JP6609727B1 - Alloy plate and manufacturing method thereof - Google Patents
Alloy plate and manufacturing method thereof Download PDFInfo
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- 239000000956 alloy Substances 0.000 title claims abstract description 137
- 229910045601 alloy Inorganic materials 0.000 title claims abstract description 131
- 238000004519 manufacturing process Methods 0.000 title claims description 35
- 239000013078 crystal Substances 0.000 claims abstract description 63
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 29
- 239000000203 mixture Substances 0.000 claims abstract description 19
- 239000000126 substance Substances 0.000 claims abstract description 14
- 239000012535 impurity Substances 0.000 claims abstract description 9
- 230000032683 aging Effects 0.000 claims description 105
- 238000005097 cold rolling Methods 0.000 claims description 52
- 229910052758 niobium Inorganic materials 0.000 abstract description 21
- 229910052719 titanium Inorganic materials 0.000 abstract description 20
- 229910052782 aluminium Inorganic materials 0.000 abstract description 16
- 229910052804 chromium Inorganic materials 0.000 abstract description 11
- 229910052759 nickel Inorganic materials 0.000 abstract description 6
- 229910052748 manganese Inorganic materials 0.000 abstract description 5
- 229910052757 nitrogen Inorganic materials 0.000 abstract description 4
- 229910052802 copper Inorganic materials 0.000 abstract description 3
- 229910052750 molybdenum Inorganic materials 0.000 abstract description 3
- 229910052698 phosphorus Inorganic materials 0.000 abstract description 3
- 229910052720 vanadium Inorganic materials 0.000 abstract description 3
- 229910052726 zirconium Inorganic materials 0.000 abstract description 3
- 229910052721 tungsten Inorganic materials 0.000 abstract description 2
- 239000002244 precipitate Substances 0.000 description 38
- 230000035882 stress Effects 0.000 description 34
- 238000001556 precipitation Methods 0.000 description 32
- 238000005728 strengthening Methods 0.000 description 29
- 239000000243 solution Substances 0.000 description 27
- 230000007774 longterm Effects 0.000 description 22
- 238000000034 method Methods 0.000 description 20
- 238000010438 heat treatment Methods 0.000 description 18
- 238000005096 rolling process Methods 0.000 description 18
- 229910001068 laves phase Inorganic materials 0.000 description 15
- 230000000052 comparative effect Effects 0.000 description 13
- 230000000694 effects Effects 0.000 description 12
- 239000002184 metal Substances 0.000 description 12
- 229910052751 metal Inorganic materials 0.000 description 12
- 238000000445 field-emission scanning electron microscopy Methods 0.000 description 8
- 150000004767 nitrides Chemical class 0.000 description 7
- 239000002994 raw material Substances 0.000 description 7
- 238000005098 hot rolling Methods 0.000 description 6
- 229910052742 iron Inorganic materials 0.000 description 6
- 238000009864 tensile test Methods 0.000 description 6
- QZPSXPBJTPJTSZ-UHFFFAOYSA-N aqua regia Chemical compound Cl.O[N+]([O-])=O QZPSXPBJTPJTSZ-UHFFFAOYSA-N 0.000 description 5
- 229910052799 carbon Inorganic materials 0.000 description 5
- 239000011159 matrix material Substances 0.000 description 5
- 150000001247 metal acetylides Chemical class 0.000 description 5
- 238000000137 annealing Methods 0.000 description 4
- 150000001875 compounds Chemical class 0.000 description 4
- 238000001816 cooling Methods 0.000 description 4
- 230000007423 decrease Effects 0.000 description 4
- 239000000463 material Substances 0.000 description 4
- 230000006911 nucleation Effects 0.000 description 4
- 238000010899 nucleation Methods 0.000 description 4
- 230000003287 optical effect Effects 0.000 description 4
- 239000011324 bead Substances 0.000 description 3
- 238000005266 casting Methods 0.000 description 3
- 238000009792 diffusion process Methods 0.000 description 3
- 238000005242 forging Methods 0.000 description 3
- 238000002844 melting Methods 0.000 description 3
- 230000008018 melting Effects 0.000 description 3
- 230000007797 corrosion Effects 0.000 description 2
- 238000005260 corrosion Methods 0.000 description 2
- 229910000816 inconels 718 Inorganic materials 0.000 description 2
- 238000011835 investigation Methods 0.000 description 2
- 238000012423 maintenance Methods 0.000 description 2
- 238000005259 measurement Methods 0.000 description 2
- 230000003647 oxidation Effects 0.000 description 2
- 238000007254 oxidation reaction Methods 0.000 description 2
- 238000005498 polishing Methods 0.000 description 2
- 238000012545 processing Methods 0.000 description 2
- 238000005204 segregation Methods 0.000 description 2
- 238000005482 strain hardening Methods 0.000 description 2
- 238000012360 testing method Methods 0.000 description 2
- 229910000604 Ferrochrome Inorganic materials 0.000 description 1
- 229910018487 Ni—Cr Inorganic materials 0.000 description 1
- 238000003723 Smelting Methods 0.000 description 1
- 239000005539 carbonized material Substances 0.000 description 1
- 230000003197 catalytic effect Effects 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 238000005520 cutting process Methods 0.000 description 1
- 230000002542 deteriorative effect Effects 0.000 description 1
- 238000005530 etching Methods 0.000 description 1
- 239000003779 heat-resistant material Substances 0.000 description 1
- 238000004881 precipitation hardening Methods 0.000 description 1
- 230000001681 protective effect Effects 0.000 description 1
- 238000004080 punching Methods 0.000 description 1
- 229910052761 rare earth metal Inorganic materials 0.000 description 1
- 238000007670 refining Methods 0.000 description 1
- 239000006104 solid solution Substances 0.000 description 1
- 239000012798 spherical particle Substances 0.000 description 1
- 230000000087 stabilizing effect Effects 0.000 description 1
- 239000010935 stainless steel Substances 0.000 description 1
- 229910001220 stainless steel Inorganic materials 0.000 description 1
- 229910000601 superalloy Inorganic materials 0.000 description 1
- 229910052718 tin Inorganic materials 0.000 description 1
- 229910052725 zinc Inorganic materials 0.000 description 1
Classifications
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C30/00—Alloys containing less than 50% by weight of each constituent
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C30/00—Alloys containing less than 50% by weight of each constituent
- C22C30/02—Alloys containing less than 50% by weight of each constituent containing copper
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
- C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
Abstract
この合金板は、質量％で、Ｃ：０．０２００％以下、Ｓｉ：０．０２〜２．００％、Ｍｎ：０．０２〜２．００％、Ｐ：０．０５０％以下、Ｓ：０．０１００％以下、Ｃｒ：１２．００〜３０．００％、Ｎｉ：３５．００〜６０．００％、Ｎ：０．０２００％以下、Ｖ：０．０２〜１．００％、Ｎｂ：２．００％超、３．５０％以下、Ａｌ：２．６０％超、４．００％以下を含有し、更に任意にＴｉ、Ｍｏ、Ｗ、Ｃｕ、Ｃｏ、Ｂ、Ｚｒ、Ｃａ、Ｍｇの１種以上を含有し、残部が１０．００％以上のＦｅ、及び不純物である化学組成を有し、ミクロ組織が、オーステナイト相を含み、前記オーステナイト相の平均結晶粒径が２０μｍ以上であり、前記オーステナイト相の結晶粒のアスペクト比であるＬ１／Ｌ２の平均値が１．０〜３．０であり、前記ミクロ組織において、γ´相の面積率が１５．０％以下であり、室温での０．２％耐力が４００ＭＰａ以上である。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. More than 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 The balance is 10.00% or more of Fe and the chemical composition is an impurity, the microstructure includes an austenite phase, the average crystal grain size of the austenite phase is 20 μm or more, and the austenite 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
本発明は、合金板及びその製造方法に関する。
本願は、２０１８年０３月２８日に、日本に出願された特願２０１８−０６２５４７号、及び２０１８年１１月２９日に日本に出願された２０１８−２２３４９３号に基づき優先権を主張し、その内容をここに援用する。The present invention relates to an alloy plate and a manufacturing method thereof.
This application claims priority based on Japanese Patent Application No. 2018-062547 filed in Japan on March 28, 2018 and 2018-223493 filed in Japan on November 29, 2018. Is hereby incorporated by reference.
自動車のエンジン排気系については、エンジンの排気ガス出口から、排気マニホールド、ターボチャージャー、触媒コンバーター、排気管などの排気部品が、それぞれフランジなどを用いて接続されている。フランジなどの接続部においては、排気ガスが外部に漏れないように、ガスケットなどを挿入して気密性を確保している。 As for the engine exhaust system of an automobile, exhaust parts such as an exhaust manifold, a turbocharger, a catalytic converter, and an exhaust pipe are connected from the exhaust gas outlet of the engine using flanges or the like. In the connection part such as a flange, a gasket or the like is inserted to ensure airtightness so that the exhaust gas does not leak outside.
上記のような排気系において用いられるガスケット（排気系ガスケット）は、自動車のエンジン排気系部材において耐熱部材として用いられるシール部品であり、板厚０．１〜０．３ｍｍ程度の金属薄板にビードと呼ばれる段差部を形成して製造される。同部品は、自動車の排気系において、フランジ等の連結部に挟み込まれて使用され、ビードの反発力によって連結部からの排気ガスの漏洩を防止する役割を果たす。排気系ガスケットは、高温の排ガスが存在する環境下に長時間曝される苛酷な条件で使用されるので、高温で長時間保持後にも高い高温強度を有することが求められる。特に、ターボチャージャーとの連結部に用いられるターボガスケットは、７００℃を超える高温環境下で使用される。このため、ターボガスケットの素材には、現在、例えばインコネル７１８といった析出硬化型のＮｉ基合金の冷延薄板が用いられている。しかしながら、近年、使用温度の高温化により、インコネル７１８のような既存の耐熱合金では、所定の高温強度が得られなくなってきた。 A gasket (exhaust system gasket) used in the exhaust system as described above is a seal part used as a heat-resistant member in an automobile engine exhaust system member, and a bead is formed on a metal thin plate having a thickness of about 0.1 to 0.3 mm. It is manufactured by forming a stepped portion called. The parts are used by being sandwiched between connecting parts such as flanges in an exhaust system of an automobile, and serve to prevent leakage of exhaust gas from the connecting parts due to the repulsive force of the beads. Since the exhaust system gasket is used under severe conditions that are exposed to an environment where high temperature exhaust gas exists for a long time, it is required to have a high high temperature strength even after being held at a high temperature for a long time. In particular, a turbo gasket used for a connecting portion with a turbocharger is used in a high temperature environment exceeding 700 ° C. For this reason, precipitation-hardened Ni-based alloy cold-rolled thin plates such as Inconel 718 are currently used as the material for the turbo gasket. However, in recent years, due to the increase in use temperature, it has become impossible to obtain a predetermined high-temperature strength with an existing heat-resistant alloy such as Inconel 718.
例えば、特許文献１には、質量％でＣ：０．１５％以下、Ｓｉ：１．０％以下、Ｍｎ：０．３％以下、Ｎｉ：３０〜４９％、Ｃｒ：１０〜１８％、Ａｌ：１．６〜３．０％を含み、ＩＶａ族とＶａ族から選ばれる１種または２種以上の元素を合計で１．５〜８．０％含有し、残部は不純物を除き本質的にＦｅからなることを特徴とするＦｅ−Ｎｉ−Ｃｒ基超耐熱合金が開示されている。 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% in total of one or more elements selected from IVa group and Va group, and the balance is essentially excluding impurities An Fe—Ni—Cr-based superalloy based on Fe is disclosed.
特許文献２には、質量％で、Ｃ：０．０２〜０．３０％、Ｓｉ：０．０２〜３．５％、Ｍｎ：０．０２〜２．５％、Ｎｉ：１０〜５０％、Ｃｒ：１２〜２５％、Ｔｉ：１．０〜５．０％、Ａｌ：０．００２〜１．０％を含有し、かつＮｂ：０．１〜３．０％、Ｂ：０．００１〜０．０１、Ｍｏ：０．１〜４．０％から選択された１種以上の元素を含有して、Ｔｉ、Ａｌ及びＮｂの合計含有量が３．０〜７．０％である耐熱ステンレス鋼であって、粒界に析出するη相［Ｎｉ３Ｔｉ：ｈｃｐ構造］と基地であるγ相（オーステナイト）結晶粒内に析出するγ´相［Ｎｉ３（Ａｌ，Ｔｉ，Ｎｂ）］との重量比率「｛η相［Ｎｉ３Ｔｉ：ｈｃｐ構造］／γ´相［Ｎｉ３（Ａｌ，Ｔｉ，Ｎｂ）］｝×１００（％）」が、０．０１％以上３０．００％以下であり、６００℃での熱間引張強さが８００Ｎ／ｍｍ２以上であり、また基地であるγ相（オーステナイト）結晶粒内に析出するγ´相［Ｎｉ３（Ａｌ，Ｔｉ，Ｎｂ）］の球状粒子の直径が１ｎｍ以上２０ｎｍ以下であることを特徴とする耐熱ステンレス鋼が開示されている。In Patent Document 2, in mass%, C: 0.02 to 0.30%, Si: 0.02 to 3.5%, Mn: 0.02 to 2.5%, Ni: 10 to 50%, Cr: 12-25%, Ti: 1.0-5.0%, Al: 0.002-1.0%, Nb: 0.1-3.0%, B: 0.001- 0.01, Mo: One or more elements selected from 0.1 to 4.0%, and the total content of Ti, Al and Nb is 3.0 to 7.0%. Η phase [Ni 3 Ti: hcp structure] precipitated in grain boundaries and γ ′ phase [Ni 3 (Al, Ti, Nb)] precipitated in the base γ phase (austenite) crystal grains 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)] There is disclosed a heat resistant stainless steel characterized in that the diameter of the spherical particles is from 1 nm to 20 nm.
特許文献３には、質量％で、Ｃ：０．０１〜０．１％、Ｓｉ：２％以下、Ｍｎ：２％以下、Ｃｒ：１２〜２５％、Ｎｂ＋Ｔａ：０．２〜２．０％、Ｔｉ：１．５％未満、Ａｌ：０．５〜３．０％、Ｎｉ：２５〜４５％、Ｃｕ：０．１〜５．０％であり、またＷ：３％以下、Ｍｏ：３％以下、Ｖ：１％以下、Ｃｏ：５％以下、Ｂ：０．００１〜０．０１％、Ｚｒ：０．００１〜０．１％、Ｃａ＋Ｍｇ：０．００１〜０．０１％の何れか１種もしくは２種以上を含み、かつＴｉ／Ａｌ＝０．１１５〜１．０、１／２Ｗ＋Ｍｏ＋Ｖ＜３％、Ｔｉ＋Ａｌ＋Ｎｂ＋Ｔａ：４．５〜７．０％を満たす冷間加工性及び過時効特性に優れた耐熱合金が開示されている。 In Patent Document 3, in mass%, 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% Ti: Less than 1.5%, Al: 0.5-3.0%, Ni: 25-45%, Cu: 0.1-5.0%, W: 3% or less, Mo: 3 %, 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% For cold workability and overaging characteristics including one or more, 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.
本発明者らの検討結果によれば、これらの先行技術に開示の合金を用いたガスケットでは、近年の排気系ガスケットに要求されるレベルの高温で長時間使用された際には、使用中に高温強度が低下する等により、十分な高温強度が確保できないことが分かった。 According to the results of the study by the present inventors, these gaskets using the alloys disclosed in the prior art are in use when used for a long time at a high temperature that is required for recent exhaust system gaskets. It was found that sufficient high temperature strength could not be secured due to a decrease in high temperature strength.
本発明は、上記の課題に鑑みてなされた。本発明は、高温で長時間使用されるガスケットとして用いた際にも十分な高温強度が確保される、合金板及びその製造方法を提供することを目的とする。 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 that can ensure sufficient high-temperature strength even when used as a gasket that is used at a high temperature for a long time.
即ち、本発明の要旨とするところは以下のとおりである。
（１）本発明の一態様に係る合金板は、質量％で、Ｃ：０．０２００％以下、Ｓｉ：０．０２〜２．００％、Ｍｎ：０．０２〜２．００％、Ｐ：０．０５０％以下、Ｓ：０．０１００％以下、Ｃｒ：１２．００〜３０．００％、Ｎｉ：３５．００〜６０．００％、Ｎ：０．０２００％以下、Ｖ：０．０２〜１．００％、Ｎｂ：２．００％超、３．５０％以下、Ａｌ：２．６０％超、４．００％以下、Ｔｉ：０〜０．８０％、Ｍｏ：０〜２．００％、Ｗ：０〜２．００％、Ｃｕ：０〜１．００％、Ｃｏ：０〜１．００％、Ｂ：０〜０．０１００％、Ｚｒ：０〜０．０１００％、Ｃａ：０〜０．００５０％、Ｍｇ：０〜０．００５０％、を含有し、残部が１０．００％以上のＦｅ、及び不純物である化学組成を有し、ミクロ組織が、オーステナイト相を含み、前記オーステナイト相の平均結晶粒径が２０μｍ以上であり、前記オーステナイト相の結晶粒の、長軸の長さをＬ１、短軸の長さをＬ２とするとき、アスペクト比であるＬ１／Ｌ２の平均値が、１．０〜３．０であり、前記ミクロ組織において、γ´相の面積率が１５．０％以下であり、室温での０．２％耐力が４００ＭＰａ以上である。That is, the gist of the present invention is as follows.
(1) The alloy plate which concerns on 1 aspect of this invention is the 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 to 30.00%, Ni: 35.00 to 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 0.0050%, Mg: 0 to 0.0050%, the balance is Fe of 10.00% or more, and the chemical composition is an impurity, and the microstructure is austenite L1 which is an aspect ratio when the average grain size of the austenite phase is 20 μm or more, the major axis length of the austenite phase crystal grains is L1, and the minor axis length is L2. / L2 has an average value of 1.0 to 3.0, and in the microstructure, the area ratio of the γ ′ phase is 15.0% or less, and the 0.2% proof stress at room temperature is 400 MPa or more. .
（２）上記（１）に記載の合金板では、前記アスペクト比が１．０〜２．０であってもよい。 (2) In the alloy plate according to (1), the aspect ratio may be 1.0 to 2.0.
（３）上記（１）または（２）に記載の合金板では、７５０℃で１００時間の時効処理を行ったとき、前記時効処理後のミクロ組織におけるγ´相の面積率が１５．０％以上、前記γ´相の平均円相当径が３０ｎｍ以上、７０ｎｍ以下、かつ、γ´相以外の析出相の面積率が５．０％以下であってもよい。 (3) In the alloy plate according to (1) or (2) above, 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 precipitated phase other than the γ ′ phase may be 5.0% or less.
（４）本発明の別の態様に係る合金板の製造方法は、（１）〜（３）のいずれかに記載の合金板を製造する方法であって、圧延加工された合金素材を１０５０〜１２００℃で、３０秒〜６００秒間溶体化熱処理する第１の工程と、板厚減少率５％以上、４５％以下で仕上げ冷延する第２の工程と、を含む。 (4) The manufacturing method of the alloy plate which concerns on another aspect of this invention is a method of manufacturing the alloy plate in any one of (1)-(3), Comprising: The alloy raw material by which the rolling process was carried out is 1500. 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% to 45% are included.
（５）本発明の別の態様に係る合金板の製造方法は、（１）〜（３）のいずれかに記載の合金板を製造する方法であって、圧延加工された合金素材を１０５０〜１２００℃で、３０秒〜６００秒間溶体化熱処理する第３の工程と、７００〜８５０℃、３０秒〜６００秒で時効処理する第４の工程と、を含む。 (5) The manufacturing method of the alloy plate which concerns on another aspect of this invention is a method of manufacturing the alloy plate in any one of (1)-(3), Comprising: The alloy raw material by which the rolling process was carried out is 1500. 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 are included.
（６）上記（５）に記載の合金板の製造方法では、前記第３の工程の後かつ前記第４の工程の前に、板厚減少率４５％以下の圧下率で仕上げ冷延する第５の工程を含んでもよい。 (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. Five steps may be included.
本発明の上記態様に係る合金板は、排気系ガスケット等に使用され、高温に長時間曝された場合でも、十分な高温強度が確保される。 The alloy plate according to the above aspect of the present invention is used for an exhaust system gasket and the like, and sufficient high-temperature strength is ensured even when exposed to a high temperature for a long time.
本発明者らは、合金の組成や製造プロセスについて詳細な調査を実施し、近年要求される厳しい使用環境に耐える、排気系ガスケットに好適な合金板を検討した。 The present inventors conducted a detailed investigation on the composition and manufacturing process of the alloy, and studied an alloy plate suitable for an exhaust system gasket that can withstand a severe use environment required in recent years.
高温で長時間使用されるガスケットは、耐熱合金薄板を用いて形成される。耐熱合金薄板は、排気ガスプロセスの高温環境にさらされることによって合金中に微細なγ´（ガンマプライム）相が析出し、このγ´相によって高温強度が確保される。γ´相は、Ｎｉ３（Ａｌ，Ｔｉ，Ｎｂ）で表現される析出物であり、合金中にＴｉを含有しない場合にはＮｉ３（Ａｌ，Ｎｂ）となる。γ´相は、母相となるオーステナイト相との界面の整合性が高く、他の析出強化相と比べて母相粒内に均一に微細分散しやすいことや、γ´相自体の強度が温度との逆依存性を有することなどから、特に高温での析出強化能が高いという特徴を有する。成分によって析出相の状態が変化するので、合金成分を適正化することにより、高温環境でγ´相の適正な析出を促すことができる。A gasket that is used for a long time at a high temperature is formed using a heat-resistant alloy thin plate. When the heat-resistant alloy thin plate is exposed to a high temperature environment of an 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 expressed 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 that is the parent phase, and is more easily and finely dispersed in the parent phase grains than other precipitation strengthening phases, and the strength of the γ' phase itself is the temperature. In particular, it has a feature of high precipitation strengthening ability at a high temperature. Since the state of the precipitated phase changes depending on the component, appropriate precipitation of the γ ′ phase can be promoted in a high temperature environment by optimizing the alloy component.
ガスケットが自動車に装着され、高温に長時間曝された（以下「高温長時間時効」ともいう。）後でも気密性保持機能を十分に維持するためには、このγ´相が高温長時間時効後も安定して存在し、それによって高温強度が維持されている必要がある。しかしながら、本発明者らの検討の結果、従来知られていた耐熱合金板をガスケットに適用した場合、特にターボガスケットのように極めて高温に長時間曝される場合には、高温強度が低下することが明らかとなった。 In order to maintain the airtightness maintenance function 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 long time aging”), this γ ′ phase is high temperature long time aging. It is necessary to exist stably afterwards and to maintain high temperature strength. 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 a very high temperature for a long time like a turbo gasket, the high-temperature strength is reduced. Became clear.
従来知られている耐熱合金において高温強度が低下する理由について、本発明者らが検討を行った結果、従来の耐熱合金では、高温長時間時効処理（高温長時間時効熱処理）において、合金中にＬａｖｅｓ相などのγ´相以外の化合物（非強化相）が析出するために高温強度が低下することがわかった。このＬａｖｅｓ相の周囲には、γ´相の存在しない無析出帯が形成されるので、その結果として、γ´相の存在比率が低下し、高温強化能が著しく劣化すると考えられる。すなわち、本発明者らの検討の結果、高温長時間時効時にＬａｖｅｓ相などのγ´相以外の析出物（非強化相）を抑制することが高温強度維持に重要であることがわかった。 As a result of investigations by the present inventors on the reason why the high temperature strength is lowered in a conventionally known heat resistant alloy, the conventional heat resistant alloy has a high temperature long time aging treatment (high temperature long time aging heat treatment). It has been found that the high-temperature strength is lowered due to precipitation of a compound other than the γ ′ phase (non-strengthened phase) such as the Laves phase. A precipitate-free zone in which no γ ′ phase is present is formed around the Laves phase. As a result, it is considered that the abundance ratio of the γ ′ phase is reduced and the high-temperature strengthening ability is significantly deteriorated. That is, as a result of the study by the present inventors, it was found that it is important for maintaining high-temperature strength to suppress precipitates (non-strengthened phase) other than the γ ′ phase such as the Laves phase during high-temperature long-term aging.
高温長時間後時効後の合金板から、強化相であるγ´相、及びＬａｖｅｓ相などの非強化相の、高温長時間時効時の析出挙動を詳細に調査した結果、γ´相が母相全面に分散して微細析出するのに対し、Ｌａｖｅｓ相は結晶粒界や転位を核生成サイトとして析出することがわかった。したがって、これらの核生成サイトをできるだけ少なくすることで、Ｌａｖｅｓ相などの非強化相の析出を抑制し、強化相であるγ´相を維持させられることが分かった。 As a result of investigating the precipitation behavior of high-strength and long-term aging of non-strengthened phases such as the strengthening phase γ 'phase and Laves phase from the alloy plate after high-temperature long-time post-aging, It was found that the Laves phase precipitates with crystal grain boundaries and dislocations as nucleation sites, while finely precipitated by being dispersed over the entire surface. Therefore, it was found that by reducing these nucleation sites as much as possible, precipitation of a non-strengthening phase such as a Laves phase can be suppressed and the γ ′ phase that is a strengthening phase can be maintained.
高温長時間時効後も高温強度を維持するためには、１）γ´相の析出を促し、Ｌａｖｅｓ相などの非強化相の析出を抑制するように成分をコントロールすること、２）熱処理（溶体化処理）によって結晶粒径を粗大化し、それによって結晶粒界を少なくし、Ｌａｖｅｓ相などの非強化相の析出サイト低減を図ること、３）最終の仕上げ冷延における板厚減少率を低下させることで合金板の結晶粒内における転位を少なくし、Ｌａｖｅｓ相などの非強化相の析出サイト低減を図ること、が重要である。本発明者らは、これらプロセス上の改善も検討し、本発明を完成させた。 In order to maintain high temperature strength even after aging for a long time at high temperature, 1) control the components to promote precipitation of γ ′ phase and suppress precipitation of non-strengthening phase such as Laves phase, and 2) heat treatment (solution) The crystal grain size is increased by reducing the grain boundaries, thereby reducing the precipitation sites of non-strengthened phases such as the Laves phase, and 3) reducing the plate thickness reduction rate in the final finish cold rolling. Therefore, it is important to reduce dislocations in the crystal grains of the alloy plate and reduce the precipitation sites of non-strengthened phases such as the Laves phase. The inventors have also studied these process improvements and completed the present invention.
以下、本発明の一実施形態に係る合金板（本実施形態に係る合金板）について説明する。
まず第１に、本実施形態に係る合金板の成分組成（化学組成）について説明する。成分の含有量に関する％は断りがない限り質量％を意味する。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 the present embodiment will be described. Unless otherwise indicated,% regarding the content of a component means the mass%.
Ｃ：０．０２００％以下
Ｃは、Ｔｉ、Ｎｂと結びついて炭化物を形成する元素である。Ｔｉ、Ｎｂが炭化物を形成すると、強化相であるγ´相の生成量が減少する。そのため、Ｃ含有量は低い方が望ましい。したがって、Ｃ含有量は０．０２００％以下とする。
一方、Ｃ含有量を０．０００２％未満にすると製造コストが著しく増加する。そのため、Ｃ含有量を０．０００２％以上としてもよい。C: 0.0200% or less C is an element that forms a carbide in combination with Ti and Nb. When Ti and Nb form carbides, the amount of γ ′ phase that is a strengthening phase is reduced. Therefore, it is desirable that the C content is low. Therefore, the C content is 0.0200% or less.
On the other hand, when the C content is less than 0.0002%, the production cost is remarkably increased. Therefore, the C content may be 0.0002% or more.
Ｓｉ：０．０２％以上、２．００％以下
Ｓｉは、精錬の際に脱酸元素として有効な元素である。また、Ｓｉは、合金の耐酸化性を改善する皮膜の保護性を向上させる元素である。これらの効果を得るため、Ｓｉ含有量を０．０２％以上とする。
一方で、Ｓｉ含有量が多すぎると、熱間加工性が顕著に劣化し、耳割れが発生したり、手入れコストが大きくなったりする。そのため、Ｓｉ含有量を２．００％以下とする。Si: 0.02% or more and 2.00% or less Si is an element effective as a deoxidizing element during refining. Si is an element that improves the protective properties 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, when there is too much Si content, hot workability will deteriorate notably, an ear crack will generate | occur | produce, or maintenance cost will become large. Therefore, the Si content is set to 2.00% or less.
Ｍｎ：０．０２％以上、２．００％以下
Ｍｎ含有量が多すぎると、熱間加工性が劣化する上、高温での耐酸化性が著しく劣化する。したがって、Ｍｎ含有量は２．００％以下とする。
一方、Ｍｎは原料スクラップなどから混入するので、Ｍｎ含有量を大きく低減させるにはスクラップの使用を減らす必要があり、コストの増大を招く。したがって、Ｍｎ含有量は０．０２％以上とする。Mn: 0.02% or more and 2.00% or less When the Mn content is too large, hot workability is deteriorated and oxidation resistance at high temperatures is remarkably deteriorated. Therefore, the Mn content is 2.00% or less.
On the other hand, since Mn is mixed from raw material scrap or the like, 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 0.02% or more.
Ｐ：０．０５０％以下
Ｐは合金の原料の一つであるフェロクロムに含まれる不純物元素である。Ｐは熱間加工性に有害であるので、Ｐ含有量を０．０５０％以下に制限する。P: 0.050% or less P is an impurity element contained in ferrochrome which is one of the raw materials of the alloy. Since P is harmful to hot workability, the P content is limited to 0.050% or less.
Ｓ：０．０１００％以下
Ｓは原料のスクラップなどに含まれる不純物元素である。Ｓは熱間加工性に有害であるので、Ｓ含有量は０．０１００％以下に制限する。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％以上、３０．００％以下
Ｃｒは耐熱合金としての耐食性を確保する観点から必須の元素である。十分な耐食性を確保する観点から、Ｃｒ含有量は１２．００％以上とする。好ましくは１４．００％以上である。
一方で、Ｃｒ含有量が多すぎると、焼鈍時にσ相などの粗大な化合物が生成し、γ´相の析出が抑制されるだけでなく、材料が脆化しやすくなる。そのため、Ｃｒ含有量は３０．００％以下とする。好ましくは、２０．００％以下である。Cr: 12.00% or more, 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 set to 12.00% or more. Preferably it is 14.00% or more.
On the other hand, if the Cr content is too large, a coarse compound such as a σ phase is generated during annealing, and not only the precipitation of the γ ′ phase is suppressed, but the material is easily embrittled. Therefore, the Cr content is 30.00% or less. Preferably, it is 20.00% or less.
Ｎｉ：３５．００％以上、６０．００％以下
Ｎｉは強力なオーステナイト安定化元素であり、ミクロ組織においてオーステナイト母相を得るために必須の元素である。また、Ｎｉは高温強度の向上に寄与する強化相であるγ´相（Ｎｉ３（Ａｌ，Ｎｂ，Ｔｉ））を得るためにも極めて重要な元素である。耐熱材料として高温での強度を確保する観点から、Ｎｉ含有量は３５．００％以上とする。好ましくは４０．００％以上である。
一方、Ｎｉ含有量が多すぎると、コストが上昇することに加え、熱間加工時の変形抵抗が高くなって、製造が困難になる。そのため、Ｎｉ含有量は６０．００％以下とする。好ましくは、５０．００％以下である。Ni: 35.00% or more and 60.00% or less Ni is a strong austenite stabilizing element and an essential element for obtaining an austenite matrix in the microstructure. Ni is also an extremely important element for obtaining a γ ′ phase (Ni 3 (Al, Nb, Ti)), which is a strengthening phase that contributes to the improvement of the high-temperature strength. From the viewpoint of securing strength at high temperature as a heat-resistant material, the Ni content is set to 35.00% or more. Preferably it is 40.00% or more.
On the other hand, if the Ni content is too large, the cost increases and the deformation resistance during hot working becomes high, which makes manufacturing difficult. Therefore, the Ni content is 60.00% or less. Preferably, it is 50.00% or less.
Ｎ：０．０２００％以下
Ｎは、Ａｌ、Ｎｂ、Ｔｉと結合して窒化物を形成する元素である。Ａｌ、Ｎｂ、Ｔｉが窒化物を形成すると、強化相であるγ´相の生成量が減少する。そのため、Ｎ含有量は低い方が望ましい。したがって、Ｎ含有量は０．０２００％以下に制限する。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 γ ′ phase that is a 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％以上、１．００％以下
Ｖは、固溶強化により高温強度の向上に寄与する元素である。この効果を得るため、Ｖ含有量は０．０２％以上とする。
一方、ＶはＣ、Ｎと結合して炭化物、窒化物を形成する。Ｖ含有量が多すぎると、粗大な炭化物、窒化物が生成し、材料の加工性が劣化する。そのため、Ｖ含有量は１．００％以下とする。V: 0.02% or more and 1.00% or less V is an element that contributes to 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. When there is too much V content, coarse carbide | carbonized_material and nitride will produce | generate and the workability of material will deteriorate. Therefore, the V content is 1.00% or less.
Ｎｂ：２．００％超、３．５０％以下
Ｎｂは強化相であるγ´相を構成する元素であり、Ｎｂを固溶したγ´相は高温強化能が高い。そのため、高温強度確保のため、Ｎｂ含有量は２．００％超とする。
一方、Ｎｂは合金の融点を下げ、高温での熱間加工を困難にする元素である。したがって、Ｎｂ含有量は３．５０％以下とする。Nb: more than 2.00% and not more than 3.50% Nb is an element constituting a γ ′ phase that is a strengthening phase, and a γ ′ phase in which Nb is dissolved has a high temperature strengthening ability. For this reason, the Nb content is more than 2.00% to ensure high temperature strength.
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：２．６０％超、４．００％以下
Ａｌは高温強度の上昇に寄与する強化相であるγ´相を構成する元素である。その強化能は、Ｔｉ、Ｎｂには及ばないが、Ｔｉ、Ｎｂに比べて、長時間安定してγ´相を維持する効果がある。長時間時効後に安定して高温強度を維持する観点から、Ａｌ含有量は２．６０％超とする。
一方、Ａｌ含有量が過剰であると、合金の融点が下がり、高温での熱間加工が困難になる。したがって、Ａｌ含有量は、４．００％以下とする。Al: more than 2.60%, 4.00% or less Al is an element constituting a γ ′ phase which is a strengthening phase contributing to an increase in high temperature strength. The strengthening ability does not reach that of Ti and Nb, but 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 high temperature strength after aging for a long time, the Al content is set to exceed 2.60%.
On the other hand, if the Al content is excessive, the melting point of the alloy decreases and hot working at high temperatures becomes difficult. Therefore, the Al content is 4.00% or less.
本実施形態に係る合金板は、さらに以下の任意元素を含有しても良い。ただし、以下の任意元素は必ずしも含有する必要がないので、その下限は０％である。 The alloy plate according to this embodiment may further contain the following optional elements. However, since the following optional elements do not necessarily need to be contained, the lower limit is 0%.
Ｔｉ：０〜０．８０％
Ｔｉは強化相であるγ´相を構成する元素であり、高温強度の向上に寄与する元素である。そのため、必要に応じて含有させてもよい。上記効果を得る場合、Ｔｉ含有量を０．２０％以上とすることが好ましい。
一方、Ｔｉを多量に含有すると、合金中のＣ、ＮとＴｉとが粗大な炭化物、窒化物を形成し、熱間加工性や冷間加工性が著しく劣化する。また、Ｔｉの含有により合金の融点が下がるので、Ｔｉ含有量が過剰であると高温での熱間加工が困難になる。しがたって、含有させる場合でもＴｉ含有量は０．８０％以下とする。好ましくは、０．６０％以下である。Ti: 0 to 0.80%
Ti is an element constituting a γ ′ phase that is a strengthening phase, and is an element that contributes to an improvement in high-temperature strength. Therefore, you may make it contain as needed. When obtaining the above effect, the Ti content is preferably 0.20% or more.
On the other hand, when Ti is contained in a large amount, C, N and Ti in the alloy form coarse carbides and nitrides, and hot workability and cold workability are remarkably deteriorated. Further, since the melting point of the alloy is lowered due to the inclusion of Ti, hot working at a high temperature becomes difficult if the Ti content is excessive. Therefore, even when it is contained, the Ti content is set to 0.80% or less. Preferably, it is 0.60% or less.
Ｍｏ：０〜２．００％
Ｍｏは母相であるオーステナイト相に固溶し、高温強度の向上に寄与する元素である。そのため、必要に応じて含有させてもよい。
一方、多量にＭｏを含有すると、熱間加工時の変形抵抗が増加し、所定の板厚に熱間圧延するのが困難になる。また、高温長時間時効時に粗大なＬａｖｅｓ相の析出が促進され、高温強度が低下する。したがって、含有させる場合でもＭｏ含有量は２．００％以下とする。Mo: 0 to 2.00%
Mo is an element that dissolves in the austenite phase that is the parent phase and contributes to the improvement of the high-temperature strength. Therefore, you may make it contain as needed.
On the other hand, when Mo is contained in a large amount, deformation resistance during hot working increases, and it becomes difficult to hot-roll to a predetermined plate thickness. Moreover, precipitation of a coarse Laves phase is promoted at the time of high temperature aging for a long time, and the high temperature strength is lowered. Therefore, even when it contains, Mo content shall be 2.00% or less.
Ｗ：０〜２．００％
Ｗはオーステナイト母相に固溶し、高温強度の向上に寄与する元素である。そのため必要に応じて含有させてもよい。しかしながら、多量にＷを含有すると、熱間加工時の変形抵抗が増加し、所定の板厚に熱間圧延するのが困難になる。また、高温長時間時効時に粗大なＬａｖｅｓ相の析出が促進され、高温強度が低下する。したがって、含有させる場合でもＷ含有量は２．００％以下とする。W: 0 to 2.00%
W is an element that dissolves in the austenite matrix and contributes to the improvement of high-temperature strength. Therefore, you may make it contain as needed. However, when a large amount of W is contained, deformation resistance during hot working increases, and it becomes difficult to hot-roll to a predetermined plate thickness. Moreover, precipitation of a coarse Laves phase is promoted at the time of high temperature aging for a long time, and the high temperature strength is lowered. Therefore, even when contained, the W content is 2.00% or less.
Ｃｕ：０〜１．００％
Ｃｕは、オーステナイト母相に固溶し、高温強度を上げる効果を有する元素である。そのため必要に応じて含有させてもよい。
一方、Ｃｕ含有量が過剰になると、熱間圧延時の耳割れが発生する場合がある。したがって、含有させる場合でもＣｕ含有量は１．００％以下とする。Cu: 0 to 1.00%
Cu is an element that has the effect of increasing the high temperature strength by dissolving in the austenite matrix. Therefore, you may make it contain as needed.
On the other hand, if the Cu content is excessive, ear cracks may occur during hot rolling. Therefore, even when it contains, Cu content shall be 1.00% or less.
Ｃｏ：０〜１．００％
ＣｏはＮｉの代替としてγ´相に固溶する元素であり、高温強度の向上に寄与する元素である。そのため、含有させてもよい。
一方で、Ｃｏを多量に含有すると、コストの増加に加えて、熱間加工時の変形抵抗が増加し、所定の板厚に熱間圧延するのが困難になる。したがって、含有させる場合でもＣｏ含有量は１．００％以下とする。Co: 0 to 1.00%
Co is an element that dissolves in the γ ′ phase as an alternative to Ni, and is an element that contributes to the improvement of high-temperature strength. Therefore, you may make it contain.
On the other hand, when Co is contained in a large amount, in addition to an increase in cost, deformation resistance during hot working increases, and it becomes difficult to hot-roll to a predetermined plate thickness. Therefore, even when it is contained, the Co content is 1.00% or less.
Ｂ：０〜０．０１００％
Ｂは結晶粒界に偏析する元素であり、結晶粒界を強化することで、粒界でのすべりを抑制し、高温強度の向上に寄与する元素である。そのため、含有させても良い。Ｂを含有することによる上述の効果を得るためには、Ｂ含有量は好ましくは０．００１０％以上である。
一方で、Ｂを多量に含有すると粒界偏析が顕著になり、熱間加工性が著しく低下する。そのため、含有させる場合でもＢ含有量は０．０１００％以下とする。B: 0 to 0.0100%
B is an element that segregates at the crystal grain boundary, and is an element that contributes to the improvement of the high-temperature strength by suppressing the slip at the grain boundary by strengthening the crystal grain boundary. Therefore, you may make it contain. In order to obtain the above-described effect by containing B, the B content is preferably 0.0010% or more.
On the other hand, when B is contained in a large amount, grain boundary segregation becomes remarkable and hot workability is remarkably lowered. Therefore, even when it contains, B content shall be 0.0100% or less.
Ｚｒ：０〜０．０１００％
Ｚｒは結晶粒界に偏析する元素であり、結晶粒界を強化することで、粒界でのすべりを抑制し、高温強度の向上に寄与する元素である。そのため、含有させても良い。Ｚｒを含有することによる上述の効果を得るためには、Ｚｒ含有量は好ましくは０．００１０％以上である。
一方で、Ｚｒを多量に添加すると粒界偏析が顕著になり、熱間加工性が著しく低下する。そのため、含有させる場合でもＺｒ含有量は０．０１００％以下とする。Zr: 0 to 0.0100%
Zr is an element that segregates at the crystal grain boundary, and is an element that contributes to the improvement of high-temperature strength by suppressing the slip at the grain boundary by strengthening the crystal grain boundary. Therefore, you may make it contain. In order to obtain the above-described effect by containing Zr, the Zr content is preferably 0.0010% or more.
On the other hand, when a large amount of Zr is added, grain boundary segregation becomes prominent and hot workability is significantly reduced. Therefore, even when it is made to contain, Zr content shall be 0.0100% or less.
Ｃａ：０〜０．００５０％
Ｃａは熱間加工性を改善する効果を有する元素である。熱間加工性が改善すると製造コストが低減できる。この効果を得るため、含有させても良い。
一方、Ｃａ含有量が多量になると、鋳造時に溶湯ノズルが詰まるなどのトラブルが生じ、製造が著しく困難になる。そのため、含有させる場合でもＣａ含有量は０．００５０％以下とする。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 acquire this effect, you may make it contain.
On the other hand, when the Ca content is large, troubles such as clogging of the molten metal nozzle occur at the time of casting, and manufacturing becomes extremely difficult. Therefore, even when it contains, Ca content shall be 0.0050% or less.
Ｍｇ：０〜０．００５０％
Ｍｇは熱間加工性改善する効果を有する元素である。熱間加工性が改善すると製造コストが低減できる。この効果を得るため、含有させても良い。
一方、Ｍｇ含有量が多量になると、鋳造時に溶湯ノズルが詰まるなどのトラブルが生じ、製造が著しく困難になる。そのため、含有させる場合でもＭｇ含有量は０．００５０％以下とする。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 acquire this effect, you may make it contain.
On the other hand, when the Mg content is large, troubles such as clogging of the molten metal nozzle occur at the time of casting, and manufacturing becomes extremely difficult. Therefore, even when it is made to contain, Mg content shall be 0.0050% or less.
以上説明した成分の含有量合計は９０％以下、好ましくは８０％以下とする。また、以上説明した成分以外の残部は１０．００％以上のＦｅ、及び不純物からなる。
残部に含まれるＦｅが合金板の化学組成の１０．００％未満であると、加工性が劣化する。そのため、本実施形態に係る合金板において、Ｆｅは１０．００％以上とする。
また、不純物は合計で２．０％以下であることが好ましい。不純物とは、上述したＰ、Ｓ以外に、例えばＲＥＭ、Ｓｎ、Ｚｎである。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 Fe and impurities of 10.00% or more.
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 set to 10.00% or more.
Moreover, it is preferable that impurities are 2.0% or less in total. Impurities are, for example, REM, Sn, and Zn other than P and S described above.
次に、合金板の結晶粒径を粗大化し、それによって結晶粒界を少なくし、Ｌａｖｅｓ相などの非強化相の析出サイト低減を図る点について説明する。 Next, the point that the crystal grain size of the alloy plate is coarsened, thereby reducing the grain boundaries and reducing the precipitation sites of non-strengthened phases such as the Laves phase will be described.
＜ミクロ組織＞
本実施形態に係る合金板は、金属組織が、母相としてのオーステナイト相を含む。母相のオーステナイト相には、少量のγ´相等が析出する場合がある。析出物を除く母相は実質的にオーステナイト相からなっていてもよい。<Microstructure>
In the alloy plate according to the present embodiment, the metal structure includes an austenite phase as a parent phase. A small amount of γ 'phase or the like may precipitate in the austenite phase of the parent phase. The parent phase excluding the precipitates may consist essentially of an austenite phase.
［オーステナイト相の平均結晶粒径が２０μｍ以上］
結晶粒界は、時効処理（時効熱処理）時にＬａｖｅｓ相などの非強化相の析出サイトになる。そのため、平均結晶粒径が小さいと、合金板の単位断面あたりの粒界存在比率が多くなり、時効処理時にこれらの非強化相の析出量が多くなる。非強化相の近傍には、γ´相が析出しない無析出相が形成されるため、その結果、強化相であるγ´相の量が少なくなる。
合金板のオーステナイト相の平均結晶粒径が２０μｍ以上であれば、非強化相の析出サイトを低減し、高温長時間時効後のγ´相の析出を確保することができる。したがって、平均結晶粒径は２０μｍ以上とする。平均結晶粒径は７５０℃程度の使用温度での熱処理では大きく変化しないため、仕上げ冷延後の合金板での値とする。合金板の母相であるオーステナイト相の平均結晶粒径を２０μｍ以上とするための製造方法（仕上げ冷延前の溶体化処理）については後述する。
平均結晶粒径の上限を規定する必要はないが、板厚に対して結晶粒が大きすぎると、プレス成型性などが著しく劣化するので、平均結晶粒径は５０μｍ以下が好ましい。[Average crystal grain size of austenite phase is 20 μm or more]
A grain boundary becomes a precipitation site of a non-strengthened phase such as a Laves phase during aging treatment (aging heat treatment). Therefore, when the average crystal grain size is small, the grain boundary existence ratio per unit cross section of the alloy plate increases, and the amount of precipitation of these non-strengthened phases increases during the aging treatment. In the vicinity of the non-strengthened phase, a non-precipitated phase in which the γ ′ phase does not precipitate is formed. As a result, the amount of the γ ′ phase that is the strengthened phase is reduced.
If the average grain size of the austenite phase of the alloy plate is 20 μm or more, the precipitation sites of the non-strengthened phase can be reduced, and the precipitation of the γ ′ phase after high-temperature and long-term aging can be ensured. Therefore, the average crystal grain size is 20 μm or more. Since the average crystal grain size does not change greatly in the heat treatment at the use temperature of about 750 ° C., the value is the value in the alloy plate after finish cold rolling. A production 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 plate, to 20 μm or more will be described later.
There is no need to define the upper limit of the average crystal grain size, but if the crystal grains are too large relative to the plate thickness, the press formability and the like are significantly deteriorated, so the average crystal grain size is preferably 50 μm or less.
平均結晶粒径は以下の方法で求めることができる。
Ｌ断面（圧延方向に平行で板厚面に平行な断面）が観察面となるように金属板からサンプルを採取し、観察面を表面研磨及びエッチングする。このサンプルを、光学顕微鏡で観察、撮影し、ＪＩＳ Ｇ０５５１：２０１３に従って単位面積当たりの結晶粒数を計数して一個あたりの平均結晶粒面積aを求め、aの平方根を平均結晶粒径とする。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 polished and etched. This sample is observed and photographed with an optical microscope, and the number of crystal grains per unit area is counted in accordance with JIS G0551: 2013 to determine the average crystal grain area a, and the square root of a is the average crystal grain size.
＜γ´相の面積率が１５．０％未満＞
本実施形態に係る合金板では、高温長時間時効処理前（ガスケット等として使用される前）のミクロ組織において、γ´相の面積率が１５．０％未満である。
高温長時間時効処理前に既にγ´相が多く析出していると、これらのγ´相が高温長時間時効処理により粗大化し、十分な高温強度を保てない。そのため、合金板のミクロ組織において、γ´相の面積率を１５．０％未満とする。好ましくは、８．０％以下、より好ましくは６．０％以下、さらに好ましくは２．０％以下である。
高温長時間時効処理前のγ´相の面積率を小さくすると、室温での０．２％耐力が低くなることが懸念されるが、後述するように板厚減少率が５％〜４５％の冷間圧延を行うことで、室温での０．２％耐力が確保できる。また、冷間圧延によれば、合金板の平坦度も向上する。<Area ratio of γ 'phase is less than 15.0%>
In the alloy plate according to the present embodiment, the area ratio of the γ ′ phase is less than 15.0% in the microstructure before high-temperature and long-term aging treatment (before being used as a gasket or the like).
If a large amount of γ ′ phase is already precipitated before the high temperature long time aging treatment, these γ ′ phases are coarsened by the high temperature long time aging treatment, and sufficient high temperature strength cannot be maintained. Therefore, the area ratio of the γ ′ phase is set to less than 15.0% in the microstructure of the alloy plate. Preferably, it is 8.0% or less, More preferably, it is 6.0% or less, More preferably, it is 2.0% or less.
When the area ratio of the γ 'phase before high temperature long-term aging treatment is reduced, there is a concern that the 0.2% proof stress at room temperature is lowered, but the plate thickness reduction rate is 5% to 45% as described later. By performing cold rolling, 0.2% proof stress at room temperature can be secured. Further, according to the cold rolling, the flatness of the alloy plate is also improved.
高温長時間時効処理前の合金板である本実施形態に係る合金板のγ´相の面積率は以下の方法で求める。
合金板の、圧延方向に平行な断面を研磨し、研磨面を王水でエッチングし、ＦＥ−ＳＥＭと同設備に付帯されるＥＤＳで観察、分析する。結晶粒内に均一に析出している微細な球状の析出物をγ´相とし、結晶粒界、結晶粒内に析出し、粗大、且つ合金板の化学組成よりも（合金板の化学組成におけるＦｅ、Ｃｒ含有量に比べて）Ｆｅ、Ｃｒが多量に含まれている析出物をγ´相以外の析出物として分類する。ＦＥ−ＳＥＭで撮影した写真から、線分法でγ´相の面積率を算出する。微細な球状の析出物とは、概ね、円相当直径で１００ｎｍ以下、アスペクト比が２以下の析出物を指す。The area ratio of the γ ′ phase of the alloy plate according to this embodiment, which is an alloy plate before high-temperature long-term 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. Fine spherical precipitates that are uniformly precipitated in the crystal grains are used as γ ′ phase, and are precipitated in the grain boundaries and in the crystal grains, which 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 (as compared to the Fe and Cr contents) is classified as a precipitate other than the γ ′ phase. From the photograph taken with the FE-SEM, the area ratio of the γ ′ phase is calculated by the line segment method. The fine spherical precipitate generally refers to a precipitate having an equivalent circle diameter of 100 nm or less and an aspect ratio of 2 or less.
［オーステナイト相の結晶粒の、アスペクト比の平均値が１．０〜３．０］
本実施形態に係る合金板は、オーステナイト相の結晶粒の、長軸の長さをＬ１、短軸の長さをＬ２とするとき、アスペクト比であるＬ１／Ｌ２の平均値が１．０〜３．０である。
結晶粒界は、時効処理時にＬａｖｅｓ相などの高温強度の向上に寄与しない非強化相の析出サイトとなる。アスペクト比は圧延によって結晶粒が伸長することにより大きくなる。圧延によって転位が多量に導入されると、その転位を核生成サイトとして非強化相の析出が促進され、強化相であるγ´相の析出が阻害される。また、結晶粒１つあたりの体積と結晶粒界の面積を考えると、アスペクト比が大きくなるにしたがって、結晶粒の体積が同一であるのに対して結晶粒界の面積のみが大きくなり、非強化相の析出サイトの割合が高くなる。したがって、非強化相の析出を抑制し、強化相であるγ´相を確保するためには、アスペクト比は小さいことが好ましい。
そのため、本実施形態に係る合金板では、非強化相の析出サイトを少なくした組織であることを示す指標の１つとして、アスペクト比の平均値を１．０〜３．０とする。アスペクト比については、１．０〜２．０であることが好ましい。[The average aspect ratio of the austenite crystal grains is 1.0 to 3.0]
The alloy plate according to the present embodiment has an aspect ratio of L1 / L2 of 1.0 to 1.0 when the major axis length of the austenite phase crystal grains is L1 and the minor axis length is L2. 3.0.
The crystal grain boundary becomes a precipitation site for a non-strengthened phase that does 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 are elongated by rolling. When a large amount of dislocations is introduced by rolling, precipitation of the non-strengthened phase is promoted using the dislocations as nucleation sites, and precipitation of the γ ′ phase, which is a strengthening phase, is inhibited. Further, 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 while the area of the crystal grain boundary increases. The proportion of the strengthening phase precipitation sites increases. Therefore, it is preferable that the aspect ratio is small in order to suppress the precipitation of the non-strengthening phase and secure the γ ′ phase that is the strengthening phase.
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 index indicating that the structure has a reduced number of precipitation sites of the non-strengthened phase. The aspect ratio is preferably 1.0 to 2.0.
アスペクト比は以下の方法で求めることができる。
Ｌ断面（圧延方向に平行で板厚面に平行な断面）が観察面となるようにサンプルを採取し、表面研磨及びエッチングする。このサンプルを、光学顕微鏡にて、観察面積が４０００００μｍ２以上（８００００μｍ２以上の領域を５か所以上）となるように金属組織を観察する。観察領域にて測定される金属組織を用い、観察された各結晶粒の長軸の長さＬ１と短軸の長さＬ２とからアスペクト比（Ｌ１／Ｌ２）を求め、これらを相加平均して算出する。測定の際、双晶界面は除くものとする。本測定は画像解析ソフトを使用してもよい。
結晶粒のＬ１、Ｌ２は、結晶粒を楕円とみなして、その楕円の長軸、短軸をそれぞれＬ１、Ｌ２と定義する。圧延材の場合、圧延方向の粒径が最も長く、板厚方向の粒径が最も短くなるので、圧延方向の結晶粒の長さをＬ１、板厚方向の結晶粒の長さをＬ２としてもよい。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 is 400000 μm 2 or more (5 or more regions of 80000 μm 2 or more). Using the metallographic structure measured in the observation region, the aspect ratio (L1 / L2) is obtained from the major axis length L1 and minor axis length L2 of each crystal grain, and these are arithmetically averaged. To calculate. In the measurement, the twin interface is excluded. This measurement may use image analysis software.
Regarding the crystal grains L1 and L2, the crystal grains are regarded as ellipses, and the major and minor axes 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, the length of crystal grains in the rolling direction is L1, and the length of crystal grains in the plate thickness direction is L2. Good.
＜室温での０．２％耐力が４００ＭＰａ以上＞
本実施形態に係る合金板では、室温における０．２％耐力が４００ＭＰａ以上である。室温における０．２％耐力が４００ＭＰａ以上であれば、合金板を加工してガスケットとして使用する際、高温長時間時効前（ガスケットとしての使用開始時）の強度を確保できる。
一方、室温での０．２％耐力が１２００ＭＰａ超であると、ガスケットへの加工の際の加工性が低下する。そのため、室温での０．２％耐力は、１２００ＭＰａ以下であることが好ましい。<0.2% proof stress at room temperature is 400 MPa or more>
In the alloy plate according to the present embodiment, the 0.2% yield strength at room temperature is 400 MPa or more. When the 0.2% proof stress at room temperature is 400 MPa or more, when the alloy plate is processed and used as a gasket, it is possible to ensure the strength before high-temperature and long-term aging (when starting use as a gasket).
On the other hand, if the 0.2% proof stress at room temperature is more than 1200 MPa, the workability when processing into a gasket is lowered. Therefore, the 0.2% proof stress at room temperature is preferably 1200 MPa or less.
室温における０．２％耐力は、試験片の長手方向が圧延方向と平行となるように採取したＪＩＳＺ２２４１：２０１１に規定される１３Ｂ号引張試験を用いて、標点間距離の変化量が３ｍｍ/ｍｉｎとなるよう引張荷重を付与する引張試験を行うことで得られる。 The 0.2% proof stress at room temperature is such that the amount of change in the distance between the gauge points is 3 mm / mm using the No. 13B tensile test defined in JISZ2241: 2011 collected so that the longitudinal direction of the test piece is parallel to the rolling direction. It can be obtained by performing a tensile test to give a tensile load so as to be min.
以上のように、本実施形態に係る合金板では、成分をコントロールしてγ´相の析出を促すとともに、溶体化処理、圧延の制御によって結晶粒界及び転位を少なくし、非強化相の析出サイト低減を図っている。
本実施形態に係る合金板は、γ´相により高温強度を確保するが、γ´相は冷間圧延を終了した合金板の段階ではほとんど析出しておらず、合金板に時効処理を施すことによって析出する。高温時効処理が長時間に及んだとき、従来の合金板であれば非強化相の析出によってγ´相の量が十分確保できないのに対し、本実施形態に係る合金板では高温長時間時効後においてもγ´相の量が確保されるので、高温長時間時効後にも十分な高温強度（例えば３８０ＭＰａ以上）が確保できる。As described above, in the alloy sheet according to the present embodiment, the component is controlled to promote the precipitation of the γ ′ phase, and the grain boundary and dislocation are reduced by the solution treatment and the control of the rolling, so that the precipitation of the non-strengthened phase is performed. The site is being reduced.
The alloy plate according to the present embodiment secures high-temperature strength by the γ ′ phase, but the γ ′ phase hardly precipitates at the stage of the alloy plate after the cold rolling, and the alloy plate is subjected to an aging treatment. To precipitate. When the high temperature aging treatment is performed for a long time, the conventional alloy plate cannot secure a sufficient amount of γ ′ phase due to the precipitation of the non-strengthened phase, whereas the alloy plate according to this embodiment has a high temperature long time aging. Since the amount of the γ ′ phase is ensured later, sufficient high-temperature strength (for example, 380 MPa or more) can be ensured even after high-temperature and long-term aging.
＜７５０℃で１００時間時効後の、ミクロ組織中のγ´相の面積率が１５．０％以上、γ´相の平均円相当径が３０ｎｍ以上、７０ｎｍ以下、γ´相以外の析出相の面積率が５．０％以下＞
ガスケットとしての使用を模擬した場合、高温長時間時効処理として７５０℃、１００時間の熱処理条件を行った後でも、高温強度が維持されていることが好ましい。熱処理後の冷却は、５００℃以下までの平均冷却速度を１０℃／秒以上とすることが望ましい。この熱処理条件は、自動車の耐熱部品の熱履歴よりも過酷であり、この条件で高温強度が確保されれば、ガスケットとして長時間使用した場合にも十分に高温強度が確保できる。本実施形態に係る合金板の時効後の特徴については、使用を模擬した高温長時間時効処理後のミクロ組織を観察することで評価ができる。
７５０℃で１００時間の熱処理条件を行った後に高温強度が維持されるためには、７５０℃で１００時間時効処理を施した後のミクロ組織において、γ´相の面積率が１５．０％以上、γ´相の平均円相当径が３０ｎｍ以上、７０ｎｍ以下、γ´相以外の析出相の面積率が５．０％以下であることが好ましい。<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 and 70 nm or less, and the precipitated phase other than the γ ′ phase Area ratio is 5.0% or less>
When simulating the use as a gasket, it is preferable that high-temperature strength is maintained even after heat treatment conditions of 750 ° C. and 100 hours as a high-temperature long-term aging treatment. As for the cooling after the heat treatment, the average cooling rate up to 500 ° C. or lower is desirably 10 ° C./second or higher. This heat treatment condition is more severe than the heat history of heat-resistant parts of automobiles. If high-temperature strength is ensured under these conditions, sufficient high-temperature strength can be ensured even when used as a gasket for a long time. The characteristics after aging of the alloy sheet according to the present embodiment can be evaluated by observing the microstructure after the high-temperature and long-term aging treatment simulating use.
In order to maintain high temperature strength after heat treatment conditions at 750 ° C. for 100 hours, the area ratio of the γ ′ phase is 15.0% or more in the microstructure after aging treatment at 750 ° C. for 100 hours. The average equivalent circle diameter of the γ ′ phase is preferably 30 nm or more and 70 nm or less, and the area ratio of the precipitated phase other than the γ ′ phase is preferably 5.0% or less.
高温強度の観点から、上記高温長時間時効処理後のγ´相の量を１５．０％以上とする。
また、高温強度はγ´相の量（面積率）のみでなく、そのサイズにも影響を受ける。γ´相が小さすぎると、転位がγ´相を簡単に切断するため、γ´相の強化能は低くなる。Ｎｉ，Ｎｂ，Ａｌ，Ｔｉの含有量が少ない場合等、γ´相の析出に時間がかかる場合に、高温長時間時効処理後もγ´相が小さくなる。一方、γ´相が大きすぎることは、γ´相の数が少ないことと同義であるので、合金板の高温強度は低くなる。したがって、高温長時間時効処理後のγ´相は、平均円相当径で３０ｎｍ以上、７０ｎｍ以下とする。
また、結晶粒界や転位上にはγ´相以外の非強化相が析出する場合がある。Ｌａｖｅｓ相に代表されるこれらの非強化相は、高温強度に寄与しないばかりか、その周囲にγ´相が欠落した無析出帯を形成し、高温強度を劣化させる場合がある。したがって、γ´相以外の析出物の面積率は５．０％以下とする。From the viewpoint of high temperature strength, the amount of the γ ′ phase after the high temperature long term aging treatment is set to 15.0% or more.
The high temperature strength is affected not only by the amount (area ratio) of the γ ′ phase but also by its size. If the γ ′ phase is too small, the dislocation easily cuts the γ ′ phase, so that the strengthening ability of the γ ′ phase is lowered. When it takes time to precipitate the γ ′ phase, such as when the content of Ni, Nb, Al, Ti is small, the γ ′ phase becomes small even after the high-temperature long-term aging treatment. On the other hand, since the γ ′ phase is too large is synonymous with the small number of γ ′ phases, the high-temperature strength of the alloy sheet is lowered. Therefore, the γ ′ phase after the high-temperature long-term 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 precipitate on the grain boundaries and dislocations. These non-strengthened phases typified by the Laves phase may not only contribute to the high temperature strength, but may form a precipitate-free zone lacking the γ ′ phase around it, thereby deteriorating the high temperature strength. Therefore, the area ratio of precipitates other than the γ ′ phase is 5.0% or less.
７５０℃で１００時間時効後のγ´相の面積率、γ´相の平均円相当径、γ´相以外の析出相の面積率は以下の方法で求める。
７５０℃で１００時間時効後の合金板の、圧延方向に平行な断面を研磨し、研磨面を王水でエッチングし、ＦＥ−ＳＥＭと同設備に付帯されるＥＤＳで観察、分析する。結晶粒内に均一に析出している微細な球状の析出物をγ´相とし、一方で、結晶粒界、結晶粒内に析出し、粗大、且つ合金板の化学組成よりもＦｅ、Ｃｒが多量に含まれている析出物をγ´相以外の析出物として分類する。ＦＥ−ＳＥＭで撮影した写真から、線分法でγ´相の面積率、γ´相以外の析出物の面積率を算出する。微細な球状の析出物とは、概ね、円相当直径で１００ｎｍ以下、アスペクト比が２．０以下の析出物を指す。
また、１つのサンプルに対して、５０個以上のγ´相について、画像解析ソフトを用いてその面積を求め、その面積と同じ面積を持つ円の直径を算出し、その平均を円相当径（円相当直径）として見積もる。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 phase other than the γ ′ phase are determined by the following methods.
A 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 with EDS attached to the same equipment as FE-SEM. Fine spherical precipitates uniformly precipitated in the crystal grains are used as the γ ′ phase. On the other hand, Fe and Cr are precipitated in the grain boundaries and in the crystal grains, and are coarser than the chemical composition of the alloy plate. A precipitate contained in a large amount is classified as a precipitate other than the γ ′ phase. From the photograph taken with FE-SEM, the area ratio of the γ ′ phase and the area ratio of precipitates other than the γ ′ phase are calculated by the line segment method. The fine spherical precipitate generally refers to a precipitate having an equivalent circle diameter of 100 nm or less and an aspect ratio of 2.0 or less.
In addition, for one sample, the area of 50 or more γ ′ phases is obtained using image analysis software, the diameter of a circle having the same area as the area is calculated, and the average is the equivalent circle diameter ( Estimated as equivalent circle diameter).
上記本実施形態に係る合金板を加工することによってガスケットを形成することができる。例えば本実施形態に係る合金板から、打ち抜きや切り出しなどの手段によってガスケットの外形を形成した上で、ビードと呼ばれる段差部を形成してガスケットとする。このようにして得られたガスケットは、７５０℃で１００時間の高温長時間時効処理後も、高い高温強度が維持できる。 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 a step portion called a bead is formed to form 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.
＜製造方法＞
次に、本実施形態に係る合金板の製造方法について説明する。
本実施形態に係る合金板は、以下の方法によって製造することができる。
（ｉ）上述した化学組成を有し、圧延加工された合金素材を、１０５０〜１２００℃、３０秒〜６００秒で溶体化熱処理する第１の工程と、板厚減少率５％以上、４５％以下で仕上げ冷延する第２の工程とを含む方法。
または、
（ｉｉ）上述した化学組成を有し、圧延加工された合金素材を、１０５０〜１２００℃、３０秒〜６００秒で溶体化熱処理する第３の工程と、７００〜８５０℃、３０秒〜６００秒で時効処理する第４の工程と、第３の工程の後かつ第４の工程の前に、任意に板厚減少率４５％以下で仕上げ冷延する第５の工程とを含む方法。<Manufacturing method>
Next, the manufacturing method of the alloy plate which concerns on this embodiment is demonstrated.
The alloy plate according to the present embodiment can be manufactured by the following method.
(I) a first step of solution heat-treating a rolled alloy material having the above-described chemical composition at 1050 to 1200 ° C. for 30 seconds to 600 seconds, and a thickness reduction rate of 5% or more and 45% And a second step of finish cold rolling.
Or
(Ii) a third step of subjecting the rolled alloy material having the above-described chemical composition to solution heat treatment at 1050 to 1200 ° C. for 30 seconds to 600 seconds, and 700 to 850 ° C. for 30 seconds to 600 seconds. A fourth step of aging treatment in step 5 and a fifth step of optionally finishing and cold rolling at a sheet thickness reduction rate of 45% or less after the third step and before the fourth step.
以下、各工程について説明する。
＜製造方法（ｉ）について＞
［圧延加工された合金素材を、１０５０℃〜１２００℃、３０秒から６００秒で溶体化熱処理する第１の工程］
本実施形態に係る合金板を得るための素材は多くの合金元素を含み、製造過程で様々な化合物が析出する。そのため、本実施形態に係る合金板のようなミクロ組織を得るためには、一度析出物を母相に溶け込ませる溶体化熱処理が必要となる。また、この溶体化処理は、析出物を溶け込ませるという役割に加え、平均結晶粒径を大きくすることも考慮した条件にする必要がある。そのため、温度を１０５０℃以上で、かつ３０秒以上の溶体化処理とする。これにより、耐熱合金薄板の平均結晶粒径を２０μｍ以上とすることができる。
一方、溶体化処理温度が高すぎたり、処理時間が長すぎたりするとコストの増加を招く。したがって、溶体化熱処理温度は１２００℃以下、時間は６００秒以下とする。溶体化処理後は、５００℃以下まで１０℃/秒以上の平均冷却速度で冷却することが好ましい。Hereinafter, each step will be described.
<About manufacturing method (i)>
[First step of solution heat-treating the rolled alloy material at 1050 ° C. to 1200 ° C. for 30 to 600 seconds]
The raw 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 like the alloy plate according to the present embodiment, solution heat treatment for once dissolving the precipitate into the matrix phase is necessary. In addition to the role of dissolving the precipitate, this solution treatment needs to be performed under conditions that also consider increasing the average crystal grain size. Therefore, the solution treatment is performed at a temperature of 1050 ° C. or higher and 30 seconds or longer. Thereby, the average crystal grain size of the heat-resistant alloy thin plate can be 20 μm or more.
On the other hand, when the solution treatment temperature is too high or the treatment time is too long, the cost increases. Accordingly, the solution heat treatment temperature is 1200 ° C. or less and the time is 600 seconds or less. After the solution treatment, it is preferable to cool to 500 ° C. or less at an average cooling rate of 10 ° C./second or more.
［板厚減少率５％以上、４５％以下で仕上げ冷延する第２の工程］
本実施形態に係る合金板は、板厚調整や強度調整のために、仕上げ冷延を施す。結晶粒内の非強化相は、転位を核生成サイトとして析出するので、仕上げ冷延後の時効処理によって結晶粒内にγ´相以外の非強化相が析出し、本来の特性を損なう場合がある。そのため、仕上げ冷延での板厚減少率（圧下率）を小さくすることにより、仕上げ冷延後の結晶粒内における転位の量を低減し、高温長時間時効における非強化相の析出を抑制する。板厚減少率の低減は、前述のように、結晶粒のアスペクト比を小さくし、非強化相の析出サイトとなる結晶粒界の面積を小さくする効果もある。
アスペクト比を１．０〜３．０、γ´相の面積率を１５．０％以下とし、７５０℃、１００時間時効後にγ´相以外の析出相を５．０％以下まで抑制し、γ´相の面積率を１５．０％以上確保できるような組織とする観点から、仕上げ冷延での板厚減少率を４５％以下とする。望ましくは、４０％以下とする。仕上げ冷延の工程内で冷延を分割して中間焼鈍を行うことはしない。また、仕上げ冷延率が高いと、転位密度が増加し、原子の拡散が促進されるため、γ´相の粒成長が促進される。γ´相の粗大化は高温強度の低下につながるので、この観点からも仕上げ冷延率は４５％以下、望ましくは、４０％以下とする。[Second step of finish cold rolling at a thickness reduction rate of 5% to 45%]
The alloy plate according to the present embodiment is subjected to finish cold rolling for thickness adjustment and strength adjustment. Since the non-strengthened phase in the crystal grains precipitates as dislocations as nucleation sites, the non-strengthened phase other than the γ 'phase may precipitate in the crystal grains due to the aging treatment after finish cold rolling, which may impair the original characteristics. is there. Therefore, by reducing the sheet thickness reduction rate (rolling 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-term aging is suppressed. . As described above, the reduction in the plate thickness reduction rate also has the effect of reducing the crystal grain aspect ratio and reducing the area of the crystal grain boundary, which is the precipitation site of 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 after aging at 750 ° C. for 100 hours, the precipitated phase other than the γ ′ phase is suppressed to 5.0% or less. The sheet thickness reduction rate in the finish cold rolling is set to 45% or less from the viewpoint of obtaining a structure that can secure the area ratio of the phase '15 .0% or more. Desirably, it is 40% or less. Intermediate annealing is not performed by dividing the cold rolling in the process of finish cold rolling. Further, when the finish cold rolling rate is high, the dislocation density increases and the diffusion of atoms is promoted, so that the γ ′ phase grain growth is promoted. Since coarsening of the γ 'phase leads to a decrease in high-temperature strength, the finish cold rolling rate is set to 45% or less, preferably 40% or less from this viewpoint.
一方、本実施形態に係る合金板は、高温環境に保持されることによってγ´相が析出し、これによって室温強度、高温強度が高く維持される。しかしながら、ガスケットとして使用される場合の使用開始直後においては、金属板にγ´相を析出させない場合、必要な強度をγ´相を析出させる以外の手段で確保する必要がある。
仕上げ冷延の加工硬化を活用することで、高温時効前の室温での強度（０．２％耐力）を確保できる。仕上げ冷延率が低いと高温時効後の高温強度には有利であるものの、高温時効前の室温での強度が不足する場合がある。そのため、時効処理を行わない場合、仕上げ圧延の板厚減少率は５％以上とする。仕上げ圧延の板厚減少率は１０％以上とすることが好ましい。On the other hand, in the alloy plate according to the present embodiment, the γ ′ phase is precipitated by being held in a high temperature environment, whereby the room temperature strength and the high temperature strength are maintained high. However, immediately after the start of use when used as a gasket, if the γ ′ phase is not precipitated on the metal plate, it is necessary to ensure the necessary strength by means other than the precipitation of the γ ′ phase.
By utilizing the work hardening of finish cold rolling, the strength at room temperature (0.2% proof stress) before high temperature aging can be secured. If the finish cold rolling rate 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 thickness reduction rate of finish rolling is set to 5% or more. The plate thickness reduction rate in finish rolling is preferably 10% or more.
＜製造方法（ｉｉ）について＞
［圧延加工された合金素材を１０５０〜１２００℃、３０秒〜６００秒で溶体化熱処理する第３の工程］
本実施形態に係る合金板を得るための素材は多くの合金元素を含み、製造過程で様々な化合物が析出する。そのため、上記のような７５０℃、１００時間時効後に適切なミクロ組織を得るためには、一度析出物を母相に溶け込ませる溶体化熱処理が必要となる。この溶体化処理は、析出物を溶け込ませるという役割に加え、平均結晶粒径を大きくすることも考慮した条件にする必要がある。そのため、温度を１０５０℃以上で、かつ３０秒以上の溶体化処理とする。これにより、耐熱合金薄板の平均結晶粒径を２０μｍ以上とすることができる。一方、溶体化処理温度が高すぎたり、処理時間が長すぎたりするとコストの増加を招く。したがって、溶体化熱処理温度は１２００℃以下、時間は６００秒以下とする。溶体化処理後は、５００℃以下まで１０℃/秒以上の平均冷却速度で冷却することが好ましい。<About production method (ii)>
[Third step of solution heat-treating the rolled alloy material at 1050 to 1200 ° C. for 30 to 600 seconds]
The raw 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, a solution heat treatment for once dissolving the precipitate in the matrix phase is required. In this solution treatment, in addition to the role of dissolving the precipitate, it is necessary to set the conditions in consideration of increasing the average crystal grain size. Therefore, the solution treatment is performed at a temperature of 1050 ° C. or higher and 30 seconds or longer. Thereby, the average crystal grain size of the heat-resistant alloy thin plate can be 20 μm or more. On the other hand, when the solution treatment temperature is too high or the treatment time is too long, the cost increases. Accordingly, the solution heat treatment temperature is 1200 ° C. or less and the time is 600 seconds or less. After the solution treatment, it is preferable to cool to 500 ° C. or less at an average cooling rate of 10 ° C./second or more.
［７００〜８５０℃、３０秒〜６００秒で時効処理する第４の工程］
仕上げ冷延の加工硬化を活用することで、高温長時間時効前の室温での強度（０．２％耐力）を確保できる。一方で、仕上げ冷延を行うことで、転位が導入され、非強化相の析出サイトが増加する。そのため、転位の導入をより抑制する場合には、仕上げ冷延を必須としない製造方法とすることが好ましい。
仕上げ冷延を行わない場合には、室温での０．２％耐力が確保できないおそれがあるが、仕上げ冷延を行わない場合でも、短時間の時効処理を行い、少量のγ´相を析出させることで、室温での０．２％耐力を確保できる。
そのため、仕上げ冷延を必須としない本実施形態に係る合金板の製造方法では、第３の工程の後に、７００〜８５０℃で、３０秒〜６００秒保持する時効処理を行う。
時効温度が７００℃未満である、または時間が３０秒未満であると、十分に強度が確保できない。
一方、時効温度が８５０℃超であると、γ´相が析出しない場合がある。また、時効時間が６００秒超である場合には、γ′相が粗大化して十分な強度が得られない場合がある。[Fourth step of aging treatment at 700 to 850 ° C. for 30 to 600 seconds]
By utilizing the work hardening of finish cold rolling, it is possible to ensure the strength at room temperature (0.2% proof stress) before aging at a high temperature for a long time. 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 finishing cold rolling.
If finish cold rolling is not performed, 0.2% proof stress at room temperature may not be ensured, but even if finish cold rolling is not performed, short-term 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 manufacturing method of the alloy plate which concerns on this embodiment which does not require finishing cold rolling, the aging treatment which hold | maintains at 700-850 degreeC for 30 second-600 second is performed after a 3rd process.
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 precipitate. In addition, when the aging time exceeds 600 seconds, the γ ′ phase may become coarse and sufficient strength may not be obtained.
［板厚減少率４５％以下で仕上げ冷延する第５の工程］
第４の工程を行うことで、室温の０．２％耐力が４００ＭＰａ以上を確保できるので、仕上げ冷延を行う必要はないが、さらに室温の０．２％耐力を高めたい場合には、第３の工程と第４の工程との間に、板厚減少率４５％以下で仕上げ冷延を行ってもよい。仕上げ冷延率が高いと、転位密度が増加し、原子の拡散が促進されるので、γ´相の粒成長が促進される。γ´相の粗大化は高温強度の低下につながるので、仕上げ冷延率は４５％以下、望ましくは、４０％以下とする。[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 ensure 400 MPa or more, so there is no need to perform finish cold rolling, but if you want to further increase the 0.2% proof stress at room temperature, Finishing and cold rolling may be performed between the step 3 and the fourth step 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 γ ′ phase grain growth is promoted. Since coarsening of the γ ′ phase leads to a decrease in high-temperature strength, the finish cold rolling rate is set to 45% or less, preferably 40% or less.
上記第１の工程、または第３の工程に至るまでの製造工程については、通常の合金板の製造方法を用いることができる。即ち、製錬工程で所定の成分に調整された合金を、連続鋳造をはじめとする鋳造工程で鋳造して鋳片とし、鋳片を熱延前加熱炉で加熱した上で熱間圧延を行う。熱間圧延の前には熱間鍛造をしてもよい。また、熱間圧延後の合金板については、熱延後焼鈍を行っても行わなくてもよい。このようにして製造した熱延板について、冷間圧延（中間冷延）を行う。中間冷延を複数回に分け、各中間冷延の間に中間焼鈍を行ってもよい。このようにして製造した中間冷延板を対象として、上記第１の工程、または第３の工程の溶体化熱処理を実施する。 For the manufacturing steps up to the first step or the third step, a normal method for manufacturing an alloy plate can be used. That is, an alloy adjusted to a predetermined component in the smelting process is cast into a slab by a casting process including continuous casting, and the slab is heated in a heating furnace before hot rolling and hot rolled. . Hot forging may be performed before hot rolling. Moreover, about the alloy plate after hot rolling, it is not necessary to perform annealing after hot rolling. The hot-rolled sheet manufactured in this way is subjected to cold rolling (intermediate cold rolling). The 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 in the first step or the third step is performed on the intermediate cold-rolled sheet thus manufactured.
以下、実施例を参照しながら、本発明を具体的に説明する。
＜実施例１＞Hereinafter, the present invention will be specifically described with reference to examples.
<Example 1>
まず、表１の化学組成を有する２５ｋｇの鋳塊を溶製した。ここで、ＡからＰは本発明の規定を満たす成分、ＱからＷは本発明を満足しない成分である。本発明範囲から外れる数値に下線を付している。
これら鋳塊を熱間鍛造により厚さ４０ｍｍにした後、熱間圧延を行って板厚５ｍｍの熱延板を得た。この熱延板に、焼鈍、冷延を施すことで種々の板厚の中間冷延板にした。この中間冷延板に表２に示す条件で溶体化処理、仕上げ冷延（冷間圧延）を施した。中間冷延板の板厚を変えることで、仕上げ冷延板の板厚はいずれも０．２０ｍｍとした。First, a 25 kg ingot having the chemical composition shown in Table 1 was melted. Here, A to P are components that satisfy the definition of the present invention, and Q to W are components that do not satisfy the present invention. Numerical values that fall outside the scope of the present invention are underlined.
These ingots were made into a thickness of 40 mm by hot forging, and then hot rolled to obtain a hot rolled sheet having a thickness of 5 mm. The hot-rolled sheet was annealed and cold-rolled to obtain intermediate cold-rolled sheets having various thicknesses. This intermediate cold-rolled sheet was subjected to solution treatment and finish cold-rolling (cold rolling) under the conditions shown in Table 2. By changing the thickness of the intermediate cold-rolled plate, the thickness of the finished cold-rolled plate was set to 0.20 mm.
この仕上げ冷延板のＬ断面が観察面となるように金属板からサンプルを採取し、観察面を表面研磨及び王水でエッチングし、光学顕微鏡で観察、撮影し、求積法によりオーステナイト相の平均結晶粒径を算出した。
また、同サンプルに対し、光学顕微鏡にて、観察面積が４０００００μｍ２以上となるように金属組織を観察した。観察された各結晶粒の長軸の長さＬ１と短軸の長さＬ２とからアスペクト比（Ｌ１／Ｌ２）を求め、これらを平均してアスペクト比を算出した。A sample is taken from the metal plate so that the L cross section of the finished cold-rolled sheet becomes the observation surface, the observation surface is etched with surface polishing and aqua regia, observed and photographed with an optical microscope, and the austenite phase is obtained by a quadrature method. The average crystal grain size was calculated.
Further, the metal structure of the sample was observed with an optical microscope so that the observation area was 400000 μm 2 or more. The aspect ratio (L1 / L2) was calculated from the observed major axis length L1 and minor axis length L2 of each crystal grain, and these were averaged to calculate the aspect ratio.
また、高温長時間時効処理前の合金板において、以下の方法で、室温の０．２％耐力を求めた。
室温における０．２％耐力は、試験片の長手方向が圧延方向と平行になるように採取したＪＩＳＺ２２４１：２０１１に規定される１３Ｂ号引張試験を用いて、標点間距離の変化量が３ｍｍ/ｍｉｎとなるように引張荷重を付与する条件で引張試験を行うことで得た。Moreover, the 0.2% yield strength of room temperature was calculated | required with the following method in the alloy plate before high temperature long-term aging treatment.
The 0.2% proof stress at room temperature is such that the amount of change in the distance between the gauge points is 3 mm / mm using the No. 13B tensile test defined in JISZ2241: 2011 taken so that the longitudinal direction of the test piece is parallel to the rolling direction. It was obtained by conducting a tensile test under the condition of applying a tensile load so as to be min.
また、高温長時間時効処理前の合金板において、以下の方法でγ´相の面積率を求めた。
高温長時間時効処理前の合金板の圧延方向に平行断面を研磨し、王水でエッチングし、ＦＥ−ＳＥＭと同設備に付帯されるＥＤＳで観察、分析した。結晶粒内に均一に析出している円相当直径で１００ｎｍ以下、アスペクト比が２以下の析出物をγ´相とした。また、結晶粒界、結晶粒内に析出し、粗大、且つ合金板の化学組成よりもＦｅ、Ｃｒが多量に含まれている析出物をγ´相以外の析出物とした。ＦＥ−ＳＥＭで撮影した写真から、線分法でγ´相の面積率を算出した。Further, 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 plate before high-temperature long-term aging treatment was polished, etched with aqua regia, and observed and analyzed with EDS attached to the same equipment as FE-SEM. A precipitate having a circle-equivalent diameter uniformly precipitated in the crystal grains and having an aspect ratio of 2 or less was defined as a γ ′ phase. In addition, a precipitate other than the γ ′ phase was defined as a precipitate that precipitated in the crystal grain boundaries and in the crystal grains and was coarse and contained a larger amount of Fe and Cr than the chemical composition of the alloy plate. From the photograph taken with the FE-SEM, the area ratio of the γ ′ phase was calculated by the line segment method.
さらに、上記の仕上げ冷延板に大気炉で７５０℃、１００時間の時効処理（高温長時間時効処理）を施した。この高温長時間時効処理前の合金板において、以下の方法でγ´相の面積率、γ´相以外の析出物の面積率を求めた。
時効処理後の金属板の圧延方向に垂直断面を研磨し、王水でエッチングし、ＦＥ−ＳＥＭと同設備に付帯されるＥＤＳで観察、分析し、結晶粒内に均一に析出している円相当直径で１００ｎｍ以下、アスペクト比が２以下の析出物をγ´相とし、結晶粒界、結晶粒内に析出し、粗大、且つＦｅ、Ｃｒが多量に（合金板の化学組成におけるＦｅ、Ｃｒ含有量よりも多く）含まれている析出物をγ´相以外の析出物として分類した。ＦＥ−ＳＥＭで撮影した写真から、線分法でγ´相の面積率、γ´相以外の析出物の面積率を算出した。
また、１つのサンプルに対して、５０個以上のγ´相について、画像解析ソフトを用いてその面積を求め、その面積と同じ面積を持つ円の直径を算出し、その平均を円相当径として見積もった。Further, the finished cold-rolled sheet was subjected to aging treatment at 750 ° C. for 100 hours (high-temperature long-time aging treatment) in an atmospheric furnace. In the alloy plate before the high-temperature long-term aging treatment, the area ratio of the γ ′ phase and the area ratio of precipitates other than the γ ′ phase were determined by the following methods.
A circle perpendicular to the rolling direction of the metal plate after aging treatment, polished with aqua regia, observed and analyzed with EDS attached to the same equipment as FE-SEM, and uniformly deposited in the crystal grains A precipitate having an equivalent diameter of 100 nm or less and an aspect ratio of 2 or less is used as a γ ′ phase, and precipitates in the grain boundaries and in the crystal grains, and is coarse and contains a large amount of Fe and Cr (Fe, Cr in the chemical composition of the alloy plate). The contained precipitates were classified as precipitates other than the γ ′ phase. From the photograph taken with FE-SEM, the area ratio of the γ ′ phase and the area ratio of precipitates other than the γ ′ phase were calculated by the line segment method.
Further, for one sample, the area of 50 or more γ ′ phases is obtained using image analysis software, the diameter of a circle having the same area as that area is calculated, and the average is defined as the equivalent circle diameter. Estimated.
さらに、７５０℃、１００時間の時効処理を施した合金板を、７５０℃の高温引張試験に供し、７５０℃における高温０．２％耐力を測定した。７５０℃、１００時間の時効処理後の７５０℃における高温０．２％耐力が３８０ＭＰａ以上であれば、高温長時間時効処理後にも十分な高温強度が確保できていると判断した。 Further, the alloy plate subjected to aging treatment at 750 ° C. for 100 hours was subjected to a high temperature tensile test at 750 ° C., and a high temperature 0.2% proof stress at 750 ° C. was measured. If the high-temperature 0.2% proof stress at 750 ° C. after aging at 750 ° C. for 100 hours was 380 MPa or more, it was judged that sufficient high-temperature strength could be secured even after high-temperature long-term aging treatment.
表２の合金板Ｎｏ．１〜２１が本発明の規定を満足する発明例、Ｎｏ．２２〜３２が本発明の規定を満たさない比較例である。 Alloy plate no. Examples Nos. 1 to 21 satisfy the provisions of the present invention. 22 to 32 are comparative examples that do not satisfy the provisions of the present invention.
発明例１〜２１はいずれも、高温長時間時効後の高温強度（０．２％耐力）が３８０ＭＰａ以上と良好な品質を確保することができた。また、板厚減少率５％以上の仕上げ冷延を行い、そのまま耐熱合金板としているので、時効処理前の室温強度（０．２％耐力）についても、いずれも４００ＭＰａ以上であった。 In each of Invention Examples 1 to 21, high temperature strength (0.2% proof stress) after aging at high temperature for a long time was 380 MPa or more, and good quality could be ensured. Moreover, since the finish cold rolling with a plate thickness reduction rate of 5% or more was performed and the heat-resistant alloy plate was used as it was, the room temperature strength (0.2% proof stress) before the aging treatment was 400 MPa or more.
一方、比較例２２は、合金の組成は本発明の規定を満足するものの、溶体化処理の温度が低く、時間が短かった。その結果、オーステナイト相の平均結晶粒径が小さく、時効後に非強化相が多量に析出し、強化相であるγ´相の量が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例２３、２４は、合金の組成は本発明の規定を満足するものの、仕上げ冷延率が大きかった。その結果、アスペクト比が３．０を超えた。その結果、時効後に非強化相が多量に析出し、高温長時間時効後に強化相であるγ´相が少なかった。また、比較例２４は、仕上げ冷延率が大きく、転位密度が高かったので、原子の拡散が速く、高温長時間時効後にγ´相が粗大化していた。その結果、高温長時間時効後の高温０．２％耐力が低かった。On the other hand, in Comparative Example 22, although the alloy composition satisfied the provisions of the present invention, the solution treatment temperature was low and the time was short. As a result, the average crystal grain size of the austenite phase was small, a large amount of non-strengthened phase was precipitated after aging, and the amount of γ ′ phase as a strengthened phase was small. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Examples 23 and 24, although the alloy composition satisfied the provisions of the present invention, the finish cold rolling rate was large. As a result, the aspect ratio exceeded 3.0. As a result, a large amount of non-strengthened phase precipitated after aging, and there were few γ ′ phases that were strengthened phases after aging at a high temperature for a long time. In Comparative Example 24, the finish cold rolling ratio was large and the dislocation density was high, so that the diffusion of atoms was fast, and the γ ′ phase was coarsened after high-temperature and long-term aging. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
比較例２５は、合金中のＣ含有量が多かった。そのため、製造段階からＮｂ、Ｔｉの炭化物が生成し、高温長時間時効後に強化相であるγ´相の生成量が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例２６は、合金中のＣｒ含有量が多かった。そのため、時効後にσ相などの非強化相が多量に析出し、高温長時間時効後に強化相であるγ´相の生成量が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例２７は、合金中のＮｉ含有量が少なかった。そのため、時効後に強化相であるγ´相の析出が極めて少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例２８は、合金中のＮ含有量が多かった。そのため、製造段階からＡｌ、Ｎｂ、の窒化物が生成し、高温長時間時効後に強化相であるγ´相の生成量が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例２９は、合金中のＶ含有量が多かった。その結果、製造段階から粗大なＶ炭化物、Ｖ窒化物が生成し、高温長時間時効後に強化相であるγ´相の生成量が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例３０は、合金中のＮｂ含有量が少なかった。そのため、高温長時間時効後にγ´相が極めて少なかった。また、γ´相のサイズも小かった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例３１は、合金中のＡｌ含有量が少なかった。そのため、高温長時間時効後にγ´相が少なかった。その結果、高温長時間時効後の高温０．２％耐力が低かった。
比較例３２は、合金の組成は本発明の規定を満足するものの溶体化温度が低すぎてオーステナイト相の平均結晶粒径が小さくなり、製造中に析出したγ´が再固溶せずに時効前の合金板に残存した。その結果、高温長時間時効中にγ´が粗大化すると共に非強化相が析出し、高温０．２％耐力が低かった。In Comparative Example 25, the C content in the alloy was high. Therefore, Nb and Ti carbides were generated from the production stage, and the amount of γ ′ phase, which is a strengthening phase, was small after high temperature aging for a long time. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 26, the Cr content in the alloy was high. Therefore, a large amount of non-strengthening phase such as σ phase was precipitated after aging, and the amount of γ ′ phase, which is a strengthening phase, was small after aging at high temperature for a long time. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 27, the Ni content in the alloy was low. Therefore, the precipitation of the γ ′ phase, which is a strengthening phase after aging, was extremely small. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 28, the N content in the alloy was high. For this reason, nitrides of Al, Nb were produced from the production stage, and the amount of γ ′ phase, which is a strengthening phase, was small after aging at a high temperature for a long time. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 29, the V content in the alloy was high. As a result, coarse V carbides and V nitrides were produced from the production stage, and the amount of γ ′ phase, which is a strengthening phase, was small after high temperature aging for a long time. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 30, the Nb content in the alloy was small. Therefore, there was very little γ 'phase after high-temperature and long-term aging. In addition, the size of the γ ′ phase was small. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 31, the Al content in the alloy was small. Therefore, there was little γ 'phase after aging at high temperature for a long time. As a result, the high temperature 0.2% proof stress after high temperature long time aging was low.
In Comparative Example 32, although the alloy composition satisfies the provisions of the present invention, the solution temperature is too low, the average crystal grain size of the austenite phase becomes small, and γ 'precipitated during production does not re-dissolve and is aged. It remained on the previous alloy plate. As a result, during the high temperature long time aging, γ ′ became coarse and a non-strengthened phase precipitated, and the high temperature 0.2% proof stress was low.
以上のように、比較例２２〜３２はいずれも、合金板の組織または化学組成が本発明を外れていた。また、その結果、高温長時間時効処理後のγ´相の量が少なく、高温長時間時効処理後の高温強度が本発明の目標に達しなかった。 As described above, in all of Comparative Examples 22 to 32, the structure or chemical composition of the alloy plate deviated from the present invention. As a result, the amount of 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.
＜実施例２＞
表１の合金Ｎｏ．Ａの成分を有する鋳塊を熱間鍛造により厚さ４０ｍｍにした後、熱間圧延を行って板厚５ｍｍの熱延板を得た。この熱延板に、焼鈍、冷延を施すことで種々の板厚の中間冷延板にした。この中間冷延板に表３に示す条件で溶体化処理、及び一部の合金板には仕上げ冷延（冷間圧延）を施した。仕上げ冷延板の板厚はいずれも０．２０ｍｍとなるようにした。
さらに一部の合金板については、表３に示す条件で時効処理（短時間時効処理）を行った。<Example 2>
Alloy No. 1 in Table 1 After the ingot having the component A was made 40 mm thick by hot forging, hot rolling was performed to obtain a hot rolled sheet having a thickness of 5 mm. The hot-rolled sheet was annealed and cold-rolled to obtain intermediate cold-rolled sheets having various thicknesses. The intermediate cold-rolled sheet was subjected to a solution treatment under the conditions shown in Table 3, and some of the alloy sheets were subjected to finish cold-rolling (cold rolling). The thickness of each finished cold-rolled sheet was set to 0.20 mm.
Further, some of the alloy plates were subjected to aging treatment (short-time aging treatment) under the conditions shown in Table 3.
得られた合金板について、実施例１と同じ方法で、オーステナイト相の平均結晶粒径、室温の０．２％耐力、γ′相の面積率、アスペクト比（Ｌ１／Ｌ２）を求めた。 For 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.
さらに、上記の仕上げ冷延板または短時間時効処理後の金属板に、大気炉で７５０℃、１００時間の時効処理（高温長時間時効処理）を施した。 Furthermore, the above-described cold-rolled sheet or the metal plate after the short-time aging treatment was subjected to aging treatment at 750 ° C. for 100 hours (high-temperature long-time aging treatment) in an atmospheric furnace.
高温長時間時効処理後の金属板について、実施例１と同じ方法でγ´相の面積率、γ´相以外の析出物の面積率を算出した。また、１つのサンプルに対して、５０個以上のγ´相について、画像解析ソフトを用いてその面積を求め、その面積と同じ面積を持つ円の直径を算出し、その平均を円相当径として見積もった。さらに、７５０℃の高温引張試験に供し、７５０℃における高温０．２％耐力を測定した。 For the metal plate after the high-temperature and long-term aging treatment, 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. Further, for one sample, the area of 50 or more γ ′ phases is obtained using image analysis software, the diameter of a circle having the same area as that area is calculated, and the average is defined as the equivalent circle diameter. Estimated. Furthermore, it used for the high temperature tensile test of 750 degreeC, and measured the high temperature 0.2% yield strength in 750 degreeC.
結果を表３に示す。
発明例である合金板Ｎｏ．Ｘ１〜Ｘ４については、本発明の第３の工程（溶体化熱処理）と、その後の第４の工程（短時間時効処理）を行っており、いずれも、室温０．２％耐力が４００ＭＰａ以上と良好であった。また、高温０．２％耐力も３８０ＭＰａ以上と高かった。
合金板Ｎｏ．Ｘ１は第５の工程（仕上げ冷延）を行わず、合金板Ｎｏ．Ｘ２は第５の工程（仕上げ冷延）における板厚減少率が４％と少ないが、第４の工程による析出硬化で室温強度が確保された。The results are shown in Table 3.
Inventive alloy plate No. For X1 to X4, the third step (solution heat treatment) of the present invention and the subsequent fourth step (short-term aging treatment) are performed, and both have a room temperature 0.2% proof stress of 400 MPa or more. It was good. Moreover, the high temperature 0.2% proof stress was as high as 380 MPa or more.
Alloy plate No. X1 does not perform the fifth step (finish cold rolling), and alloy plate No. Although X2 has a small thickness reduction rate of 4% in the fifth step (finish cold rolling), room temperature strength was ensured by precipitation hardening in the fourth step.
一方、比較例である合金板Ｎｏ．Ｘ５は第５の工程（仕上げ冷延）と第４の工程（短時間時効処理）とをともに行わなかった。また、合金板Ｎｏ．Ｘ６は第５の工程（仕上げ冷延）の板厚減少率が３％であり、かつ第４の工程（短時間時効処理）を行わなかった。そのため、いずれも室温０．２％耐力が４００ＭＰａ未満であった。
合金板Ｎｏ．Ｘ７は第５の工程（仕上げ冷延）の板厚減少率が３％であり、かつ第４の工程（短時間時効処理）の温度が低かった。また、合金板Ｎｏ．Ｘ８は第５の工程（仕上げ冷延）の板厚減少率が３％であり、かつ第４の工程（短時間時効処理）の温度が高かった。そのため、いずれも室温０．２％耐力が４００ＭＰａ未満であった。On the other hand, the alloy plate No. X5 did not perform both the fifth step (finish cold rolling) and the fourth step (short-term aging treatment). Alloy plate No. X6 had a thickness reduction rate of 3% in the fifth step (finish cold rolling), and the fourth step (short-term aging treatment) was not performed. Therefore, in all cases, the room temperature 0.2% proof stress was less than 400 MPa.
Alloy plate No. X7 had a thickness reduction rate of 3% in the fifth step (finish cold rolling), and the temperature in the fourth step (short-term aging treatment) was low. Alloy plate No. In X8, the thickness reduction rate in the fifth step (finish cold rolling) was 3%, and the temperature in the fourth step (short-term aging treatment) was high. Therefore, in all cases, the room temperature 0.2% proof stress was less than 400 MPa.
本発明の合金板は、排気系ガスケット等に使用され、高温に長時間曝された場合でも、十分な高温強度が確保される。そのため、本発明の合金板を用いたガスケットでは、高温長時間使用後でも高温強度が確保される。そのため、産業上の利用可能性が高い。 The alloy plate of the present invention is used for an exhaust system gasket or the like, and a sufficient high temperature strength is ensured even when exposed to a high temperature for a long time. Therefore, in the gasket using the alloy plate of the present invention, high-temperature strength is ensured even after high-temperature and long-time use. Therefore, industrial applicability is high.
Claims (6)
Ｃ：０．０２００％以下、
Ｓｉ：０．０２〜２．００％、
Ｍｎ：０．０２〜２．００％、
Ｐ：０．０５０％以下、
Ｓ：０．０１００％以下、
Ｃｒ：１２．００〜３０．００％、
Ｎｉ：３５．００〜６０．００％、
Ｎ：０．０２００％以下、
Ｖ：０．０２〜１．００％、
Ｎｂ：２．００％超、３．５０％以下、
Ａｌ：２．６０％超、４．００％以下、
Ｔｉ：０〜０．８０％、
Ｍｏ：０〜２．００％、
Ｗ：０〜２．００％、
Ｃｕ：０〜１．００％、
Ｃｏ：０〜１．００％、
Ｂ：０〜０．０１００％、
Ｚｒ：０〜０．０１００％、
Ｃａ：０〜０．００５０％、
Ｍｇ：０〜０．００５０％、
を含有し、残部が１０．００％以上のＦｅ、及び不純物である化学組成を有し、
ミクロ組織が、オーステナイト相を含み、前記オーステナイト相の平均結晶粒径が２０μｍ以上であり、
前記オーステナイト相の結晶粒の、長軸の長さをＬ１、短軸の長さをＬ２とするとき、アスペクト比であるＬ１／Ｌ２の平均値が、１．０〜３．０であり、
前記ミクロ組織において、γ´相の面積率が１５．０％以下であり、
室温での０．２％耐力が４００ＭＰａ以上である
ことを特徴とする合金板。% By mass
C: 0.0200% or less,
Si: 0.02-2.00%,
Mn: 0.02 to 2.00%
P: 0.050% or less,
S: 0.0100% or less,
Cr: 12.00 to 30.00%,
Ni: 35.00-60.00%,
N: 0.0200% or less,
V: 0.02 to 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,
The microstructure includes an austenite phase, and the average crystal grain size of the austenite phase is 20 μm or more,
When the major axis length of the austenite phase crystal grains is L1, and the minor axis length 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% yield strength at room temperature of 400 MPa or more.
ことを特徴とする請求項１に記載の合金板。The alloy plate according to claim 1, wherein the aspect ratio is 1.0 to 2.0.
前記時効処理後のミクロ組織におけるγ´相の面積率が１５．０％以上、
前記γ´相の平均円相当径が３０ｎｍ以上、７０ｎｍ以下、かつ
γ´相以外の析出相の面積率が５．０％以下である、
ことを特徴とする請求項１または２に記載の合金板。When aging treatment was performed 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 equivalent circle diameter of the γ ′ phase is 30 nm or more and 70 nm or less, and the area ratio of the precipitated phase other than the γ ′ phase is 5.0% or less.
The alloy plate according to claim 1 or 2, characterized by the above.
圧延加工された合金素材を１０５０〜１２００℃で、３０秒〜６００秒間溶体化熱処理する第１の工程と、
板厚減少率５％以上、４５％以下で仕上げ冷延する第２の工程と、
を含む
ことを特徴とする合金板の製造方法。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 thickness reduction rate of 5% or more and 45% or less;
The manufacturing method of the alloy plate characterized by including.
圧延加工された合金素材を１０５０〜１２００℃で、３０秒〜６００秒間溶体化熱処理する第３の工程と、
７００〜８５０℃、３０秒〜６００秒で時効処理する第４の工程と、
を含む
ことを特徴とする合金板の製造方法。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;
The manufacturing method of the alloy plate characterized by including.
ことを特徴とする請求項５に記載の合金板の製造方法。6. The alloy according to claim 5, further comprising a fifth step of finishing and cold rolling after the third step and before the fourth step at a reduction ratio of a sheet thickness reduction rate of 45% or less. A manufacturing method of a board.
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Citations (5)
| Publication number | Priority date | Publication date | Assignee | Title |
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| JP2001073053A (en) * | 1999-06-30 | 2001-03-21 | Sumitomo Metal Ind Ltd | Ni BASE HEAT RESISTANT ALLOY |
| JP2011127204A (en) * | 2009-12-21 | 2011-06-30 | Hitachi Metals Ltd | Hydrogen embrittlement-resistant high strength austenitic alloy |
| WO2013141030A1 (en) * | 2012-03-23 | 2013-09-26 | 株式会社クボタ | Cast product having alumina barrier layer, and method for manufacturing same |
| JP2017053006A (en) * | 2015-09-10 | 2017-03-16 | 新日鐵住金株式会社 | MANUFACTURING METHOD OF Ni-BASED HEAT RESISTANT ALLOY TUBE |
| WO2017119415A1 (en) * | 2016-01-05 | 2017-07-13 | 新日鐵住金株式会社 | Austenitic heat-resistant alloy and method for manufacturing same |
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| Publication number | Priority date | Publication date | Assignee | Title |
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| JP2011127204A (en) * | 2009-12-21 | 2011-06-30 | Hitachi Metals Ltd | Hydrogen embrittlement-resistant high strength austenitic alloy |
| WO2013141030A1 (en) * | 2012-03-23 | 2013-09-26 | 株式会社クボタ | Cast product having alumina barrier layer, and method for manufacturing same |
| JP2017053006A (en) * | 2015-09-10 | 2017-03-16 | 新日鐵住金株式会社 | MANUFACTURING METHOD OF Ni-BASED HEAT RESISTANT ALLOY TUBE |
| WO2017119415A1 (en) * | 2016-01-05 | 2017-07-13 | 新日鐵住金株式会社 | Austenitic heat-resistant alloy and method for manufacturing same |
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