which is one of alloy raw materials. Since P is harmful to hot workability, the P content is limited to 0.050% or less.

S: 0.0100% or less S is an impurity element contained in raw material scrap or the like. Since S is harmful to hot workability, the S content is limited to 0.0100% or less.

Cr: 12.00% or more and 30.00% or less Cr is an essential element from the viewpoint of ensuring corrosion resistance as a heat resistant alloy. From the viewpoint of ensuring sufficient corrosion resistance, the Cr content is 12.00% or more. It is preferably 14.00% or more.
On the other hand, if the Cr content is too high, a coarse compound such as a σ phase is generated during annealing, not only the precipitation of the γ ′ phase is suppressed, but also the material is easily embrittled. Therefore, the Cr content is 30.00% or less. It is preferably 20.00% or less.

Ni: 35.00% or more and 60.00% or less Ni is a strong austenite stabilizing element and is an essential element for obtaining an austenite matrix in the microstructure. Further, Ni is an extremely important element for obtaining a γ ′ phase (Ni ₃ (Al, Nb, Ti)) which is a strengthening phase contributing to the improvement of high temperature strength. From the viewpoint of securing strength at high temperature as a heat resistant material, the Ni content is 35.00% or more. It is preferably at least 40.00%.
On the other hand, if the Ni content is too high, the cost increases, and the deformation resistance during hot working increases, making manufacturing difficult. Therefore, the Ni content is 60.00% or less. Preferably, it is 50.00% or less.

N: 0.0200% or less N is an element that combines with Al, Nb, and Ti to form a nitride. When Al, Nb, and Ti form a nitride, the amount of the γ ′ phase, which is the strengthening phase, is reduced. Therefore, it is desirable that the N content is low. Therefore, the N content is limited to 0.0200% or less.

V: 0.02% or more and 1.00% or less V is an element that contributes to the improvement of high temperature strength by solid solution strengthening. In order to obtain this effect, the V content is 0.02% or more.
On the other hand, V combines with C and N to form carbides and nitrides. If the V content is too large, coarse carbides and nitrides are generated, and the workability of the material deteriorates. Therefore, the V content is 1.00% or less.

Nb: more than 2.00% and 3.50% or less Nb is an element that constitutes the γ'phase which is a strengthening phase, and the γ'phase in which Nb is solid-soluted has a high temperature strengthening ability. Therefore, in order to secure high temperature strength, the Nb content is set to more than 2.00%.
On the other hand, Nb is an element that lowers the melting point of the alloy and makes hot working at high temperatures difficult. Therefore, the Nb content is 3.50% or less.

Al: over 2.60%, 4.00% or less Al is an element constituting the γ'phase, which is a strengthening phase contributing to the increase in high temperature strength. Although its strengthening ability does not reach that of Ti and Nb, it has an effect of maintaining the γ'phase stably for a long time as compared with Ti and Nb. From the viewpoint of stably maintaining the high temperature strength after aging for a long time, the Al content is more than 2.60%.
On the other hand, if the Al content is excessive, the melting point of the alloy is lowered, and hot working at high temperature becomes difficult. Therefore, the Al content is 4.00% or less.

  The alloy plate according to this embodiment may further contain the following arbitrary elements. However, since the following optional elements are not necessarily contained, the lower limit thereof is 0%.

Ti: 0 to 0.80%
Ti is an element that constitutes the γ ′ phase that is a strengthening phase, and is an element that contributes to the improvement of high temperature strength. Therefore, you may make it contain as needed. To obtain the above effect, the Ti content is preferably 0.20% or more.
On the other hand, when a large amount of Ti is contained, C, N and Ti in the alloy form coarse carbides and nitrides, and hot workability and cold workability are significantly deteriorated. Further, since the melting point of the alloy decreases due to the inclusion of Ti, if the Ti content is excessive, hot working at high temperature becomes difficult. Therefore, the Ti content is 0.80% or less even when it is contained. It is preferably 0.60% or less.

Mo: 0-2.00%
Mo is an element that forms a solid solution in the austenite phase, which is the parent phase, and contributes to the improvement of high temperature strength. Therefore, you may make it contain as needed.
On the other hand, when a large amount of Mo is contained, the deformation resistance at the time of hot working increases, and it becomes difficult to hot roll to a predetermined plate thickness. In addition, the precipitation of coarse Laves phase is promoted during high temperature aging for a long time, and the high temperature strength decreases. Therefore, the Mo content is 2.00% or less even when it is contained.

W: 0-2.00%
W is an element that forms a solid solution in the austenite matrix and contributes to the improvement of high temperature strength. Therefore, you may contain as needed. However, if W is contained in a large amount, the deformation resistance at the time of hot working increases, and it becomes difficult to perform hot rolling to a predetermined plate thickness. In addition, the precipitation of coarse Laves phase is promoted during high temperature aging for a long time, and the high temperature strength decreases. Therefore, the W content is 2.00% or less even when it is contained.

Cu: 0 to 1.00%
Cu is an element that has a solid solution in the austenite matrix and has the effect of increasing the high temperature strength. Therefore, you may contain as needed.
On the other hand, if the Cu content is excessive, edge cracking may occur during hot rolling. Therefore, the Cu content is 1.00% or less even when it is contained.

Co: 0-1.00%
Co is an element that forms a solid solution in the γ ′ phase as a substitute for Ni, and is an element that contributes to the improvement of high temperature strength. Therefore, it may be contained.
On the other hand, when a large amount of Co is contained, not only the cost increases but also the deformation resistance during hot working increases, which makes it difficult to hot-roll to a predetermined plate thickness. Therefore, the Co content is 1.00% or less even when it is contained.

B: 0 to 0.0100%
B is an element that segregates at the crystal grain boundaries, and by strengthening the crystal grain boundaries, it suppresses slippage at the grain boundaries and contributes to the improvement of high temperature strength. Therefore, it may be contained. In order to obtain the above-mentioned effects by containing B, the B content is preferably 0.0010% or more.
On the other hand, if B is contained in a large amount, the grain boundary segregation becomes remarkable and the hot workability is significantly deteriorated. Therefore, even if it is contained, the B content is 0.0100% or less.

Zr: 0 to 0.0100%
Zr is an element that segregates at the crystal grain boundaries, and by strengthening the crystal grain boundaries, it suppresses slippage at the grain boundaries and contributes to the improvement of high temperature strength. Therefore, it may be contained. In order to obtain the above-mentioned effects by containing Zr, the Zr content is preferably 0.0010% or more.
On the other hand, if a large amount of Zr is added, segregation of grain boundaries becomes remarkable, and hot workability is significantly reduced. Therefore, the Zr content is 0.0100% or less even when it is contained.

Ca: 0 to 0.0050%
Ca is an element having an effect of improving hot workability. If the hot workability is improved, the manufacturing cost can be reduced. In order to obtain this effect, it may be contained.
On the other hand, when the Ca content is large, troubles such as clogging of the molten metal nozzle occur during casting, and manufacturing becomes extremely difficult. Therefore, the Ca content is 0.0050% or less even when it is contained.

Mg: 0 to 0.0050%
Mg is an element having an effect of improving hot workability. If the hot workability is improved, the manufacturing cost can be reduced. In order to obtain this effect, it may be contained.
On the other hand, if the Mg content is large, troubles such as clogging of the molten metal nozzle occur during casting, and manufacturing becomes extremely difficult. Therefore, the Mg content is 0.0050% or less even when it is contained.

The total content of the components described above is 90% or less, preferably 80% or less. The balance other than the components described above is composed of 10.00% or more Fe and impurities.
If Fe contained in the balance is less than 10.00% of the chemical composition of the alloy plate, workability deteriorates. Therefore, in the alloy plate according to this embodiment, Fe is 10.00% or more.
Further, the total amount of impurities is preferably 2.0% or less. The impurities are, for example, REM, Sn, and Zn in addition to P and S described above.

  Next, the point that the grain size of the alloy plate is coarsened to reduce the grain boundaries to reduce the precipitation sites of the non-strengthened phase such as the Laves phase will be described.

<Microstructure>
The metal structure of the alloy plate according to this embodiment includes an austenite phase as a matrix phase. A small amount of γ'phase or the like may be precipitated in the austenite phase of the mother phase. The mother phase excluding the precipitate may substantially consist of an austenite phase.

[Average grain size of austenite phase is 20 μm or more]
The grain boundaries serve as precipitation sites for non-strengthened phases such as the Laves phase during the aging treatment (aging heat treatment). Therefore, when the average crystal grain size is small, the existence ratio of grain boundaries per unit cross section of the alloy plate increases, and the amount of precipitation of these non-reinforced phases increases during the aging treatment. In the vicinity of the non-reinforced phase, a non-precipitation phase in which the γ'phase does not precipitate is formed, and as a result, the amount of the γ'phase, which is the strengthening phase, decreases.
When the average crystal grain size of the austenite phase of the alloy plate is 20 μm or more, the precipitation site of the non-strengthened phase can be reduced and the precipitation of the γ ′ phase after aging at high temperature for a long time can be secured. Therefore, the average crystal grain size is set to 20 μm or more. Since the average crystal grain size does not change significantly by heat treatment at a working temperature of about 750 ° C., the value is the value for the alloy sheet after finish cold rolling. The manufacturing method (solution treatment before finish cold rolling) for setting the average crystal grain size of the austenite phase, which is the parent phase of the alloy sheet, to 20 μm or more will be described later.
Although it is not necessary to specify the upper limit of the average crystal grain size, if the crystal grain size is too large relative to the plate thickness, press moldability and the like will be significantly deteriorated, so the average crystal grain size is preferably 50 μm or less.

The average crystal grain size can be determined by the following method.
A sample is taken from the metal plate so that the L cross section (cross section parallel to the rolling direction and parallel to the plate thickness surface) becomes the observation surface, and the observation surface is surface-polished and etched. This sample is observed and photographed with an optical microscope, the number of crystal grains per unit area is counted according to JIS G0551: 2013, the average crystal grain area a per one is determined, and the square root of a is taken as the average crystal grain size.

<Area ratio of γ'phase is less than 15.0%>
In the alloy sheet according to this embodiment, the area ratio of the γ'phase is less than 15.0% in the microstructure before the high temperature long-term aging treatment (before being used as a gasket or the like).
If a large amount of γ'phase has already been precipitated before the high temperature long time aging treatment, these γ'phases become coarse due to the high temperature long time aging treatment and sufficient high temperature strength cannot be maintained. Therefore, in the microstructure of the alloy plate, the area ratio of the γ'phase is set to less than 15.0%. It is preferably 8.0% or less, more preferably 6.0% or less, still more preferably 2.0% or less.
If the area ratio of the γ'phase before the high temperature long-time aging treatment is reduced, the 0.2% proof stress at room temperature may decrease, but as described later, the reduction rate of the plate thickness is 5% to 45%. By performing cold rolling, 0.2% proof stress at room temperature can be secured. Further, cold rolling also improves the flatness of the alloy sheet.

The area ratio of the γ'phase of the alloy plate according to the present embodiment, which is the alloy plate before the high temperature long time aging treatment, is obtained by the following method.
The cross section of the alloy plate parallel to the rolling direction is polished, the polished surface is etched with aqua regia, and observed and analyzed by EDS attached to the same equipment as FE-SEM. The fine spherical precipitates that are uniformly precipitated in the crystal grains are defined as the γ'phase, and are precipitated in the grain boundaries and crystal grains, and are coarser than the chemical composition of the alloy plate (in the chemical composition of the alloy plate. A precipitate containing a large amount of Fe and Cr (compared to the Fe and Cr contents) is classified as a precipitate other than the γ'phase. The area ratio of the γ'phase is calculated by the line segment method from the photograph taken by the FE-SEM. The fine spherical precipitate generally means a precipitate having a circle equivalent diameter of 100 nm or less and an aspect ratio of 2 or less.

[Average value of aspect ratio of crystal grains of austenite phase is 1.0 to 3.0]
In the alloy plate according to the present embodiment, when the major axis length and the minor axis length of the austenite phase crystal grains are L1 and L2, respectively, the average value of L1 / L2, which is the aspect ratio, is 1.0 to It is 3.0.
The grain boundaries serve as precipitation sites for non-reinforced phases that do not contribute to the improvement of high temperature strength such as the Laves phase during the aging treatment. The aspect ratio increases as the crystal grains expand due to rolling. When a large amount of dislocations is introduced by rolling, the dislocations are used as nucleation sites to promote the precipitation of the non-strengthened phase and inhibit the precipitation of the γ ′ phase which is the strengthened phase. Also, considering the volume per crystal grain and the area of the crystal grain boundary, as the aspect ratio increases, the volume of the crystal grain is the same, but only the area of the crystal grain boundary increases. The percentage of precipitation sites in the strengthening phase increases. Therefore, in order to suppress the precipitation of the non-strengthened phase and secure the γ ′ phase which is the strengthened phase, the aspect ratio is preferably small.
Therefore, in the alloy plate according to the present embodiment, the average value of the aspect ratio is set to 1.0 to 3.0 as one of the indexes indicating that the structure has less precipitation sites of the non-reinforced phase. The aspect ratio is preferably 1.0 to 2.0.

The aspect ratio can be obtained by the following method.
A sample is taken so that the L cross section (cross section parallel to the rolling direction and parallel to the plate thickness surface) becomes the observation surface, and surface polishing and etching are performed. The metal structure of this sample is observed with an optical microscope so that the observation area becomes 400000 μm ² or more (5 or more regions having 80000 μm ² or more). Using the metallographic structure measured in the observation region, the aspect ratio (L1 / L2) was calculated from the length L1 of the major axis and the length L2 of the minor axis of each observed crystal grain, and these were arithmetically averaged. To calculate. The twin interface is excluded during the measurement. This measurement may use image analysis software.
Regarding the crystal grains L1 and L2, the crystal grain is regarded as an ellipse, and the major axis and the minor axis of the ellipse are defined as L1 and L2, respectively. In the case of a rolled material, since the grain size in the rolling direction is the longest and the grain size in the plate thickness direction is the shortest, even if the length of the crystal grain in the rolling direction is L1 and the length of the crystal grain in the plate thickness direction is L2, Good.

<0.2% proof stress at room temperature is 400 MPa or more>
In the alloy plate according to this embodiment, the 0.2% proof stress at room temperature is 400 MPa or more. When the 0.2% proof stress at room temperature is 400 MPa or more, the strength before aging at high temperature for a long time (at the start of use as a gasket) can be secured when the alloy plate is processed and used as a gasket.
On the other hand, if the 0.2% proof stress at room temperature is more than 1200 MPa, the workability during processing into a gasket is deteriorated. Therefore, the 0.2% proof stress at room temperature is preferably 1200 MPa or less.

  The 0.2% proof stress at room temperature was determined by using the No. 13B tensile test specified in JIS Z2241: 2011, which was taken so that the longitudinal direction of the test piece was parallel to the rolling direction. It can be obtained by performing a tensile test in which a tensile load is applied so that it becomes min.

As described above, in the alloy plate according to the present embodiment, while controlling the components to promote the precipitation of the γ ′ phase, the solution treatment and rolling are controlled to reduce the grain boundaries and dislocations, and the precipitation of the non-reinforced phase. We are trying to reduce the number of sites.
The alloy plate according to the present embodiment secures high-temperature strength by the γ'phase, but the γ'phase is hardly precipitated at the stage of the alloy plate after cold rolling, and the alloy plate is subjected to an aging treatment. Is deposited by. When the high temperature aging treatment is performed for a long time, the conventional alloy plate cannot secure a sufficient amount of the γ ′ phase due to the precipitation of the non-reinforced phase, whereas the alloy plate according to the present embodiment has a high temperature long time aging. Since the amount of the γ'phase is secured afterwards, sufficient high temperature strength (for example, 380 MPa or more) can be secured even after aging at high temperature for a long time.

After aging at <750 ° C. for 100 hours, the area ratio of the γ ′ phase in the microstructure is 15.0% or more, the average equivalent circle diameter of the γ ′ phase is 30 nm or more, 70 nm or less, and the precipitation phase other than the γ ′ phase Area ratio is 5.0% or less>
When simulating the use as a gasket, it is preferable that the high temperature strength is maintained even after the heat treatment condition of 750 ° C. for 100 hours as the high temperature long time aging treatment. For cooling after the heat treatment, it is desirable that the average cooling rate up to 500 ° C. or less is 10 ° C./second or more. This heat treatment condition is more severe than the heat history of heat-resistant parts of automobiles, and if high temperature strength is secured under this condition, sufficient high temperature strength can be secured even when used as a gasket for a long time. The characteristics of the alloy sheet according to the present embodiment after aging can be evaluated by observing the microstructure after the high-temperature long-time aging treatment that simulates the use.
In order to maintain the high temperature strength after heat treatment at 750 ° C. for 100 hours, the area ratio of γ ′ phase is 15.0% or more in the microstructure after aging treatment at 750 ° C. for 100 hours. It is preferable that the average equivalent circle diameter of the γ ′ phase is 30 nm or more and 70 nm or less, and the area ratio of the precipitation phase other than the γ ′ phase is 5.0% or less.

From the viewpoint of high temperature strength, the amount of the γ'phase after the high temperature long-term aging treatment is 15.0% or more.
Further, the high temperature strength is affected not only by the amount (area ratio) of the γ'phase, but also by its size. When the γ'phase is too small, the dislocation easily cuts the γ'phase, so that the strengthening ability of the γ'phase becomes low. When it takes time to precipitate the γ ′ phase, such as when the content of Ni, Nb, Al, and Ti is small, the γ ′ phase becomes small even after the high temperature long-term aging treatment. On the other hand, the fact that the γ'phase is too large has the same meaning as the number of γ'phases is small, so the high temperature strength of the alloy sheet becomes low. Therefore, the γ ′ phase after the high temperature long-time aging treatment has an average equivalent circle diameter of 30 nm or more and 70 nm or less.
In addition, non-strengthened phases other than the γ'phase may be precipitated on the grain boundaries and dislocations. These non-strengthened phases represented by the Laves phase may not contribute to the high temperature strength, but may form a precipitation-free zone around which the γ'phase is lacking and deteriorate the high temperature strength. Therefore, the area ratio of the precipitate other than the γ'phase is set to 5.0% or less.

The area ratio of the γ ′ phase after aging at 750 ° C. for 100 hours, the average equivalent circle diameter of the γ ′ phase, and the area ratio of the precipitated phases other than the γ ′ phase are determined by the following methods.
The cross section parallel to the rolling direction of the alloy plate after aging at 750 ° C. for 100 hours is polished, the polished surface is etched with aqua regia, and observed and analyzed by EDS attached to the same equipment as FE-SEM. The fine spherical precipitates that are uniformly precipitated in the crystal grains are the γ'phase, while they are precipitated in the crystal grain boundaries and in the crystal grains and are coarser, and Fe and Cr are more than the chemical composition of the alloy plate. The precipitate contained in a large amount is classified as a precipitate other than the γ'phase. From the photograph taken by FE-SEM, the area ratio of the γ'phase and the area ratio of the precipitates other than the γ'phase are calculated by the line segment method. A fine spherical precipitate generally means a precipitate having an equivalent circle diameter of 100 nm or less and an aspect ratio of 2.0 or less.
For one sample, the area of 50 or more γ'phases is calculated using image analysis software, the diameter of a circle having the same area as that area is calculated, and the average is calculated as the equivalent circle diameter ( Estimate as circle equivalent diameter).

  A gasket can be formed by processing the alloy plate according to the present embodiment. For example, the outer shape of the gasket is formed from the alloy plate according to the present embodiment by means such as punching or cutting, and then a step portion called a bead is formed to obtain a gasket. The gasket thus obtained can maintain high high-temperature strength even after high-temperature long-term aging treatment at 750 ° C. for 100 hours.

<Manufacturing method>
Next, a method for manufacturing the alloy plate according to this embodiment will be described.
The alloy plate according to this embodiment can be manufactured by the following method.
(I) A first step of solution heat treating the rolled alloy material having the above-described chemical composition at 1050 to 1200 ° C. for 30 seconds to 600 seconds, and a plate thickness reduction rate of 5% or more, 45% A second step of finish cold rolling below.
Or
(Ii) The third step of solution heat treating the rolled alloy material having the above-mentioned chemical composition at 1050 to 1200 ° C for 30 seconds to 600 seconds, and 700 to 850 ° C for 30 seconds to 600 seconds. A method including a fourth step of aging treatment according to 1., and a fifth step of finishing cold-rolling at a sheet thickness reduction rate of 45% or less after the third step and before the fourth step.

Hereinafter, each step will be described.
<About manufacturing method (i)>
[First step of solution heat treating the rolled alloy material at 1050 to 1200 ° C. for 30 to 600 seconds]
The material for obtaining the alloy plate according to the present embodiment contains many alloy elements, and various compounds are precipitated during the manufacturing process. Therefore, in order to obtain a microstructure such as the alloy plate according to this embodiment, solution heat treatment for once dissolving the precipitate in the matrix is necessary. In addition, the solution treatment must be made under the condition of increasing the average crystal grain size in addition to the role of dissolving the precipitate. Therefore, the solution treatment is performed at a temperature of 1050 ° C. or higher for 30 seconds or longer. Thereby, the average crystal grain size of the heat-resistant alloy thin plate can be set to 20 μm or more.
On the other hand, if the solution heat treatment temperature is too high or the treatment time is too long, the cost will increase. Therefore, the solution heat treatment temperature is 1200 ° C. or lower and the time is 600 seconds or shorter. After the solution treatment, it is preferable to cool to 500 ° C. or less at an average cooling rate of 10 ° C./sec or more.

[Second step of finish cold rolling at a sheet thickness reduction rate of 5% or more and 45% or less]
The alloy plate according to the present embodiment is subjected to finish cold rolling in order to adjust the plate thickness and the strength. The non-strengthened phase in the crystal grains precipitates dislocations as nucleation sites, so non-strengthened phases other than the γ ′ phase are precipitated in the crystal grains by the aging treatment after finish cold rolling, which may impair the original properties. is there. Therefore, by reducing the sheet thickness reduction rate (rolling reduction rate) in finish cold rolling, the amount of dislocations in the crystal grains after finish cold rolling is reduced, and the precipitation of non-strengthened phases during high temperature long time aging is suppressed. . As described above, the reduction of the plate thickness reduction rate also has the effect of reducing the aspect ratio of the crystal grains and the area of the crystal grain boundaries serving as precipitation sites for the non-strengthened phase.
The aspect ratio is 1.0 to 3.0, the area ratio of the γ ′ phase is 15.0% or less, and the precipitation phase other than the γ ′ phase is suppressed to 5.0% or less after aging at 750 ° C. for 100 hours. The sheet thickness reduction rate in finish cold rolling is set to 45% or less from the viewpoint of forming a structure capable of ensuring an area ratio of the 1'phase of 15.0% or more. It is preferably 40% or less. The intermediate annealing is not performed by dividing the cold rolling in the finish cold rolling step. Further, when the finish cold rolling rate is high, the dislocation density is increased and the diffusion of atoms is promoted, so that the grain growth of the γ ′ phase is promoted. Since the coarsening of the γ'phase leads to a decrease in high-temperature strength, the finish cold rolling rate is also 45% or less, and preferably 40% or less, also from this viewpoint.

On the other hand, in the alloy plate according to the present embodiment, the γ'phase is precipitated by being kept in a high temperature environment, whereby the room temperature strength and the high temperature strength are maintained high. However, when the γ ′ phase is not deposited on the metal plate immediately after the start of use in the case of being used as a gasket, it is necessary to secure the necessary strength by means other than depositing the γ ′ phase.
By utilizing the work hardening of finish cold rolling, it is possible to secure the strength (0.2% yield strength) at room temperature before high temperature aging. When the finish cold rolling ratio is low, it is advantageous for high temperature strength after high temperature aging, but the strength at room temperature before high temperature aging may be insufficient. Therefore, when the aging treatment is not performed, the sheet thickness reduction rate of finish rolling is set to 5% or more. The plate thickness reduction rate of finish rolling is preferably 10% or more.

<About manufacturing method (ii)>
[Third step of solution heat treating the rolled alloy material at 1050 to 1200 ° C. for 30 to 600 seconds]
The material for obtaining the alloy plate according to the present embodiment contains many alloy elements, and various compounds are precipitated during the manufacturing process. Therefore, in order to obtain an appropriate microstructure after aging at 750 ° C. for 100 hours as described above, solution heat treatment for once dissolving the precipitate in the mother phase is necessary. This solution treatment needs to have a condition that the average crystal grain size is increased in addition to the role of dissolving the precipitate. Therefore, the solution treatment is performed at a temperature of 1050 ° C. or higher for 30 seconds or longer. Thereby, the average crystal grain size of the heat-resistant alloy thin plate can be set to 20 μm or more. On the other hand, if the solution heat treatment temperature is too high or the treatment time is too long, the cost will increase. Therefore, the solution heat treatment temperature is 1200 ° C. or lower and the time is 600 seconds or shorter. After the solution treatment, it is preferable to cool to 500 ° C. or less at an average cooling rate of 10 ° C./sec or more.

[Fourth step of aging treatment at 700 to 850 ° C. for 30 to 600 seconds]
By utilizing the work hardening of finish cold rolling, the strength (0.2% yield strength) at room temperature before high temperature long time aging can be secured. On the other hand, by performing finish cold rolling, dislocations are introduced and the precipitation sites of the non-strengthened phase increase. Therefore, in order to further suppress the introduction of dislocations, it is preferable to use a manufacturing method that does not require finish cold rolling.
If finish cold rolling is not performed, 0.2% proof stress at room temperature may not be secured, but even if finish cold rolling is not performed, a short aging treatment is performed to precipitate a small amount of γ'phase. By doing so, 0.2% proof stress at room temperature can be secured.
Therefore, in the method for manufacturing an alloy sheet according to the present embodiment, which does not require finish cold rolling, an aging treatment of holding at 700 to 850 ° C for 30 seconds to 600 seconds is performed after the third step.
If the aging temperature is less than 700 ° C. or the time is less than 30 seconds, sufficient strength cannot be secured.
On the other hand, if the aging temperature is higher than 850 ° C, the γ'phase may not be precipitated in some cases. If the aging time is longer than 600 seconds, the γ'phase may be coarsened and sufficient strength may not be obtained.

[Fifth step of finish cold rolling at a sheet thickness reduction rate of 45% or less]
By performing the fourth step, the 0.2% proof stress at room temperature can be secured at 400 MPa or more, so it is not necessary to perform finish cold rolling, but if it is desired to further increase the 0.2% proof stress at room temperature, Between the step 3 and the fourth step, finish cold rolling may be performed at a sheet thickness reduction rate of 45% or less. When the finish cold rolling rate is high, the dislocation density increases and the diffusion of atoms is promoted, so that the grain growth of the γ ′ phase is promoted. Since coarsening of the γ'phase leads to a decrease in high temperature strength, the finish cold rolling rate is 45% or less, and preferably 40% or less.

  For the manufacturing steps up to the first step or the third step, a usual alloy plate manufacturing method can be used. That is, an alloy adjusted to a predetermined composition in the smelting process is cast into a slab by a casting process including continuous casting, and the slab is heated in a pre-hot rolling heating furnace and then hot-rolled. . Hot forging may be performed before hot rolling. Further, the alloy sheet after hot rolling may or may not be annealed after hot rolling. Cold rolling (intermediate cold rolling) is performed on the hot-rolled sheet thus manufactured. Intermediate cold rolling may be divided into a plurality of times, and intermediate annealing may be performed between each intermediate cold rolling. The solution heat treatment of the first step or the third step is performed on the intermediate cold-rolled sheet thus manufactured.

Hereinafter, the present invention will be specifically described with reference to Examples.
<Example 1>

First, a 25 kg ingot having the chemical composition shown in Table 1 was melted. Here, A to P are components that satisfy the requirements of the present invention, and Q to W are components that do not satisfy the present invention. Numerical values outside the scope of the present invention are underlined.
These ingots were hot-forged to have a thickness of 40 mm and then hot-rolled to obtain a hot-rolled sheet having a sheet thickness of 5 mm. The hot rolled sheet was annealed and cold rolled to obtain intermediate cold rolled sheets having various sheet thicknesses. This intermediate cold-rolled sheet was subjected to solution heat treatment and finish cold rolling (cold rolling) under the conditions shown in Table 2. By changing the plate thickness of the intermediate cold-rolled sheet, the plate thickness of the finish cold-rolled sheet was set to 0.20 mm.

A sample was taken from the metal plate so that the L cross section of this finish cold-rolled sheet would be the observation surface, the observation surface was surface-polished and etched with aqua regia, observed with an optical microscope, photographed, and the austenite phase was observed by the quadrature method. The average crystal grain size was calculated.
Further, the metal structure of the same sample was observed with an optical microscope so that the observation area was 400000 μm ² or more. The aspect ratio (L1 / L2) was determined from the length L1 of the major axis and the length L2 of the minor axis of each observed crystal grain, and these were averaged to calculate the aspect ratio.

Further, 0.2% proof stress at room temperature was obtained by the following method in the alloy plate before high temperature long time aging treatment.
The 0.2% proof stress at room temperature was determined by using the No. 13B tensile test specified in JIS Z2241: 2011, which was taken so that the longitudinal direction of the test piece was parallel to the rolling direction, and the amount of change in gauge length was 3 mm / It was obtained by conducting a tensile test under the condition that a tensile load is applied so as to be min.

In addition, the area ratio of the γ'phase was determined by the following method in the alloy plate before the high temperature long time aging treatment.
A cross section parallel to the rolling direction of the alloy sheet before high temperature long time aging treatment was polished, etched with aqua regia, and observed and analyzed by EDS attached to the same equipment as FE-SEM. A γ'phase is a precipitate having a circle-equivalent diameter of 100 nm or less and an aspect ratio of 2 or less, which is uniformly precipitated in the crystal grains. In addition, the precipitates that were precipitated in the crystal grain boundaries and in the crystal grains, were coarse, and contained Fe and Cr in a larger amount than the chemical composition of the alloy plate were defined as precipitates other than the γ'phase. The area ratio of the γ'phase was calculated by the line segment method from the photograph taken by the FE-SEM.

Furthermore, the finish cold-rolled sheet was subjected to an aging treatment (high temperature long-term aging treatment) at 750 ° C. for 100 hours in an atmospheric furnace. The area ratio of the γ'phase and the area ratio of the precipitates other than the γ'phase in the alloy plate before the high temperature long-time aging treatment were determined by the following methods.
A circle perpendicular to the rolling direction of the metal plate after aging treatment is polished, etched with aqua regia, and observed and analyzed by EDS attached to the same equipment as FE-SEM. Precipitates having an equivalent diameter of 100 nm or less and an aspect ratio of 2 or less are defined as γ'phases and are precipitated in the grain boundaries and in the crystal grains, and are coarse and have a large amount of Fe and Cr (Fe, Cr in the chemical composition of the alloy plate). The precipitate contained (more than the content) was classified as a precipitate other than the γ'phase. The area ratio of the γ'phase and the area ratio of the precipitate other than the γ'phase were calculated by the line segment method from the photograph taken by the FE-SEM.
For one sample, the area of 50 or more γ'phases was calculated using image analysis software, the diameter of a circle having the same area as that area was calculated, and the average was taken as the equivalent circle diameter. I estimated.

  Further, the alloy plate that had been subjected to an aging treatment at 750 ° C. for 100 hours was subjected to a high temperature tensile test at 750 ° C., and the high temperature 0.2% proof stress at 750 ° C. was measured. When the high temperature 0.2% proof stress at 750 ° C. after aging treatment at 750 ° C. for 100 hours was 380 MPa or more, it was judged that sufficient high temperature strength could be secured even after the high temperature long time aging treatment.

  Alloy plate No. of Table 2 Invention examples Nos. 1 to 21 satisfying the requirements of the present invention, No. 22 to 32 are comparative examples that do not satisfy the requirements of the present invention.

  In each of Inventive Examples 1 to 21, good quality with high temperature strength (0.2% proof stress) after aging at high temperature for a long time of 380 MPa or more could be secured. Further, since the finish cold-rolling with a plate thickness reduction rate of 5% or more was performed as it was as a heat-resistant alloy plate, the room temperature strength (0.2% proof stress) before aging treatment was 400 MPa or more in all cases.

On the other hand, in Comparative Example 22, although the composition of the alloy satisfied the requirements of the present invention, the solution treatment temperature was low and the time was short. As a result, the average grain size of the austenite phase was small, a large amount of the non-reinforced phase was precipitated after aging, and the amount of the γ ′ phase that was the strengthened phase was small. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
In Comparative Examples 23 and 24, the alloy compositions satisfied the regulations of the present invention, but the finish cold rolling rate was large. As a result, the aspect ratio exceeded 3.0. As a result, a large amount of non-reinforcing phase was precipitated after aging, and the γ'phase, which is a strengthening phase, was few after aging at high temperature for a long time. Further, in Comparative Example 24, the finish cold rolling rate was high and the dislocation density was high, so that the diffusion of atoms was fast and the γ ′ phase was coarsened after aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.

Comparative Example 25 had a large C content in the alloy. Therefore, carbides of Nb and Ti were generated from the manufacturing stage, and the amount of the γ ′ phase, which is a strengthening phase, was small after aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
Comparative Example 26 had a high Cr content in the alloy. Therefore, a large amount of non-reinforcing phase such as σ phase was precipitated after aging, and the amount of γ ′ phase that was a strengthening phase was small after aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
In Comparative Example 27, the Ni content in the alloy was small. Therefore, the precipitation of the γ ′ phase, which is the strengthening phase, was extremely small after aging. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
Comparative Example 28 had a large N content in the alloy. Therefore, nitrides of Al and Nb were generated from the manufacturing stage, and the amount of the γ'phase, which is a strengthening phase, was small after aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
Comparative Example 29 had a large V content in the alloy. As a result, coarse V carbides and V nitrides were produced from the manufacturing stage, and the amount of the γ'phase, which is the strengthening phase, was small after the aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
Comparative Example 30 had a small Nb content in the alloy. Therefore, the γ'phase was extremely small after aging at high temperature for a long time. The size of the γ'phase was also small. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
Comparative Example 31 had a small Al content in the alloy. Therefore, the γ'phase was small after aging at high temperature for a long time. As a result, the 0.2% proof stress at high temperature after aging at high temperature for a long time was low.
In Comparative Example 32, the composition of the alloy satisfies the requirements of the present invention, but the solution temperature is too low, the average crystal grain size of the austenite phase becomes small, and γ'precipitated during the production does not re-dissolve and is aged. It remained on the previous alloy plate. As a result, γ ′ was coarsened during high temperature long time aging, a non-reinforced phase was precipitated, and the high temperature 0.2% proof stress was low.

  As described above, in Comparative Examples 22 to 32, the structure or chemical composition of the alloy plate was outside the scope of the present invention. Further, as a result, the amount of the γ'phase after the high temperature long time aging treatment was small, and the high temperature strength after the high temperature long time aging treatment did not reach the target of the present invention.

<Example 2>
Alloy No. of Table 1 The ingot having the component A was hot-forged to a thickness of 40 mm and then hot-rolled to obtain a hot-rolled sheet having a sheet thickness of 5 mm. The hot rolled sheet was annealed and cold rolled to obtain intermediate cold rolled sheets having various sheet thicknesses. This intermediate cold-rolled sheet was subjected to solution treatment under the conditions shown in Table 3, and some alloy sheets were subjected to finish cold rolling (cold rolling). The thickness of the finished cold-rolled sheet was 0.20 mm.
Further, some of the alloy sheets were subjected to an aging treatment (short-time aging treatment) under the conditions shown in Table 3.

  With respect to the obtained alloy plate, the average crystal grain size of the austenite phase, the 0.2% proof stress at room temperature, the area ratio of the γ'phase, and the aspect ratio (L1 / L2) were determined in the same manner as in Example 1.

  Furthermore, the finish cold-rolled sheet or the metal sheet after the short-time aging treatment was subjected to an aging treatment (high-temperature long-time aging treatment) at 750 ° C. for 100 hours in an atmospheric furnace.

  The area ratio of the γ ′ phase and the area ratio of precipitates other than the γ ′ phase were calculated in the same manner as in Example 1 for the metal plate after the high temperature and long time aging treatment. For one sample, the area of 50 or more γ'phases was calculated using image analysis software, the diameter of a circle having the same area as that area was calculated, and the average was taken as the equivalent circle diameter. I estimated. Further, it was subjected to a high temperature tensile test at 750 ° C. and the high temperature 0.2% proof stress at 750 ° C. was measured.

The results are shown in Table 3.
Inventive alloy plate No. Regarding X1 to X4, the third step (solution heat treatment) of the present invention and the subsequent fourth step (short-time aging treatment) are performed, and both have room temperature 0.2% proof stress of 400 MPa or more. It was good. The high temperature 0.2% proof stress was as high as 380 MPa or more.
Alloy plate No. The alloy plate No. X1 was not subjected to the fifth step (finish cold rolling). Although X2 had a small reduction rate of plate thickness in the fifth step (finish cold rolling) of 4%, room temperature strength was secured by precipitation hardening in the fourth step.

On the other hand, alloy plate No. For X5, neither the fifth step (finish cold rolling) nor the fourth step (short-time aging treatment) was performed. In addition, the alloy plate No. For X6, the sheet thickness reduction rate in the fifth step (finish cold rolling) was 3%, and the fourth step (short-time aging treatment) was not performed. Therefore, the room temperature 0.2% proof stress was less than 400 MPa.
Alloy plate No. For X7, the reduction rate of the plate thickness in the fifth step (finish cold rolling) was 3%, and the temperature in the fourth step (short-time aging treatment) was low. In addition, the alloy plate No. X8 had a plate thickness reduction rate of 3% in the fifth step (finish cold rolling), and had a high temperature in the fourth step (short-time aging treatment). Therefore, the room temperature 0.2% proof stress was less than 400 MPa.

  INDUSTRIAL APPLICABILITY The alloy sheet of the present invention is used for an exhaust system gasket and the like, and ensures sufficient high temperature strength even when exposed to high temperatures for a long time. Therefore, in the gasket using the alloy plate of the present invention, high temperature strength is ensured even after use at high temperature for a long time. Therefore, it has high industrial applicability.

Claims (6)

In mass%,
C: 0.0200% or less,
Si: 0.02-2.00%,
Mn: 0.02-2.00%,
P: 0.050% or less,
S: 0.0100% or less,
Cr: 12.00-30.00%,
Ni: 35.00-60.00%,
N: 0.0200% or less,
V: 0.02-1.00%,
Nb: over 2.00%, 3.50% or less,
Al: over 2.60%, 4.00% or less,
Ti: 0 to 0.80%,
Mo: 0-2.00%,
W: 0 to 2.00%,
Cu: 0 to 1.00%,
Co: 0 to 1.00%,
B: 0 to 0.0100%,
Zr: 0 to 0.0100%,
Ca: 0 to 0.0050%,
Mg: 0 to 0.0050%,
And the balance is 10.00% or more of Fe, and a chemical composition that is an impurity,
The microstructure includes an austenite phase, the austenite phase has an average crystal grain size of 20 μm or more,
When the length of the major axis of the crystal grains of the austenite phase is L1 and the length of the minor axis is L2, the average value of the aspect ratio L1 / L2 is 1.0 to 3.0,
In the microstructure, the area ratio of the γ'phase is 15.0% or less,
An alloy plate having a 0.2% proof stress at room temperature of 400 MPa or more. The alloy plate according to claim 1, wherein the aspect ratio is 1.0 to 2.0. When aged at 750 ° C for 100 hours,
The area ratio of the γ'phase in the microstructure after the aging treatment is 15.0% or more,
The average circle equivalent diameter of the γ'phase is 30 nm or more and 70 nm or less, and the area ratio of the precipitation phase other than the γ'phase is 5.0% or less.
The alloy plate according to claim 1 or 2, characterized in that. A method for producing the alloy plate according to any one of claims 1 to 3,
A first step of solution heat treating the rolled alloy material at 1050 to 1200 ° C. for 30 seconds to 600 seconds;
A second step of finish cold rolling at a sheet thickness reduction rate of 5% or more and 45% or less;
A method for manufacturing an alloy plate, comprising: A method for producing the alloy plate according to any one of claims 1 to 3,
A third step of solution heat treating the rolled alloy material at 1050 to 1200 ° C. for 30 seconds to 600 seconds;
A fourth step of aging treatment at 700 to 850 ° C. for 30 to 600 seconds;
A method for manufacturing an alloy plate, comprising: The alloy according to claim 5, further comprising a fifth step of finish cold rolling at a reduction rate of a sheet thickness reduction rate of 45% or less after the third step and before the fourth step. Method of manufacturing a plate.