JP3803694B2 - Nitrogen-containing sintered hard alloy - Google Patents
Nitrogen-containing sintered hard alloy Download PDFInfo
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Description
【０００１】
【産業上の利用分野】
この発明は耐熱衝撃性、耐摩耗性に優れ、特に切削工具に適応した場合に極めて優れた性能を発揮する窒素含有焼結硬質合金に関するものである。
【０００２】
【従来の技術】
窒素含有焼結硬質合金は、従来から切削工具として利用されているが、切削条件によっては、熱衝撃性に対する弱さから熱亀裂の発生を招き、その亀裂を起点として大破することがあった。熱亀裂の発生原因は窒素含有焼結硬質合金の熱伝導率の低さにある。切削中、工具の表面にのみ熱がこもり温度は上昇するが、内部は温度があまり上昇しない。よって表面部は内部と比較して熱による膨張が大きくなる。この状態で切削を終了し工具が冷やされると表面部は収縮しようとするが、内部に拘束されることで十分収縮しきれず、結果として熱亀裂が発生する。
【０００３】
【発明が解決しようとする課題】
本発明は窒素含有焼結硬質合金の耐熱衝撃性を向上させた窒素含有焼結硬質合金を提供するものである。
【０００４】
【課題を解決するための手段】
前述の如く熱亀裂の発生は、表面部と内部との温度差によって生じる。考えられる手段としては、窒素含有焼結硬質合金そのものの熱伝導率を向上させる手段が挙げられるが、窒素含有焼結硬質合金の熱伝導率向上にはおのずと限界がある。しかしながら窒素含有焼結合金の表面部に熱伝導率の高いＷＣリッチ（残部Ｃｏ、Ｎｉを主成分とする金属結合相）な層を配置させると切削中に発生する熱を合金全体に伝導させ、熱発散（フィン）効果があることが研究の結果明らかとなった。
【０００５】
この発明は上記研究結果にもとづいてなされたものであって、ＷＣを必須とし周期律表の４ａ、５ａ、６ａ族から選ばれた少なくとも１種類の遷移金属の炭化物、窒化物、炭窒化物あるいはこれらの複合炭窒化物からなる硬質相と残部がＮｉ及びＣｏならびに不可避的不純物を含む結合相とからなる窒素含有焼結硬質合金において、
（１）合金表面部にＮｉ、Ｃｏが主成分である金属結合相とＷＣから構成されるシミダシ層１が存在し（図１参照）、層１は内部方向に３層に分かれ、最外層はＷＣが０体積％以上３０体積％以下（好ましくは、０〜５体積％）で残部がＣｏ、Ｎｉを主成分とする金属結合相で構成され、中間層はＷＣが５０体積％以上１００体積％以下（好ましくは、８０〜１００体積％）で残部がＣｏ、Ｎｉを主成分とする金属結合相で構成され、最下層はＷＣが０体積％以上３０体積％以下（好ましくは、０〜５体積％）で残部がＣｏ、Ｎｉを主成分とする金属結合相で構成され、かつ最外層、最下層の厚みがそれぞれ０．１μｍ〜１０μｍ以下（好ましくは、０．１μｍ〜０．５μｍ）で構成され、中間層の厚みが０．５μｍ以上１０μｍ以下（好ましくは、０．５μｍ〜５μｍ）であることを特徴とするものである。表面部に前述の層を構成している窒素含有焼結硬質合金は、実施例でも明らかなように耐熱衝撃性が著しく向上する。尚、最外層と最下層は実質Ｎｉ、Ｃｏを主成分である金属結合相リッチになってしまうが、製法上避けられない生成層であり、規定内の厚みであると性能上問題はない。
【０００６】
（２）合金表面部にＮｉ、Ｃｏが主成分である金属結合相とＷＣから構成されるシミダシ層１の直下に全く金属結合相を含まないか、含んだとしても２体積％以下であり、かつその領域がシミダシ層１の直下から内部方向に２μｍ以上１００μｍ以内（好ましくは、２μｍ〜５０μｍ）であれば、この領域が非常に高硬度な領域となり、耐摩耗性と耐熱衝撃性を両立させうることができる。また、
【０００７】
（３）合金表面部にＮｉ、Ｃｏを主成分とする金属結合相とＷＣから構成されるシミダシ層１の直下にＷＣを全く含まないか、含んだとしても２体積％以下である領域が１μｍ以上、５００μｍ以下（好ましくは、２０μｍ〜１００μｍ）であるか、又は該条件下でかつ該領域直下から内部方向へＷＣの体積％が漸増し、シミダシ層１の直下から１ｍｍ以内（好ましくは、０．３〜０．７ｍｍ）でＷＣ体積％が合金平均ＷＣ体積％になることで構成される窒素含有焼結硬質合金は、ＷＣが存在することで合金全体のヤング率が向上し、機械的強度に対して非常に強くなる。よって表面部にはＷＣを存在させずに内部にのみ存在させることで耐熱衝撃性と耐欠損性を両立させることが可能となる。
【０００８】
（４）（２）と（３）の特徴を合わせると耐熱衝撃性、耐摩耗性、耐欠損性の３特性の向上が可能となる。
以上４つの知見を得た。
【０００９】
次に数値限定の理由について述べる。まず（１）において中間層が５０体積％以上１００体積％以下と規定したのは、ＷＣが５０体積％以下（残部Ｃｏ、Ｎｉを主成分とした金属結合相）であると所望の熱伝導率が得られずに熱発散層としての役割を果たさない為である。又この中間層の厚みを０．５μｍ以上１０μｍ以下と規定したのは、０．５μｍ未満だと所望の熱伝導が得られず、１０μｍより厚くなると著しい耐摩耗性の劣化を招く為である。
尚、最外層および最下層は最も重要な層である中間層を得る為には必然的に形成される層であり、少なくとも０．１μｍは必要であるが、厚みが１０μｍを越えると切削時に被削材の主成分である鉄と溶着を起こし、欠損にいたる場合がある。１０μｍ以下であると切削性能に対し影響はないことが研究結果より判明した為である。
【００１０】
次に（２）においてＣｏ、Ｎｉを主成分とする金属結合相が表面部に２体積％以下と規定しているのは、それ以上の比率で金属結合相が存在すると耐摩耗性の著しい向上が認められない為である。また、その領域が２μｍ未満では耐摩耗性の向上は認められず、１００μｍを越えると逆に硬くなりすぎて脆くなり耐欠損性の劣化を招くためである。
【００１１】
さらに（３）において、ＷＣを含まないか含んだとしても２体積％以下の領域が１μｍ未満であるとＷＣによる硬度低下の影響を受け耐摩耗性の劣化を招き、５００μｍを越えるとＷＣによる合金そのものの靱性強化の恩恵にこうむれなくなるということが種々の試験の結果判明した為である。
尚、本発明合金の構造は、規定の組成において焼結温度を１３５０℃〜１７００℃とし、焼結雰囲気、冷却速度を操作することで得られる。層１内の３層の厚みは焼結雰囲気、冷却速度で制御することが可能である。
【００１２】
尚、ＷＣの体積％については次のような測定手法を用いた。ＷＣ体積％が既知であるWC-Co の超硬合金の断面をラッピングし、４８００倍のＳＥＭ写真を撮る。その写真を画像解析装置により、写真内のＷＣの占める面積を算出し、ＷＣ体積％と画像解析装置によるＷＣの占める面積との検量線を引く。本発明の合金については、観察したい部位を断面ラッピングし、４８００倍のＳＥＭ写真からＷＣの占める面積を画像解析装置より算出、検量線よりＷＣの体積％を求めた。
【００１３】
【実施例】
次にこの発明の切削工具を実施例により具体的に説明する。原料粉末として、平均粒径１．５μｍのＴｉＣＮ粉末、ＷＣ粉末、ＴａＣ粉末、ＮｂＣ粉末、Ｍｏ2 Ｃ粉末、ＶＣ粉末、（Ｔｉ0.5 Ｗ0.3 Ｔａ0.1 Ｎｂ0.1 ）Ｃ0.5 Ｎ0.5 粉末、Ｃｏ粉末、Ｎｉ粉末を用意し、これら原料粉末を表１−１に示される組成に配合し、湿式アトライターにて１２時間混合した後、１．５ton/cm2 の圧力でＣＮＭＧ４３２形状の圧粉体を作成し、この圧粉体にホーニング処理を施した後、表１−２に示す焼結条件で表２−１〜表２−９の構造をした焼結硬質合金を作成した。これらの表で、「シミダシ層直下から内部への構造」とは、シミダシ層直下を０とし合金内部に向かっての深さに応じて変化する硬質相や結合相の組成割合を示す。例えば、試料番号ａ−７では、シミダシ層直下のＷＣ量は、シミダシ層直下から内部まで合金平均ＷＣ体積％になり、一方結合相量は、２．５μm まで１．８体積％で、２．５μm 〜６０μm の間で漸増し、６０μm より内部では合金平均結合相体積％となる。各深さでの残部の硬質相量は、硬質相量＝１００─（合金平均結合相体積％）─（合金平均ＷＣ体積％）で表される。
尚、焼結条件は表１−２に示す。
【００１４】
【表１】
【表３】
【表４】
【表５】
【表６】
【表７】
【表８】
【表９】
【表１０】
【表１１】
（実施例１）
ａ−１からａ−１５の試料について（Ａ）耐熱衝撃性試験及び（Ｂ）耐摩耗試験を行った。結果を表３に示す。
【表１２】
ＴｉＣＮとＷＣが硬質相である焼結硬質合金において、規定通りのシミダシ層を兼ね備えていれば、従来以上の耐熱衝撃性が得られていることが分かる。さらに規定内の結合相分布があれば耐摩耗性が、規定内のＷＣ分布があればさらに耐熱衝撃性が向上することが分かる。
【００２７】
（実施例２）
ｂ−１からｂ−１５の試料について（Ｃ）耐熱衝撃性試験および（Ｄ）耐摩耗試験を行った。結果を表４に示す。
【表１３】
４ａ，５ａ，６ａ族が硬質相である焼結硬質合金において規定通りのシミダシ層を兼ね備えていれば、従来以上の耐熱衝撃性が得られることが分かる。さらに規定内の結合相分布があれば耐摩耗性が、規定内のＷＣ分布があればさらに耐熱衝撃性が向上することが分かる。
【００３０】
（実施例３）
ｃ−１からｃ−１５の試料について（Ｅ）耐熱衝撃性試験および（Ｆ）耐摩耗試験を行った。結果を表５に示す。
【表１４】
４ａ，５ａ，６ａ族の固溶体硬質層を原料とした焼結硬質合金において規定通りのシミダシ層を兼ね備えていれば、従来以上の耐熱衝撃性が得られることが分かる。さらに規定内の結合相分布があれば耐摩耗性が、規定内のＷＣ分布があればさらに耐熱衝撃性が向上することが分かる。
【００３３】
（実施例４）
ａ−１とａ−２とｄ−１の試料について（Ｇ）耐熱衝撃性試験を行った。結果を表６に示す。
【表１５】
シミダシ層を兼ね備えていても、ＷＣを主成分とする層が存在しなければ、耐熱衝撃性の向上は認められないことが分かる。
【００３６】
【発明の効果】
表面部のシミダシ層にＷＣを主成分とする層が存在すれば、耐熱衝撃性が各段に向上し、かつシミダシ層直下に規定の結合相分布、ＷＣ分布を配置させれば、耐摩耗性、さらなる耐熱衝撃性を向上させることが可能である。このような焼結硬質合金工具は従来焼結硬質合金では不可能であった湿式断続加工への適用が可能となる。
【図面の簡単な説明】
【図１】図１（Ａ）は、シミダシ層が３層に分かれ、最外層、最下層にＣｏ，Ｎｉの結合層、中間層にＷＣ層があることを示す合金組織の顕微鏡写真（ＳＥＭ写真）であり、（Ｂ）及び（Ｃ）は夫々同組織中のＣｏ及びＮｉ元素の分布を示す顕微鏡写真（ＥＤＸ分析）である。[0001]
[Industrial application fields]
The present invention relates to a nitrogen-containing sintered hard alloy that is excellent in thermal shock resistance and wear resistance, and exhibits particularly excellent performance when applied to a cutting tool.
[0002]
[Prior art]
Nitrogen-containing sintered hard alloys have been conventionally used as cutting tools. However, depending on cutting conditions, thermal cracking may occur due to weakness against thermal shock, and the crack may be used as a starting point. The cause of thermal cracking is the low thermal conductivity of the nitrogen-containing sintered hard alloy. During cutting, heat is accumulated only on the surface of the tool and the temperature rises, but the temperature does not rise so much inside. Therefore, the surface portion is more thermally expanded than the inside. When cutting is finished in this state and the tool is cooled, the surface portion tends to shrink, but it is not fully contracted by being restrained inside, and as a result, a thermal crack occurs.
[0003]
[Problems to be solved by the invention]
The present invention provides a nitrogen-containing sintered hard alloy having improved thermal shock resistance of the nitrogen-containing sintered hard alloy.
[0004]
[Means for Solving the Problems]
As described above, the occurrence of thermal cracks is caused by the temperature difference between the surface portion and the inside. Possible means include means for improving the thermal conductivity of the nitrogen-containing sintered hard alloy itself, but there is a natural limit to improving the thermal conductivity of the nitrogen-containing sintered hard alloy. However, if a layer with a high WC-rich (metal binding phase composed mainly of Co and Ni) is placed on the surface of the nitrogen-containing sintered alloy, heat generated during cutting is conducted to the entire alloy, Research has revealed that there is a heat dissipation (fin) effect.
[0005]
The present invention has been made based on the above research results, and requires WC as a carbide, nitride, carbonitride of at least one transition metal selected from groups 4a, 5a, and 6a of the periodic table. In a nitrogen-containing sintered hard alloy composed of a hard phase composed of these composite carbonitrides and the balance of Ni and Co and a binder phase containing inevitable impurities in the balance,
(1) On the surface of the alloy, there is a squeeze layer 1 composed of a metallic binder phase mainly composed of Ni and Co and WC (see FIG. 1). The layer 1 is divided into three layers in the inner direction, and the outermost layer is WC is 0% by volume or more and 30% by volume or less (preferably 0 to 5% by volume), and the balance is composed of a metallic binder phase mainly composed of Co and Ni. The intermediate layer has a WC of 50% by volume to 100% by volume. In the following (preferably 80 to 100% by volume), the balance is composed of a metal binder phase mainly composed of Co and Ni, and the lowermost layer has a WC of 0 to 30% by volume (preferably 0 to 5% by volume). %) And the balance is composed of a metallic binder phase mainly composed of Co and Ni, and the outermost layer and the lowermost layer each have a thickness of 0.1 μm to 10 μm or less (preferably 0.1 μm to 0.5 μm). And the thickness of the intermediate layer is 0.5 μm or more and 10 μm or less (preferably , It is characterized in that it is 0.5 m to 5 m). The nitrogen-containing sintered hard alloy having the above-mentioned layer on the surface part has a markedly improved thermal shock resistance, as is apparent from the examples. Although the outermost layer and the lowermost layer are substantially rich in a metal bonded phase mainly composed of Ni and Co, they are generation layers that are unavoidable in terms of the production method, and there is no problem in performance if the thickness is within the specified range.
[0006]
(2) The alloy surface portion does not contain any metal binder phase immediately below the stagnation layer 1 composed of a metal binder phase and WC mainly composed of Ni and Co, and even if included, it is 2% by volume or less. And if the region is 2 μm or more and within 100 μm (preferably 2 μm to 50 μm) in the inner direction from directly under the stigmatic layer 1, this region becomes a very hard region, and both wear resistance and thermal shock resistance are achieved. Can be obtained. Also,
[0007]
(3) The area of the alloy surface portion containing no or only WC immediately below the stagnation layer 1 composed of a metallic binder phase mainly composed of Ni and Co and WC, or 2% by volume or less is 1 μm. As described above, it is 500 μm or less (preferably 20 μm to 100 μm), or the volume percentage of WC gradually increases from the region directly below to the inside under the conditions, and within 1 mm (preferably 0 3 to 0.7 mm), and the nitrogen-containing sintered hard alloy constituted by the WC volume% being the average WC volume% of the alloy improves the Young's modulus of the whole alloy due to the presence of WC, and the mechanical strength It becomes very strong against. Therefore, it is possible to achieve both thermal shock resistance and fracture resistance by allowing the surface portion to exist only inside without WC.
[0008]
(4) When the characteristics of (2) and (3) are combined, it is possible to improve the three characteristics of thermal shock resistance, wear resistance, and fracture resistance.
The above four findings were obtained.
[0009]
Next, the reason for the numerical limitation will be described. First, in (1), the intermediate layer is defined as 50% by volume or more and 100% by volume or less because the WC is 50% by volume or less (metal bonding phase mainly composed of the remaining Co and Ni) and the desired thermal conductivity. This is because it does not serve as a heat dissipating layer. The reason why the thickness of the intermediate layer is specified to be not less than 0.5 μm and not more than 10 μm is that if it is less than 0.5 μm, desired heat conduction cannot be obtained, and if it exceeds 10 μm, the wear resistance is remarkably deteriorated.
The outermost layer and the lowermost layer are inevitably formed in order to obtain an intermediate layer which is the most important layer, and at least 0.1 μm is necessary. In some cases, welding occurs with iron, which is the main component of the cutting material, leading to defects. This is because it has been found from research results that the cutting performance is not affected when the thickness is 10 μm or less.
[0010]
Next, in (2), the metal binder phase containing Co and Ni as the main component is defined as 2% by volume or less on the surface part. When the metal binder phase is present in a higher ratio, the wear resistance is remarkably improved. Is not allowed. In addition, if the area is less than 2 μm, no improvement in wear resistance is observed, and if it exceeds 100 μm, the area becomes too hard and brittle, leading to deterioration of chipping resistance.
[0011]
Further, in (3), even if WC is not included or not, if the area of 2% by volume or less is less than 1 μm, the wear resistance deteriorates due to the effect of hardness reduction by WC, and if it exceeds 500 μm, the alloy by WC It is because it became clear as a result of various tests that it would not be able to suffer from the benefits of its toughness enhancement.
The structure of the alloy of the present invention can be obtained by setting the sintering temperature to 1350 ° C. to 1700 ° C. in a specified composition and operating the sintering atmosphere and the cooling rate. The thickness of the three layers in the layer 1 can be controlled by the sintering atmosphere and the cooling rate.
[0012]
The following measurement method was used for the volume percent of WC. Wrap a cross section of a WC-Co cemented carbide with a known WC volume% and take a SEM photograph at 4800x. The area occupied by the WC in the photograph is calculated for the photograph by an image analyzer, and a calibration curve is drawn between the WC volume% and the area occupied by the WC by the image analyzer. For the alloy of the present invention, the part to be observed was lapped in cross section, the area occupied by WC was calculated from an image analysis apparatus from a 4800 times SEM photograph, and the volume% of WC was obtained from a calibration curve.
[0013]
【Example】
Next, the cutting tool of the present invention will be specifically described with reference to examples. As raw material powder, an average particle size of 1.5 μm TiCN powder, WC powder, TaC powder, NbC powder, Mo 2 C powder, VC powder, (Ti 0.5 W 0.3 Ta 0.1 Nb 0.1 ) C 0.5 N 0.5 powder, Co powder, Prepare Ni powder, mix these raw material powders with the composition shown in Table 1-1, mix for 12 hours in a wet attritor, and then create a CNMG432 shaped green compact with a pressure of 1.5 ton / cm 2 And after performing the honing process to this green compact, the sintered hard alloy which made the structure of Table 2-1-Table 2-9 on the sintering conditions shown in Table 1-2 was created. In these tables, “the structure from directly under the stigma layer to the inside” indicates the composition ratio of the hard phase and the binder phase that changes according to the depth toward the inside of the alloy, with 0 directly under the stigma layer. For example, in Sample No. a-7, the amount of WC immediately under the stigma layer is the average WC volume% of the alloy from directly under the stigma layer to the inside, while the amount of the binder phase is 1.8% by volume up to 2.5 μm. It gradually increases between 5 μm and 60 μm, and within 60 μm, it becomes the alloy average binder phase volume%. The amount of the remaining hard phase at each depth is expressed as hard phase amount = 100— (alloy average bond phase volume%) — (alloy average WC volume%).
The sintering conditions are shown in Table 1-2.
[0014]
[Table 1]
[Table 3]
[Table 4]
[Table 5]
[Table 6]
[Table 7]
[Table 8]
[Table 9]
[Table 10]
[Table 11]
Example 1
Samples a-1 to a-15 were subjected to (A) thermal shock resistance test and (B) wear resistance test. The results are shown in Table 3.
[Table 12]
In a sintered hard alloy in which TiCN and WC are hard phases, it can be seen that thermal shock resistance higher than that of the prior art can be obtained if it has a regular stagnation layer. Further, it can be seen that if the binder phase distribution is within the specified range, the wear resistance is improved, and if the WC distribution is within the specified range, the thermal shock resistance is further improved.
[0027]
(Example 2)
Samples b-1 to b-15 were subjected to (C) thermal shock resistance test and (D) abrasion resistance test. The results are shown in Table 4.
[Table 13]
It can be seen that if the sintered hard alloy in which the groups 4a, 5a, and 6a are the hard phase also has a regular stagnation layer, the thermal shock resistance higher than the conventional one can be obtained. Further, it can be seen that if the binder phase distribution is within the specified range, the wear resistance is improved, and if the WC distribution is within the specified range, the thermal shock resistance is further improved.
[0030]
Example 3
Samples c-1 to c-15 were subjected to (E) thermal shock resistance test and (F) abrasion resistance test. The results are shown in Table 5.
[Table 14]
It can be seen that if the sintered hard alloy using the solid solution hard layer of the 4a, 5a, 6a group as a raw material also has a normal stagnation layer, the thermal shock resistance higher than the conventional one can be obtained. Further, it can be seen that if the binder phase distribution is within the specified range, the wear resistance is improved, and if the WC distribution is within the specified range, the thermal shock resistance is further improved.
[0033]
Example 4
(G) Thermal shock resistance test was performed on samples a-1, a-2, and d-1. The results are shown in Table 6.
[Table 15]
It can be seen that, even if it has a stagnation layer, an improvement in thermal shock resistance is not recognized unless a layer mainly composed of WC exists.
[0036]
【The invention's effect】
If there is a layer composed mainly of WC in the surface layer, the thermal shock resistance is improved in each step, and if the specified binder phase distribution and WC distribution are arranged directly under the layer, the wear resistance is improved. It is possible to further improve the thermal shock resistance. Such sintered hard alloy tools can be applied to wet interrupted machining, which was impossible with conventional sintered hard alloys.
[Brief description of the drawings]
FIG. 1 (A) is a micrograph (SEM photograph) of an alloy structure showing that there are three layers of the stagnation layer, the outermost layer, the Co and Ni bonding layers in the lowermost layer, and the WC layer in the intermediate layer. (B) and (C) are photomicrographs (EDX analysis) showing the distribution of Co and Ni elements in the same structure, respectively.
Claims (4)
Priority Applications (7)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP04929095A JP3803694B2 (en) | 1995-02-15 | 1995-02-15 | Nitrogen-containing sintered hard alloy |
| EP95107670A EP0687744B1 (en) | 1994-05-19 | 1995-05-18 | Nitrogen-containing sintered hard alloy |
| DE69523342T DE69523342T2 (en) | 1994-05-19 | 1995-05-18 | Hard sintered alloy containing nitrogen |
| EP97115279A EP0822265B1 (en) | 1994-05-19 | 1995-05-18 | Nitrogen-containing sintered hard alloy |
| DE69513086T DE69513086T2 (en) | 1994-05-19 | 1995-05-18 | Hard sintered alloy containing nitrogen |
| KR1019950012885A KR0180522B1 (en) | 1994-05-19 | 1995-05-19 | Nitrogen containing sintered hard alloy |
| US08/709,176 US6057046A (en) | 1994-05-19 | 1996-09-06 | Nitrogen-containing sintered alloy containing a hard phase |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP04929095A JP3803694B2 (en) | 1995-02-15 | 1995-02-15 | Nitrogen-containing sintered hard alloy |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JPH08225877A JPH08225877A (en) | 1996-09-03 |
| JP3803694B2 true JP3803694B2 (en) | 2006-08-02 |
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| Application Number | Title | Priority Date | Filing Date |
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| JP04929095A Expired - Lifetime JP3803694B2 (en) | 1994-05-19 | 1995-02-15 | Nitrogen-containing sintered hard alloy |
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Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2011218481A (en) * | 2010-04-08 | 2011-11-04 | Mitsubishi Materials Corp | Cutting insert made from titanium-carbonitride-based cermet, and manufacturing method thereof |
Families Citing this family (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| KR100996838B1 (en) * | 2005-03-28 | 2010-11-26 | 쿄세라 코포레이션 | Super hard alloy and cutting tool |
| JP5185032B2 (en) * | 2007-09-14 | 2013-04-17 | 住友電気工業株式会社 | Cutting tools |
| JP5462549B2 (en) | 2009-08-20 | 2014-04-02 | 住友電気工業株式会社 | Cemented carbide |
| JP6066365B2 (en) * | 2014-06-17 | 2017-01-25 | 住友電気工業株式会社 | Cemented carbide and cutting tools |
-
1995
- 1995-02-15 JP JP04929095A patent/JP3803694B2/en not_active Expired - Lifetime
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2011218481A (en) * | 2010-04-08 | 2011-11-04 | Mitsubishi Materials Corp | Cutting insert made from titanium-carbonitride-based cermet, and manufacturing method thereof |
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| JPH08225877A (en) | 1996-09-03 |
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