JPWO2007066555A1 - Co-based alloy and manufacturing method thereof - Google Patents
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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Abstract
Ｆｅ：０．０１〜１０％，Ｎｉ：０．０１〜３０％，Ｍｎ：０．０１〜２５％の一種又は二種以上を含むＣｏ基合金であり、熱誘起又は応力誘起変態で生じたｈ．ｃ．ｐ．構造のε相が１０体積％以上生成した金属組織を有する。必要に応じＡｌ：０．０１〜１０％，Ｃｒ：０．０１〜３５％，Ｖ：０．０１〜２０％，Ｔｉ：０．０１〜１５％，Ｍｏ：０．０１〜３０％，Ｎｂ：０．０１〜１０％，Ｚｒ：０．０１〜３％，Ｗ：０．０１〜３０％，Ｔａ：０．０１〜１０％，Ｈｆ：０．０１〜５％，Ｓｉ：０．０１〜８％，Ｃ：０．００１〜３％，Ｂ：０．００１〜３％，Ｐ：０．００１〜３％，ミッシュメタル：０．００１〜３％の一種又は二種以上を添加しても良い。高弾性変形能を示し、延性，加工性も良好であり、磁場印加により変位制御できるアクチュエータ，センサー等の機能材料として使用される。Co-based alloy containing one or more of Fe: 0.01 to 10%, Ni: 0.01 to 30%, Mn: 0.01 to 25%, and h generated by heat-induced or stress-induced transformation . c. p. It has a metal structure in which the ε phase of the structure is formed by 10% by volume or more. If necessary, Al: 0.01 to 10%, Cr: 0.01 to 35%, V: 0.01 to 20%, Ti: 0.01 to 15%, Mo: 0.01 to 30%, Nb: 0.01 to 10%, Zr: 0.01 to 3%, W: 0.01 to 30%, Ta: 0.01 to 10%, Hf: 0.01 to 5%, Si: 0.01 to 8 %, C: 0.001 to 3%, B: 0.001 to 3%, P: 0.001 to 3%, Misch metal: 0.001 to 3%, or one or more of them may be added. . It exhibits high elastic deformability, has good ductility and workability, and is used as a functional material for actuators, sensors, etc. that can control displacement by applying a magnetic field.
Description
本発明は、弾性変形能が高く、磁場印加で変位制御できる強磁性Ｃｏ基合金及びその製造方法に関する。 The present invention relates to a ferromagnetic Co-based alloy having high elastic deformability and capable of displacement control by application of a magnetic field, and a method for producing the same.
低ヤング率で高弾性変形能を示す材料としてチタン合金が注目されており、人工歯根，人工骨，眼鏡フレーム等に使用されている。たとえば、文献１，２では、４Ａ族や５Ａ族元素を含むチタン合金が低ヤング率で高弾性変形能材料になることを紹介している。
材料の変形をもたらす外部因子には、文献１，２で説明されている外部応力の外に温度，磁場等がある。
温度による変位制御には形状記憶合金が知られており、数％レベルの寸法変化が得られている。形状記憶効果は、変形した材料を所定温度以上に加熱したときに生じるマルテンサイト逆変態を利用して元の形状が回復する現象である。形状記憶効果を用いると熱駆動型のアクチュエータとして利用できるが、温度制御が必要なことに加え、冷却時の形状変化が熱拡散で律速されるため応答性が悪い。
強磁性の形状記憶合金もアクチュエータ材料として注目されている。強磁性の形状記憶合金では、従来の磁歪材料を超える数％の寸法変化が外部磁場の印加で得られ、熱駆動型形状記憶合金の欠点である低応答性も解消されている。強磁性型形状記憶合金としては、たとえばＮｉ−Ｍｎ−Ｇａ系があり、磁場印加で形状変化を引き起こすアクチュエータ材料が文献３に紹介されている。Ｃｏ−Ｎｉ−Ａｌ系も文献４，５に、Ｃｏ系合金も文献６に紹介されている。
しかし、Ｎｉ−Ｍｎ−Ｇａ系の材料は延性に劣り、機械部品に要求される複雑で精密な形状を付与し難い。Ｃｏ−Ｎｉ−Ａｌ系合金は、γ相を第二相として利用して延性を改善しているが、磁化の強さが小さい。また、Ｃｏ系合金は、延性，磁化の強さは比較的優れ、熱駆動型の形状記憶効果が得られるものの、磁歪特性が不十分で超弾性特性も得られない。
文献１：特開２００２−３３２５３１号公報
文献２：特開２００２−２４９８３６号公報
文献３：米国特許５，９５８，１５４号明細書
文献４：特開２００２−１２９２７３号公報
文献５：特開２００４−２７７８６５号公報
文献６：特開２００４−２３８７２０号公報Titanium alloys are attracting attention as a material having a low Young's modulus and high elastic deformability, and are used for artificial tooth roots, artificial bones, eyeglass frames, and the like. For example, Documents 1 and 2 introduce that a titanium alloy containing a 4A group element or a 5A group element becomes a highly elastic deformable material with a low Young's modulus.
In addition to the external stress described in Documents 1 and 2, external factors that cause deformation of the material include temperature and magnetic field.
A shape memory alloy is known for displacement control by temperature, and a dimensional change of several percent level is obtained. The shape memory effect is a phenomenon in which the original shape is recovered by utilizing a martensite reverse transformation that occurs when a deformed material is heated to a predetermined temperature or higher. If the shape memory effect is used, it can be used as a heat-driven actuator, but in addition to requiring temperature control, the shape change during cooling is limited by thermal diffusion, so the response is poor.
Ferromagnetic shape memory alloys are also attracting attention as actuator materials. In a ferromagnetic shape memory alloy, a dimensional change of several percent exceeding that of a conventional magnetostrictive material can be obtained by applying an external magnetic field, and the low response, which is a drawback of the heat-driven shape memory alloy, is also eliminated. As the ferromagnetic shape memory alloy, for example, there is a Ni—Mn—Ga system, and an actuator material that causes a shape change by applying a magnetic field is introduced in Reference 3. Co-Ni-Al system is also introduced in References 4 and 5, and Co-based alloy is also introduced in Reference 6.
However, Ni—Mn—Ga-based materials are inferior in ductility, and it is difficult to give complex and precise shapes required for mechanical parts. The Co—Ni—Al-based alloy uses the γ phase as the second phase to improve ductility, but has a low magnetization strength. Co-based alloys are relatively excellent in ductility and magnetization strength, and provide a heat-driven shape memory effect, but have insufficient magnetostriction characteristics and do not provide superelastic characteristics.
Document 1: Japanese Patent Application Laid-Open No. 2002-325331 Document 2: Japanese Patent Application Laid-Open No. 2002-249836 Document 3: US Patent No. 5,958,154 Document 4: Japanese Patent Application Laid-Open No. 2002-129273 Document 5: Japanese Patent Application Laid-Open No. 2004-294273 Publication No. 277865 Publication 6: JP 2004-238720 A
本発明者等は、従来のチタン合金やＮｉ−Ｍｎ−Ｇａ系合金の欠点に鑑み、高弾性変形能を維持しながら強磁性で磁場駆動でき、延性も良好な材料を種々調査・検討した。その結果、Ｃｏをベースとし、合金成分，組成の適切な選択及びｈ．ｃ．ｐ．構造を有するε相の適量生成により、高弾性変形能を有するＣｏ基合金が得られることを見出した。
本発明は、かかる知見をベースとし、Ｆｅ，Ｎｉ，Ｍｎの一種又は二種以上を添加した成分系においてε相の生成量を制御することにより、高弾性変形能を有し延性，加工性も良好なＣｏ基合金を提供することを目的とする。
本発明のＣｏ基合金は、質量比でＦｅ：０．０１〜１０％，Ｎｉ：０．０１〜３０％，Ｍｎ：０．０１〜２５％の一種又は二種以上を含んでいる。Ｆｅ，Ｎｉ，Ｍｎの二種又は三種を添加する場合、好ましくは合計含有量を０．０２〜５０％の範囲に設定する。以下、質量比を単に％と表記する。
更に、Ｆｅ，Ｎｉ，Ｍｎ以外に、Ａｌ：０．０１〜１０％，Ｃｒ：０．０１〜３５％，Ｖ：０．０１〜２０％，Ｔｉ：０．０１〜１５％，Ｍｏ：０．０１〜３０％，Ｎｂ：０．０１〜１０％，Ｚｒ：０．０１〜３％，Ｗ：０．０１〜３０％，Ｔａ：０．０１〜１０％，Ｈｆ：０．０１〜５％，Ｓｉ：０．０１〜８％，Ｃ：０．００１〜３％，Ｂ：０．００１〜３％，Ｐ：０．００１〜３％，ミッシュメタル：０．００１〜３％から選ばれた一種又は二種以上を合計含有量：０．００１〜５０％の範囲で含むこともできる。
提案したＣｏ基合金は少なくとも常温で強磁性であり、熱誘起又は応力誘起で生成したｈ．ｃ．ｐ．構造のε相が分布している。ε相の割合は１０体積％以上とされ、成分調整，製造条件等で制御できる。
ｈ．ｃ．ｐ．構造のε相が分布した金属組織は、所定組成のＣｏ基合金を９００〜１４００℃で溶体化処理することにより形成される。溶体化処理後、加工率：１０％以上で加工しても良く、更に３００〜８００℃で時効処理しても良い。In view of the shortcomings of conventional titanium alloys and Ni—Mn—Ga alloys, the present inventors have investigated and examined various materials that can be driven by a magnetic field with ferromagnetism and have good ductility while maintaining high elastic deformability. As a result, based on Co, appropriate selection of alloy components, composition and h. c. p. It has been found that a Co-based alloy having a high elastic deformability can be obtained by generating an appropriate amount of an ε phase having a structure.
The present invention is based on such knowledge, and by controlling the amount of ε phase generated in a component system to which one or more of Fe, Ni, and Mn are added, it has high elastic deformability and ductility and workability. An object is to provide a good Co-based alloy.
The Co-based alloy of the present invention contains one or more of Fe: 0.01 to 10%, Ni: 0.01 to 30%, and Mn: 0.01 to 25% by mass ratio. When adding 2 or 3 types of Fe, Ni and Mn, the total content is preferably set in the range of 0.02 to 50%. Hereinafter, the mass ratio is simply expressed as%.
Furthermore, in addition to Fe, Ni and Mn, Al: 0.01 to 10%, Cr: 0.01 to 35%, V: 0.01 to 20%, Ti: 0.01 to 15%, Mo: 0. 01-30%, Nb: 0.01-10%, Zr: 0.01-3%, W: 0.01-30%, Ta: 0.01-10%, Hf: 0.01-5%, Si: 0.01 to 8%, C: 0.001 to 3%, B: 0.001 to 3%, P: 0.001 to 3%, Misch metal: 0.001 to 3% Or 2 or more types can also be included in total content: 0.001 to 50% of range.
The proposed Co-based alloy is ferromagnetic at least at room temperature, and is formed by heat induction or stress induction. c. p. The ε phase of the structure is distributed. The ratio of the ε phase is 10% by volume or more, and can be controlled by adjusting the components, manufacturing conditions, and the like.
h. c. p. The metal structure in which the ε phase of the structure is distributed is formed by solution treatment of a Co-based alloy having a predetermined composition at 900 to 1400 ° C. After the solution treatment, the processing rate may be processed at 10% or more, and further, an aging treatment may be performed at 300 to 800 ° C.
図１は、Ｃｏ−Ｆｅ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼすＦｅ含有量の影響を表したグラフ
図２は、Ｃｏ−Ｎｉ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼすＮｉ含有量の影響を表したグラフ
図３は、Ｃｏ−Ｍｎ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼすＭｎ含有量の影響を表したグラフ
図４は、ｈ．ｃ．ｐ．構造のε相が生成した金属組織をもつＣｏ−３．８０％Ｆｅ合金の顕微鏡写真
図５は、同Ｃｏ−３．８０％Ｆｅ合金の応力−歪線図
図６は、Ｃｏ−Ｆｅ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす溶体化処理条件の影響を表したグラフ
図７は、Ｃｏ−Ｆｅ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす冷間加工率の影響を表したグラフ
図８は、Ｃｏ−Ｆｅ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす時効処理条件の影響を表したグラフ
図９は、Ｃｏ−Ｎｉ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす溶体化処理条件の影響を表したグラフ
図１０は、Ｃｏ−Ｎｉ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす冷間加工率の影響を表したグラフ
図１１は、Ｃｏ−Ｎｉ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす時効処理条件の影響を表したグラフ
図１２は、Ｃｏ−Ｍｎ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす溶体化処理条件の影響を表したグラフ
図１３は、Ｃｏ−Ｍｎ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす冷間加工率の影響を表したグラフ
図１４は、Ｃｏ−Ｍｎ二元合金の金属組織に占めるε相の体積分率，回復歪量，磁化の強さに及ぼす時効処理条件の影響を表したグラフ
図１５は、Ｃｏ−３．８０％Ｆｅ合金の試験温度に応じた応力−歪線図FIG. 1 is a graph showing the influence of Fe content on the volume fraction of ε phase, the recovery strain amount, and the strength of magnetization in the metal structure of the Co—Fe binary alloy. FIG. 3 is a graph showing the effect of Ni content on the volume fraction, recovery strain amount, and strength of magnetization of the ε phase in the metal structure of the alloy. FIG. 3 shows the ε phase in the metal structure of the Co—Mn binary alloy. FIG. 4 is a graph showing the influence of the Mn content on the volume fraction, the recovery strain amount, and the strength of magnetization. c. p. FIG. 5 is a stress-strain diagram of the Co-3.80% Fe alloy. FIG. 6 is a Co—Fe binary. FIG. 7 is a graph showing the effect of solution treatment conditions on the volume fraction, recovery strain amount, and magnetization strength of the ε phase in the alloy metal structure. FIG. 7 shows the ε phase in the Co—Fe binary alloy metal structure. FIG. 8 is a graph showing the influence of the cold work rate on the volume fraction, recovery strain amount, and strength of magnetization. FIG. 8 shows the volume fraction of ε phase and the recovery strain amount in the metal structure of the Co—Fe binary alloy. FIG. 9 is a graph showing the influence of aging treatment conditions on the strength of magnetization. FIG. 9 is a solution treatment on the volume fraction of ε phase, recovery strain amount, and magnetization strength in the metal structure of a Co—Ni binary alloy. FIG. 10 is a graph showing the influence of processing conditions. FIG. 10 shows the volume fraction and recovery of the ε phase in the metal structure of the Co—Ni binary alloy. FIG. 11 is a graph showing the effect of the cold work rate on the amount and the strength of magnetization. FIG. 11 shows the effect on the volume fraction of ε phase, the recovery strain amount, and the strength of magnetization in the metal structure of the Co—Ni binary alloy. FIG. 12 shows the influence of solution treatment conditions on the volume fraction of ε phase, recovery strain amount, and strength of magnetization in the metal structure of the Co—Mn binary alloy. Graph FIG. 13 is a graph showing the influence of the cold work rate on the volume fraction of ε phase, the recovery strain amount, and the strength of magnetization in the metal structure of the Co—Mn binary alloy. FIG. 15 is a graph showing the influence of aging treatment conditions on the volume fraction of ε phase, recovery strain amount, and strength of magnetization in the metal structure of the binary alloy. FIG. 15 shows the test temperature of Co-3.80% Fe alloy. Corresponding stress-strain diagram
本発明者等は、弾性変形能が高く、磁場勾配の印加・除去により変位制御でき加工可能な材料を開発するため、強磁性元素であるＣｏに種々の合金元素を添加し、組織と弾性変形能，磁気特性との関係を調査・検討した。その結果、適量のＦｅ，Ｎｉ，Ｍｎの少なくとも一種又は二種以上をＣｏと合金化させると、加工性が改善された強磁性高弾性変形合金となることを解明した。
強磁性材料は、Ｆ＝−Ｍ（ｄＨ／ｄｘ）の関係式に従い磁場勾配に応じた力を受けて変位することが知られている。この関係式は、所定の磁場勾配ｄＨ／ｄｘで得られる力Ｆが磁化の強さＭに比例することを意味する。材料に与える力Ｆが大きいほど得られる歪量が大きくなるため、歪量の増大には磁化の強さＭが大きい方がよく、変形に要する応力が高いと磁場による歪量が小さくなるためヤング率は低いことが望ましい。しかし、弾性変形能が小さいと磁場勾配の除去により歪が残留する。そのため、磁場勾配の印加・除去により大きな可逆歪を得るには、大きな弾性変形能が必要となる。
Ｃｏは高いキュリー温度を有する強磁性元素であるが、純Ｃｏは加工性に乏しく、磁化率も低い。Ｃｏは高温でｆ．ｃ．ｃ．構造のγ相であり、冷却中に４２０℃付近でｈ．ｃ．ｐ．構造のε相にマルテンサイト変態する。Ｃｏに種々の合金元素を添加してγ相とε相の相安定性を調査した結果、γ相を安定化し、加工性を向上させる種々の元素を特定できた。
γ相は不規則構造のｆ．ｃ．ｃ．のため加工性が良好であるが、ε相は弾性変形量を大きくする効果を有する相であるものの加工硬化しやすく、加工割れ等の欠陥を発生させやすい。そこで、γ相に固溶して安定化させる元素を添加すると常温でもγ相が残留しＣｏ基合金の加工性向上を期待できる。なかでも、磁気モーメントをもつＦｅ，Ｎｉ，Ｍｎは、加工性向上の外に、Ｃｏ基合金の磁気特性改善にも有効である。
成分調整，製造条件等でε相の体積分率を１０体積％以上に調整すると、高弾性変形の効果が顕著になる。ただし、ε相の過剰分布は残留γ相の加工性向上効果を相殺するので、ε相の上限を好ましくは９９体積％とする。基本的にはｈ．ｃ．ｐ．構造のε相と残りがｆ．ｃ．ｃ．構造のγ相の複相組織を有するＣｏ基合金であるが、高弾性特性，磁気特性，加工性等に悪影響を及ぼさない限り異相の存在を許容する。異相には、第三成分の添加で生じる金属間化合物，炭化物，α相（ｂ．ｃ．ｃ．構造）等がある。
Ｆｅを添加したＣｏ合金は特に優れた磁気特性を示す。磁場印加による変位制御素子としての用途では比較的低い磁場で大きな磁化の強さを示すことが好ましいが、本発明Ｃｏ基合金の磁化曲線を測定すると、０．２Ｔ（テスラ）の外部磁場に対し３０ｅｍｕ／ｇ以上の磁化の強さを示し、変位制御素子として有望な材料といえる。Ｎｉ，Ｍｎは、磁化の強さを弱める作用を呈するものの、依然として磁化の強さが高レベルに保たれるので変位制御素子としての要求特性が付与される。
一方、組織と弾性変形挙動との関係を調査した結果、次のようなことが解明された。
焼鈍されたＦｅやＣｏ等の金属材料の弾性限歪は通常０．３％程度であり、冷間加工を施すと加工硬化して辷り変形が抑制され、硬度，引張り強さの上昇と共に降伏応力，弾性限が高くなる。他方、本発明Ｃｏ基合金は、たとえば１％の曲げ変形に対し通常の弾性歪量を超える０．４％以上の歪が回復する。しかも、液体窒素温度（−１９６℃）から４００℃までの広い温度範囲で適用できる。
しかし、変形に要する応力が高いと磁場による変位が得にくくなるためヤング率は低い方が望ましい。ヤング率は原子間の凝集力に関わる物性値であり、加工や熱処理では制御し難いと考えられている。この点、本発明では、マルテンサイト変態時の格子軟化，マルテンサイトバリアントの方位性等を制御することにより、低ヤング率を達成している。
マルテンサイト変態では、変態に伴う前駆現象の一つとしてマルテンサイト変態温度から数１０〜１００℃程度の幅をもった温度域で弾性定数が低下する格子軟化（ｌａｔｔｉｃｅ ｓｏｆｔｅｎｉｎｇ）現象がみられる。格子軟化をＣｏ基合金に利用すると、マルテンサイト変態温度Ｍｓの近傍でヤング率の低下が予測される。
純Ｃｏではマルテンサイト変態による格子軟化が４２０℃近傍で生じるが、Ｆｅ，Ｎｉ，Ｍｎの添加で変態温度が低下する。変態温度の低下に伴い、所望の温度範囲において低いヤング率が得られる。更に、Ｃｏ−Ｘ（Ｘ：Ｆｅ，Ｎｉ，Ｍｎ）系のマルテンサイト変態は非熱弾性型であるためヒステリシスが約１５０℃であり、加工によりヒステリシスが更に広がるので、広い温度幅で低ヤング率を実現するのに適している。
Ｃｏ基合金のマルテンサイト変態では、Ｆｅ合金等と同様に十分にＭｓ温度以下に冷却しても試料全体がマルテンサイト相に変態するわけではなく、母相がある程度残留する。そのため、溶体化処理後の加工によってマルテンサイト相が応力誘起されるが、応力に対してＳｃｈｍｉｄ因子の最も大きなバリアントが優先的に生成される。また、熱的に誘起されたマルテンサイト相に応力を加えると、マルテンサイト相の一部が応力方向に優先的なバリアントに再配列される。これらの優先的バリアントが応力方向に対し低い変形応力を示すことも、Ｃｏ基合金が低ヤング率を示す原因の一つと考えられる。
Ｃｏ−Ｎｉ−Ａｌ系合金（文献４，５）は、β相又はβ＋γ相の組織を有し、β相のマルテンサイト変態・逆変態を形状記憶特性付与に利用しているが、γ相がε相に変態することはない。延性に乏しいＢ２構造のβ相が存在することは加工性低下の原因にもなるが、本発明のＣｏ基合金はβ相が存在しない点でも延性に優れた材料といえる。しかも、強磁性であるものの磁化の強さが小さいＣｏ−Ｎｉ−Ａｌ系合金に比較し、本発明のＣｏ基合金は磁化の強さが極めて大きいため、磁場勾配による変位制御を行う上で優れた性能を発揮する。
形状記憶特性を有する文献６のＣｏ基合金では、形状記憶効果に優れているが超弾性が得られず、極めて小さな磁歪のために磁場による変位制御も困難である。これは、形状記憶効果，超弾性，磁場による変位制御がそれぞれ異なる現象に由来する結果である。すなわち、形状記憶効果は、マルテンサイト相状態で変形を与え、母相にマルテンサイト逆変態させることにより形状回復する現象である。超弾性は、母相状態で応力を与えると応力誘起のマルテンサイト変態をし、応力を除荷すると母相へマルテンサイト逆変態するために形状回復する現象である。他方、強磁性形状記憶合金における大きな磁歪は、均一磁場によるマルテンサイトバリアントの再配列や磁場誘起マルテンサイト変態により得られる現象である。
本特許のＣｏ基合金は、主に加工硬化により大きな弾性変形能を実現したものであり、応力の負荷・除荷によるマルテンサイト相の可逆的な変態・逆変態が明確に検出されない点で超弾性と本質的に異なり、マルテンサイト相（ε相）はヤング率を低く保ちながら弾性変形能を補助するのみである。更に、加工性が改善され、磁化の強さも増大しているので、磁場勾配の印加・除去による変位制御が可能な材料として優れた機能を発揮する。
本発明のＣｏ基合金は、Ｆｅ：０．０１〜１０％，Ｎｉ：０．０１〜３０％，Ｍｎ：０．０１〜２５％から選ばれた一種又は二種以上を含むＣｏ合金を基本とする。
Ｆｅ，Ｎｉ，Ｍｎは、マルテンサイト変態温度を低下させ、延性，加工性の向上に寄与し、磁化率を高める効果を奏する。このような効果は、０．０１％以上のＦｅ，Ｎｉ又はＭｎ添加で顕著になる。しかし、過剰添加はマルテンサイト及び磁気変態温度を室温以下に下げ、ε相の生成抑制，磁気特性の低下を招くので、Ｆｅ：１０％，Ｎｉ：３０％，Ｍｎ：２５％を上限とした。Ｆｅ，Ｎｉ，Ｍｎの二種又は二種を添加する場合、合計含有量を０．０２〜５０％の範囲に収めることが好ましい。Ｆｅ，Ｎｉ，Ｍｎそれぞれの含有量は、好ましくは１〜８％，１〜２５％，１〜２０％、更に好ましくは２〜６％，５〜２０％，５〜１５％の範囲で定める。
Ａｌ，Ｃｒ，Ｖ，Ｔｉ，Ｍｏ，Ｎｂ，Ｚｒ，Ｗ，Ｔａ，Ｈｆ，Ｓｉ，Ｃ，Ｂ，Ｐ，ミッシュメタルから選ばれた一種又は二種以上の第三成分を必要に応じてＣｏ−（Ｆｅ，Ｍｎ，Ｎｉ）系に添加できる。複数の第三成分を添加する場合には、０．００２〜５０％（好ましくは、０．００５〜３０％）の範囲で合計含有量を選定する。
Ａｌ，Ｖ，Ｔｉは、マルテンサイト変態温度を低下させる成分である。しかし、過剰添加はγ相を安定化させε相の体積分率を低下させるので、添加する場合にはＡｌ：０．０１〜１０％，Ｖ：０．０１〜２０％，Ｔｉ：０．０１〜１５％の範囲でＡｌ，Ｖ，Ｔｉの含有量を定める。
Ｃｒ，Ｍｏは耐食性の向上に有効な成分であるが、過剰添加は延性の著しい劣化を招く。添加する場合、Ｃｒ：０．０１〜３５％，Ｍｏ：０．０１〜３０％の範囲で含有量を選定する。
Ｎｂ，Ｚｒ，Ｗ，Ｔａ，Ｈｆは材料の強化に有効な成分であるが、過剰添加は延性の著しい劣化を招くので、添加する場合にはＮｂ：０．０１〜１０％，Ｚｒ：０．０１〜３％，Ｗ：０．０１〜３０％，Ｔａ：０．０１〜１０％，Ｈｆ：０．０１〜５％の範囲で含有量を選定する。
Ｓｉはマルテンサイト変態温度を上昇させる成分であるが、過剰添加は延性の著しい劣化を招くので、添加する場合にはＳｉ：０．０１〜８％の範囲で含有量を選定する。
Ｃ，Ｂ，Ｐ，ミッシュメタルは結晶粒微細化に有効な成分であるが、過剰添加は延性の著しい劣化を招く。そこで、Ｃ，Ｂ，Ｐ，ミッシュメタルを添加する場合にはＣ：０．００１〜３％，Ｂ：０．００１〜３％，Ｐ：０．００１〜３％，ミッシュメタル：０．００１〜３％の範囲で含有量を選定する。
所定組成に調整されたＣｏ基合金を溶解した後、鋳造，鍛造，熱間圧延等を経て圧延，引抜き，鍛造等の冷間加工によって目標サイズの板材，線材，管材等に成形される。冷間加工されたＣｏ基合金を温度：９００〜１４００℃で溶体化処理すると、冷間加工までの工程で導入された歪が除去され材料が均質化される。溶体化温度は、十分に再結晶温度以上である必要があるため９００℃以上、融点以下（具体的には、１４００℃以下）が必要であり、好ましくは１０００〜１２５０℃の範囲に設定される。
溶体化温度から室温に冷却する過程で、ｆ．ｃ．ｃ．構造のγ相からｈ．ｃ．ｐ．構造のε相にマルテンサイト変態する。マルテンサイト変態点（Ｍｓ温度）が室温より高い場合でも、Ｆｅ，Ｎｉ，Ｍｎによりγ相が安定化されるので冷却後の組織がε相に単相化することはない。
溶体化処理されたＣｏ基合金に圧延，鍛造，曲げ加工，絞り等の加工を施してもよい。加工温度は常温が通常であるが、７００℃以下にも設定できる。マルテンサイト変態には、Ｍｓ温度以下に冷却することにより生じる熱誘起変態の他に応力誘起変態もあり、加工はε相の体積分率を増加させる有効な手段である。
応力誘起変態はＭｓ温度以上の環境で応力を負荷したときに生じる変態であるが、加工に伴う動的再結晶や析出等が懸念される熱間加工は好ましくない。熱間加工が通常０．６ＴＭ（ＴＭ：融点）以上の温度での加工と定義されていることを考慮し、加工温度を０．６ＴＭ以下，具体的には７００℃以下に設定する。ε相の誘起効果は加工率の増加に応じて顕著になるので、冷間加工時の加工率を１０％以上に設定する。加工設備の能力に応じて加工率の上限が定められるが、過剰な加工率では加工設備の負担が大きくなるので９０％を上限とすることが好ましい。
弾性変形能の向上，ヤング率の低下に及ぼすε相の影響は、金属組織全体にε相が占める体積分率を１０体積％以上とすることにより顕著になる。ただし、ε相の過剰分布はヤング率低下の効果を弱め、却って加工性向上に有効なγ相の体積分率を相対的に減らすことになるので、ε相の上限を好ましくは９９体積％とする。
９００〜１４００℃での溶体化，７００℃以下の温度で１０％以上の加工の後に、３００〜８００℃（好ましくは、４００〜７００℃）で時効処理してもよい。時効処理を施すと、歪時効効果や回復・再結晶により強度を上昇又は低下させることができる。コットレル効果，鈴木効果等が起こる場合に強度が上昇し、回復や再結晶が起こる場合に強度が低下する。時効処理には少なくとも原子の短距離拡散が必要なため時効温度を３００℃以上とするが、８００℃を超える高温加熱では十分な弾性歪が得られない。
以下、実施例により本発明の作用・効果を具体的に説明するが、実施例は本発明の具体的な理解を助けるものであり本発明の技術的範囲に何ら影響を及ぼすものでないことは勿論である。In order to develop a material that has high elastic deformability and can be processed by applying and removing a magnetic field gradient, the inventors can add various alloying elements to Co, which is a ferromagnetic element. Investigated and investigated the relationship between performance and magnetic properties. As a result, it has been clarified that when at least one or two or more of Fe, Ni and Mn in an appropriate amount are alloyed with Co, a ferromagnetic highly elastic deformation alloy having improved workability is obtained.
It is known that a ferromagnetic material is displaced by receiving a force corresponding to a magnetic field gradient according to a relational expression of F = −M (dH / dx). This relational expression means that the force F obtained with a predetermined magnetic field gradient dH / dx is proportional to the magnetization intensity M. The greater the force F applied to the material, the greater the amount of strain that can be obtained. Therefore, it is better to increase the amount of strain by increasing the magnetization strength M. If the stress required for deformation is high, the amount of strain due to the magnetic field decreases. It is desirable that the rate be low. However, if the elastic deformability is small, strain remains due to the removal of the magnetic field gradient. Therefore, in order to obtain a large reversible strain by applying / removing the magnetic field gradient, a large elastic deformability is required.
Co is a ferromagnetic element having a high Curie temperature, but pure Co has poor workability and a low magnetic susceptibility. Co is f. c. c. Γ phase of the structure, h. c. p. It undergoes martensitic transformation to the ε phase of the structure. As a result of investigating the phase stability of the γ phase and the ε phase by adding various alloying elements to Co, it was possible to identify various elements that stabilize the γ phase and improve the workability.
The γ phase has an irregular structure f. c. c. Therefore, the workability is good, but the ε phase is a phase having an effect of increasing the amount of elastic deformation, but it is easy to work harden and easily generate defects such as work cracks. Therefore, when an element that is solid-solved and stabilized in the γ phase is added, the γ phase remains even at room temperature, and improvement in workability of the Co-based alloy can be expected. Among these, Fe, Ni, and Mn having a magnetic moment are effective not only for improving workability but also for improving magnetic properties of a Co-based alloy.
When the volume fraction of the ε phase is adjusted to 10% by volume or more by component adjustment, production conditions, etc., the effect of high elastic deformation becomes remarkable. However, since the excessive distribution of the ε phase cancels out the effect of improving the workability of the residual γ phase, the upper limit of the ε phase is preferably 99% by volume. Basically h. c. p. Ε phase of structure and the rest is f. c. c. Although it is a Co-based alloy having a γ-phase double-phase structure, the presence of a different phase is allowed as long as it does not adversely affect high elastic properties, magnetic properties, workability, and the like. Examples of the different phase include intermetallic compounds, carbides, α-phase (bcc structure) generated by the addition of the third component.
The Co alloy to which Fe is added exhibits particularly excellent magnetic properties. In applications as a displacement control element by applying a magnetic field, it is preferable to show a large magnetization intensity with a relatively low magnetic field, but when measuring the magnetization curve of the Co-based alloy of the present invention, it is possible to obtain an external magnetic field of 0.2 T (Tesla). It exhibits a magnetization strength of 30 emu / g or more, and can be said to be a promising material for a displacement control element. Although Ni and Mn exhibit an action of weakening the magnetization strength, the required strength as a displacement control element is imparted because the magnetization strength is still maintained at a high level.
On the other hand, as a result of investigating the relationship between the structure and the elastic deformation behavior, the following was clarified.
Annealed metal materials such as Fe and Co usually have an elastic limit strain of about 0.3%. When cold-worked, they work harden and curl deformation is suppressed. Yield stress increases with increasing hardness and tensile strength. , The elastic limit becomes high. On the other hand, the Co-based alloy of the present invention recovers a strain of 0.4% or more exceeding the normal elastic strain amount with respect to, for example, 1% bending deformation. Moreover, it can be applied in a wide temperature range from liquid nitrogen temperature (−196 ° C.) to 400 ° C.
However, if the stress required for deformation is high, it is difficult to obtain displacement due to a magnetic field, so a lower Young's modulus is desirable. Young's modulus is a physical property value related to the cohesive force between atoms, and is considered difficult to control by processing or heat treatment. In this regard, the present invention achieves a low Young's modulus by controlling the lattice softening during martensitic transformation, the orientation of martensitic variants, and the like.
In the martensitic transformation, a lattice softening phenomenon in which the elastic constant decreases in a temperature range having a range of several tens to 100 ° C. from the martensitic transformation temperature is seen as one of the precursory phenomena accompanying the transformation. When lattice softening is used for a Co-based alloy, a decrease in Young's modulus is predicted near the martensitic transformation temperature Ms.
In pure Co, lattice softening due to martensitic transformation occurs in the vicinity of 420 ° C., but the transformation temperature decreases with the addition of Fe, Ni, and Mn. As the transformation temperature decreases, a low Young's modulus is obtained in the desired temperature range. Furthermore, since the Co-X (X: Fe, Ni, Mn) -based martensitic transformation is a non-thermoelastic type, the hysteresis is about 150 ° C., and the hysteresis further spreads by processing, so the Young's modulus is low over a wide temperature range. It is suitable for realizing.
In the martensitic transformation of the Co-based alloy, the entire sample does not transform into the martensite phase even if it is sufficiently cooled below the Ms temperature as in the case of the Fe alloy and the like, and the matrix phase remains to some extent. Therefore, although the martensite phase is stress-induced by processing after the solution treatment, the largest variant of the Schmid factor is preferentially generated with respect to the stress. Further, when stress is applied to the thermally induced martensite phase, a part of the martensite phase is rearranged into a variant preferential in the stress direction. The fact that these preferential variants exhibit a low deformation stress with respect to the stress direction is also considered to be one of the causes that the Co-based alloy exhibits a low Young's modulus.
Co-Ni-Al alloys (References 4 and 5) have a β-phase or β + γ-phase structure, and use β-phase martensite transformation / reverse transformation for imparting shape memory characteristics. There is no transformation to the ε phase. Although the presence of a β phase having a B2 structure with poor ductility also causes a decrease in workability, the Co-based alloy of the present invention can be said to be a material with excellent ductility in that no β phase is present. In addition, the Co-based alloy of the present invention has an extremely large magnetization compared to a Co-Ni-Al-based alloy that is ferromagnetic but has a small magnetization strength, so it is excellent in controlling displacement by a magnetic field gradient. Demonstrate performance.
The Co-based alloy of Reference 6 having shape memory characteristics is excellent in shape memory effect, but cannot obtain superelasticity, and it is difficult to control displacement by a magnetic field due to extremely small magnetostriction. This is a result of the shape memory effect, superelasticity, and displacement control by a magnetic field originating from different phenomena. That is, the shape memory effect is a phenomenon in which shape is recovered by applying deformation in the martensite phase state and reversely transforming the parent phase into martensite. Superelasticity is a phenomenon in which shape recovery occurs because stress-induced martensitic transformation occurs when stress is applied in the parent phase state, and when the stress is unloaded, martensite reversely transforms to the parent phase. On the other hand, large magnetostriction in a ferromagnetic shape memory alloy is a phenomenon obtained by rearrangement of martensite variants by a uniform magnetic field and magnetic field induced martensitic transformation.
The Co-based alloy of this patent achieves a large elastic deformability mainly by work hardening, and it is superb in that reversible transformation / reverse transformation of the martensite phase due to stress loading / unloading is not clearly detected. Essentially different from elasticity, the martensite phase (ε phase) only assists the elastic deformability while keeping the Young's modulus low. Furthermore, since the workability is improved and the strength of magnetization is also increased, it exhibits an excellent function as a material capable of displacement control by applying and removing a magnetic field gradient.
The Co-based alloy of the present invention is based on a Co alloy containing one or more selected from Fe: 0.01 to 10%, Ni: 0.01 to 30%, and Mn: 0.01 to 25%. To do.
Fe, Ni, and Mn have the effect of lowering the martensite transformation temperature, contributing to the improvement of ductility and workability, and increasing the magnetic susceptibility. Such an effect becomes remarkable when 0.01% or more of Fe, Ni, or Mn is added. However, excessive addition lowers the martensite and magnetic transformation temperature to room temperature or lower, and suppresses the formation of the ε phase and lowers the magnetic properties, so the upper limit was set to Fe: 10%, Ni: 30%, and Mn: 25%. When adding 2 types or 2 types of Fe, Ni, and Mn, it is preferable to keep total content in the range of 0.02 to 50%. The content of each of Fe, Ni, and Mn is preferably set in the range of 1 to 8%, 1 to 25%, and 1 to 20%, more preferably 2 to 6%, 5 to 20%, and 5 to 15%.
If necessary, one or more third components selected from Al, Cr, V, Ti, Mo, Nb, Zr, W, Ta, Hf, Si, C, B, P, and misch metal may be added to Co- It can be added to the (Fe, Mn, Ni) system. When a plurality of third components are added, the total content is selected in the range of 0.002 to 50% (preferably 0.005 to 30%).
Al, V, and Ti are components that lower the martensitic transformation temperature. However, excessive addition stabilizes the γ phase and lowers the volume fraction of the ε phase. Therefore, when added, Al: 0.01 to 10%, V: 0.01 to 20%, Ti: 0.01 The content of Al, V, Ti is determined in the range of ˜15%.
Cr and Mo are effective components for improving the corrosion resistance, but excessive addition causes remarkable deterioration of ductility. When adding, content is selected in the range of Cr: 0.01-35%, Mo: 0.01-30%.
Nb, Zr, W, Ta, and Hf are effective components for strengthening the material. However, excessive addition causes significant deterioration of ductility. Therefore, when added, Nb: 0.01 to 10%, Zr: 0. The content is selected in the range of 01 to 3%, W: 0.01 to 30%, Ta: 0.01 to 10%, Hf: 0.01 to 5%.
Si is a component that raises the martensite transformation temperature. However, excessive addition causes significant deterioration of ductility. Therefore, when Si is added, the content is selected in the range of Si: 0.01 to 8%.
C, B, P, and misch metal are effective components for refining crystal grains, but excessive addition causes a significant deterioration in ductility. Therefore, when adding C, B, P, Misch metal, C: 0.001 to 3%, B: 0.001 to 3%, P: 0.001 to 3%, Misch metal: 0.001 The content is selected in the range of 3%.
After melting the Co-based alloy adjusted to a predetermined composition, it is formed into a target-size plate, wire, tube, etc. by cold working such as rolling, drawing, forging, etc. through casting, forging, hot rolling and the like. When the cold-worked Co-based alloy is subjected to a solution treatment at a temperature of 900 to 1400 ° C., the strain introduced in the steps up to the cold working is removed and the material is homogenized. Since the solution temperature needs to be sufficiently higher than the recrystallization temperature, it must be 900 ° C. or higher and the melting point or lower (specifically, 1400 ° C. or lower), and is preferably set in the range of 1000 to 1250 ° C. .
In the process of cooling from solution temperature to room temperature, f. c. c. From the gamma phase of the structure h. c. p. It undergoes martensitic transformation to the ε phase of the structure. Even when the martensitic transformation point (Ms temperature) is higher than room temperature, the γ phase is stabilized by Fe, Ni, and Mn, so that the structure after cooling does not become a single phase in the ε phase.
The solution-treated Co-based alloy may be subjected to processing such as rolling, forging, bending, and drawing. The processing temperature is usually room temperature, but can be set to 700 ° C. or lower. In the martensitic transformation, there is a stress-induced transformation in addition to a heat-induced transformation caused by cooling below the Ms temperature, and machining is an effective means for increasing the volume fraction of the ε phase.
The stress-induced transformation is a transformation that occurs when stress is applied in an environment at or above the Ms temperature, but hot working that is concerned about dynamic recrystallization, precipitation, etc. associated with working is not preferred. Considering that hot working is normally defined as processing at a temperature of 0.6 T M (T M : melting point) or higher, the processing temperature is set to 0.6 T M or lower, specifically, 700 ° C. or lower. . Since the induction effect of the ε phase becomes significant as the working rate increases, the working rate during cold working is set to 10% or more. Although the upper limit of the processing rate is determined according to the capacity of the processing facility, it is preferable to set the upper limit to 90% because an excessive processing rate increases the burden on the processing facility.
The effect of the ε phase on the improvement of the elastic deformability and the decrease in the Young's modulus becomes significant when the volume fraction of the ε phase in the entire metal structure is 10% by volume or more. However, the excessive distribution of the ε phase weakens the effect of lowering the Young's modulus and, on the contrary, relatively reduces the volume fraction of the γ phase effective for improving the workability, so the upper limit of the ε phase is preferably 99% by volume. To do.
After solution treatment at 900 to 1400 ° C. and processing of 10% or more at a temperature of 700 ° C. or less, an aging treatment at 300 to 800 ° C. (preferably 400 to 700 ° C.) may be performed. When the aging treatment is performed, the strength can be increased or decreased by the strain aging effect or recovery / recrystallization. When the Cottrell effect, the Suzuki effect, etc. occurs, the strength increases, and when recovery or recrystallization occurs, the strength decreases. Since the aging treatment requires at least short-distance diffusion of atoms, the aging temperature is set to 300 ° C. or higher. However, sufficient elastic strain cannot be obtained by high-temperature heating exceeding 800 ° C.
Hereinafter, the operation and effect of the present invention will be described in detail by way of examples. However, it should be understood that the examples are intended to assist the specific understanding of the present invention and do not affect the technical scope of the present invention. It is.
表１〜３のＣｏ基合金を溶解し、鋳造，熱間圧延を経て板厚：０．５ｍｍまで冷間圧延し、更に１２００℃×１５分で溶体化した。
表１のＦ１〜Ｆ８はＣｏ−Ｆｅ系，表２のＮ１〜Ｎ８はＣｏ−Ｎｉ系，表３のＭ１〜Ｍ８はＣｏ−Ｍｎ系を基本とする合金設計である。
表１〜３には、各Ｃｏ基合金の室温におけるε相の体積分率，回復歪量，磁化の強さを調査した結果を併せ示す。比較のため、純Ｃｏ，純Ｆｅ，ＳＵＳ３１６Ｌの同様な特性を表４に示す。また、Ｆｅ，Ｎｉ，Ｍｎのε相の体積分率，回復歪量，磁化の強さに及ぼす影響を図１〜３にグラフ化した。
回復歪量は、三点曲げ試験で１％の曲げ歪量を与えた後、除荷したときに戻る形状歪量とした。
磁化の強さは、振動試料型磁力計を用いて磁場を印加し、０．２Ｔのときの磁化の強さとした。
表１〜３，図１〜３にみられるように、Ｃｏ−Ｆｅ，Ｃｏ−Ｎｉ，Ｃｏ−Ｍｎ何れの合金系においても、金属組織全体の１０体積％以上をε相で占めていたが、Ｆｅ，Ｎｉ，Ｍｎの増量に応じてε相の体積分率，回復歪量が低下する傾向にあった。
Ｆｅ：３．８０％のＣｏ−Ｆｅ系合金（Ｆ２）では図４にみられるように、ε相がバンド状に存在していた。回復歪量はε相が多いほど大きくなる傾向にあり、１０体積％以上のε相で０．４％以上の回復歪量が得られた。
Ｃｏ−Ｆｅ系では、磁化の強さがＦｅの増量に伴い増加し、７．６１％（Ｎｏ．Ｆ４）で最大値を示し、優れた磁気特性であった。Ｃｏ−Ｎｉ系，Ｃｏ−Ｍｎ系では、Ｎｉ，Ｍｎの増量に応じ磁化の強さに低下がみられたが、依然として４３．９ｅｍｕ／ｇ以上の大きな磁化の強さが維持されていた。
一方、比較例Ｃ１（純Ｃｏ）は全体の９３体積％をε相で占めているが、回復歪量は本発明材よりも低くなっていた。比較例Ｃ２（純Ｆｅ）は磁化の強さは大きいものの、ε相が存在しておらず、回復歪量も更に低くなっていた。常磁性材料である比較例Ｃ３（ＳＵＳ３１６Ｌ）は、０．２Ｔでの磁化の強さはほぼ０に近く、回復歪量も低い値であった。
Ｃｏ−Ｆｅ，Ｃｏ−Ｎｉ，Ｃｏ−Ｍｎ何れの合金系においても、冷間圧延を施していない材料（表１〜３）に比較してε相の体積分率が増加していた。また、合金Ｆ２の応力−歪線図（図５）にみられるように、１％の歪印加に対して除荷時に０．６１％の回復歪を示している。超弾性材料に比較してより線形的である点も特長である。他の合金系Ｃｏ−Ｎｉ，Ｃｏ−Ｍｎでも冷間圧延により回復歪量は増加し、０．６％前後の回復歪量が得られた。冷間圧延による磁気特性の変化は磁化の強さが数ｅｍｕ／ｇ程度低下したに留まり、依然として優れた磁気特性を有していた。
一方、比較例Ｃ１，Ｃ２では冷間圧延によるε相分率の変化がみられず、回復歪量は発明材に比較して小さかった。Ｃ３は大きな回復歪量が得られたが、０．２Ｔの外部磁場を印加したときの磁化の強さがほぼゼロであった。
F1 to F8 in Table 1 are based on Co—Fe, N1 to N8 in Table 2 are based on Co—Ni, and M1 to M8 in Table 3 are based on Co—Mn based.
Tables 1 to 3 also show the results of investigating the volume fraction of the ε phase at room temperature, the amount of recovery strain, and the strength of magnetization of each Co-based alloy. For comparison, the same characteristics of pure Co, pure Fe, and SUS316L are shown in Table 4. Also, the effects of Fe, Ni, Mn on the volume fraction of the ε phase, the recovery strain amount, and the strength of magnetization are graphed in FIGS.
The amount of recovery strain was defined as the amount of shape strain that returned when unloading after giving a bending strain amount of 1% in a three-point bending test.
The strength of magnetization was the strength of magnetization at 0.2 T when a magnetic field was applied using a vibrating sample magnetometer.
As can be seen in Tables 1 to 3 and FIGS. 1 to 3, in any alloy system of Co—Fe, Co—Ni, and Co—Mn, 10% by volume or more of the entire metal structure was occupied by the ε phase. There was a tendency for the volume fraction of the ε phase and the recovery strain to decrease with increasing amounts of Fe, Ni and Mn.
In the Fe: 3.80% Co—Fe-based alloy (F2), as shown in FIG. 4, the ε phase was present in a band shape. The amount of recovery strain tended to increase as the amount of ε phase increased, and a recovery strain amount of 0.4% or more was obtained with 10 volume% or more of ε phase.
In the Co—Fe system, the strength of magnetization increased with the increase of Fe, and showed a maximum value at 7.61% (No. F4), which was excellent magnetic characteristics. In the Co—Ni system and Co—Mn system, the magnetization strength decreased with the increase of Ni and Mn, but the large magnetization strength of 43.9 emu / g or more was still maintained.
On the other hand, Comparative Example C1 (pure Co) accounted for 93% by volume of the total in the ε phase, but the recovery strain was lower than that of the material of the present invention. In Comparative Example C2 (pure Fe), although the magnetization strength was large, the ε phase did not exist and the recovery strain amount was further reduced. Comparative Example C3 (SUS316L), which is a paramagnetic material, had a magnetization intensity at 0.2T close to 0 and a low recovery strain.
In any of the alloy systems of Co—Fe, Co—Ni, and Co—Mn, the volume fraction of the ε phase was increased as compared with materials that were not subjected to cold rolling (Tables 1 to 3). Further, as seen in the stress-strain diagram of the alloy F2 (FIG. 5), the recovery strain of 0.61% is shown at the time of unloading with respect to the application of 1% strain. Another advantage is that it is more linear than superelastic materials. Even in other alloy-based Co—Ni and Co—Mn, the amount of recovery strain was increased by cold rolling, and a recovery strain amount of about 0.6% was obtained. The change in magnetic properties due to cold rolling was that the strength of magnetization was reduced by about several emu / g and still had excellent magnetic properties.
On the other hand, in Comparative Examples C1 and C2, no change in the ε phase fraction due to cold rolling was observed, and the amount of recovery strain was smaller than that of the inventive material. Although a large amount of recovery strain was obtained for C3, the strength of magnetization when an external magnetic field of 0.2 T was applied was almost zero.
表１のＦ２，表２のＮ４，表３のＭ３をそれぞれＣｏ−Ｆｅ系，Ｃｏ−Ｎｉ系，Ｃｏ−Ｍｎ系合金の代表として選択し、溶体化後に冷間加工，時効処理を施した。
製造条件と物性値との関係を表９に示す。また、Ｃｏ−Ｆｅ系，Ｃｏ−Ｎｉ系，Ｃｏ−Ｍｎ系それぞれについて、ε相の体積分率，回復歪量，磁化の強さに及ぼす溶体化処理条件の影響（冷間加工率：４０％，時効処理条件：７００℃×２時間に固定），冷間加工率の影響（溶体化処理条件：１２００℃×１５分，時効処理条件：７００℃×２時間に固定），時効処理条件の影響（溶体化処理条件：１２００℃×１５分，冷間加工率：４０％又は８０％に固定）を図６〜１４にグラフ化した。
試験Ｎｏ．４，６，７では、冷間加工率、時効処理条件を固定し、溶体化温度を変化させたが、ε相の体積分率、回復歪量、磁化の強さの何れにも大きな変化はみられなかった。この関係は、試験Ｎｏ．１２，１３及び試験Ｎｏ．１９，２０においても同様であり、図６，９，１２からも理解できる。また、溶体化温度が高すぎる試験Ｎｏ．２では液相が出現して部分溶融してしまった。
試験Ｎｏ．１１，１２，１４，１５では、溶体化処理条件，時効処理条件を固定し、冷間加工率を変化させた。冷間加工率が高いほどε相の体積分率が高く、回復歪量も大きくなっていた。磁化の強さは僅かに減少するが、大きな減少ではなかった。この関係は、試験Ｎｏ．４，５及び試験Ｎｏ．１９，２１でもみられ、図７，１０，１３からも理解できる。
試験Ｎｏ．１６〜１９，２２，２３では、溶体化処理条件、冷間加工率を固定し、時効処理条件を変化させた。時効温度の上昇に伴いε相の体積分率，回復歪量が減少する傾向にあった。回復歪量は冷間加工まま材に比較して小さく、試験Ｎｏ．１７〜１９，２２では溶体化まま材と溶体化＋冷間加工材で得られる回復歪量の中間の値であった。磁化の強さには顕著な変化がみられなかった。この関係は、試験Ｎｏ．１，３，４及び試験Ｎｏ．８，１０，１２においても同様である。
加工率が８０％と大きな試験Ｎｏ．９，１５では、時効処理により回復歪量が高くなっている。一方、試験Ｎｏ．２３は冷間加工後の時効による回復歪量が表３に示す加工前のものと同等であり、顕著な時効効果が得られなかった。
以上の結果は、適切な時効により回復歪量の調整が可能なことを意味し、図８，１１，１４によっても支持される。
Table 9 shows the relationship between the manufacturing conditions and the physical property values. In addition, for each of the Co—Fe system, Co—Ni system, and Co—Mn system, the influence of the solution treatment conditions on the volume fraction of ε phase, the recovery strain amount, and the strength of the magnetization (cold working rate: 40%). , Aging treatment conditions: 700 ° C x 2 hours fixed), Effect of cold working rate (Solution treatment conditions: 1200 ° C x 15 minutes, Aging treatment conditions: 700 ° C x 2 hours fixed), Effects of aging treatment conditions (Solution treatment conditions: 1200 ° C. × 15 minutes, cold working rate: fixed at 40% or 80%) are graphed in FIGS.
Test No. In 4, 6, and 7, the cold working rate and the aging treatment conditions were fixed and the solution temperature was changed, but there was a big change in any of the volume fraction of ε phase, the recovery strain amount, and the strength of magnetization. It was not seen. This relationship is related to test no. 12, 13 and test no. The same applies to 19, 20 and can be understood from FIGS. In addition, Test No. with a solution temperature too high. In 2, the liquid phase appeared and partially melted.
Test No. In 11, 12, 14, and 15, the solution treatment conditions and the aging treatment conditions were fixed, and the cold working rate was changed. The higher the cold working rate, the higher the volume fraction of the ε phase and the greater the recovery strain amount. Although the magnetization intensity decreased slightly, it was not a large decrease. This relationship is related to test no. 4, 5 and test no. 19 and 21 and can also be understood from FIGS.
Test No. In 16-19, 22, and 23, the solution treatment conditions and the cold working rate were fixed, and the aging treatment conditions were changed. As the aging temperature increased, the volume fraction of ε phase and the amount of recovery strain tended to decrease. The amount of recovery strain is smaller than that of the cold-worked material. In 17-19 and 22, it was an intermediate value of the amount of recovery strain obtained with the material in solution and the solution plus cold worked material. There was no significant change in the strength of magnetization. This relationship is related to test no. 1, 3, 4 and test no. The same applies to 8, 10, and 12.
Test No. with a processing rate as large as 80%. In Nos. 9 and 15, the recovery strain amount is increased by the aging treatment. On the other hand, test no. In No. 23, the amount of recovery strain due to aging after cold working was equivalent to that before working shown in Table 3, and a remarkable aging effect was not obtained.
The above results mean that the amount of recovery strain can be adjusted by appropriate aging, which is also supported by FIGS.
Ｃｏ−２．０５％Ｆｅ，Ｃｏ−１０％Ｎｉ，Ｃｏ−５％ＭｎをそれぞれＣｏ−Ｆｅ系，Ｃｏ−Ｎｉ系，Ｃｏ−Ｍｎ系の基本組成とし、第三成分を添加して種々のＣｏ基合金を調製した。溶解後、実施例１と同様に鋳造，熱間圧延を経て板厚：０．５ｍｍに冷間圧延し、溶体化処理，冷間圧延，時効処理を施した。
得られた各Ｃｏ基合金についてε相の体積分率，回復歪量，磁化の強さを測定した結果を表１０（Ｃｏ−Ｆｅ系），表１１（Ｃｏ−Ｎｉ系），表１２（Ｃｏ−Ｍｎ系）に示す。
表１０〜１２の調査結果にみられるように、第三成分の添加で延性，磁性，耐食性，強度等が改善されたＣｏ基合金は、何れの溶体化材でもε相が５０体積％以上存在し、１％歪印加に対し０．４１％以上の回復歪量が得られた。また、０．２Ｔの磁場印加に対し６９．９ｅｍｕ／ｇ以上の高い磁化の強さも有していた。４０％の冷間圧延によりε相の体積分率が何れも増加し、磁化の強さを大きく損なうことなく０．６％前後の高い回復歪量が得られるようになった。更に時効処理を施すことにより、回復歪量を調節できた。
Table 10 (Co-Fe series), Table 11 (Co-Ni series), Table 12 (Co series) show the results of measuring the volume fraction of ε phase, the amount of recovery strain, and the strength of magnetization for each of the obtained Co-based alloys. -Mn system).
As can be seen from the results of the investigations in Tables 10 to 12, the Co-based alloy whose ductility, magnetism, corrosion resistance, strength, etc. have been improved by the addition of the third component is present in 50% by volume or more of the ε phase in any solution material. In addition, a recovery strain amount of 0.41% or more was obtained with respect to 1% strain application. Further, it had a high magnetization intensity of 69.9 emu / g or more with respect to 0.2 T magnetic field application. The 40% cold rolling increased the volume fraction of the ε phase, and a high recovery strain amount of about 0.6% was obtained without greatly deteriorating the magnetization strength. Furthermore, the recovery strain could be adjusted by applying an aging treatment.
表１のＦ２合金を選択し、鋳造，熱間圧延を経て板厚：０．３３ｍｍまで冷間圧延し、更に１２００℃×１５分で溶体化し、最後に圧下率：２０％で冷間圧延した。
得られたＣｏ−Ｆｅ系合金について、−５０℃，２５℃，１００℃，２００℃の各温度におけるε相の体積分率，回復歪量，磁化の強さを求めた。回復歪量は、各温度での引張り試験で１％の歪量を与えた後、除荷したときに戻る形状歪量とした。ε相の体積分率，回復歪量，磁化の強さについては、各温度において実施例１と同じ方法で求めた。
表１３の調査結果にみられるように、ε相の体積分率は−５０℃から２００℃までの温度範囲で大きく変化していない。回復歪量は試験温度が上昇すると低下する傾向にあったが、２００℃においても依然大きな回復歪を示した。
形状記憶合金では、変形応力が温度に対して大きく変化し、たとえばＴｉ−Ｎｉ系合金では見掛け上の降伏応力の温度依存性は約５ＭＰａ／℃である。一方、本発明のＣｏ−Ｆｅ系合金は応力−歪線図（図１５）にみられるように応力の温度依存性が小さく、約０．５ＭＰａ／℃とＴｉ−Ｎｉ系合金の約１０分の１程度であるので、室温以下から高温までの広い温度範囲にわたる利用にも適している。また、キュリー温度が非常に高いため、２００℃でも大きな磁化の強さを示した。
同じ条件下で製造した表２のＮ４合金冷延材，表３のＭ３合金冷延材についても同じ試験で温度による影響を調査し、調査結果を表１３に併せ示した。Ｃｏ−Ｎｉ系，Ｃｏ−Ｍｎ系でも大半をε相で占める金属組織を有し、回復歪量，磁化の強さ共に大きな値を示した。
About the obtained Co-Fe type | system | group alloy, the volume fraction of the epsilon phase in each temperature of -50 degreeC, 25 degreeC, 100 degreeC, and 200 degreeC, the recovery | restoration distortion amount, and the strength of magnetization were calculated | required. The amount of recovery strain was defined as the amount of geometric strain that returned when unloading after giving a strain amount of 1% in a tensile test at each temperature. The volume fraction of ε phase, the recovery strain amount, and the strength of magnetization were obtained by the same method as in Example 1 at each temperature.
As seen in the investigation results in Table 13, the volume fraction of the ε phase does not change significantly in the temperature range from −50 ° C. to 200 ° C. The amount of recovery strain tended to decrease as the test temperature increased, but still showed a large recovery strain even at 200 ° C.
In the shape memory alloy, the deformation stress changes greatly with respect to the temperature. For example, in the Ti—Ni alloy, the temperature dependence of the apparent yield stress is about 5 MPa / ° C. On the other hand, as shown in the stress-strain diagram (FIG. 15), the Co—Fe based alloy of the present invention has a low temperature dependence of stress, about 0.5 MPa / ° C., about 10 minutes of that of Ti—Ni based alloy. Since it is about 1, it is suitable for use over a wide temperature range from below room temperature to high temperature. In addition, since the Curie temperature is very high, a large magnetization strength was exhibited even at 200 ° C.
The N4 alloy cold-rolled material in Table 2 and the M3 alloy cold-rolled material in Table 3 manufactured under the same conditions were examined for the influence of temperature in the same test, and the results of the investigation are also shown in Table 13. The Co—Ni system and Co—Mn system also have a metal structure that occupies most of the ε phase, and both recovery strain and magnetization strength are large.
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| US10844465B2 (en) | 2017-08-09 | 2020-11-24 | Garrett Transportation I Inc. | Stainless steel alloys and turbocharger kinematic components formed from stainless steel alloys |
| WO2019099719A1 (en) * | 2017-11-16 | 2019-05-23 | Arconic Inc. | Cobalt-chromium-aluminum alloys, and methods for producing the same |
| KR102136455B1 (en) * | 2018-03-16 | 2020-07-21 | 서울대학교산학협력단 | Self-healable trip superalloys and manufacturing method for the same |
| DE102018208737A1 (en) * | 2018-06-04 | 2019-12-05 | Siemens Aktiengesellschaft | Y, Y` hardened cobalt-nickel base alloy, powder, component and process |
| DE102018208736A1 (en) * | 2018-06-04 | 2019-12-05 | Siemens Aktiengesellschaft | Y, Y 'hardened cobalt-nickel base alloy, powder, component and process |
| CN110592432B (en) * | 2019-09-25 | 2020-09-04 | 北京北冶功能材料有限公司 | Cobalt-based wrought superalloy and preparation method thereof |
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| JP3822573B2 (en) * | 2003-03-18 | 2006-09-20 | 本田技研工業株式会社 | Shape memory alloy and manufacturing method thereof |
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2006
- 2006-11-22 EP EP06833680A patent/EP1959024A4/en not_active Withdrawn
- 2006-11-22 WO PCT/JP2006/323877 patent/WO2007066555A1/en active Application Filing
- 2006-11-22 JP JP2007549087A patent/JPWO2007066555A1/en active Pending
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2008
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| JPS6314864A (en) * | 1986-07-08 | 1988-01-22 | Ulvac Corp | Co alloy sputtering target and its production |
| JPH0251839A (en) * | 1988-06-11 | 1990-02-21 | Degussa Ag | Target for magnetron cathode sputtering apparatus |
| JP2004238720A (en) * | 2003-02-10 | 2004-08-26 | Kiyohito Ishida | Shape memory alloy |
Also Published As
| Publication number | Publication date |
|---|---|
| EP1959024A4 (en) | 2009-12-23 |
| US20080289730A1 (en) | 2008-11-27 |
| WO2007066555A1 (en) | 2007-06-14 |
| EP1959024A1 (en) | 2008-08-20 |
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