JP2008266688A - Mn-cu damping alloy and producing method therefor - Google Patents
Mn-cu damping alloy and producing method therefor Download PDFInfo
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Abstract
Description
本発明は、Ｍｎ−Ｃｕ系制振合金及びその製造方法に関し、更に詳しくは、加工性及び制振特性に優れたＭｎ−Ｃｕ系制振合金及びその製造方法に関する。 The present invention relates to a Mn—Cu based damping alloy and a method for producing the same, and more particularly to a Mn—Cu based damping alloy having excellent workability and damping characteristics and a method for producing the same.
制振合金は、外部からの振動エネルギーを内部摩擦によって熱に変換することにより振動を吸収する合金であり、その制振機能により、複合型（例えば、片状黒鉛鋳鉄）、強磁性型（例えば、Ｆｅ−Ｃｒ合金）、転位型（例えば、Ｍｇ−Ｚｒ合金）、及び、双晶型（例えば、Ｍｎ−Ｃｕ合金）に分類される。
双晶型に分類されるＭｎ−Ｃｕ系制振合金は、双晶の運動により振動を吸収し、高い減衰能を発揮する。この双晶は、Ｍｎ−Ｃｕ系合金をオーステナイト（γ）相領域から徐冷することによって形成される。Ｍｎ−Ｃｕ系制振合金を得るにはこの熱処理が欠かせないが、熱処理方法によって制振特性に違いが出るため、より良い制振特性を得るために種々の熱処理方法が提案されている。
A damping alloy is an alloy that absorbs vibration by converting vibration energy from outside into heat by internal friction. Due to its damping function, a composite type (for example, flake graphite cast iron), a ferromagnetic type (for example, , Fe—Cr alloy), dislocation type (for example, Mg—Zr alloy), and twin type (for example, Mn—Cu alloy).
A Mn—Cu vibration damping alloy classified as a twin type absorbs vibrations due to the movement of twins and exhibits high damping ability. This twin is formed by slowly cooling the Mn—Cu alloy from the austenite (γ) phase region. This heat treatment is indispensable for obtaining a Mn—Cu-based vibration damping alloy, but various heat treatment methods have been proposed in order to obtain better vibration damping characteristics since the vibration damping characteristics differ depending on the heat treatment method.
例えば、特許文献１には、マンガン基双晶型制振合金を８００〜１１００℃で加熱した後、２５０℃〜４７５℃の温度範囲に入るまで０．８５〜１．６５℃／分で定速徐冷し（γ相領域から定速徐冷）、更に１０℃／分以上で急冷するマンガン基双晶型制振合金の熱処理方法が開示されている。 For example, in Patent Document 1, a manganese-based twin-type vibration damping alloy is heated at 800 to 1100 ° C., and then at a constant speed of 0.85 to 1.65 ° C./min until it enters a temperature range of 250 ° C. to 475 ° C. A method for heat treatment of a manganese-based twin-type vibration-damping alloy that is gradually cooled (constant-speed gradually cooling from the γ phase region) and further rapidly cooled at 10 ° C./min or more is disclosed.
しかしながら、特許文献１に記載のＭｎ−Ｃｕ系合金は、高周波誘導加熱炉を用いて、アルゴン雰囲気下で鋼塊（Ｍｎ：ｂａｌ．、Ｃｕ：２０ａｔ％、Ｎｉ：４ａｔ％、Ｆｅ：２ａｔ％）を作製し、熱間鍛造、熱間圧延及び冷間圧延を経て半製品（棒材、板材等）とし、その後で、上記熱処理（すなわち、８００〜１１００℃で加熱した後、２５０℃〜４７５℃の温度範囲に入るまで０．８５〜１．６５℃／分で定速徐冷し（γ相領域から定速徐冷）、更に１０℃／分以上で急冷する熱処理）を行うものである。
このように、鋳込み後の鋼塊を熱間鍛造等の加工により半製品（棒材、板材等）とする場合、ＭｎとＣｕのマクロ偏析が存在するため、加工の際、内部には不均一な応力が発生し、割れ（特に表面への）が極めて発生しやすいという問題があった。更に、鋳込み後の鋼塊は、ＭｎとＣｕのマクロ偏析が大きいため、熱間鍛造等の加工による割れ発生が顕著となり、歩留まり低下の要因になるという問題もあった。また、ＭｎとＣｕのマクロ偏析が大きいと、溶融開始温度が９００℃〜１０００℃以上に達するＭｎ−Ｃｕ組成比（濃度比率）が局所的に存在し、必然的に熱間鍛造ができる温度範囲が狭くなり、熱間加工の際に割れが発生しやすいという問題があった。
更に、鋳込み後の鋼塊は、結晶粒が十分に微細化されておらず、強度が低く、割れ発生の原因になるという問題があった。
However, the Mn—Cu alloy described in Patent Document 1 is a steel ingot (Mn: bal., Cu: 20 at%, Ni: 4 at%, Fe: 2 at%) using a high frequency induction heating furnace in an argon atmosphere. Is made into a semi-finished product (bar material, plate material, etc.) through hot forging, hot rolling and cold rolling, and then heated at 800-1100 ° C. and then 250 ° C.-475 ° C. Until the temperature falls within a temperature range of 0.85 to 1.65 ° C./min, heat treatment is performed at a constant rate of slow cooling (constant rate slow cooling from the γ-phase region), and a rapid cooling at 10 ° C./min or higher.
Thus, when a steel ingot after casting is made into a semi-finished product (bar material, plate material, etc.) by processing such as hot forging, macro segregation of Mn and Cu exists, and therefore, the inside is not uniform during processing. There was a problem that a large amount of stress was generated and cracks (particularly to the surface) were very likely to occur. Furthermore, since the steel ingot after casting has large macrosegregation of Mn and Cu, there is a problem that cracking due to processing such as hot forging becomes remarkable, which causes a decrease in yield. Moreover, when the macrosegregation of Mn and Cu is large, the Mn—Cu composition ratio (concentration ratio) where the melting start temperature reaches 900 ° C. to 1000 ° C. or more exists locally, and the temperature range in which hot forging can be inevitably performed. There is a problem that cracks tend to occur during hot working.
Further, the steel ingot after casting has a problem that the crystal grains are not sufficiently refined, the strength is low, and cracks are caused.
本発明は、上記事情に鑑みてなされたものであり、その目的は、加工時における割れ発生が抑制され、良好な制振特性が得られるＭｎ−Ｃｕ系制振合金及びその製造方法を提供することにある。 The present invention has been made in view of the above circumstances, and an object of the present invention is to provide a Mn—Cu-based vibration damping alloy in which generation of cracks during processing is suppressed and good vibration damping characteristics can be obtained, and a method for producing the same. There is.
上記課題を解決するために、本発明に係るＭｎ−Ｃｕ系制振合金は、Ｃｕ：１５〜２５ａｔ％を含み、更に、Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％を含み、残部がＭｎ及び不可避的不純物からなることを要旨とする。この場合に、更に、Ｎｉ：２〜８ａｔ％及びＦｅ：１〜３ａｔ％を含むものでもよく、これに更に、Ａｌ：２〜５ａｔ％を含むものでもよい。 In order to solve the above problems, the Mn—Cu vibration damping alloy according to the present invention contains Cu: 15-25 at%, and further Ca: 0.001-0.007 at% or Mg: 0.005-0. The content is .23 at%, with the balance being Mn and inevitable impurities. In this case, it may further contain Ni: 2 to 8 at% and Fe: 1 to 3 at%, and may further contain Al: 2 to 5 at%.
上記課題を解決するために、本発明に係るＭｎ−Ｃｕ系制振合金の製造方法は、Ｃｕ：１５〜２５ａｔ％を含み、更に、Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％を含み、残部がＭｎ及び不可避的不純物からなるＭｎ−Ｃｕ系合金の鋼塊を８００℃以上固相線温度以下の固溶化温度に加熱し、前記固溶化温度で保持した後、空冷以上の冷却速度で冷却する固溶化処理工程を備えたことを要旨とする。前記固溶化処理工程は、前記固溶化温度での保持時間が８時間以上であることが望ましい。
そして、前記固溶化処理工程の後に、前記合金を最終製品形状又はこれより大きい形状まで加工する加工工程と、前記加工工程の後に、前記合金を８００℃以上固相線温度以下の焼鈍温度に加熱し、前記焼鈍温度で０．５時間以上８時間以下保持した後、前記合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷する高温焼鈍工程と、前記高温焼鈍工程の後に、前記合金を１０℃／分以上で冷却する急冷工程とを備えてもよい。尚、前記Ｍｎ−Ｃｕ系合金は、更に、Ｎｉ：２〜８ａｔ％及びＦｅ：１〜３ａｔ％を含むものでもよく、これに更に、Ａｌ：２〜５ａｔ％を含むものでもよい。
In order to solve the above problems, a method for producing a Mn—Cu vibration damping alloy according to the present invention includes Cu: 15 to 25 at%, and Ca: 0.001 to 0.007 at% or Mg: 0.00. A steel ingot of Mn—Cu alloy containing 005 to 0.23 at% and the balance consisting of Mn and inevitable impurities was heated to a solid solution temperature not lower than 800 ° C. and not higher than the solidus temperature, and held at the solid solution temperature. Then, the gist is that a solution treatment step of cooling at a cooling rate higher than air cooling is provided. In the solution treatment step, the retention time at the solution temperature is preferably 8 hours or more.
Then, after the solution treatment step, a processing step for processing the alloy into a final product shape or a shape larger than this, and after the processing step, the alloy is heated to an annealing temperature of 800 ° C. or higher and a solidus temperature or lower. Then, after holding at the annealing temperature for 0.5 hours or more and 8 hours or less, until the temperature of the alloy enters any of the ranges of 250 ° C. to 450 ° C., a constant speed of 0.85 to 1.65 ° C./min You may provide the high temperature annealing process which anneals, and the rapid cooling process which cools the said alloy at 10 degree-C / min or more after the said high temperature annealing process. The Mn—Cu based alloy may further contain Ni: 2-8 at% and Fe: 1-3 at%, and may further contain Al: 2-5 at%.
本発明に係るＭｎ−Ｃｕ系制振合金は、上記構成を備えたものであるから、鋼塊にした後の結晶粒の微細化が促進され、これにより、合金強度が高まるため、加工時における割れ発生が抑制される。また、本発明に係るＭｎ−Ｃｕ系制振合金は、上記構成を備えたものであるから、適切な熱処理を行うことにより、Ｍｎ及びＣｕの偏析が低減されるため、加工時における不均一応力の発生が低減される。従って、より一層、加工時における割れ発生が抑制される。 Since the Mn—Cu based vibration damping alloy according to the present invention has the above-described configuration, the refinement of crystal grains after being made into a steel ingot is promoted, thereby increasing the alloy strength. Cracking is suppressed. In addition, since the Mn—Cu vibration damping alloy according to the present invention has the above-described configuration, by performing an appropriate heat treatment, segregation of Mn and Cu is reduced, and therefore, non-uniform stress during processing. Is reduced. Therefore, the generation of cracks during processing is further suppressed.
本発明に係るＭｎ−Ｃｕ系制振合金の製造方法によれば、所定の組成を備えたＭｎ−Ｃｕ系合金の鋼塊が８００℃以上固相線温度以下の固溶化温度に加熱され、その固溶化温度で保持された後、空冷以上の冷却速度で冷却されるため、ＭｎとＣｕのマクロ偏析が低減されるとともに、結晶粒が微細化される。
従って、固溶化処理工程後の合金に熱間加工がなされても、不均一応力の発生による割れや、低強度に起因する割れの発生が抑制される。
高温焼鈍工程においては、加工された合金が８００℃以上固相線温度以下の焼鈍温度に加熱され、その焼鈍温度で保持された後、合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷されるため、αＭｎ相の生成や熱膨張によるひずみを回避しつつ、十分な双晶が形成され、良好な制振特性が得られる。
According to the method for producing a Mn—Cu vibration damping alloy according to the present invention, a steel ingot of a Mn—Cu alloy having a predetermined composition is heated to a solution temperature of 800 ° C. or more and a solidus temperature or less, After being held at the solution temperature, it is cooled at a cooling rate equal to or higher than that of air cooling, so that macrosegregation of Mn and Cu is reduced and crystal grains are refined.
Therefore, even when hot working is performed on the alloy after the solution treatment step, cracks due to generation of non-uniform stress and cracks due to low strength are suppressed.
In the high-temperature annealing step, the processed alloy is heated to an annealing temperature of 800 ° C. or more and a solidus temperature or less and held at the annealing temperature, and then the temperature of the alloy is in any of the ranges of 250 ° C. to 450 ° C. Since it is cooled at a constant rate of 0.85 to 1.65 ° C./min until it enters, sufficient twins are formed while avoiding distortion due to the formation of αMn phase and thermal expansion, and good damping characteristics can get.
以下に、本発明の一実施の形態について詳細に説明する。
（組成及びその限定理由）
本発明の一実施形態に係るＭｎ−Ｃｕ系制振合金は、γ相領域から冷却することによって形成される双晶の運動により振動を吸収する双晶型制振合金のうち、Ｍｎ−Ｃｕ−Ｎｉ−Ｆｅ系のものが好ましく、Ｃｕ：１５〜２５ａｔ％、Ｎｉ：２〜８ａｔ％、Ｆｅ：１〜３ａｔ％を含有し、更に、Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％を含有し、残部がＭｎ及び不可避的不純物からなる。
本発明の一実施形態に係るＭｎ−Ｃｕ系制振合金は、更に、Ａｌ：２〜５ａｔ％を含有してもよい。
Hereinafter, an embodiment of the present invention will be described in detail.
(Composition and reasons for limitation)
The Mn—Cu-based damping alloy according to an embodiment of the present invention is a Mn—Cu— of twin type damping alloys that absorb vibrations by the movement of twins formed by cooling from the γ phase region. Ni-Fe-based materials are preferred, including Cu: 15 to 25 at%, Ni: 2 to 8 at%, Fe: 1 to 3 at%, and Ca: 0.001 to 0.007 at% or Mg: 0. 005 to 0.23 at%, with the balance being Mn and inevitable impurities.
The Mn—Cu vibration damping alloy according to one embodiment of the present invention may further contain Al: 2 to 5 at%.
（１）Ｃｕ：１５〜２５ａｔ％、Ｎｉ：２〜８ａｔ％及びＦｅ：１〜３ａｔ％。
Ｍｎをベースとして、この組成にすると後述する高温焼鈍工程を行うことにより双晶が形成され、良好な減衰特性が得られるためである。Ｃｕを１５〜２５ａｔ％としたのは、成形加工性を高め、常温近傍に変態点を移動させるためである。Ｎｉを２〜８ａｔ％及びＦｅ：１〜３ａｔ％としたのは、良好な減衰特性を得るためである。
(1) Cu: 15 to 25 at%, Ni: 2 to 8 at%, and Fe: 1 to 3 at%.
This is because, when Mn is used as a base and this composition is used, twins are formed by performing a high-temperature annealing process, which will be described later, and good damping characteristics can be obtained. The reason why the Cu content is set to 15 to 25 at% is to improve the molding processability and move the transformation point to around room temperature. The reason why Ni is set to 2 to 8 at% and Fe: 1 to 3 at% is to obtain good damping characteristics.
（２）Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％。
Ｃａ及びＭｇは、結晶粒の微細化、及び、その微細化によって強度を高めて、熱間加工時の割れを抑制するために含有させる。これらは、いずれか一方を含有させればよい。これらの効果を得るために、Ｃａを０．００１ａｔ％以上、Ｍｇを０．００５ａｔ％以上とした。一方、Ｃａ又はＭｇの量が多すぎると、酸化物又は介在物が形成され加工性が悪化するとともに、特に過剰な介在物は、結晶粒界や双晶界面に析出し、双晶界面の移動が困難となり、減衰特性を劣化させる。また、Ｍｇの場合、沸点が低いので量が多すぎるとスプラッシュが発生してしまう。そこで、これらを回避するために、Ｃａを０．００７ａｔ％以下、Ｍｇを０．２３ａｔ％以下とした。
(2) Ca: 0.001 to 0.007 at% or Mg: 0.005 to 0.23 at%.
Ca and Mg are contained in order to increase the strength by refinement of crystal grains and the refinement, and to suppress cracking during hot working. Any one of these may be contained. In order to obtain these effects, Ca is set to 0.001 at% or more, and Mg is set to 0.005 at% or more. On the other hand, if the amount of Ca or Mg is too large, oxides or inclusions are formed and the workability deteriorates. In particular, excessive inclusions precipitate at the grain boundaries and twin interfaces, and the migration of twin interfaces Becomes difficult and deteriorates the attenuation characteristics. In the case of Mg, since the boiling point is low, splashing occurs when the amount is too large. Therefore, in order to avoid these, Ca is set to 0.007 at% or less and Mg is set to 0.23 at% or less.
（３）残部がＭｎ及び不可避的不純物。
Ｍｎは、制振材料として周知であり、本発明に係るＭｎ−Ｃｕ系制振合金のベースである。具体的には、Ｍｎは、７０〜７５ａｔ％である。
不可避的不純物には、Ｃ、Ｏ、Ｎ等がある。
(3) The balance is Mn and inevitable impurities.
Mn is well known as a vibration damping material and is the base of the Mn—Cu based vibration damping alloy according to the present invention. Specifically, Mn is 70 to 75 at%.
Inevitable impurities include C, O, N, and the like.
（４）Ａｌ：２〜５ａｔ％
Ａｌは、必要に応じて含有させればよく、含有させる場合には、２〜５ａｔ％含有させる。Ａｌを２〜５ａｔ％としたのは、上記組成にＡｌをこの割合で添加すれば、剛性を高めつつ、良好な減衰特性が得られるためである。
(4) Al: 2 to 5 at%
Al may be contained if necessary, and if contained, 2 to 5 at% is contained. The reason why Al is 2 to 5 at% is that if Al is added to the above composition in this proportion, good damping characteristics can be obtained while increasing rigidity.
（製造方法）
本発明の一実施形態に係るＭｎ−Ｃｕ系制振合金は、鋳造工程、固溶化処理工程、加工工程、高温焼鈍工程、急冷工程を行うことにより製造される。
以下これらの各工程について説明する。
(Production method)
The Mn—Cu vibration damping alloy according to an embodiment of the present invention is manufactured by performing a casting process, a solution treatment process, a processing process, a high-temperature annealing process, and a rapid cooling process.
Each of these steps will be described below.
（１）鋳造工程は、Ｃｕ：１５〜２５ａｔ％を含み、更に、Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％を含み、残部がＭｎ及び不可避的不純物からなるＭｎ−Ｃｕ系合金を溶解・鋳造して鋼塊を得る工程である。
「Ｍｎ−Ｃｕ系合金」は、上記組成を備えたものであれば特に限定されず、後述する高温焼鈍工程においてγ相領域から冷却することによって双晶が形成されるものであればよく、更に、Ｎｉ：２〜８ａｔ％及びＦｅ：１〜３ａｔ％を含むものや、更に、Ｎｉ：２〜８ａｔ％、Ｆｅ：１〜３ａｔ％及びＡｌ：２〜５ａｔ％を含むものでもよい。
「鋼塊」は、所定の組成に配合され、溶解・鋳造されたものであればよく、加工の有無は問わない。
(1) The casting process includes Cu: 15 to 25 at%, and further includes Ca: 0.001 to 0.007 at% or Mg: 0.005 to 0.23 at%, and the balance is from Mn and inevitable impurities. This is a step of obtaining a steel ingot by melting and casting the Mn-Cu alloy.
The “Mn—Cu-based alloy” is not particularly limited as long as it has the above composition, and may be any one in which twins are formed by cooling from the γ-phase region in the high-temperature annealing step described below. Ni: 2 to 8 at% and Fe: 1 to 3 at%, and further, Ni: 2 to 8 at%, Fe: 1 to 3 at%, and Al: 2 to 5 at% may be included.
The “steel ingot” is not particularly limited as long as it is blended in a predetermined composition and melted and cast.
（２）固溶化処理工程は、上記（１）で得られた鋼塊を８００℃以上固相線温度以下の固溶化温度に加熱し、その固溶化温度で保持した後、空冷以上の冷却速度で冷却する工程である。固溶化処理工程は、ＭｎとＣｕのマクロ偏析を解消するために行う。
「固溶化温度に加熱する」のは、γ相単相となる温度にするためである。固溶化温度は、より短時間でＭｎとＣｕのマクロ偏析を解消するには高い方が好ましく、具体的には、８００℃以上が好ましく、８５０℃以上が更に好ましく、８７０℃以上がより更に好ましい。一方、固溶化温度が高くなるほど結晶粒が粗大化し、固溶化温度が高すぎると液相が生成する。従って、固溶化温度は、固相線温度以下が好ましい。Ｃｕを２０±５ａｔ％程度含むＭｎ−Ｃｕ系制振合金の固相線温度は、約１０００℃であるため、具体的には、固溶化温度は、１０００℃以下が好ましく、９５０℃以下が更に好ましく、９２５℃以下がより更に好ましい。
(2) In the solution treatment step, the steel ingot obtained in (1) above is heated to a solid solution temperature not lower than 800 ° C. and not higher than the solidus temperature, held at the solid solution temperature, and then cooled by air cooling or higher. It is the process of cooling at. The solution treatment step is performed to eliminate macro segregation of Mn and Cu.
The reason for “heating to the solid solution temperature” is to set the temperature to a single phase of γ phase. The solution temperature is preferably higher in order to eliminate macro segregation of Mn and Cu in a shorter time, specifically, 800 ° C or higher is preferable, 850 ° C or higher is more preferable, and 870 ° C or higher is even more preferable. . On the other hand, as the solution temperature increases, the crystal grains become coarser, and when the solution temperature is too high, a liquid phase is generated. Therefore, the solid solution temperature is preferably equal to or lower than the solidus temperature. Since the solidus temperature of the Mn—Cu vibration damping alloy containing about 20 ± 5 at% of Cu is about 1000 ° C., specifically, the solution temperature is preferably 1000 ° C. or less, and more preferably 950 ° C. or less. Preferably, 925 ° C. or lower is even more preferable.
「固溶化温度で保持する」のは、ＭｎとＣｕのマクロ偏析を解消するためである。また、Ｍｇ又はＣａを含有させたことにより促進される結晶粒の微細化を図るためである。保持時間は、特に限定されないが、ＭｎとＣｕのマクロ偏析を解消するには長い方が好ましい。ＭｎとＣｕのマクロ偏析を解消するのは、（１）熱間加工時に内部に不均一な応力が発生して割れ発生の原因になったり、（２）Ｍｎ−Ｃｕ濃度比率が低融点（９００℃）以下に達する個所が局所的に存在し、熱間鍛造ができる温度範囲が狭くなり、割れ発生の原因になるからである。そして、割れ発生は、歩留まり低下の要因にもなる。そこで、割れ発生を抑制するためには、固溶化温度での保持時間は、４時間以上が好ましく、８時間以上が更に好ましい。また、ＭｎとＣｕのマクロ偏析を解消すれば損失係数を大きくすることができ、より高い制振特性が得られる。 “Hold at the solution temperature” is to eliminate macro segregation of Mn and Cu. Moreover, it is for aiming at the refinement | miniaturization of the crystal grain accelerated | stimulated by containing Mg or Ca. The holding time is not particularly limited, but a longer one is preferable for eliminating macro segregation of Mn and Cu. The reason for eliminating macro segregation of Mn and Cu is that (1) non-uniform stress is generated inside during hot working, causing cracks, or (2) the Mn—Cu concentration ratio is low melting point (900 This is because there are local portions that reach below (° C.), the temperature range in which hot forging can be performed is narrowed, and cracking occurs. The occurrence of cracks also causes a decrease in yield. Therefore, in order to suppress the occurrence of cracks, the retention time at the solution temperature is preferably 4 hours or more, and more preferably 8 hours or more. Further, if macro segregation of Mn and Cu is eliminated, the loss factor can be increased, and higher vibration damping characteristics can be obtained.
「空冷以上の冷却速度で冷却する」のは、冷却速度が遅すぎると冷却時にαＭｎ相が生成するのでこれを回避しつつ、冷却時のＭｎとＣｕのマクロ偏析を抑制するためである。 The reason for “cooling at a cooling rate higher than air cooling” is to prevent macro segregation of Mn and Cu during cooling while avoiding this because an αMn phase is generated during cooling if the cooling rate is too slow.
尚、「固溶化処理時の雰囲気」は、特に限定されないが、不活性雰囲気中（例えば、アルゴン雰囲気中）、還元雰囲気中（例えば、水素雰囲気中）、あるいは大気中のいずれであってもよい。大気中で固溶化処理を行うと、製造コストを低減できるという利点がある。大気中で固溶化処理を行うと、合金表面に酸化被膜が形成されるが、後述する加工工程の際に酸化被膜を除去すれば制振特性を劣化させることがない。 The “atmosphere during the solution treatment” is not particularly limited, and may be any of an inert atmosphere (for example, an argon atmosphere), a reducing atmosphere (for example, a hydrogen atmosphere), or the air. . When the solution treatment is performed in the air, there is an advantage that the manufacturing cost can be reduced. When the solution treatment is performed in the atmosphere, an oxide film is formed on the surface of the alloy. However, if the oxide film is removed during the processing step described later, the vibration damping characteristics are not deteriorated.
（３）加工工程は、上記（２）で固溶化処理された合金を最終製品形状又はこれより大きい形状まで加工する工程である。
「固溶化処理された合金を最終製品形状又はこれより大きい形状まで加工する」のは、所望の最終製品を得るために必要だからである。これにより、鋳造時における鋳造組織の破壊、鋳造欠陥の消滅、偏析の解消等の効果も得られる。この加工は、熱間鍛造又は熱間圧延等の熱間加工により行われ、合金は、最終製品又はそれより大きい所定の形状（例えば、棒材や板材）に加工される。この後、更に、冷間圧延等の冷間加工を行ってもよい。また、必要に応じて固溶化処理の際に合金表面に形成された酸化被膜の除去を行ってもよい。「最終製品形状より大きい形状」にするのは、後述する高温焼鈍工程の条件によっては、表面に酸化被膜が生成したり、寸法変化を生ずる場合があるため、精加工仕上げのための削り代を見込んでおくためである。
(3) The processing step is a step of processing the alloy that has been subjected to the solution treatment in (2) to a final product shape or a shape larger than this.
“Solubilized alloy is processed into a final product shape or a larger shape” because it is necessary to obtain a desired final product. Thereby, effects, such as destruction of a cast structure at the time of casting, disappearance of a casting defect, elimination of segregation, are also acquired. This processing is performed by hot processing such as hot forging or hot rolling, and the alloy is processed into a final product or a predetermined shape larger than the final product (for example, a bar or a plate). Thereafter, cold working such as cold rolling may be further performed. Moreover, you may remove the oxide film formed in the alloy surface in the case of a solution treatment as needed. The “larger shape than the final product shape” means that depending on the conditions of the high-temperature annealing process described later, an oxide film may be formed on the surface or dimensional changes may occur. This is to keep in mind.
（４）高温焼鈍工程は、上記（３）で加工された合金を８００℃以上固相線温度以下の焼鈍温度に加熱し、その焼鈍温度で保持した後、その合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷する工程である。 (4) In the high-temperature annealing step, the alloy processed in the above (3) is heated to an annealing temperature not lower than 800 ° C. and not higher than the solidus temperature, held at the annealing temperature, and then the temperature of the alloy is 250 ° C. to 450 ° C. This is a step of slow cooling at a constant rate of 0.85 to 1.65 ° C./min until entering any of the range of ° C.
「焼鈍温度に加熱し、その焼鈍温度で保持する」のは、合金をγ相状態にして、固溶化処理での冷却により生じたひずみを解消するためである。更に、固溶化処理で解消されなかったＭｎとＣｕのマクロ偏析を消滅させるためである。
従って、焼鈍温度は、８００℃以上が好ましく、８５０℃以上が更に好ましく、８７０℃以上がより更に好ましい。できるだけ短時間でひずみや偏析を消滅させるには、焼鈍温度は高い方が好ましいが、焼鈍温度が高すぎると、結晶粒が粗大化し、材料が脆化する。更に温度が高い場合には、液相が生成してしまう。また、焼鈍温度が高くなるほど、合金表面からＭｎが揮発しやすくなる。従って、焼鈍温度は、固相線温度以下が好ましく、９５０℃以下が更に好ましく、９２５℃以下がより更に好ましい。
焼鈍温度での保持時間は、特に限定されないが、ひずみや偏析を消滅させるためには、０．５時間以上が好ましく、１時間以上が更に好ましい。一方、必要以上長く保持すると、合金表面からＭｎが蒸発する。従って、焼鈍温度での保持時間は、加工された合金に十分な削り代がある場合には、８時間以下が好ましく、４時間以下が更に好ましく、また、加工された合金の削り代が少ない場合には、４時間以下が好ましく、３時間以下が更に好ましい。加工された合金が最終製品形状に近くなるほど、焼鈍温度での保持時間を短くするとよい。
The reason for “heating to the annealing temperature and holding at the annealing temperature” is to eliminate the strain caused by cooling in the solution treatment by bringing the alloy into a γ phase state. Furthermore, this is to eliminate macro segregation of Mn and Cu that has not been eliminated by the solution treatment.
Accordingly, the annealing temperature is preferably 800 ° C. or higher, more preferably 850 ° C. or higher, and even more preferably 870 ° C. or higher. In order to eliminate strain and segregation in as short a time as possible, the annealing temperature is preferably high, but if the annealing temperature is too high, the crystal grains become coarse and the material becomes brittle. Further, when the temperature is high, a liquid phase is generated. Further, as the annealing temperature becomes higher, Mn is more easily volatilized from the alloy surface. Therefore, the annealing temperature is preferably not more than the solidus temperature, more preferably not more than 950 ° C, and still more preferably not more than 925 ° C.
The holding time at the annealing temperature is not particularly limited, but is preferably 0.5 hours or longer and more preferably 1 hour or longer in order to eliminate strain and segregation. On the other hand, if it is kept longer than necessary, Mn evaporates from the alloy surface. Accordingly, the holding time at the annealing temperature is preferably 8 hours or less, more preferably 4 hours or less when the machined alloy has a sufficient machining allowance, and when the machining alloy has a less machining allowance. Is preferably 4 hours or shorter, more preferably 3 hours or shorter. The closer the processed alloy is to the final product shape, the shorter the holding time at the annealing temperature.
「その合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷する」のは、定速で、かつ、この冷却速度でスピノーダル分解を起こす温度（４００℃近傍）を通過させると、スピノーダル分解により双晶が形成され、これにより、制振特性が備えられるからである。一方、冷却速度が遅すぎると双晶形成に寄与しないαＭｎ相が生成し、逆に、冷却速度が速すぎると双晶形成が不十分になる上、合金内部に熱膨張によるひずみが発生する。冷却速度にばらつきがあっても、同様である。従って、αＭｎ相の生成を回避しつつ双晶を十分に形成させるには、定速徐冷速度は、０．８５℃／分以上が好ましく、１．３３℃／分以上が更に好ましい。また、熱膨張によるひずみを回避しつつ双晶を十分に形成させるには、定速徐冷速度は、１．６５℃／分以下が好ましく、１．６０℃／分以下が更に好ましい。
定速徐冷の終了温度は、特に限定されないが、スピノーダル分解により、双晶が十分に形成された後であればよく、具体的には、合金が２５０℃〜４５０℃に達するいずれかの温度である。残留ひずみを回避し十分な制振特性を得るには、定速徐冷の終了温度は、４５０℃以下が好ましく、４００℃以下が更に好ましく、３５０℃以下がより更に好ましい。αＭｎ相の生成を抑制するには、定速徐冷の終了温度は、２５０℃以上が好ましく、２７５℃以上が更に好ましく、３００℃以上がより更に好ましい。
“Slow cooling at a constant rate of 0.85 to 1.65 ° C./min until the temperature of the alloy falls in any of the ranges of 250 ° C. to 450 ° C.” means that the cooling rate is constant. This is because if a temperature causing spinodal decomposition (near 400 ° C.) is passed, twins are formed by spinodal decomposition, thereby providing damping characteristics. On the other hand, if the cooling rate is too slow, an αMn phase that does not contribute to twin formation is generated. Conversely, if the cooling rate is too fast, twin formation is insufficient and strain due to thermal expansion occurs inside the alloy. The same applies if the cooling rate varies. Therefore, in order to sufficiently form twins while avoiding the formation of αMn phase, the constant slow cooling rate is preferably 0.85 ° C./min or more, and more preferably 1.33 ° C./min or more. In order to sufficiently form twins while avoiding strain due to thermal expansion, the constant slow cooling rate is preferably 1.65 ° C./min or less, and more preferably 1.60 ° C./min or less.
The end temperature of constant-speed slow cooling is not particularly limited, but may be after the twins are sufficiently formed by spinodal decomposition. Specifically, any temperature at which the alloy reaches 250 ° C. to 450 ° C. It is. In order to avoid residual strain and obtain sufficient vibration damping characteristics, the constant temperature slow cooling end temperature is preferably 450 ° C. or lower, more preferably 400 ° C. or lower, and even more preferably 350 ° C. or lower. In order to suppress the formation of αMn phase, the constant temperature slow cooling end temperature is preferably 250 ° C. or higher, more preferably 275 ° C. or higher, and still more preferably 300 ° C. or higher.
尚、「焼鈍工程の雰囲気」は、特に限定されないが、不活性雰囲気（例えば、アルゴン雰囲気、窒素雰囲気）、又は、還元雰囲気（例えば、水素雰囲気）が好ましい。大気中で焼鈍すると、表面に酸化被膜が形成されるので、これを回避するためである。 The “atmosphere of the annealing step” is not particularly limited, but an inert atmosphere (for example, an argon atmosphere or a nitrogen atmosphere) or a reducing atmosphere (for example, a hydrogen atmosphere) is preferable. This is to avoid an oxide film formed on the surface when annealing is performed in the atmosphere.
（５）急冷工程は、上記（４）で制振特性が備えられた合金を１０℃／分以上で冷却する工程である。
「１０℃／分以上で冷却する」のは、αＭｎ相の生成を抑制するためであり、２０℃／分以上で冷却するのが更に好ましい。尚、冷却方法としては、衝風冷却、水冷、油冷を用いることができ、高速冷却をする場合は、水冷及び油冷を行うとよい。
急冷工程の終了温度は、１００℃以下の温度であればよく、例えば、室温であればよい。
(5) The rapid cooling step is a step of cooling the alloy having the vibration damping characteristics in (4) above at 10 ° C./min or more.
“Cooling at 10 ° C./min or more” is to suppress the formation of αMn phase, and it is more preferable to cool at 20 ° C./min or more. In addition, as a cooling method, blast cooling, water cooling, and oil cooling can be used, and when performing high speed cooling, it is good to perform water cooling and oil cooling.
The end temperature of the rapid cooling process may be a temperature of 100 ° C. or lower, for example, room temperature.
尚、上記（４）の高温焼鈍工程、又は、上記（５）の急冷工程後を行った後の合金は、そのまま最終製品として販売・使用等しても良く、又は、精仕上げ加工（焼鈍時に生成した酸化被膜の除去、寸法変化の矯正、Ｍｎの蒸発により生じた表面変質層の除去等）を行ってもよい。 The alloy after the high-temperature annealing step (4) or after the rapid cooling step (5) may be sold and used as a final product as it is, or a fine finishing process (during annealing) The generated oxide film may be removed, the dimensional change may be corrected, or the surface altered layer generated by evaporation of Mn may be removed).
（作用）
次に、本発明の一実施形態に係るＭｎ−Ｃｕ系制振合金及びその製造方法の作用について説明する。
固溶化処理工程において、所定の組成からなるＭｎ−Ｃｕ系合金の鋼塊が８００℃以上固相線温度以下の固溶化温度に加熱され、その固溶化温度で保持された後、空冷以上の冷却速度で冷却されると、ＭｎとＣｕのマクロ偏析が低減されるとともに、Ｍｇ又はＣａを含有させたことにより、結晶粒が微細化される。
そのため、加工工程においては、合金は、割れ発生が抑制されつつ、最終製品形状又はこれより大きい形状まで加工される。この工程において割れ発生が抑制されることで、最終的に得られるＭｎ−Ｃｕ系制振合金の割れ発生が抑制される。
次に、合金は、焼鈍温度に加熱され、その焼鈍温度で保持された後、合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷される。従って、得られたＭｎ−Ｃｕ系制振合金は、αＭｎ相の生成や熱膨張によるひずみを回避しつつ、十分な双晶が形成され、高い制振特性が備えられる。
(Function)
Next, the operation of the Mn—Cu vibration damping alloy and the manufacturing method thereof according to one embodiment of the present invention will be described.
In the solution treatment step, the Mn-Cu alloy ingot having a predetermined composition is heated to a solid solution temperature not lower than 800 ° C and not higher than the solidus temperature, and maintained at the solid solution temperature, and then cooled beyond air cooling. When cooled at a speed, macrosegregation of Mn and Cu is reduced, and the crystal grains are refined by containing Mg or Ca.
Therefore, in the processing step, the alloy is processed to a final product shape or a shape larger than this while suppressing the occurrence of cracks. By suppressing the generation of cracks in this step, the generation of cracks in the finally obtained Mn—Cu vibration damping alloy is suppressed.
The alloy is then heated to the annealing temperature, held at that annealing temperature, and then at 0.85 to 1.65 ° C./min until the temperature of the alloy enters any of the ranges of 250 ° C. to 450 ° C. Cooled at a constant speed. Therefore, the obtained Mn—Cu-based damping alloy has sufficient twinning properties and high damping characteristics while avoiding strain due to the formation of αMn phase and thermal expansion.
また、得られたＭｎ−Ｃｕ系制振合金は、固溶化処理により、ＭｎとＣｕのマクロ偏析が低減されているため、Ｃｕに対する室温での固溶限が０％であるＣ（不可避的に含まれる）の双晶界面又は結晶粒界への析出が抑制され、これにより、制振特性の経時劣化が抑制される。 In addition, since the obtained Mn—Cu vibration-damping alloy has reduced the macrosegregation of Mn and Cu by the solution treatment, C (inevitable) has a solid solubility limit of 0% with respect to Cu at room temperature. Precipitation) at the twin interface or the grain boundary is suppressed, thereby suppressing the deterioration of the damping characteristics with time.
以下、本発明の実施例及び比較例について説明する。尚、以下の実施例及び比較例においては、実施例１のＮｉを２ａｔ％、８ａｔ％とした組成、実施例１のＦｅを１ａｔ％、３ａｔ％とした組成については実施例１と同等の性能を有するため省略した。また、実施例６のＡｌを２ａｔ％とした組成も、実施例６と同等の性能を有するため省略した。更に、実施例５のＣｕを１５ａｔ％とした組成も、実施例５と同等の性能を有するため省略した。
（実施例１〜８及び比較例１〜４）
（割れについて）
実施例１〜８及び比較例１〜４について、表１に示すＡ〜Ｊの組成となるように原料を高周波誘導加熱炉に投入し、アルゴン雰囲気下で溶解・鋳造し、それぞれの組成について、１５０ｋｇ鋼塊（直径１７０ｍｍ〜直径１９０ｍｍ×長さ８００ｍｍ）を作製した。
Examples of the present invention and comparative examples will be described below. In the following examples and comparative examples, the composition of Example 1 with Ni of 2 at% and 8 at% and the composition of Example 1 with Fe of 1 at% and 3 at% are equivalent to those of Example 1. It was omitted because it has The composition of Example 6 with Al of 2 at% was also omitted because it had the same performance as in Example 6. Further, the composition of Example 5 with Cu at 15 at% is omitted because it has the same performance as Example 5.
(Examples 1-8 and Comparative Examples 1-4)
(About cracking)
For Examples 1 to 8 and Comparative Examples 1 to 4, the raw materials were put into a high-frequency induction heating furnace so as to have the compositions of A to J shown in Table 1, and melted and cast in an argon atmosphere. A 150 kg steel ingot (diameter 170 mm-diameter 190 mm x length 800 mm) was produced.
次に、作製した各鋼塊に対して固溶化処理を行った（比較例２を除く）。固溶化処理は、表２に示す条件、すなわち、大気雰囲気下、９００℃に加熱し、その温度で８時間（実施例３については２４時間）保持した後、空冷することにより行った。 Next, a solution treatment was performed on each steel ingot produced (except for Comparative Example 2). The solution treatment was performed by heating to 900 ° C. under the conditions shown in Table 2, that is, in an air atmosphere, holding at that temperature for 8 hours (24 hours for Example 3), and then air cooling.
次に、得られた各鋼塊に対して熱間加工を行った。熱間加工は、８５０℃にて２時間加熱後、プレス鍛造することにより行った。これにより、８０ｍｍ角の角片を作製した。
また、実施例１〜８及び比較例１〜４の８０ｍｍ角の角片の割れ長さの計測を行った。この計測は、以下の手順で行った。８０ｍｍ角の横断面より、任意の個所から１５ｍｍ角（１辺は表層）を切り出し、これを顕微鏡観察用試料として調整し、鏡面研磨及び腐食後、光学顕微鏡にて表層からの割れ個所を観察した。存在する割れ個所に対して、最も長いものから５個所の長さを計測し、その平均値を「割れ長さ」と定義した。
その結果を表２に示す。実施例１〜８は、いずれも割れ長さが０．７０ｍｍ以下となったのに対して、比較例１〜４は、いずれも割れ長さが０．７０ｍｍを超えた。表２では、割れ長さが０．７０ｍｍ以下のものを「○」で示し、割れ長さが０．２５ｍｍ以下のものを「◎」で示し、割れ長さが０．７０ｍｍを超えたものを「×」で示した。
また、図１に、実施例２及び比較例２の割れ個所を光学顕微鏡で撮影した組織写真（５０倍）を示す。同図に示したように、実施例２は、比較例２に比して、結晶粒が微細化され、割れ長さが低減されていることがわかった。
Next, hot working was performed on each obtained steel ingot. Hot working was performed by press forging after heating at 850 ° C. for 2 hours. Thereby, an 80 mm square piece was produced.
Moreover, the measurement of the crack length of the 80 mm square piece of Examples 1-8 and Comparative Examples 1-4 was performed. This measurement was performed according to the following procedure. From a cross section of 80 mm square, a 15 mm square (one side is the surface layer) was cut out from an arbitrary location, this was adjusted as a sample for microscope observation, and after the mirror polishing and corrosion, the cracked portion from the surface layer was observed with an optical microscope. . From the longest one, the length of the five existing cracks was measured, and the average value was defined as the “crack length”.
The results are shown in Table 2. In all of Examples 1 to 8, the crack length was 0.70 mm or less, while in Comparative Examples 1 to 4, the crack length exceeded 0.70 mm. In Table 2, those having a crack length of 0.70 mm or less are indicated by “◯”, those having a crack length of 0.25 mm or less are indicated by “◎”, and those having a crack length exceeding 0.70 mm. Indicated by “x”.
Moreover, the structure photograph (50 times) which image | photographed the crack part of Example 2 and Comparative Example 2 with the optical microscope in FIG. 1 is shown. As shown in the figure, it was found that in Example 2, the crystal grains were refined and the crack length was reduced as compared with Comparative Example 2.
まず、固溶化処理の有無のみが異なる比較例１と２とを比較すると、固溶化処理を行った方が割れ長さが短かったことから、固溶化処理が熱間鍛造時における割れ抑制に有効であることが確認できた。その理由は、固溶化処理によって、組織が均質化され、ＭｎとＣｕのマクロ偏析が低減されるためと考えられる。
次に、Ｃａの有無が異なる実施例１，２，４と比較例１とを比較すると、Ｃａを添加した方が割れ長さが短かったことから、Ｍｎ−Ｃｕ−Ｎｉ−Ｆｅを基本組成とする合金にＣａを添加すると、熱間鍛造時における割れ抑制に有効であることがわかった。これは、Ｃａ添加により結晶粒の微細化が促進されたためと考えられる（Ｃａの有無が異なる実施例５と比較例３との比較（但し、これらは、Ｍｎ−Ｃｕを基本組成とする）及びＣａの有無が異なる実施例６と比較例４との比較についても同様である）。
次に、Ｍｇの有無が異なる実施例７、８と比較例１とを比較すると、Ｍｇを添加した方が割れ長さが短かったことから、Ｍｎ−Ｃｕ−Ｎｉ−Ｆｅを基本組成とする合金にＭｇを添加すると、熱間鍛造時における割れ抑制に有効であることがわかった。これは、Ｍｇ添加により結晶粒の微細化が促進されたためと考えられる。
次に、保持時間のみが異なる実施例２と実施例３とを比較すると、保持時間が長い方が割れ長さが短かった。これは、保持時間が長い方が、ＭｎとＣｕのマクロ偏析が低減されることによる不均一応力の低減や、結晶粒の微細化による強度の改善が図られたためと考えられる。
First, comparing Comparative Examples 1 and 2 that differ only in the presence or absence of the solution treatment, the crack length was shorter when the solution treatment was performed, so the solution treatment was effective in suppressing cracks during hot forging. It was confirmed that. The reason is considered to be that the structure is homogenized by the solution treatment and the macrosegregation of Mn and Cu is reduced.
Next, when Examples 1, 2, 4 and Comparative Example 1 with different presence or absence of Ca were compared with each other, since the crack length was shorter when Ca was added, Mn—Cu—Ni—Fe was used as the basic composition. It was found that adding Ca to the alloy to be used is effective in suppressing cracking during hot forging. This is considered to be because the refinement of crystal grains was promoted by the addition of Ca (comparison between Example 5 and Comparative Example 3 in which Ca is different or not (however, these are based on Mn—Cu) and The same applies to the comparison between Example 6 and Comparative Example 4 in which the presence or absence of Ca is different.
Next, when Examples 7 and 8 differing in the presence or absence of Mg and Comparative Example 1 were compared, the crack length was shorter when Mg was added, so an alloy having a basic composition of Mn—Cu—Ni—Fe. It has been found that the addition of Mg to is effective in suppressing cracking during hot forging. This is presumably because the refinement of crystal grains was promoted by the addition of Mg.
Next, when Example 2 and Example 3 in which only the holding time is different are compared, the longer the holding time, the shorter the crack length. This is probably because the longer the holding time, the reduction of non-uniform stress due to the reduction of Mn and Cu macrosegregation and the improvement of the strength due to the refinement of crystal grains.
（損失係数について）
次に、実施例１〜８及び比較例１〜４の８０ｍｍ角の角片の中心部から厚さ２０ｍｍ×幅２０ｍｍ×長さ２００ｍの角片を切り出し、各角片を水素雰囲気下、９００℃に加熱し、この温度で３時間保持した後、冷却速度１．５℃／分で各角片が３００℃になるまで定速徐冷（空冷）した（表３参照）。更に、各角片が１００℃以下になるまで水冷により急冷した。
(About loss factor)
Next, square pieces having a thickness of 20 mm, a width of 20 mm, and a length of 200 m were cut out from the center of the 80 mm square pieces of Examples 1 to 8 and Comparative Examples 1 to 4, and each square piece was 900 ° C. in a hydrogen atmosphere. And kept at this temperature for 3 hours, and then cooled at a constant rate of 1.5 ° C./min until each square piece reached 300 ° C. (air cooling) (see Table 3). Furthermore, it cooled rapidly by water cooling until each square piece became 100 degrees C or less.
次に、各角片から放電加工により厚さ１ｍｍ×幅１０ｍｍ×長さ１６０ｍｍの角片を試験片として切り出した。切り出した試験片を用いて、「ＪＩＳ Ｇ０６０２」に準拠した中央加振法による減衰特性を測定した。減衰特性の測定は次のようにして行った。
まず、試験片の１次共振周波数を測定し、その周波数において振幅ひずみが１×１０−３となるバースト正弦波を加振した場合の振動減衰波形を測定した。次に、得られた減衰波形をフーリエ変換し、周波数分布を求め、半値幅法により、ピーク高さが半分となる範囲Δｆとピーク周波数ｆ０とにより損失係数＝Δｆ／（１．７３２ｆ０）を求めた。尚、加振にはＥＭＩＣ社製の電磁型加振器を用い、振幅ひずみの測定には小野測器社製のＣＦ−５２００型ＦＦＴアナライザーを用いた。その結果を表３に示す。比較例２を除き、いずれも損失係数が０．１２以上となり、良好な結果が得られた。表２では、損失係数が０．１２以上のものを「○」で示し、損失係数が０．１６以上のものを「◎」で示し、損失係数が０．１１以下のものを「×」で示した。
Next, a square piece having a thickness of 1 mm, a width of 10 mm, and a length of 160 mm was cut out from each square piece as a test piece by electric discharge machining. Using the cut out test piece, the attenuation characteristic by the central excitation method based on “JIS G0602” was measured. The attenuation characteristics were measured as follows.
First, the primary resonance frequency of the test piece was measured, and a vibration attenuation waveform was measured when a burst sine wave having an amplitude distortion of 1 × 10 −3 was applied at that frequency. Next, the obtained attenuation waveform is subjected to Fourier transform to obtain a frequency distribution, and a loss factor = Δf / (1.732f 0 ) by a half-width method with a range Δf in which the peak height is halved and a peak frequency f 0. Asked. An electromagnetic exciter manufactured by EMIC was used for excitation, and a CF-5200 type FFT analyzer manufactured by Ono Sokki Co., Ltd. was used for measurement of amplitude distortion. The results are shown in Table 3. Except for Comparative Example 2, the loss factor was 0.12 or more, and good results were obtained. In Table 2, “○” indicates that the loss factor is 0.12 or more, “◎” indicates that the loss coefficient is 0.16 or more, and “×” indicates that the loss coefficient is 0.11 or less. Indicated.
まず、固溶化処理の有無のみが異なる比較例１と２とを比較すると、固溶化処理を行った方が損失係数が高かったことから、固溶化処理を行っておくと、良い制振特性が得られることが確認できた。その理由は、固溶化処理で組織が均質化され、ＭｎとＣｕのマクロ偏析が低減され、高温焼鈍で双晶が十分に形成されたためと考えられる。
次に、Ｃａの有無が異なる実施例１〜２と比較例１とを比較すると損失係数に差がなかった。同様にＣａの有無が異なる実施例４と比較例１とを比較すると、逆に、比較例１の方が損失係数が高かった。固溶化処理での保持時間に差がある実施例２と実施例３とを比較すると、実施例３の方が損失係数が高かった。これらのことから、Ｃａ量を調整し、保持時間を長くすれば、損失係数が高くなり、制振特性に効果があることがわかった。
一方、Ｃａの有無が異なる実施例５と比較例３との比較や、Ｍｇの有無が異なる実施例７と比較例１との比較では損失係数に差がなかった。また、Ｃａの有無が異なる実施例６と比較例４との比較や、Ｍｇの有無が異なる実施例８と比較例１との比較では、これらの組成では、Ｍｇ又はＣａを添加しても損失係数に影響を与えない（損失係数を低下させない）ことがわかった。
First, comparing Comparative Examples 1 and 2 that differ only in the presence or absence of the solution treatment, the loss factor was higher when the solution treatment was performed. Therefore, when the solution treatment was performed, good damping characteristics were obtained. It was confirmed that it was obtained. The reason is considered to be that the structure is homogenized by the solution treatment, macro segregation of Mn and Cu is reduced, and twins are sufficiently formed by high-temperature annealing.
Next, when Examples 1 and 2 and Comparative Example 1 with different presence or absence of Ca were compared, there was no difference in loss factor. Similarly, when Example 4 and Comparative Example 1 with different presence or absence of Ca were compared, the loss factor of Comparative Example 1 was higher. When Example 2 and Example 3 having a difference in retention time in the solution treatment were compared, Example 3 had a higher loss factor. From these facts, it was found that if the Ca amount is adjusted and the holding time is lengthened, the loss factor increases and the damping characteristics are effective.
On the other hand, there was no difference in loss factor between the comparison between Example 5 and Comparative Example 3 with different presence or absence of Ca, and the comparison between Example 7 and Comparative Example 1 with different presence or absence of Mg. Moreover, in the comparison between Example 6 and Comparative Example 4 in which the presence or absence of Ca is different, or in the comparison between Example 8 and Comparative Example 1 in which the presence or absence of Mg is different, in these compositions, even if Mg or Ca is added, loss is caused. It has been found that the coefficient is not affected (the loss coefficient is not reduced).
表２及び表３の結果を総合すると、実施例１〜８は、割れ長さ及び損失係数の両者が良好な値を示したが、比較例１〜４は、特に割れ長さの点で実施例に劣ることがわかった。このことから、Ｃａ又はＭｇを含有させることにより、制振性を維持しつつ、成形加工性を高めることに効果があることがわかった。 When the results of Tables 2 and 3 are combined, Examples 1 to 8 show good values for both the crack length and loss factor, but Comparative Examples 1 to 4 are particularly implemented in terms of crack length. I found it inferior to the examples. From this, it was found that inclusion of Ca or Mg is effective in improving the moldability while maintaining the vibration damping property.
以上、本発明の実施の形態について詳細に説明したが、本発明は上記実施の形態に何ら限定されるものではなく、本発明の要旨を逸脱しない範囲内で種々の改変が可能である。 Although the embodiments of the present invention have been described in detail above, the present invention is not limited to the above embodiments, and various modifications can be made without departing from the scope of the present invention.
本発明に係るＭｎ−Ｃｕ系制振合金及びその製造方法は、制振特性が要求される機器に用いられる各種の部品（例えば、ネジ、ワッシャー、インシュレータ、台座、バネ、バイトホルダー、軸受等）の材料やその製造方法として用いることができる。 The Mn-Cu vibration damping alloy and the manufacturing method thereof according to the present invention are various parts (for example, screws, washers, insulators, pedestals, springs, tool holders, bearings, etc.) used in equipment that requires vibration damping characteristics. It can be used as a material and a manufacturing method thereof.
Claims (8)
更に、Ｃａ：０．００１〜０．００７ａｔ％又はＭｇ：０．００５〜０．２３ａｔ％を含み、
残部がＭｎ及び不可避的不純物からなることを特徴とするＭｎ−Ｃｕ系制振合金。 Cu: 15-25 at%,
Furthermore, Ca: 0.001-0.007 at% or Mg: 0.005-0.23 at%,
A Mn-Cu vibration damping alloy, characterized in that the balance consists of Mn and inevitable impurities.
前記加工工程の後に、前記合金を８００℃以上固相線温度以下の焼鈍温度に加熱し、前記焼鈍温度で０．５時間以上８時間以下保持した後、前記合金の温度が２５０℃〜４５０℃の範囲のいずれかに入るまで、０．８５〜１．６５℃／分で定速徐冷する高温焼鈍工程と、
前記高温焼鈍工程の後に、前記合金を１０℃／分以上で冷却する急冷工程とを備えたことを特徴とする請求項４又は５に記載のＭｎ−Ｃｕ系制振合金の製造方法。 After the solution treatment step, a processing step for processing the alloy to a final product shape or a shape larger than this,
After the processing step, the alloy is heated to an annealing temperature of 800 ° C. or more and a solidus temperature or less and held at the annealing temperature for 0.5 hours or more and 8 hours or less, and then the temperature of the alloy is 250 ° C. to 450 ° C. Until it enters any of the range of high temperature annealing step of constant cooling at 0.85 to 1.65 ° C./min,
The method for producing a Mn—Cu vibration damping alloy according to claim 4, further comprising a quenching step of cooling the alloy at 10 ° C./min or more after the high-temperature annealing step.
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