JP3878232B2 - Coated cemented carbide - Google Patents
Coated cemented carbide Download PDFInfo
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- JP3878232B2 JP3878232B2 JP00202895A JP202895A JP3878232B2 JP 3878232 B2 JP3878232 B2 JP 3878232B2 JP 00202895 A JP00202895 A JP 00202895A JP 202895 A JP202895 A JP 202895A JP 3878232 B2 JP3878232 B2 JP 3878232B2
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Description
【０００１】
【産業上の利用分野】
この発明は、切削工具に適する、耐摩耗性、耐欠損性に優れる被覆超硬合金、特に、ＷＣ基超硬合金を母材とする被覆超硬合金に関するものである。
【０００２】
【従来の技術及びその課題】
切削工具に使用される被覆超硬合金は、耐摩耗性と耐欠損性という相反する性質を満足する必要がある。
【０００３】
ところで、ＷＣ基超硬合金母材の最表面に、ＷＣと結合金属のみからなる層（脱β層）を設け、その内側の領域のＢ１型結晶を小さくすることにより、耐摩耗性と耐欠損性の向上を図ることが従来から提案されている。
【０００４】
ところが、脱β層の内側部分のＢ１型結晶の粒度を小さくすると、耐摩耗性は向上するが、強度、即ち耐欠損性が低下するという問題がある。これは、亀裂が脱β層を通り抜けた後に、その直下で急激に亀裂進展が促進されるため、脱β層の効果が十分発揮されないことによる。
【０００５】
一方、Ｂ１型結晶の粒度を大きくすれば、強度は向上するが、耐塑性変形性が低下するという問題がある。
【０００６】
そこで、この発明は、脱β層の内側部分のＢ１型結晶の粒度の分布等を工夫することにより、耐摩耗性と耐欠損性をうまくバランスさせた、切削工具に好適なＷＣ基超硬合金を母材とする被覆超硬合金を得ようとするものである。
【０００７】
【課題を解決するための手段とその作用】
この発明におけるＷＣ基超硬合金母材には、Ｂ１型結晶が５〜４０ｗｔ％含まれる。５ｗｔ％未満の場合には、脱β層の内側の０〜１００μｍにおけるＢ１型結晶が、さらにその内側の十分内部のＢ１型結晶に比べ、半価幅の比率が９０％より大きくなってその差が小さくなり、脱β層をとおり抜けた亀裂の伝播を阻止できず、強度の向上効果がほとんどない。一方、４０ｗｔ％以上の場合には、合金内部の強度低下が顕著になる。
【０００８】
また、この発明におけるＷＣ基超硬合金母材には、結合相として５〜１５ｗｔ％の鉄族金属が含まれている。５ｗｔ％未満の場合には、焼結性が低下して強度が低下し、１５ｗｔ％以上の場合には、耐塑性変形性が低下する。
【０００９】
また、この発明におけるＷＣ基超硬合金母材の表面付近には、５〜５０μｍ、好ましくは１０〜３０μｍの脱β層を有する。５μｍよりも薄い場合には、亀裂の導入に対する抵抗が不足するため耐欠損性の向上が十分得られず、５０μｍ以上の場合には、耐塑性変形性が低下する。
【００１０】
また、この発明におけるＷＣ基超硬合金母材は、脱β層の内側１０〜２００μｍの領域を、Ｂ１型結晶が、この領域のさらに内側部分のＢ１型結晶と比較し、格子定数の差が０〜０．０１、好ましくは０〜０．００５であり、かつＸ線回折の（４２２）回折面における半価幅の比率が７０〜９０％となるようにしている。この領域の厚みが１０μｍよりも薄いと強度向上効果が小さくなり、２００μｍを超えると耐摩耗性の低下につながる。そして、上記格子定数の差が０〜０．０１で半価幅の比率を７０〜９０％とするのは、次の理由からである。格子定数の差が０．０１を超えると、この領域と、この領域の内側部分とで、Ｂ１型結晶の組成の差が大きくなりすぎるので、耐摩耗性と強度とのバランスが崩れることになる。また、半価幅の比率が７０％より小さいと、この領域と、この領域の内側部分とで粒度の差が大きくなりすぎ、即ち、脱β層の直下の領域においてＢ１型結晶の粒度が粗すぎたり、あるいは逆に、この領域の内側部分のＢ１型結晶の粒度が細かくなりすぎるので、強度と耐摩耗性とのバランスが崩れてしまうことになる。逆に、半価幅の比率が９０％より大きくなると、脱β層の直下の領域におけるＢ１型結晶の粗粒化による強度向上効果が十分得られなくなる。
【００１１】
次に、脱β層の直下の領域の内側部分における半価幅の絶対値は０．４〜１．０であることが好ましい。０．４より大きいことにより、上記領域の内側部分の耐摩耗性向上と上記領域の強度向上による両性能バランスの向上効果が大きくなる。但し、上記領域の内側部分のＢ１型結晶の半価幅が１．０を超えると、上記領域のＢ１型結晶の半価幅もそれにつられて大きくなるので、上記両性能のバランスが崩れる傾向を示す。即ち、耐摩耗性は大きく向上するが、耐欠損性が低下する傾向を示す。
【００１２】
また、上記Ｂ１型結晶は、IVａ族元素とＶａ族元素を主成分とする金属の炭化物、窒化物、炭窒化物及びこれらの固溶体からなり、IVａ族金属成分の量（Ｍ１「ｍｏｌ」）とＶａ族金属成分の量（Ｍ２「ｍｏｌ」）をｍｏｌ比で表した場合に、以下の関係にあるようにすると、耐摩耗性と耐欠損性のバランスがとりやすくなる。
【００１３】
０．２≦Ｍ１／（Ｍ１＋Ｍ２）≦０．９（好ましくは０．７）
特に、IVａ族金属がＺｒ、Ｖａ族金属がＴａとＮｂの少なくとも一方を主成分とした場合、より効果的である。
【００１４】
ところで、この発明の被覆超硬合金は、次のようにして製造することができる。
【００１５】
まず、ＷＣ、鉄族金属、Ｂ１型結晶炭化物の原料をこの発明の範囲で配合、成形し、１３５０〜１５００℃の範囲で焼結する。但し、合金内部の格子定数は原料の配合により調整する。また、内部の半価幅は配合原料の粒度と、焼結温度、時間により調整する。この際、昇温中の１２５０℃以上から焼結キープの間で、温度上昇に伴う平衡窒素分圧の上昇にあわせ、窒素量を徐々に変化（上昇）させ、常に一定の弱脱窒雰囲気に保つ。そしてこの雰囲気で焼結キープ中も保ち、毎分５℃以上の速度で冷却する。
【００１６】
上記平衡窒素分圧は合金系（Ｂ１型結晶の種類）によって異なるがＢ１型結晶が（ＴｉＴａＮｂ）Ｎ系の場合、経験的に１５００℃で約１０ｔｏｒｒ、１２５０℃で約１ｔｏｒｒ程度と考えられる。
【００１７】
従って、脱β層の形成は、脱窒による超硬合金表面付近でのＢ１型結晶の結合層への溶解と脱β層直下付近での再析出による物であるため、これらの溶解析出速度の制御で脱β層自体とのその直下のＢ１型結晶を制御できる。
【００１８】
上記方法で制御した脱β層、脱β層直下のＢ１型結晶は、冷却速度を遅くすると冷却中に変化してしまい、焼結キープ中までに保たれていたこの発明の構造を保てなくなる。従って、５℃／ｍｉｎ以上の冷却速度が必要となる。
【００１９】
そして、脱β層の厚みと脱β層直下の半価幅の調整は上記方法で窒素分圧の調整及び焼結キープ時間により行うことができる。
【００２０】
【実施例】
（実施例１）
原料粉末として、ＷＣ、ＴｉＣ、ＴｉＮ、ＴａＣ、ＴａＮ、ＮｂＣ、ＮｂＮ、ＶＣ、ＶＮ及びＣｏ、Ｎｉを用意して、表１に示す組成からなる完粉をＣＮＭＧ１２０４０８の形状でチップにプレス後、１２５０℃から１４５０℃まで５℃／ｍｉｎで減圧Ｎ2 雰囲気中で、温度上昇とともにＮ2 分圧を０．０５ｔｏｒｒから１ｔｏｒｒまで連続的に上昇させながら昇温後、１時間２ｔｏｒｒ窒素雰囲気中で保持後、５℃／ｍｉｎで冷却した。
【００２１】
ついで、この合金を基体として、通常のＣＶＤで内層に５μｍＴｉＣ、ＴｉＣＮ、ＴｉＮ、中間層にＴｉＢＮ、ＴｉＣＮＯ、外層に酸化アルミニウム、酸化ジルコニウム、酸化ハフニウムの複層を２μｍ被覆して下記の切削条件で切削テストを行った。
【００２２】
また、比較の為に試料Ｎo.２、３と同じ組成で、従来の焼結法である真空焼結（昇温、キープ、冷却ともに真空）を行った試料２’、３’も作製し、上記と同じコーティング膜を被覆した。
【００２３】
なお、内部については、試料を表面から５００μｍ以上平研を用いて粗研削した後♯１５００及び♯３０００のダイヤモンドペーストによるバフ研磨により表面を鏡面仕上げ、脱β層直下の領域については♯１５００より粗い番手のダイヤモンドペーストによるバフ研磨で表面から順次追い込んでいき、脱β層がなくなるのを光学顕微鏡で確認した後♯３０００のダイヤモンドペーストバフ研磨で鏡面仕上げする。この面についてＸ線回折装置を用いて、（４２２）面の回折ピークで格子定数及び半価幅を測定する。
【００２４】
これらのサンプルを用い、下に示すテスト１及びテスト２の条件で切削し、その時の逃げ面摩耗量（ｍｍ）と欠損率（％）を測定した。
【００２５】
テスト １ （耐摩耗性テスト）
切削速度 ２２０ｍ／ｍｉｎ 被削材 ＳＣＭ４３５
送り ０．４５ｍｍ／ｒｅｖ 切削時間 ２０ｍｉｎ
切込み ２．０ｍｍ
テスト ２ （靱性テスト）
切削速度 １００ｍ／ｍｉｎ 被削材 ＳＣＭ４３５ ４溝材
送り ０．４５ｍｍ／ｒｅｖ 切削時間 ３０秒
切込み ２．０ｍｍ ８回繰り返し
なお、表１中にIVａ族金属成分の量をＭ１ｍｏｌ、Ｖａ族金属成分の量をＭ２ｍｏｌとした時のＭ１／（Ｍ１＋Ｍ２）の値を同時に示した。
【００２６】
テスト結果及びＢ１型結晶の格子定数、半価幅等を比較品と合わせ表２に示す。
【００２７】
なお、表２中のＡ、Ｂ、Ｃ、Ｄは、
｜Ａ−Ｂ｜：脱β層の内側部分（Ａ）と十分内部（Ｂ）でのＢ１型結晶の格子定
数の差
Ｃ：脱β層の内側におけるＢ１型結晶の（４２２）回折面における半価幅
Ｄ：十分内部におけるＢ１型結晶の（４２２）回折面における半価幅
１００＊Ｃ／Ｄ：脱β層の内側部分と十分内部でのＢ１型結晶の（４２２）回折
面における半価幅の比（％）
Ｅ：脱β層の内側の領域の厚み（μｍ）
Ｆ：脱β層の厚み（μｍ）
【００２８】
【表１】
【表２】
（実施例２）
原料粉末として、ＷＣ、ＺｒＣ、ＺｒＮ、ＴａＣ、ＴａＮ、ＮｂＣ、ＮｂＮ、ＶＣ、ＶＮ及びＣＯ、Ｎｉを用意して、表３に示す組成からなる完粉をＣＮＭＧ１２０４０８の形状でチップにプレス後、１２５０℃から１４５０℃まで２℃／ｍｉｎでＮ２＋Ｈ２の減圧雰囲気中で昇温し、温度上昇とともにＮ２分圧を０．０１ｔｏｒｒから２ｔｏｒｒまで連続的に上昇させながら昇温（Ｈ２分圧は５ｔｏｒｒで一定）後、１時間２ｔｏｒｒの窒素ｐ雰囲気中で保持し、１０℃／ｍｉｎで冷却した。
【００３１】
ついで、この合金を基本として、通常のＣＶＤで内層に５μｍＴｉＣ、ＴｉＣＮ、ＴｉＮ、中間層にＴｉＢＮ、ＴｉＣＮＯ、外層に酸化アルミニウム、酸化ジルコニウム、酸化ハフニウムの複層を２μｍ被覆して下記の切削条件で切削テストを行った。また、試料１２、１３の組成で従来焼結である真空焼結（昇温、キープ、冷却とも）を真空度を変えて行った従来比較品１２’、１３’及び１２″、１３″を表４中に併せてのせた。
【００３２】
これらの試料を用い、実施例１のテスト１とテスト２と同条件で切削評価を行った。試料Ｎo.１２及び１３はＢ１型結晶の量、結合層量、Ｍ１／（Ｍ１＋Ｍ２）値ともに実施例１のＮo.３及び４とほぼ同じであり、Ｂ１型結晶組成のみが１２及び１３ではＺｒとＴａを主成分とする金属の炭窒化合物等としたものであるが、これらの比較により、この発明の効果は一段と顕著になっていることが判る。
【００３３】
【表３】
【表４】
【発明の効果】
以上のように、この発明は、ＷＣ基超硬合金母材の脱β層の直下部分に、内部との組成の差が小さく、かつ内部よりも粗い亀裂伝播抵抗に優れるＢ１型結晶を配置して、強度（耐欠損性）と耐摩耗性を向上させつつ、さらに内部はそれよりも細いＢ１型結晶を配置することにより、耐摩耗性を向上させたものであるから、耐摩耗性と耐欠損性のバランスを大きく向上させた、切削工具に好適な被覆超硬合金が得られる。
【図面の簡単な説明】
【図１】この発明に係る被覆超硬合金の断面組織を示す模式図[0001]
[Industrial application fields]
The present invention relates to a coated cemented carbide suitable for a cutting tool and excellent in wear resistance and fracture resistance, and more particularly to a coated cemented carbide based on a WC-based cemented carbide.
[0002]
[Prior art and problems]
The coated cemented carbide used for cutting tools must satisfy the conflicting properties of wear resistance and fracture resistance.
[0003]
By the way, on the outermost surface of the WC-base cemented carbide base material, a layer (de-β layer) consisting only of WC and a binding metal is provided, and by reducing the B1 type crystal in the inner region, wear resistance and fracture resistance are reduced. It has been proposed to improve the performance.
[0004]
However, when the grain size of the B1-type crystal in the inner part of the de-β layer is reduced, the wear resistance is improved, but there is a problem that the strength, that is, the fracture resistance is lowered. This is because after the crack has passed through the de-β layer, the crack growth is rapidly promoted immediately below the crack, so that the effect of the de-β layer is not sufficiently exhibited.
[0005]
On the other hand, if the grain size of the B1 type crystal is increased, the strength is improved, but there is a problem that the plastic deformation resistance is lowered.
[0006]
In view of this, the present invention provides a WC-based cemented carbide suitable for a cutting tool, in which wear resistance and fracture resistance are well balanced by devising the distribution of the grain size of the B1 type crystal inside the de-β layer. It is intended to obtain a coated cemented carbide with a base material.
[0007]
[Means for solving the problems and their functions]
The WC-based cemented carbide base material according to the present invention contains 5 to 40 wt% of B1-type crystals. In the case of less than 5 wt%, the B1 type crystal at 0 to 100 μm inside the de-β layer is more than 90% of the ratio of the half-value width compared to the B1 type crystal inside sufficiently further inside. , The propagation of cracks passing through the de-β layer cannot be prevented, and there is almost no effect of improving the strength. On the other hand, in the case of 40 wt% or more, the strength reduction inside the alloy becomes remarkable.
[0008]
Further, the WC-based cemented carbide base material in the present invention contains 5 to 15 wt% of an iron group metal as a binder phase. If it is less than 5 wt%, the sinterability is lowered and the strength is lowered, and if it is 15 wt% or more, the plastic deformation resistance is lowered.
[0009]
Moreover, it has 5-50 micrometers, Preferably it is 10-30 micrometers de-beta layer near the surface of the WC group cemented carbide base material in this invention. When the thickness is less than 5 μm, the resistance to the introduction of cracks is insufficient, so that the fracture resistance cannot be sufficiently improved. When the thickness is 50 μm or more, the plastic deformation resistance is lowered.
[0010]
Further, in the WC-based cemented carbide base material according to the present invention, the area of 10-200 μm inside the deβ layer is compared with the B1 type crystal in the further inner part of this area, and the difference in lattice constant is It is 0 to 0.01, preferably 0 to 0.005 , and the ratio of the half width on the (422) diffraction plane of X-ray diffraction is 70 to 90%. If the thickness of this region is thinner than 10 μm, the effect of improving the strength is reduced, and if it exceeds 200 μm, the wear resistance is reduced. The reason why the difference in lattice constant is 0 to 0.01 and the ratio of the half width is 70 to 90% is as follows. If the difference in lattice constant exceeds 0.01, the difference between the composition of the B1 type crystal between this region and the inner part of this region becomes too large, and the balance between wear resistance and strength is lost. . Also, if the half width ratio is less than 70%, the difference in grain size between this region and the inner part of this region becomes too large, that is, the grain size of the B1 type crystal is coarse in the region directly under the de-β layer. To the contrary, or conversely, since the grain size of the B1 type crystal in the inner portion of this region becomes too fine, the balance between strength and wear resistance is lost. On the other hand, when the ratio of the half width exceeds 90%, the effect of improving the strength due to the coarsening of the B1 type crystal in the region immediately below the deβ layer is not obtained.
[0011]
Next, it is preferable that the absolute value of the half width in the inner portion of the region immediately below the de-β layer is 0.4 to 1.0. When the ratio is larger than 0.4, the effect of improving both the performance balance by improving the wear resistance of the inner portion of the region and improving the strength of the region is increased. However, when the half width of the B1 type crystal in the inner portion of the region exceeds 1.0, the half width of the B1 type crystal in the region also increases accordingly. Show. That is, the wear resistance is greatly improved, but the fracture resistance tends to decrease.
[0012]
The B1 type crystal is composed of a metal carbide, nitride, carbonitride, and solid solution thereof, each of which includes a group IVa element and a group Va element as a main component, and the amount of the group IVa metal component (M1 “mol”) When the amount of the Va group metal component (M2 “mol”) is expressed in terms of a mole ratio, it is easy to balance the wear resistance and the fracture resistance if the following relationship is satisfied.
[0013]
0.2 ≦ M1 / (M1 + M2) ≦ 0.9 (preferably 0.7)
In particular, it is more effective when the IVa group metal is mainly composed of Zr and the Va group metal is composed mainly of at least one of Ta and Nb.
[0014]
By the way, the coated cemented carbide of the present invention can be manufactured as follows.
[0015]
First, raw materials of WC, iron group metal, and B1 type crystal carbide are blended and molded within the scope of the present invention, and sintered at 1350 to 1500 ° C. However, the lattice constant inside the alloy is adjusted by the blending of raw materials. Also, the internal half width is adjusted by the particle size of the blended raw material, the sintering temperature, and the time. At this time, the amount of nitrogen is gradually changed (increased) in accordance with the increase in the equilibrium nitrogen partial pressure accompanying the temperature rise from 1250 ° C. or higher during the temperature increase to the sintering keep, and a constant weak denitrification atmosphere is always obtained. keep. And it keeps also in a sintering keep in this atmosphere, and cools at a speed | rate of 5 degreeC or more per minute.
[0016]
The equilibrium nitrogen partial pressure varies depending on the alloy system (type of B1 type crystal), but when the B1 type crystal is a (TiTaNb) N system, it is empirically considered to be about 10 torr at 1500 ° C. and about 1 torr at 1250 ° C.
[0017]
Therefore, the formation of the de-β layer is due to the dissolution of the B1 type crystal in the bonding layer near the cemented carbide surface by denitrification and the re-precipitation near the de-β layer. Control can control the B1-type crystal directly below the de-beta layer itself.
[0018]
The B1-type crystal directly under the de-β layer and the de-β layer controlled by the above method changes during cooling when the cooling rate is slowed down, and the structure of the present invention maintained until the sintering keep cannot be maintained. . Therefore, a cooling rate of 5 ° C./min or more is required.
[0019]
The thickness of the de-β layer and the half width immediately below the de-β layer can be adjusted by adjusting the nitrogen partial pressure and sintering keep time by the above method.
[0020]
【Example】
Example 1
As raw material powders, WC, TiC, TiN, TaC, TaN, NbC, NbN, VC, VN, Co, and Ni are prepared, and a complete powder having the composition shown in Table 1 is pressed into a chip in the form of CNMG120408, and then 1250 in vacuo N 2 atmosphere ° C. from 5 ° C. / min up to 1450 ° C., while continuously increased from 0.05torr to 1torr the N 2 partial pressure with increasing temperature after heating, after maintained in 1 hour 2torr nitrogen atmosphere Cooled at 5 ° C./min.
[0021]
Then, using this alloy as a base, the inner layer is coated with 5 μm TiC, TiCN, TiN, the intermediate layer is TiBN, TiCNO, the outer layer is coated with 2 μm of aluminum oxide, zirconium oxide, and hafnium oxide by the following cutting conditions. A cutting test was performed.
[0022]
For comparison, Samples 2 ′ and 3 ′ having the same composition as Samples No. 2 and 3 and subjected to vacuum sintering (vacuum for heating, keeping and cooling), which is a conventional sintering method, are also produced. The same coating film as above was coated.
[0023]
As for the inside, the sample was roughly ground from the surface by using a plane polishing of 500 μm or more, and then the surface was mirror-finished by buffing with # 1500 and # 3000 diamond paste, and the region directly under the de-β layer was rougher than # 1500. The surface is sequentially driven from the surface by buffing with a count diamond paste, and it is confirmed with an optical microscope that the β-free layer disappears, and then mirror-finished by # 3000 diamond paste buffing. With respect to this surface, the lattice constant and the half width are measured at the diffraction peak of the (422) surface using an X-ray diffractometer.
[0024]
Using these samples, cutting was performed under the conditions of Test 1 and Test 2 shown below, and the flank wear amount (mm) and the defect rate (%) at that time were measured.
[0025]
Test 1 (Abrasion resistance test)
Cutting speed 220m / min Work material SCM435
Feed 0.45mm / rev Cutting time 20min
Cutting depth 2.0mm
Test 2 (Toughness test)
Cutting speed 100 m / min Work material SCM435 Four-groove material feed 0.45 mm / rev Cutting time 30 seconds infeed 2.0 mm Repeated 8 times In Table 1, the amount of IVa group metal component is M1 mol, the amount of Va group metal component The value of M1 / (M1 + M2) when M was 2 mol was simultaneously shown.
[0026]
Table 2 shows the test results and the lattice constant, half-value width, etc. of the B1-type crystal together with the comparative product.
[0027]
In Table 2, A, B, C and D are
| A−B |: Difference in lattice constant of the B1 type crystal between the inner part (A) and the sufficiently inner part (B) of the deβ layer. C: Half of the (422) diffraction plane of the B1 type crystal inside the deβ layer. Valence width D: Half width at the (422) diffraction plane of the B1 type crystal sufficiently inside 100 * C / D: Half width at the (422) diffraction plane of the B1 type crystal inside the inner part of the de-beta layer and inside sufficiently Ratio (%)
E: Thickness (μm) of the area inside the de-β layer
F: De-β layer thickness (μm)
[0028]
[Table 1]
[Table 2]
(Example 2)
As raw material powders, WC, ZrC, ZrN, TaC, TaN, NbC, NbN, VC, VN and CO, Ni are prepared. After pressing the finished powder having the composition shown in Table 3 into the shape of CNMG120408, 1250 The temperature is raised in a reduced pressure atmosphere of N2 + H2 at a rate of 2 ° C./min from 1 ° C. to 1450 ° C., and the temperature is raised while the N 2 partial pressure is continuously increased from 0.01 torr to 2 torr as the temperature rises (H 2 partial pressure is constant at 5 torr) Thereafter, it was kept in a nitrogen p atmosphere of 2 torr for 1 hour and cooled at 10 ° C./min.
[0031]
Then, based on this alloy, the inner layer is coated with 5 μm TiC, TiCN, TiN, the intermediate layer is TiBN, TiCNO, the outer layer is coated with 2 μm of aluminum oxide, zirconium oxide, and hafnium oxide by the following cutting conditions. A cutting test was performed. Further, conventional comparative products 12 ′, 13 ′ and 12 ″, 13 ″ obtained by changing the degree of vacuum by vacuum sintering (temperature rising, keeping, cooling), which is the conventional sintering with the composition of Samples 12 and 13, are shown. 4 was put together.
[0032]
Using these samples, cutting evaluation was performed under the same conditions as in Test 1 and Test 2 of Example 1. Samples No. 12 and No. 13 are almost the same as No. 3 and No. 4 in Example 1 in terms of the amount of B1 type crystal, the amount of the bonding layer, and the M1 / (M1 + M2) value. It is understood that the effects of the present invention are more prominent by comparing these with a carbon carbonitride compound of a metal mainly containing Ta and Ta.
[0033]
[Table 3]
[Table 4]
【The invention's effect】
As described above, according to the present invention, a B1 type crystal having a small composition difference from the inside and having a coarser crack propagation resistance than the inside is disposed in a portion immediately below the de-β layer of the WC-based cemented carbide base material. Thus, while improving the strength (breakage resistance) and wear resistance, the wear resistance is improved by arranging a B1 type crystal that is thinner than that inside. A coated cemented carbide suitable for a cutting tool with a greatly improved balance of chipping properties can be obtained.
[Brief description of the drawings]
FIG. 1 is a schematic diagram showing a cross-sectional structure of a coated cemented carbide according to the present invention.
Claims (4)
０．２≦Ｍ１／（Ｍ１＋Ｍ２）≦０．９The B1 type crystal is composed of a metal carbide, nitride, carbonitride, and solid solution of a metal having a group IVa element and a group Va element as main components, and the amount of the group IVa metal component (M1 “mol”) and the group Va metal component The coated cemented carbide according to claim 1, wherein the amount of M2 (M2 “mol”) is expressed in terms of a mole ratio, and has the following relationship.
0.2 ≦ M1 / (M1 + M2) ≦ 0.9
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP00202895A JP3878232B2 (en) | 1995-01-10 | 1995-01-10 | Coated cemented carbide |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP00202895A JP3878232B2 (en) | 1995-01-10 | 1995-01-10 | Coated cemented carbide |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JPH08188846A JPH08188846A (en) | 1996-07-23 |
| JP3878232B2 true JP3878232B2 (en) | 2007-02-07 |
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| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP00202895A Expired - Lifetime JP3878232B2 (en) | 1995-01-10 | 1995-01-10 | Coated cemented carbide |
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| Country | Link |
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| JP (1) | JP3878232B2 (en) |
Families Citing this family (8)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US6638474B2 (en) * | 2000-03-24 | 2003-10-28 | Kennametal Inc. | method of making cemented carbide tool |
| SE529590C2 (en) * | 2005-06-27 | 2007-09-25 | Sandvik Intellectual Property | Fine-grained sintered cemented carbides containing a gradient zone |
| BRPI0618596B1 (en) | 2005-11-17 | 2018-03-06 | Boehlerit Gmbh & Co. Kg | PROCESS FOR COATING A TOOL OR PART OF A TOOL, APPLIED TO AN OBJECT, WITH AT LEAST ONE LAYER OF TITANIUM CARBONITRET, AND TOOL OR PART OF THE TOOL |
| AT503050B1 (en) * | 2005-11-17 | 2007-09-15 | Boehlerit Gmbh & Co Kg | Coating a tool with titanium, zirconium, hafnium, vanadium, niobium, tantalum or chromium carbonitride by chemical vapor deposition comprises increasing the temperature during deposition |
| JP5460301B2 (en) * | 2009-12-24 | 2014-04-02 | 京セラ株式会社 | Cutting tools |
| JP7098969B2 (en) * | 2018-03-09 | 2022-07-12 | 住友電気工業株式会社 | Cemented Carbide, Cutting Tools Containing It, Cemented Carbide Manufacturing Method and Cutting Tool Manufacturing Method |
| CN109097654B (en) * | 2018-08-22 | 2020-07-31 | 株洲欧科亿数控精密刀具股份有限公司 | Numerical control blade for heavy-load machining and preparation method thereof |
| CN113667874B (en) * | 2020-05-15 | 2022-02-11 | 四川大学 | YA-class gradient hard alloy material and preparation method and application thereof |
-
1995
- 1995-01-10 JP JP00202895A patent/JP3878232B2/en not_active Expired - Lifetime
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| Publication number | Publication date |
|---|---|
| JPH08188846A (en) | 1996-07-23 |
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