JP4177467B2 - High toughness hard alloy and manufacturing method thereof - Google Patents
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【０００１】
【発明の属する技術分野】
本発明は、高靱性の硬質合金とその製造方法に関するものである。
【従来の技術】
【０００２】
一般的に硬質相をＷＣを主体とし結合相をＣｏ、Ｎｉなどの鉄族金属とする硬質合金はＷＣ基超硬合金と呼ばれ、硬質相をＴｉＣ（Ｎ）を主体とし結合相を鉄族金属とする硬質合金はＴｉＣ（Ｎ）基サーメットと呼ばれる。
【０００３】
これらの硬質合金は一般的に1350℃以上1600℃以下の温度で１時間ほど真空中で無加圧で保持されて焼結が行われる。場合によってはその後、焼結温度よりも低い温度でＨＩＰ（熱間静水圧プレス処理）がなされることもある。そして、このような焼結条件下では液相が生成し、ＷＣ粒は溶解再析出現象により焼結中に粒成長を起こしやすいことが知られている。
【０００４】
これらの方法によって作製された硬質合金の断面組織写真を図４に示す。図における白い斑点部分が結合相だが、どの切断面でも硬質相の粒界や粒同士の間隙に存在するだけである。その形状も結合相粒と呼べるような明確なものは少なく、存在してもアスペクト比が３より小さい形状のものである。このような組織となる理由は、焼結中に結合相が溶解して液相が発生すると、ＷＣ基超硬合金やＴｉＣ（Ｎ）基サーメットでは硬質相と結合相間の濡れ性が高いことにより、液相が硬質相粒界や間隙に流動するためである。
【０００５】
一方、ＷＣ基超硬合金やサーメットは高い硬度を有する材料を作製するため、微粒原料を使用し、粒成長抑制のためＶＣ、Ｃｒ3 Ｃ2 、ＮｂＣ、ＴａＣ、ＴｉＣなどの化合物を添加し、緻密化できる限界の低温で焼結することで超微粒の硬質合金を作製する努力がなされてきた。特にＷＣ基超硬合金ではＶＣやＣｒ3 Ｃ2 の添加により、平均粒径が約 0.5μmのＷＣ組織を有する硬質合金を作製（特開平1-215947号公報、同4-289146号公報、同5-98385 号公報参照）している。
【０００６】
【発明が解決しようとする課題】
しかし、上記の組織を有する硬質合金は靱性が十分とはいえない。また、平均粒径が約 0.5μm以下の細かい粒径を有するＷＣ基超硬合金を工業的に製造することはできていない。そのため、硬質合金を利用できる分野にも制限があった。
従って、本発明の主目的は、より高靱性で硬度や強度にも優れる硬質合金とその製造方法を提供することにある。
【０００７】
【課題を解決するための手段】
本発明は上記の目的を達成するため、ＷＣ，ＴｉＣ，ＴｉＮおよびＴｉＣＮから選択された少なくとも１種からなる硬質相と、鉄族金属からなる結合相とを具え、結合相が、硬質相粉末の３倍以上の平均粒径を有する結合相粉末で通電加圧焼結されて形成され、通電加圧焼結を行う焼結装置の加圧軸と平行な硬質合金の断面においてアスペクト比が５〜２０となる扁平な形状で、加圧軸と垂直な方向に伸びるように配列に方向性を有している結合相組織を含むことを特徴とする。特に、アスペクト比が５〜20となる形状の結合相組織は、その最大長さが硬質相の平均結晶粒径の３倍以上であることが好ましい。
【０００８】
ここで、アスペクト比とは結合相組織の最大長さと平均厚みの比をいう。後述するように、本発明の硬質合金は通電加圧焼結により得られ、その際の加圧により結合相組織が押し潰される。その結果、上記アスペクト比の組織が形成されるのである。このような結合相組織は焼結時の加圧により形成されるため、加圧軸と垂直な方向に伸びるように配列に方向性を有し、扁平な形状となる。なお、言うまでもないが、本発明合金中には不可避的不純物を含んでいてもよい。不可避的不純物には、例えばＡｌ，Ｂａ，Ｃａ，Ｃｕ，Ｆｅ，Ｍｇ，Ｍｎ，Ｎｉ，Ｓｉ，Ｓｒ，Ｓ，Ｏ，Ｎ，Ｍｏ，Ｓｎ，Ｃｒ等が挙げられる。
【０００９】
従来、このような不均一な結合相組織を有する硬質合金は、液相焼結材料としては好ましくないと考えられた。しかし、後に実施例で示すように、本発明硬質合金の破壊靱性は従来の合金のそれよりも高い値を示す。これは、亀裂進展の際に亀裂がこの扁平な結合相を必ず横切ることになるため、この際の結合相の塑性変形に亀裂進展エネルギーが費やされるからと考えられる。
【００１０】
このような硬質合金は、さらに次の各要件を単独で、または複合して具えることが好適である。
(1) 硬質相の平均粒径が0.01〜１μmである。
ＷＣ基超硬合金やＴｉＣ（Ｎ）基サーメットはその硬質相粒子を微粒とすることで焼結体の硬度を向上できる。しかし、硬質相粒度を小さくすると硬度が高くなる反面、靱性が低下し、非常に脆くなる。特に硬質相粒径が１μmよりも小さくなるとその傾向は顕著になる。
【００１１】
後述する本発明製造方法によれば、上記限定範囲の硬質相を有する硬質合金を得ることができる。しかもこの方法によれば、結合金属の形状を扁平な形状とすることができるため靱性を向上でき、従来、相反する関係にあった硬度と靱性を両立させることができる。特に硬質相の粒径が0.01〜１μmのときにその効果が大きい。0.01μmよりも小さい原料を使用することは工業的に高コストであり、１μmより大きい粒度を有する合金では靱性の向上効果が小さい。特に好ましい粒径は 0.1〜１μmである。
【００１２】
ここで得られる合金の硬質相粒径は使用する原料の硬質相粒径に主に依存する。現状技術では直接炭化法で作製したＷＣ粉末や粉砕工程で微細化したＷＣ粉末を用いればよい。今後さらに微粒のＷＣ粉末が開発された際にも本発明を適用することで一層微粒のＷＣを有する超硬合金を作製できる。
【００１３】
(2) Ｃｒ、Ｖ、Ｃｒの炭化物、Ｖの炭化物から選択された少なくとも１種を含有し、その合計含有量が結合相量に対して１wt％以下である。
このような粒成長抑制材の添加量を限定することにより、得られる合金の硬質相粒径を微細にでき、かつ抗折力も向上できる。通電加圧焼結法により急速昇温し、低温で短時間の焼結を行えば、硬質相の粒成長を抑制でき、その微粒化が実現できる。しかし、粒成長抑制材（ＶＣやＣｒの炭化物）を添加した焼結では、抗折力が従来技術で作製した合金よりも低強度のものしか得られない。そして、この原因はＶＣやＣｒの炭化物の凝集体が要因と思われる異常な組織にあることが判明した。
【００１４】
そのため、粒成長抑制材の添加量を結合相量に対して１wt％以下とすることで硬質相が微粒の合金を作製でき、しかも抗折力が従来よりも高い合金を作製できる。なお、Ｃｒ、Ｖ、Ｃｒの炭化物、Ｖの炭化物の添加量は無添加が望ましいが、原料粉末中に不可避不純物の形でＣｒやＶが混入することが考えられる。そのため、これら粒成長抑制材の合計含有量を結合相量に対して 0.3wt％以下とすることがより好ましい。
【００１５】
( ３−１ ) 加圧軸と平行の断面においてアスペクト比が５〜２０となる形状の結合相組織が、硬質合金の通電加圧焼結を行う際の加圧方向となる厚さ方向の一方側に有し、厚さ方向の他方側には有さないように、厚さ方向に組織が変化している。
( ３−２ ) 硬質合金の前記厚さ方向の一方側における硬質相の平均粒径が０．０１〜１μmの範囲にあり、前記厚さ方向の他方側における硬質相の平均粒径が同範囲にないように厚さ方向に硬質相の粒径が変化している。
( ３−３ ) 硬質合金の前記厚さ方向一方側から前記厚さ方向他方側に向かって結合相量が変化している。
【００１６】
前記厚さ方向一方側と厚さ方向他方側とで硬度や靱性の異なる硬質合金は従来より提案されているが、従来の焼結法では硬質相の粒成長および焼結中の液相の移動が激しく、狙いとする合金を作製することが難しかった。後述する本発明方法では、焼結中の硬質相の粒成長、液相の移動が少ないため、厚さ方向に硬質相粒径や結合相組織、結合相量の異なる硬質合金を作製することができる。なお、これらの組成の変化の仕方には、段階的なものと実質上連続的なものとの双方を含む。
【００１７】
(4) 金属材料の基体上に接合されている。
従来、焼結体と金属基体とのろう付けによる接合では接合強度が不十分だったが、焼結接合することにより、高い接合強度が得られると共に、ろう付け工程を省略することができる。
【００１８】
上記硬質合金の製造方法は、ＷＣ，ＴｉＣ，ＴｉＮおよびＴｉＣＮから選択された少なくとも１種からなる硬質相粉末と、硬質相粉末の３倍以上の平均粒径を有する鉄族金属からなる結合相粉末とを混合する工程と、この混合粉末から構成される原料部材を通電加熱装置に配置する工程と、この原料部材を1100℃〜1350℃、5〜200MPaで通電加圧焼結する工程とを具えることを特徴とする。結合相粉末の平均粒径は好ましくは５倍以上である。
【００１９】
ここで、原料部材には、各原料の粉末自体や、予めプレスした圧粉体、中間焼結体、これらの積層体などが含まれる。また、必要に応じて原料粉末を混合する際に高融点化合物などの粒成長抑制材を加えればよい。高融点化合物としては、IVa,Va,VIa属元素の炭化物，窒化物，炭窒化物を挙げることができる。粒成長抑制材は無添加が最も好ましい。添加する場合には極力少なくする。特に、Ｃｒ、Ｖ、Ｃｒの炭化物、Ｖの炭化物の合計含有量は結合相量に対して１wt％以下とする。より好ましい含有量は0.3wt ％以下である。
【００２０】
この方法により急速昇温して低温で短時間の焼結を行えば、結合相が移動する時間が十分でないため、結合相粒子が加圧軸の方向に押しつぶされた扁平な形状に形成される。
【００２１】
焼結は液相の存在下で行うことが望ましい。焼結は固相焼結で行うと、扁平な結合相組織が生成しやすいが、液相を出現させて焼結を行うことで緻密化が促進される。これにより、焼結体の強度は向上する。従って、扁平な結合相組織が消失しない程度の温度で液相を生成させて、短時間で焼結することにより、緻密で強度、靱性、硬度に優れた合金を作製することができる。
【００２２】
上記各焼結条件の限定理由は次の通りである。
焼結温度は、1100℃未満では緻密化が進行しにくく、1350℃を越えると液相のシミ出しが生じやすくなるためである。なお、ここでいう焼結温度は焼結炉を制御するときの黒鉛型表面の温度のことを指す。実際の試料温度はこの温度よりも
150℃〜300 ℃程度高い温度になっているものと思われる。
【００２３】
また、加圧力は、５MPa 以下では加圧焼結の効果が見られず、200MPaより加圧力を大きくすることは設備的に難しく、コストアップの要因となるためである。特に好ましい圧力は10〜50MPa である。その理由は安価な黒鉛型が利用できるためである。
【００２４】
さらに、焼結時間は10分以内であることが好ましい。焼結時間を短くすることで硬質相の粒成長および焼結中の液相の移動を抑制し、厚さ方向にＷＣの粒径や結合相量の異なる硬質合金を作製することができる。より好ましくは５分以内である。なお、焼結雰囲気は0.1torr以下の真空が好ましい。
【００２５】
硬質相粒径や組織が厚さ方向に変化する硬質合金を製造するには、原料粉末の配合を変えることで硬質相（結合相）粒径や結合相量の異なる複数種の混合粉末を準備しておけばよい。結合相粒径の異なる複数種の混合粉末を準備する場合、これらの混合粉末のいずれかに硬質相粉末の３倍以上の粒径の結合相を含むようにする。そして、原料部材を通電加熱装置に配置する工程において、これら複数種の混合粉末を硬質相（結合相）粒子の粒径順または結合相量順に積層して配置する。準備された混合粉末の種類が少なければ、厚さ方向に段階的に硬質相（結合相）粒径や結合相量の異なる硬質合金を得ることができ、この種類を多くして積層される各層の厚みを薄くすれば実質上連続的に硬質相（結合相）粒径や結合相量の変化する硬質合金を得ることができる。本発明の方法では、焼結中の硬質相の粒成長、液相の移動が少ないため、このような構成の焼結体を安定して製造することができる。
【００２６】
また、このような傾斜構造の硬質合金を基体上に接合するには、基体と共に原料部材を通電加圧装置に配置すればよい。その際、接合面側の硬質相（結合相）の粒径を大きく、その反対面側の粒径を小さくすることが望ましい。
【００２７】
以下、発明の実施の形態について説明する。
（実施例１）
平均粒径１μmのＷＣ粉末、平均粒径１μmのＣｏ粉末、平均粉径 1.5μmのＴｉＣＮ粉末、平均粒径２μmのＴｉＣ粉末、平均粒径１μmのＮｉ粉末を準備し、表１に記載した組成に配合し、ボールミルで20時間混合粉砕して原料粉末（No.1-1〜1-7）を作製した。また、Ｃｏ粉末とＮｉ粉末をそれぞれ平均粒径３μmと５μmの粗い原料に変えた原料粉末No.2-2〜2-7も同様にして作製した。
【００２８】
【表１】
次に、これらの粉末を黒鉛型に装入し、通電加熱焼結装置を用いて、50MPa の圧力を上下方向から負荷しながら昇温スピード 190℃／min となるように黒鉛型に電流を通じ、1130℃に達した時点で５分間キープし、約 100℃／min の速度で冷却を行うことによって25×８×５mmの形状の焼結体（試料No. １〜14）を得た。
【００３０】
これらの焼結体を加圧軸に平行な面で切断して断面を平面研削し、鏡面研磨後、光学顕微鏡により任意の３視野の組織写真撮影を1500倍にて行い、この写真を用いて、結合相金属のアスペクト比を算出した。ここで、アスペクト比は結合相粒の最大長さを平均厚みで割ることにより算出した。また、ダイヤモンド製ヴィッカース圧子を用いて、50kgの荷重でインデンテーション法により、硬度と破壊靱性を測定した。さらに３点曲げ試験により、曲げ強度も測定した。これらの測定結果を表２に示す。
【００３１】
【表２】
表２より、主体となる硬質相粒径に対して３倍以上大きな結合相金属粉末を用いた原料No.2-1〜2-7の焼結体の結合相のアスペクト比は硬質相と結合相の粒子径の差がない原料No.1-1〜1-7を焼結した合金のそれよりも大きく、約５〜15の値となっている。また、これら試料No.2,4,6,8,10,12,14 の破壊靱性は試料No.1,3,5,7,9,11,13の破壊靱性よりも大幅に優れることが確認できた。
【００３３】
焼結体を加圧軸に平行な面で切断し、平面研削・鏡面研磨した断面の光学顕微鏡写真を図１に示す。同写真の白い部分が結合相である。この写真から明らかなように、一部の結合相は加圧軸に対して垂直な方向に長く伸びた形状、すなわちアスペクト比で５〜20の形状となっている。また、この組織写真中に確認できた結合相金属は、焼結装置の加圧軸に垂直な方向に伸びるように配列しており、方向性を有していることがわかる。
【００３４】
次に、焼結体を加圧軸に垂直な面で切断し、平面研削・鏡面研磨した断面の光学顕微鏡写真を図２に示す。ここでも白い部分が結合相を表している。写真中央部に示される結合相粒は円形であり、この結合相金属は扁平な形状であることがわかる。
【００３５】
さらに、原料粉末No.1-1,2-1,1-5,2-5,1-6,2-6を用いて作製した試料No.1,2,9,10,11,12 の曲げ強度を比較した。ＶＣ、Ｃｒ3 Ｃ2 の添加量が結合相量に対して５wt％であるNo.9,10 の試料の曲げ強度はNo.1,2と比較して低強度となっていた。しかし、１wt％以下である試料No.11,12の曲げ強度はＶＣ、Ｃｒ3 Ｃ2 無添加のNo.1,2と同程度であった。中でも結合相金属のアスペクト比が５〜20の範囲にある試料12の破壊靱性はアスペクト比が３の試料No.11 の合金よりも高く、曲げ強度、破壊靱性、硬度が高レベルとなっていることがわかる。
【００３６】
（実施例２）
実施例１で作製した原料粉末No.1-1と同一の組成で、ＷＣ粉末を平均粒径0.25μmのものに変更した原料粉末No.3-1を作製した。また、実施例１で作製した原料粉末No.1-2と同一の組成で、ＴｉＣＮ粉末を平均粒径 0.5μm、Ｎｉ粉末を平均粒径1.5μm のものに変更した原料粉末No.3-2を作製した。
【００３７】
これらの粉末を実施例１と同様にして、通電加熱焼結装置で最高キープ温度での保持時間を表３に記載したように変化させて焼結し、試料No. 15〜No. 22を得た。これらの試料の鏡面研磨した断面組織をＦＥ−ＳＥＭで写真撮影後、写真を用いて主体となる硬質相の平均粒子径をフルマンの式から算出した。また、これらの試料の結合相金属のアスペクト比、Ｈｖ硬度および破壊靱性を実施例１と同様にして測定し、その結果を表３中に記載した。
【００３８】
【表３】
表３の結果より、通電加圧焼結法により平均粒子径が１μm以下のＷＣ基超硬合金、ＴｉＣ（Ｎ）基サーメットが作製できたことがわかる。このため、これらの合金は高い硬度を示した。また、これらの合金はアスペクト比５〜20の結合相金属粒を有していることから、破壊靱性値も非常に大きく、硬度と靱性が高レベルで両立していることが判明した。
【００４０】
なお、本実施例で行った通電加圧焼結法での実際の試料温度はＰＲ熱電対による測定の結果、約1380℃であることが判明した。この温度はＷＣ基超硬合金の共晶組成の融点1320℃を上回っており、少なくとも部分的には液相が出現していたものと考えられた。
【００４１】
（実施例３）
平均粒径0.25μmのＷＣ粉末、平均粒径1.3μm のＣｏ粉末、平均粒径１μmのＶＣ粉末、平均粒径 1.5μmのＣｒ3 Ｃ2 粉末を準備し、表４に記載した組成に配合し、アトライターで10時間混合粉砕して原料粉末（原料粉末No.3-3〜3-11 ）を作製した。
【００４２】
【表４】
これらの原料粉末を１ ton／cm2 の圧力で金型プレスし、プレス体を焼結炉にセットして、0.01Torr以下の真空中で昇温速度10℃／min 、最高キープ温度1350℃、キープ時間１時間、冷却速度５℃／min の条件（従来の液相焼結法条件）で焼結した。さらにその後、1320度、キープ時間１時間、アルゴン中で100MPaの条件でＨＩＰ処理を行い、25×８×５mmの形状の焼結体（試料No.23 〜31：表５参照）を得た。
【００４４】
これらの焼結体は平面研削、鏡面研磨後、ＦＥ−ＳＥＭにより組織写真撮影を行い、撮影した写真を用いてフルマンの式により、ＷＣの平均粒子径を算出した。また、20mmスパンの３点曲げ試験で曲げ強度も測定した。これらの測定結果を表５に示す。
【００４５】
【表５】
次に、原料No.3-3〜3-11 を用いて、通電加熱焼結装置を用いて、50MPa の圧力を上下方向から負荷しながら昇温スピード 190℃／min となるように黒鉛型に電流を通じ、1130℃に達した時点で５分間キープし、約 100℃／min の速度で冷却を行うことによって硬質合金（試料No.32 〜40）を作製した。これらの試料No.32〜40も同様にして、ＷＣの平均粒度、曲げ強度と結合相のアスペクト比を測定した。その結果も表５に記載した。
【００４７】
その結果、従来焼結法で得られた最小のＷＣ平均粒度はＶＣを添加した場合の 0.5μmであるのに対して、通電加圧焼結法では原料組成に関わらず、平均粒度 0.3μmの微粒ＷＣ合金が作製でき、しかも優れた破壊靱性を有することが確認できた。ところが、ＶＣやＣｒ3 Ｃ2 を結合相量に対して１wt％を越える量添加した合金については、ＷＣ粒度が微細にも関わらず曲げ強度が著しく低下していることが判明した。ただ、ＶＣやＣｒ3 Ｃ2 を結合相量に対して１wt％以下の含有量とした試料No.34,35,38,39,40 の合金は非常に優れた曲げ強度を実現し、従来焼結法以上の曲げ強度を実現できることが判明した。
【００４８】
なお、本実施例で行った通電加圧焼結法での実際の試料温度はＰＲ熱電対による測定の結果、約1380℃であることが判明した。この温度はＷＣ基超硬合金の共晶組成の融点1320℃を上回っており、少なくとも部分的には液相が出現していたものと考えられた。
【００４９】
（実施例４）
硬質相としてＷＣ、結合相として平均粒径 3.0μmのＣｏを10wt％、Ｎｉを２wt％配合し、10時間の混合、粉砕をアトライターで行った粉末をＷＣ粒径の大きさに分けて２種類用意した。そして、黒鉛型中にＷＣ粒径の大きい粉末（平均粒径 2.5μm）が下部層、ＷＣ粒径の小さい粉末（平均粒径0.25μm）が上部層となるように層状にプレスして充填し、41MPa の圧力を上下方向から負荷しながら昇温スピード 300℃／分となるように黒鉛型に電流を通じ、1200℃に達した時点で３分間キープし、 100℃／分の条件で冷却を行うことによって硬質合金を作製した。
【００５０】
得られた直径30mm、厚み８mmの円板状焼結体の加圧軸に平行な断面を＃250 の砥石で平面研削後、鏡面研磨して光学顕微鏡により観察したところ、上部層にはアスペクト比が約８の扁平な形状をしたＣｏが部分的に見られた。また、ダイヤモンド製ヴィッカース圧子を用いた硬度、破壊靱性測定でも高硬度、高靱性を示した。これは、この層が約 0.3μmの微粒ＷＣを主体とすることで高硬度となり、扁平なＣｏが存在するため亀裂進展エネルギーを吸収し、微粒ＷＣによる靱性の低下を抑制できたためと思われる。
【００５１】
また、ＥＰＭＡにて組成分析を行ったが、各層間でのＣｏ、Ｎｉ元素の移動は比較的少なく、従来の製造法による焼結体で問題があった層間の成分の拡散が抑制され、各層間には亀裂の発生もなくしっかりと接合されていた。
【００５２】
ＷＣ基超硬合金はＷＣ粒径が小さいほど硬度が高く、ＷＣ粒径が大きいほど靱性が高くなることから、本構造の焼結体は上部側で比較的高靱性で耐摩耗性に優れ、下部側では上部層よりもさらに靱性に優れるため、通常相反する両特性を両立することのできる材料となっている。
【００５３】
（実施例５）
硬質相として平均粒径0.25μmのＷＣ、結合相として平均粒径 0.5μmのＣｏを12wt％配合し、10時間の混合粉砕を行った粉末Ａと、硬質相として平均粒径0.25μmのＷＣ、結合相として平均粒径 3.0μmのＣｏを12wt％配合し、10時間の混合粉砕を行った粉末Ｂを用意した。そして、粉末Ａが上部層となるようにそれらを層状にプレスして、黒鉛型に充填し、30MPa の圧力を上下方向から負荷しながら昇温スピード 190℃／分となるように黒鉛型に電流を通じ、1250℃に達した時点で２分間キープし、 200℃／分の速度で冷却を行うことによって硬質合金を作製した。
【００５４】
得られた直径30mm、厚み８mmの円板状焼結体の加圧軸に平行な断面を＃250 の砥石で平面研削後、鏡面研磨して光学顕微鏡により観察したところ、上部層にはアスペクト比が約２のＣｏ相、下部層にはアスペクト比が約８のＣｏ相が部分的に見られた。さらに、ダイヤモンド製ヴィッカース圧子を用いた破壊靱性測定でも下部層ほど高靱性を示した。これは扁平なＣｏが亀裂進展エネルギーを吸収することによって高い靱性を示したと思われる。
【００５５】
また、ＥＰＭＡにて組成分析を行ったが、各層間でのＣｏ元素の移動は比較的少なく、従来の製造法による焼結体で問題があった層間の成分の拡散が抑制され、各層間には亀裂の発生もなくしっかりと接合されていた。
【００５６】
結合相金属のアスペクト比が大きな硬質合金ではアスペクト比が小さな合金よりも靱性が高くなる。しかし、結合相金属のアスペクト比が大きな硬質合金ではミクロ的にはＣｏの分散が不均一であり、微視的な耐摩耗性が要求される用途ではアスペクト比の小さな合金の方が耐摩耗性に優れるケースもある。そのようなケースでは、本構造の焼結体は上部側で耐摩耗性に優れ、下部側で靱性に優れるため、通常相反する両特性を両立することのできる材料となっている。なお、下部層の結合相量を上部層よりも多くすることでさらに靱性に優れた合金とすることもできる。
【００５７】
（実施例６）
平均粒径0.25μmのＷＣ粉末と平均粒径３μmのＣｏ粉末を20wt％配合し、10時間混合粉砕した粉末を黒鉛型中で鋼の基体上に配置した。そして、60MPa の圧力を上下方向から負荷しながら昇温スピードを 190℃／分となるように黒鉛型に電流を通じ、1300℃に達した時点で１分間キープし、 100℃／分の速度で冷却を行うことによって硬質合金を鋼上に接合した。
【００５８】
得られた直径50mm、厚み20mmの円板状焼結体の加圧軸に平行な断面を＃250 の砥石で平面研削後、鏡面研磨して光学顕微鏡により観察したところ、上部層（焼結体表面側）にはアスペクト比が約８のＣｏが部分的に見られた。さらに、ダイヤモンド製ヴィッカース圧子を用いた硬度、破壊靱性測定でも高硬度、高靱性を示した。これは通電加圧焼結により、ＷＣ粒径が約 0.3μmの微粒組織とできたことで高硬度を実現し、扁平なＣｏが亀裂進展エネルギーを吸収することで微粒ＷＣによる靱性の低下を抑制できたためと思われる。
【００５９】
また、ＥＰＭＡにて組成分析を行ったが、各層間でのＣｏ元素の移動は比較的少なく、従来の製造法による焼結体で問題があった層間の成分の拡散が抑制され、各層間には亀裂の発生もなくしっかりと接合されていた。
【００６０】
本構造の焼結体は上部層は粒径の細かいＷＣからなっているため高耐摩耗性、下部層は鋼としたことによる高強度、高靱性を得ることができ、通常相反する両特性を両立することのできる材料となっている。
【００６１】
（実施例７）
平均粒径 0.5μmのＴｉＣＮと平均粒径 3.0μmのＮｉ粉末を12wt％配合し、10時間ボールミルで混合粉砕した粉末Ａと、平均粒径２μmのＴｉＣＮ粉末と平均粒径２μmのＮｉ粉末を12wt％配合し、ボールミルで10時間混合粉砕した粉末Ｂを用意し、この二つの粉末を表６に示す割合に配合して５種類の原料粉末No.4-1〜4-5を作製した。
【００６２】
【表６】
図３に示すように、これらの原料粉末１をNo.4-1が上部側、No.4-5が下部側（基体表面側）となるように順に黒鉛型２内で鋼の基体３上に配置した。そして、上下部加圧ラム４，５により50MPa の圧力を上下方向から負荷しながら昇温スピードを 150℃／分となるように黒鉛型に電流を通じ、1150℃に達した時点で３分間キープし、約 200℃／min の速度で冷却を行うことによって硬質合金を鋼上に接合した。この図では原料粉末１の断面構造を省略化しているが、実際には積層構造となっている。なお、上下部加圧ラム４，５に接続されているのは電源６、黒鉛型２に設置されているのは熱電対７である。
【００６４】
得られた直径30mm、厚み10mmの円板状焼結体の加圧軸に平行な断面を＃250 の砥石で平面研削後、鏡面研磨して光学顕微鏡により観察したところ、上部層にはアスペクト比が約10のＮｉ相が部分的に見られ、下部の組織になるほどアスペクト比は減少し、下部層にはアスペクト比が約３のＮｉ相が部分的に見られた。ダイヤモンド製ヴィッカース圧子を用いた硬度、破壊靱性測定では、上部層は高硬度、高靱性を示し、下部層では上部層よりもさらに高靱性を示した。これは上部層ではＴｉＣＮ粒径が約 0.6μmの微粒組織とできたことで高硬度を実現し、扁平なＣｏが亀裂進展エネルギーを吸収することで微粒ＷＣによる靱性の低下を抑制でたものと思われる。さらに下部層では結合相のアスペクト比は小さいもののＴｉＣＮ粒径が大きくなっているため、上部層よりも高靱性を示したものと思われる。
【００６５】
また、ＥＰＭＡにて組成分析を行ったが、各層間でのＮｉ元素の移動は比較的少なく、従来の製造法による焼結体で問題があった層間の成分の拡散が抑制され、各層間には亀裂の発生もなくしっかりと接合されていた。
【００６６】
ＴｉＣＮ−ＮｉのサーメットはＴｉＣＮ粒径が小さいほど硬度が高くなることから、本構造の焼結体は上部側で耐摩耗性に優れ下部側で靱性に優れるため、通常相反する両特性を両立することのできる材料となっている。
【００６７】
（実施例８）
平均粒径 0.5μmのＴｉＣＮ粉末に平均粒径 3.0μmのＮｉ粉末を20wt％添加した後、アトライターで30時間混合粉砕後した粉末Ａと、平均粒径 0.5μmのＴｉＣＮ粉末に平均粒径 3.0μmのＮｉ粉末を20wt％添加した後、３時間混合粉砕した粉末Ｂを用意した。そして、実施例７で行ったのと同様にして、これらの混合割合の異なる５種類の粉末の積層を黒鉛型内で鋼の基体上に行い、40MPa の圧力を上下方向から負荷しながら昇温スピード 200℃／分となるように黒鉛型に電流を通じ、1120℃に達した時点で５分間キープし、 100℃／分の速度で冷却を行うことによって硬質合金を鋼上に接合した。
【００６８】
得られた直径50mm、厚み30mmの円板状焼結体の加圧軸に平行な断面を＃250 の砥石で平面研削後、鏡面研磨して光学顕微鏡により観察したところ、上部層にはアスペクト比が約６のＮｉ相が見られ、下部層（接合面側）にはアスペクト比が約12のＮｉ相が部分的に見られた。さらに、ダイヤモンド製ヴィッカース圧子を用いた破壊靱性測定でも下部層の方が高靱性を示していた。これは粉末Ａに比べて、粉末Ｂでは混合粉砕時間が短いことによって原料時点でのＮｉの粒径が大きく、分散が不均一であることから、その後の通電加圧焼結によって、アスペクト比の大きい扁平な形状をしたＮｉ相が下部層で多く存在し、これらが亀裂進展エネルギーを吸収したため、下部層の破壊靱性は大きくなったものと思われる。
【００６９】
また、ＥＰＭＡにて組成分析を行ったが、各層間でのＮｉ元素の移動は比較的少なく、従来の製造法による焼結体で問題があった層間の成分の拡散が抑制され、各層間には亀裂の発生もなくしっかりと接合されていた。
【００７０】
本構造の焼結体は上部層はＮｉの分散を均一にすることによって高耐摩耗性、下部層はＮｉの分散が不均一で高靱性、そして高強度、高靱性の鋼層となっていることによって、通常相反する両特性を両立することのできる材料となっている。
【００７１】
なお、本実施例の中で、粉末Ｂの硬質相として平均粒径５μmのＷＣ、結合相金属として平均粒径３μmのＣｏ粉末を用いれば、扁平な結合相組織の効果のみでなく、ＷＣ粒がＴｉＣＮ粒子に比べて高靱性であること、硬質相粒子が粗大化した効果などにより、さらに優れた靱性を下部層に保持させることができる。
【００７２】
【発明の効果】
以上説明したように、本発明硬質合金はアスペクト比が５〜20となる形状の結合相組織を含む断面を有し、高い靱性を具えている。また、硬質相粒子が微粒のものは同時に高い硬度も具える。そのため、高靱性が要求される切削工具や耐摩耗部材、耐衝撃用工具などに利用することができる。特に、厚さ方向に組成の異なる合金とすることで、合金の厚さ方向一方側と厚さ方向他方側とで相反する特性を有する合金とできる。
【００７３】
また、本発明製造方法は、本発明硬質合金を製造するのに最適な方法で、短時間による焼結が可能なため、コストダウンに寄与できる。
【図面の簡単な説明】
【図１】本発明硬質合金を加圧軸と平行な面で切断した断面の組織を示す顕微鏡写真である。
【図２】本発明硬質合金を加圧軸と垂直な面で切断した断面の組織を示す顕微鏡写真である。
【図３】本発明硬質合金を製造する装置の概略図である。
【図４】従来の硬質合金の断面組織を示す顕微鏡写真である。
【符号の説明】
１ 原料粉末 ２ 黒鉛型 ３ 基体 ４ 上部加圧ラム
５ 下部加圧ラム ６ 電源 ７ 熱電対[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a tough hard alloy and a method for producing the same.
[Prior art]
[0002]
In general, a hard alloy whose hard phase is mainly WC and whose binder phase is an iron group metal such as Co or Ni is called a WC-based cemented carbide, and its hard phase is mainly TiC (N) and the binder phase is an iron group. The hard alloy used as a metal is called a TiC (N) -based cermet.
[0003]
These hard alloys are generally sintered at a temperature of 1350 ° C. or higher and 1600 ° C. or lower for 1 hour in a vacuum without pressure. In some cases, HIP (hot isostatic pressing) may be performed at a temperature lower than the sintering temperature. It is known that a liquid phase is generated under such sintering conditions, and the WC grains are likely to undergo grain growth during sintering due to the dissolution and reprecipitation phenomenon.
[0004]
The cross-sectional structure photograph of the hard alloy produced by these methods is shown in FIG. The white spots in the figure are the binder phase, but any cut surface exists only at the grain boundaries of the hard phase or between the grains. There are few clear shapes that can be called bonded phase grains, and even if they exist, they have shapes with an aspect ratio of less than 3. The reason for this structure is that when the binder phase dissolves and a liquid phase is generated during sintering, WC-based cemented carbide and TiC (N) -based cermet have high wettability between the hard phase and the binder phase. This is because the liquid phase flows into the hard phase grain boundaries and gaps.
[0005]
On the other hand, WC-based cemented carbide and cermet use fine raw materials to produce materials with high hardness, and VC, Cr to suppress grain growthThree C2 Efforts have been made to produce ultrafine hard alloys by adding compounds such as NbC, TaC, and TiC and sintering them at a low temperature that can be densified. Especially for WC-based cemented carbide, VC and CrThree C2 Is added to produce a hard alloy having a WC structure with an average particle size of about 0.5 μm (see Japanese Patent Application Laid-Open Nos. 1-215947, 4-289146, and 5-98385).
[0006]
[Problems to be solved by the invention]
However, the hard alloy having the above structure cannot be said to have sufficient toughness. In addition, a WC-based cemented carbide having an average particle size of about 0.5 μm or less cannot be industrially produced. For this reason, there is a limit to the field in which hard alloys can be used.
Accordingly, a main object of the present invention is to provide a hard alloy having higher toughness and excellent hardness and strength and a method for producing the same.
[0007]
[Means for Solving the Problems]
In order to achieve the above object, the present invention provides at least one selected from WC, TiC, TiN and TiCN.Consist ofHard phase and iron group metalsConsist ofWith a binder phase,A hard alloy in which the binder phase is formed by current-pressure-sintering with a powder-phase powder having an average particle size of three times or more that of the hard-phase powder, and is parallel to the pressure axis of a sintering apparatus that performs current-pressure sintering In cross sectionAspect ratio of 5-20FlatIn shape,The array has directionality so as to extend in the direction perpendicular to the pressure axis.Bond phase structureincludingIt is characterized by that. In particular, it is preferable that the maximum length of the binder phase structure having an aspect ratio of 5 to 20 is three times or more the average crystal grain size of the hard phase.
[0008]
Here, the aspect ratio refers to the ratio between the maximum length of the binder phase structure and the average thickness. As will be described later, the hard alloy of the present invention is obtained by electric pressure sintering, and the binder phase structure is crushed by the pressure applied at that time. As a result, a structure having the aspect ratio is formed. Since such a binder phase structure is formed by pressing during sintering,To extend in the direction perpendicular to the pressure axisIt has directionality in the arrangement and becomes a flat shapeThe NaNeedless to say, the alloy of the present invention may contain inevitable impurities. Inevitable impurities include, for example, Al, Ba, Ca, Cu, Fe, Mg, Mn, Ni, Si, Sr, S, O, N, Mo, Sn, Cr and the like.
[0009]
Conventionally, it has been considered that a hard alloy having such a non-uniform bond phase structure is not preferable as a liquid phase sintered material. However, as will be shown later in Examples, the fracture toughness of the hard alloy of the present invention is higher than that of the conventional alloy. This is presumably because crack propagation energy is expended on plastic deformation of the binder phase at this time because the crack always crosses the flat binder phase during the crack extension.
[0010]
Such a hard alloy preferably further comprises the following requirements alone or in combination.
(1) The average particle size of the hard phase is 0.01 to 1 μm.
WC-based cemented carbide and TiC (N) -based cermet can improve the hardness of the sintered body by making the hard phase particles fine. However, if the hard phase particle size is reduced, the hardness increases, but the toughness decreases and the brittleness becomes very brittle. In particular, when the hard phase particle size is smaller than 1 μm, the tendency becomes remarkable.
[0011]
According to the production method of the present invention to be described later, a hard alloy having a hard phase in the above-mentioned limited range can be obtained. In addition, according to this method, the shape of the bonding metal can be made flat, so that the toughness can be improved, and both the hardness and the toughness that have conventionally been in a contradictory relationship can be achieved. In particular, the effect is large when the particle size of the hard phase is 0.01 to 1 μm. Use of a raw material smaller than 0.01 μm is industrially expensive, and an alloy having a particle size larger than 1 μm has a small effect of improving toughness. A particularly preferred particle size is 0.1 to 1 μm.
[0012]
The hard phase particle size of the alloy obtained here mainly depends on the hard phase particle size of the raw material used. In the current technology, a WC powder produced by a direct carbonization method or a WC powder refined by a grinding process may be used. Even when finer WC powder is developed in the future, a cemented carbide having a finer WC can be produced by applying the present invention.
[0013]
(2) It contains at least one selected from Cr, V, a carbide of Cr, and a carbide of V, and the total content thereof is 1 wt% or less with respect to the amount of the binder phase.
By limiting the addition amount of such a grain growth inhibitor, the hard phase grain size of the obtained alloy can be made fine and the bending strength can be improved. If the temperature is rapidly raised by an electric pressure sintering method and sintering is performed at a low temperature for a short time, the grain growth of the hard phase can be suppressed and the atomization can be realized. However, sintering with the addition of a grain growth inhibitor (VC or Cr carbide) can only provide a bending strength lower than that of an alloy produced by the prior art. And it has been found that this cause is in an abnormal structure that seems to be caused by an aggregate of carbides of VC and Cr.
[0014]
For this reason, an alloy having a fine hard phase can be produced by setting the addition amount of the grain growth inhibitor to 1 wt% or less with respect to the amount of the binder phase, and an alloy having a higher bending strength than the conventional one can be produced. The addition amount of Cr, V, Cr carbide, and V carbide is preferably not added, but it is considered that Cr and V are mixed in the raw material powder in the form of inevitable impurities. Therefore, it is more preferable that the total content of these grain growth inhibitors is 0.3 wt% or less with respect to the amount of the binder phase.
[0015]
( 3-1. ) In a cross section parallel to the pressure axisA binder phase structure having an aspect ratio of 5 to 20But,Hard alloyOne side in the thickness direction, which is the pressing direction when conducting electrical pressure sinteringHaveThe other side in the thickness directionThe structure changes in the thickness direction so that it does not exist.
( 3-2 ) Hard alloyOne side of the thickness directionThe average particle size of the hard phase in the range of 0.01 to 1 μm,The other side in the thickness directionInAverage particle size of hard phaseAre not in the same range, the particle size of the hard phase changes in the thickness direction.
( 3-3 ) Hard alloyOne side in the thickness directionFromThe other side in the thickness directionThe amount of the binder phase is changing toward.
[0016]
One side in the thickness directionWhenThe other side in the thickness directionHard alloys with different hardness and toughness have been proposed in the past, but with the conventional sintering method, the grain growth of the hard phase and the movement of the liquid phase during the sintering are severe, making it possible to produce the target alloy. was difficult. In the method of the present invention to be described later, since there is little grain growth of the hard phase and liquid phase movement during the sintering, the hard phase particle size and bonding in the thickness directionphaseHard alloys with different structures and binder phases can be produced. In addition, the method of changing these compositions includes both stepwise and substantially continuous.
[0017]
(4) Bonded on a metallic substrate.
Conventionally, the joining strength by the brazing of the sintered body and the metal substrate has been insufficient, but by joining the sintered body, a high joining strength can be obtained and the brazing step can be omitted.
[0018]
The method for producing the hard alloy is at least one selected from WC, TiC, TiN and TiCN.Consist ofHard phase powder and more than 3 times hard phase powderaverageIron group metal with particle sizeConsist ofA step of mixing the binder phase powder, a step of placing a raw material member composed of the mixed powder in an electric heating device, and a step of applying current and pressure sintering the raw material member at 1100 ° C. to 1350 ° C. and 5 to 200 MPa. It is characterized by comprising. Of binder phase powderaverageThe particle size is preferably 5 times or more.
[0019]
Here, the raw material member includes each raw material powder itself, a previously pressed green compact, an intermediate sintered body, a laminate of these, and the like. Moreover, what is necessary is just to add grain growth suppression materials, such as a high melting point compound, when mixing raw material powder as needed. Examples of the high melting point compound include carbides, nitrides, and carbonitrides of group IVa, Va, and VIa elements. It is most preferable that the grain growth inhibitor is not added. When adding, minimize as much as possible. In particular, the total content of Cr, V, Cr carbide, and V carbide is 1 wt% or less based on the amount of the binder phase. A more preferable content is 0.3 wt% or less.
[0020]
If the temperature is rapidly raised by this method and sintering is performed at a low temperature for a short time, the binder phase particles are formed in a flat shape that is crushed in the direction of the pressure axis because the time for the binder phase to move is not sufficient. .
[0021]
Sintering is desirably performed in the presence of a liquid phase. When sintering is performed by solid phase sintering, a flat binder phase structure is likely to be generated, but densification is promoted by performing sintering by causing a liquid phase to appear. Thereby, the intensity | strength of a sintered compact improves. Therefore, a dense alloy having excellent strength, toughness, and hardness can be produced by generating a liquid phase at a temperature at which the flat binder phase structure is not lost and sintering in a short time.
[0022]
The reasons for limiting the above sintering conditions are as follows.
This is because if the sintering temperature is less than 1100 ° C., densification hardly proceeds, and if it exceeds 1350 ° C., the liquid phase tends to stain. The sintering temperature here refers to the temperature of the graphite mold surface when the sintering furnace is controlled. Actual sample temperature is lower than this temperature
It seems that the temperature is higher by 150 ° C to 300 ° C.
[0023]
Moreover, if the applied pressure is 5 MPa or less, the effect of pressure sintering is not observed, and it is difficult to increase the applied pressure from 200 MPa, which causes a cost increase. A particularly preferred pressure is 10 to 50 MPa. The reason is that an inexpensive graphite mold can be used.
[0024]
Furthermore, the sintering time is preferably within 10 minutes. By shortening the sintering time, the grain growth of the hard phase and the movement of the liquid phase during the sintering are suppressed.Bonded phaseDifferent amounts of hard alloys can be made. More preferably, it is within 5 minutes. The sintering atmosphere is preferably a vacuum of 0.1 torr or less.
[0025]
In order to produce a hard alloy whose hard phase particle size and structure change in the thickness direction, multiple types of mixed powders with different hard phase (binding phase) particle size and binding phase amount are prepared by changing the composition of the raw material powder. You just have to. When preparing a plurality of types of mixed powders having different binder phase particle sizes, one of these mixed powders includes a binder phase having a particle size three times larger than that of the hard phase powder. And in the process of arrange | positioning a raw material member to an electrical heating apparatus, these multiple types of mixed powder are laminated | stacked and arrange | positioned in order of the particle size order of the hard phase (binding phase) particle | grains or the amount of binding phases. If there are few types of mixed powders prepared, hard alloys with different hard phase (binding phase) particle sizes and bonding phase amounts can be obtained stepwise in the thickness direction. If the thickness is reduced, it is possible to obtain a hard alloy whose hard phase (binding phase) particle size and amount of binding phase change substantially continuously. In the method of the present invention, since the hard phase grain growth and liquid phase movement during sintering are small, a sintered body having such a configuration can be stably produced.
[0026]
In addition, in order to join such a hard alloy having an inclined structure on a substrate, the raw material member may be disposed in an energizing and pressing device together with the substrate. At that time, it is desirable to increase the particle size of the hard phase (binding phase) on the bonding surface side and to decrease the particle size on the opposite surface side.
[0027]
Hereinafter, embodiments of the present invention will be described.
Example 1
A WC powder having an average particle size of 1 μm, a Co powder having an average particle size of 1 μm, a TiCN powder having an average particle size of 1.5 μm, a TiC powder having an average particle size of 2 μm, and a Ni powder having an average particle size of 1 μm were prepared. And mixed and pulverized with a ball mill for 20 hours to prepare raw material powders (No. 1-1 to 1-7). Further, raw material powders Nos. 2-2 to 2-7 were prepared in the same manner by changing the Co powder and Ni powder into coarse raw materials having average particle diameters of 3 μm and 5 μm, respectively.
[0028]
[Table 1]
Next, these powders were charged into a graphite mold, and an electric heating and sintering apparatus was used to pass a current through the graphite mold so that the heating rate was 190 ° C / min while applying a pressure of 50 MPa from the vertical direction. When the temperature reached 1130 ° C., it was kept for 5 minutes and cooled at a rate of about 100 ° C./min to obtain a sintered body (sample Nos. 1 to 14) having a shape of 25 × 8 × 5 mm.
[0030]
These sintered bodies are cut along a plane parallel to the pressure axis, the cross-section is ground, and after mirror polishing, an arbitrary three-field structure photograph is taken with an optical microscope at a magnification of 1500 times. The aspect ratio of the binder phase metal was calculated. Here, the aspect ratio was calculated by dividing the maximum length of the binder phase grains by the average thickness. Further, hardness and fracture toughness were measured by an indentation method with a load of 50 kg using a diamond Vickers indenter. Furthermore, the bending strength was also measured by a three-point bending test. These measurement results are shown in Table 2.
[0031]
[Table 2]
From Table 2, the aspect ratio of the binder phase of the sintered bodies of raw materials No. 2-1 to 2-7 using binder phase metal powder that is three times larger than the main hard phase particle size is the same as that of the hard phase. It is larger than that of the alloy obtained by sintering the raw materials No. 1-1 to 1-7 having no difference in phase particle diameter, and is a value of about 5 to 15. It was also confirmed that the fracture toughness of these sample Nos. 2, 4, 6, 8, 10, 12, and 14 is significantly superior to the fracture toughness of sample Nos. 1, 3, 5, 7, 9, 11, and 13. did it.
[0033]
FIG. 1 shows an optical micrograph of a cross-section obtained by cutting the sintered body along a plane parallel to the pressing axis and performing surface grinding and mirror polishing. The white part of the photo is the binder phase. As is apparent from this photograph, some of the binder phases have a shape extending long in the direction perpendicular to the pressure axis, that is, a shape having an aspect ratio of 5 to 20. In addition, it can be seen that the binder phase metals confirmed in the structure photograph are arranged so as to extend in a direction perpendicular to the pressure axis of the sintering apparatus, and have a directivity.
[0034]
Next, FIG. 2 shows an optical micrograph of a cross section obtained by cutting the sintered body along a plane perpendicular to the pressure axis and performing surface grinding and mirror polishing. Again, the white part represents the binder phase. It can be seen that the binder phase grains shown in the center of the photograph are circular, and this binder phase metal has a flat shape.
[0035]
Furthermore, bending of sample Nos. 1, 2, 9, 10, 11, and 12 produced using raw powder Nos. 1-1, 2-1, 1-5, 2-5, 1-6, and 2-6 The strength was compared. VC, CrThree C2 The bending strength of the samples No. 9 and 10 in which the added amount of 5% by weight with respect to the amount of the binder phase was lower than those of No. 1 and 2. However, the bending strength of sample Nos. 11 and 12, which is 1 wt% or less, is VC, CrThree C2 It was the same as No. 1 and 2 with no additive. In particular, the fracture toughness of sample 12 in which the aspect ratio of the binder phase metal is in the range of 5 to 20 is higher than that of the sample No. 11 alloy having an aspect ratio of 3, and the bending strength, fracture toughness and hardness are high. I understand that.
[0036]
(Example 2)
A raw material powder No. 3-1 having the same composition as that of the raw material powder No. 1-1 produced in Example 1 and having an WC powder having an average particle size of 0.25 μm was produced. In addition, the raw material powder No. 3-2 having the same composition as the raw material powder No. 1-2 produced in Example 1, with the TiCN powder changed to an average particle size of 0.5 μm and the Ni powder to an average particle size of 1.5 μm. Was made.
[0037]
These powders were sintered in the same manner as in Example 1 by changing the holding time at the maximum keeping temperature as described in Table 3 using an electric heating and sintering apparatus to obtain samples No. 15 to No. 22. It was. The cross-sectional structures of these samples subjected to mirror polishing were photographed with FE-SEM, and the average particle size of the main hard phase was calculated from the Fullman equation using photographs. Further, the aspect ratio, Hv hardness and fracture toughness of the binder phase metal of these samples were measured in the same manner as in Example 1, and the results are shown in Table 3.
[0038]
[Table 3]
From the results of Table 3, it can be seen that a WC-based cemented carbide and a TiC (N) -based cermet having an average particle size of 1 μm or less can be produced by an electric pressure sintering method. For this reason, these alloys showed high hardness. Moreover, since these alloys have binder phase metal grains having an aspect ratio of 5 to 20, it has been found that the fracture toughness value is very large and the hardness and toughness are compatible at a high level.
[0040]
In addition, as a result of measurement with a PR thermocouple, the actual sample temperature in the electric current pressure sintering method performed in this example was found to be about 1380 ° C. This temperature was higher than the melting point 1320 ° C. of the eutectic composition of the WC-base cemented carbide, and it was considered that a liquid phase had appeared at least partially.
[0041]
(Example 3)
WC powder with an average particle size of 0.25 μm, Co powder with an average particle size of 1.3 μm, VC powder with an average particle size of 1 μm, Cr with an average particle size of 1.5 μmThree C2 Powders were prepared, blended in the composition described in Table 4, and mixed and ground for 10 hours with an attritor to prepare raw material powders (raw material powder Nos. 3-3 to 3-11).
[0042]
[Table 4]
1 ton / cm of these raw material powders2 The mold was pressed at a pressure of 1, and the press body was set in a sintering furnace. In a vacuum of 0.01 Torr or less, the heating rate was 10 ° C / min, the maximum keeping temperature was 1350 ° C, the keeping time was 1 hour, the cooling rate was 5 ° C / min. Sintering was performed under min conditions (conventional liquid phase sintering method conditions). Thereafter, the HIP process was performed under conditions of 1320 degrees, 1 hour, and 100 MPa in argon, to obtain a sintered body having a shape of 25 × 8 × 5 mm (sample Nos. 23 to 31: see Table 5).
[0044]
These sintered bodies were subjected to surface grinding and mirror polishing, and then a structure photograph was taken by FE-SEM, and the average particle diameter of WC was calculated by the Fullman equation using the taken pictures. The bending strength was also measured by a three-point bending test with a span of 20 mm. These measurement results are shown in Table 5.
[0045]
[Table 5]
Next, using raw materials No. 3-3 to 3-11, using an electric heating and sintering machine, the graphite mold was made so that the heating rate was 190 ° C / min while applying a pressure of 50 MPa from the vertical direction. A hard alloy (samples Nos. 32 to 40) was produced by keeping the current for 5 minutes when the temperature reached 1130 ° C. and cooling at a rate of about 100 ° C./min. These sample Nos. 32 to 40 were similarly measured for WC average particle size, bending strength, and binder phase aspect ratio. The results are also shown in Table 5.
[0047]
As a result, the minimum WC average particle size obtained by the conventional sintering method is 0.5 μm when VC is added, whereas the current pressure sintering method has an average particle size of 0.3 μm regardless of the raw material composition. It was confirmed that a fine-grained WC alloy could be produced and that it had excellent fracture toughness. However, VC and CrThree C2 It was found that the bending strength of the alloy added with an amount exceeding 1 wt% with respect to the amount of the binder phase was remarkably lowered despite the fine WC grain size. Just VC or CrThree C2 Sample No. 34, 35, 38, 39, 40 alloy with a content of 1 wt% or less of the binder phase content realizes a very good bending strength and a bending strength higher than the conventional sintering method. It turns out that you can.
[0048]
In addition, as a result of measurement with a PR thermocouple, the actual sample temperature in the electric current pressure sintering method performed in this example was found to be about 1380 ° C. This temperature was higher than the melting point 1320 ° C. of the eutectic composition of the WC-base cemented carbide, and it was considered that a liquid phase had appeared at least partially.
[0049]
Example 4
WC as the hard phase, 10 wt% of Co with an average particle size of 3.0 μm as the binder phase, and 2 wt% of Ni, mixed for 10 hours and pulverized with an attritor, divided into two WC particle sizes Kind prepared. The graphite mold is pressed and packed in layers so that a powder with a large WC particle size (average particle size 2.5 μm) is the lower layer and a powder with a small WC particle size (average particle size 0.25 μm) is the upper layer. , While applying a pressure of 41 MPa from the top and bottom, the current is passed through the graphite mold so that the heating rate is 300 ° C / min. When the temperature reaches 1200 ° C, the temperature is kept for 3 minutes, and cooling is performed at 100 ° C / min. Thus, a hard alloy was produced.
[0050]
A cross section parallel to the pressure axis of the disk-shaped sintered body with a diameter of 30 mm and a thickness of 8 mm was ground with a # 250 grindstone, then mirror-polished and observed with an optical microscope. The upper layer had an aspect ratio Co having a flat shape of about 8 was partially seen. In addition, hardness and fracture toughness measurement using a diamond Vickers indenter showed high hardness and high toughness. This is presumably because this layer is mainly composed of fine WC particles of about 0.3 μm and has high hardness, and because flat Co exists, it absorbs crack propagation energy and suppresses toughness deterioration due to fine WC particles.
[0051]
In addition, although composition analysis was performed by EPMA, the movement of Co and Ni elements between layers was relatively small, and diffusion of components between layers that had a problem with a sintered body by a conventional manufacturing method was suppressed. The layers were firmly joined without cracks.
[0052]
Since the WC-based cemented carbide has a higher hardness as the WC particle size is smaller and a toughness becomes higher as the WC particle size is larger, the sintered body of this structure has relatively high toughness and excellent wear resistance on the upper side. Since the lower side is more tougher than the upper layer, it is a material that can satisfy both conflicting characteristics.
[0053]
(Example 5)
WC with an average particle size of 0.25 μm as the hard phase, 12 wt% of Co with an average particle size of 0.5 μm as the binder phase, mixed and ground for 10 hours, and WC with an average particle size of 0.25 μm as the hard phase, As a binder phase, 12 wt% Co having an average particle diameter of 3.0 μm was blended, and powder B was prepared by mixing and grinding for 10 hours. Then, they are pressed into layers so that the powder A becomes the upper layer, filled in the graphite mold, and the current is applied to the graphite mold so that the temperature rise rate is 190 ° C./min while applying a pressure of 30 MPa from the vertical direction. Then, when the temperature reached 1250 ° C., the hard alloy was produced by keeping for 2 minutes and cooling at a rate of 200 ° C./min.
[0054]
A cross section parallel to the pressure axis of the disk-shaped sintered body with a diameter of 30 mm and a thickness of 8 mm was ground with a # 250 grindstone, then mirror-polished and observed with an optical microscope. The upper layer had an aspect ratio Was partially observed in the Co phase, and the Co layer having an aspect ratio of approximately 8 was partially observed in the lower layer. Furthermore, in the fracture toughness measurement using a diamond Vickers indenter, the lower layer showed higher toughness. It seems that flat Co exhibited high toughness by absorbing crack propagation energy.
[0055]
In addition, the composition analysis was performed by EPMA. However, the movement of Co element between the layers was relatively small, and the diffusion of the components between the layers, which was problematic in the sintered body by the conventional manufacturing method, was suppressed. Was firmly joined without cracks.
[0056]
A hard alloy with a large aspect ratio of the binder phase metal has higher toughness than an alloy with a small aspect ratio. However, a hard alloy with a high binder phase metal aspect ratio has a non-uniform Co dispersion microscopically, and an alloy with a lower aspect ratio is more resistant to wear in applications that require microscopic wear resistance. In some cases, it is excellent. In such a case, the sintered body of this structure is excellent in wear resistance on the upper side and excellent in toughness on the lower side, and thus is a material that can satisfy both conflicting characteristics. In addition, it can also be set as the alloy which was further excellent in toughness by making the amount of binder phases of a lower layer larger than an upper layer.
[0057]
(Example 6)
20 wt% of WC powder having an average particle size of 0.25 μm and Co powder having an average particle size of 3 μm were mixed and pulverized for 10 hours, and placed on a steel substrate in a graphite mold. Then, while applying a pressure of 60MPa from the top and bottom, current was passed through the graphite mold so that the heating rate was 190 ° C / min. When the temperature reached 1300 ° C, the temperature was kept for 1 minute and cooled at a rate of 100 ° C / min. The hard alloy was joined on the steel by
[0058]
The cross section parallel to the pressure axis of the disk-shaped sintered body having a diameter of 50 mm and a thickness of 20 mm was ground with a # 250 grindstone, mirror-polished and observed with an optical microscope. Co having an aspect ratio of about 8 was partially observed on the surface side. Furthermore, the hardness and fracture toughness measurement using a diamond Vickers indenter showed high hardness and high toughness. This is achieved by forming a fine grain structure with a WC grain size of approximately 0.3μm by applying electrical pressure and sintering, and flat Co absorbs crack growth energy, thereby suppressing toughness degradation due to fine grain WC. It seems that it was made.
[0059]
In addition, the composition analysis was performed by EPMA. However, the movement of Co element between the layers was relatively small, and the diffusion of the components between the layers, which was problematic in the sintered body by the conventional manufacturing method, was suppressed. Was firmly joined without cracks.
[0060]
The sintered body of this structure has high wear resistance because the upper layer is made of WC with a small particle size, and the lower layer can have high strength and high toughness due to the steel, both of which are contradictory to each other. It is a material that can be compatible.
[0061]
(Example 7)
12 wt% of TiCN having an average particle size of 0.5 μm and Ni powder having an average particle size of 3.0 μm, mixed and ground for 10 hours by a ball mill, 12 wt% of TiCN powder having an average particle size of 2 μm and Ni powder having an average particle size of 2 μm A powder B prepared by mixing and pulverizing with a ball mill for 10 hours was prepared, and these two powders were blended in the proportions shown in Table 6 to prepare five kinds of raw material powders Nos. 4-1 to 4-5.
[0062]
[Table 6]
As shown in FIG. 3, these raw material powders 1 are placed on the steel substrate 3 in the
[0064]
The cross section parallel to the pressure axis of the disk-shaped sintered body having a diameter of 30 mm and a thickness of 10 mm was ground with a # 250 grindstone, then mirror-polished and observed with an optical microscope. The upper layer had an aspect ratio However, the Ni phase having an aspect ratio of about 3 was partially observed in the lower layer. In the hardness and fracture toughness measurement using a diamond Vickers indenter, the upper layer showed high hardness and high toughness, and the lower layer showed higher toughness than the upper layer. This is because the upper layer has a fine structure with a TiCN particle size of about 0.6 μm, and high hardness is achieved, and flat Co absorbs crack growth energy, thereby suppressing toughness degradation due to fine WC. Seem. Furthermore, although the TiCN grain size is large in the lower layer although the binder phase has a small aspect ratio, it seems to have exhibited higher toughness than the upper layer.
[0065]
Moreover, although composition analysis was performed by EPMA, the movement of Ni element between each layer was relatively small, and the diffusion of the components between the layers, which was problematic in the sintered body by the conventional manufacturing method, was suppressed, and between each layer Was firmly joined without cracks.
[0066]
Since the TiCN-Ni cermet has a higher hardness as the TiCN particle size is smaller, the sintered body of this structure is superior in wear resistance on the upper side and excellent in toughness on the lower side, and thus has both conflicting properties. It has become a material that can.
[0067]
(Example 8)
After adding 20 wt% of Ni powder with an average particle size of 3.0 μm to TiCN powder with an average particle size of 0.5 μm, the powder A was mixed and ground for 30 hours with an attritor, and TiCN powder with an average particle size of 0.5 μm had an average particle size of 3.0. After adding 20 wt% of μm Ni powder, powder B was prepared by mixing and grinding for 3 hours. Then, in the same manner as in Example 7, these five kinds of powders having different mixing ratios were laminated on the steel substrate in the graphite mold, and the temperature was raised while applying a pressure of 40 MPa from the vertical direction. The hard alloy was bonded onto the steel by passing an electric current through the graphite mold at a speed of 200 ° C / min, keeping it for 5 minutes when it reached 1120 ° C, and cooling at a rate of 100 ° C / min.
[0068]
The cross section parallel to the pressure axis of the disk-shaped sintered body with a diameter of 50 mm and a thickness of 30 mm was ground with a # 250 grindstone, then mirror-polished and observed with an optical microscope. The upper layer had an aspect ratio A Ni phase having an aspect ratio of about 12 was partially observed in the lower layer (bonding surface side). Furthermore, in the fracture toughness measurement using a diamond Vickers indenter, the lower layer showed higher toughness. Compared with powder A, powder B has a shorter mixing and grinding time, so the Ni particle size at the raw material is large and the dispersion is non-uniform. It seems that the fracture toughness of the lower layer has increased because a large amount of Ni phase having a flat shape exists in the lower layer and these absorbed the crack propagation energy.
[0069]
Moreover, although composition analysis was performed by EPMA, the movement of Ni element between each layer was relatively small, and the diffusion of the components between the layers, which was problematic in the sintered body by the conventional manufacturing method, was suppressed, and between each layer Was firmly joined without cracks.
[0070]
The sintered body of this structure has high wear resistance by making Ni dispersion uniform in the upper layer, and the lower layer is steel layer with non-uniform Ni dispersion and high toughness, and high strength and high toughness. Therefore, it is a material capable of satisfying both characteristics which are usually contradictory.
[0071]
In this example, when WC having an average particle size of 5 μm is used as the hard phase of the powder B and Co powder having an average particle size of 3 μm is used as the binder phase metal, not only the effect of the flat binder phase structure but also WC particles Is more tougher than TiCN particles, and due to the effect of coarsening of the hard phase particles, it is possible to retain even better toughness in the lower layer.
[0072]
【The invention's effect】
As described above, the hard alloy of the present invention has a cross section including a binder phase structure having an aspect ratio of 5 to 20, and has high toughness. In addition, fine particles of hard phase particles have high hardness at the same time. Therefore, it can be used for cutting tools, wear-resistant members, impact-resistant tools and the like that require high toughness. In particular, by using alloys with different compositions in the thickness direction,One side in the thickness directionWhenThe other side in the thickness directionAnd an alloy having the opposite characteristics.
[0073]
Further, the production method of the present invention is an optimal method for producing the hard alloy of the present invention, and can be sintered in a short time, thus contributing to cost reduction.
[Brief description of the drawings]
FIG. 1 is a photomicrograph showing the structure of a cross section of a hard alloy of the present invention cut along a plane parallel to a pressure axis.
FIG. 2 is a photomicrograph showing the structure of a cross section of the hard alloy of the present invention cut along a plane perpendicular to the pressure axis.
FIG. 3 is a schematic view of an apparatus for producing the hard alloy of the present invention.
FIG. 4 is a photomicrograph showing the cross-sectional structure of a conventional hard alloy.
[Explanation of symbols]
1
5 Lower pressure ram 6 Power supply 7 Thermocouple
Claims (16)
鉄族金属からなる結合相とを具え、
結合相が、硬質相粉末の３倍以上の平均粒径を有する結合相粉末で通電加圧焼結されて形成され、通電加圧焼結を行う焼結装置の加圧軸と平行な硬質合金の断面においてアスペクト比が５〜２０となる扁平な形状で、加圧軸と垂直な方向に伸びるように配列に方向性を有している結合相組織を含むことを特徴とする高靱性硬質合金。A hard phase consisting of at least one selected from WC, TiC, TiN and TiCN;
Comprising a binding phase consisting of iron group metals,
A hard alloy in which the binder phase is formed by current-pressure-sintering with a powder-phase powder having an average particle size of three times or more that of the hard-phase powder, and is parallel to the pressure axis of a sintering apparatus that performs current-pressure sintering A high toughness hard alloy comprising a flat phase having an aspect ratio of 5 to 20 in a cross section and including a binder phase structure having orientation in an arrangement so as to extend in a direction perpendicular to the pressure axis .
その合計含有量が結合相量に対して１wt％以下であることを特徴とする請求項１記載の高靱性硬質合金。Further, it contains at least one selected from Cr, V, Cr carbide, V carbide,
2. The high toughness hard alloy according to claim 1, wherein the total content is 1 wt% or less based on the amount of the binder phase.
この混合粉末から構成される原料部材を通電加熱装置に配置する工程と、
この原料部材を1100℃〜1350℃、5〜200MPaで通電加圧焼結する工程とを具えることを特徴とする高靱性硬質合金の製造方法。Mixing WC, TiC, and hard phase powder composed of at least one selected from TiN and TiCN, the raw material powder containing a binder phase powder of iron group metals having an average particle size of more than 3 times the hard phase powder Process,
Arranging the raw material member composed of the mixed powder in an electric heating device;
A method for producing a high toughness hard alloy comprising the step of subjecting the raw material member to electric current pressure sintering at 1100 ° C. to 1350 ° C. and 5 to 200 MPa.
粒成長抑制材の合計含有量が結合相量に対して１wt％以下であることを特徴とする請求項９記載の高靱性硬質合金の製造方法。Furthermore, at least one kind of grain growth inhibitor powder selected from Cr, V, Cr carbide, V carbide is mixed with the hard phase powder and the binder phase powder,
The method for producing a high toughness hard alloy according to claim 9 , wherein the total content of the grain growth inhibitor is 1 wt% or less with respect to the amount of the binder phase.
これら複数種の混合粉末を積層して硬質相粒径が厚さ方向に変化された原料部材を通電加圧焼結することを特徴とする請求項９記載の高靱性硬質合金の製造方法。 Prepare multiple types of raw material powder mixed hard phase powder and binder phase powder so that the hard phase particle size is different,
The method for producing a high toughness hard alloy according to claim 9 , wherein a plurality of kinds of mixed powders are laminated and a raw material member whose hard phase particle size is changed in the thickness direction is subjected to current-pressure sintering.
これらの混合粉末のいずれかは硬質相粉末の３倍以上の粒径を有する結合相粉末を含み、
これら複数種の混合粉末を積層して結合相粒径が厚さ方向に変化された原料部材を通電加圧焼結することを特徴とする請求項９記載の高靱性硬質合金の製造方法。 Prepare multiple types of raw material powders mixed with hard phase powder and binder phase powder so that the particle size of binder phase is different,
Any of these mixed powders includes a binder phase powder having a particle size three times larger than that of the hard phase powder,
The method for producing a high toughness hard alloy according to claim 9 , wherein a plurality of kinds of mixed powders are laminated and a raw material member whose binder phase particle size is changed in the thickness direction is subjected to current-pressure sintering.
これらの混合粉末のいずれかは硬質相粉末の３倍以上の粒径を有する結合相粉末を含み、
これら複数種の混合粉末を積層して結合相量が厚さ方向に変化された原料部材を通電加圧焼結することを特徴とする請求項９記載の高靱性硬質合金の製造方法。 Prepare multiple types of raw material powder mixed hard phase powder and binder phase powder so that the binder phase content is different,
Any of these mixed powders includes a binder phase powder having a particle size three times larger than that of the hard phase powder,
The method for producing a high toughness hard alloy according to claim 9 , wherein a plurality of kinds of mixed powders are laminated and a raw material member whose binder phase amount is changed in the thickness direction is subjected to current-pressure sintering.
通電加圧装置には原料部材と基体との複合体を配置し、
この複合体を通電加圧焼結して、基体に原料部材の焼結体を焼結接合することを特徴とする請求項９記載の高靱性硬質合金の製造方法。After mixing each powder of the hard phase and the binder phase, comprising the step of placing a raw material member made of this mixed powder on the base of the metal material,
In the energizing and pressing apparatus, a composite of the raw material member and the substrate is arranged
10. The method for producing a high toughness hard alloy according to claim 9, wherein the composite is subjected to current and pressure sintering, and the sintered body of the raw material member is sintered and joined to the substrate.
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| CN105177390B (en) * | 2015-08-28 | 2017-03-22 | 北京科技大学 | Metal ceramics and preparation method thereof |
| JP7307394B2 (en) * | 2019-01-09 | 2023-07-12 | 三菱マテリアル株式会社 | WC-based cemented carbide cutting tools and surface-coated WC-based cemented carbide cutting tools with excellent plastic deformation resistance and chipping resistance |
| JP7161677B2 (en) * | 2019-02-18 | 2022-10-27 | 三菱マテリアル株式会社 | WC-Based Cemented Carbide Cutting Tool and Surface-Coated WC-Based Cemented Carbide Cutting Tool with Excellent Fracture Resistance |
| CN115961199B (en) * | 2022-10-26 | 2024-05-17 | 四川轻化工大学 | High-strength high-wear-resistance titanium-based high-manganese steel-bonded ceramic alloy and preparation method thereof |
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