The present invention relates to a sintered carboni- tride alloy having titanium as main component and intended preferably for metal cutting applications.

Titanium based carbonitride alloys, usually named cermets, are well established as cutting tool material often used for finishing at high cutting speed. More recently, the area of application has been widened towards more toughness demanding applications which has been made possible by, e.g., increased binder phase and nitride contents in these alloys compared to tungsten based, more brittle hard materials. Another way of ob¬ taining increased toughness is disclosed in Swedish Patent Application No. 9004121-1 in which extremely fine-grained alloys are made using melt metallurgically made intermetallic prealloys. During sintering of titanium based carbonitride alloys a solution-precipitation process takes place. As a result, a very common structure in such alloys is hard constituent grains with a core-rim structure. An early patent in this area is U.S. Patent 3,971,656 which discloses Ti- and N-rich cores and rims rich in Mo, and C. By a suitable combination of core-rim-structures in well balanced proportions optimal properties re¬ garding wear resistance, toughness behaviour and/or plastic deformation can be obtained as disclosed in Swedish Patent Application No. 8902306-3.

Swedish Patent Application No. 9101385-4 discloses a titanium based carbonitride alloy with coarse grains with core-rim structure in a more finegrained matrix. EP-A-447388 and EP-A-464396 disclose manufacturing of carbonitrides directly by carbonitriding of the oxides of the metals or the metals themselves.

According to the present invention it has now been found that it is possible to further improve the properties of a material according to SE 9004121-1 regarding toughness behaviour and/or wear resistance. By adding additional more coarsegrained hard constituents to original prealloyed powder a material with improved wear resistance and/or a significant improvement in toughness behaviour relative to the original material is obtained. Figure 1 shows in about 5000X magnification the microstructure of an alloy according to the invention.

Figure 2 shows in about 5000X magnification the microstructure of an alternative embodiment of an alloy according to the invention in which A - agglomerate, B - matrix and C - transition zone.

According to the invention there is now provided a titanium based carbonitride alloy with 70 - 97 % by volume hard constituents, in which titanium is the dominating hard constituent former i.e. more than 50 mole-% of the metallic elements is titanium. Additional metals present are Zr, Hf, V, Nb, Ta, Cr, Mo and/or W. Small additions of Al may be present usually in the binder phase which is based on Fe, Ni and/or Co, prefe¬ rably Ni and/or Co. The material according to the invention is manufactured in two steps. In the first step, disclosed in the above mentioned patent application SE 9004121-1, a powder is made by a method comprising casting a pre- alloy of hard constituent forming and binder phase forming metals without intentional additions of C, N, B and/or 0 to form a cast prealloy of brittle intermetallic phases of hard constituent forming metals and binder phase forming metals mixed in atomic scale. The alloy is then disintegrated into a powder of a grain size <50 μm. This powder is carbonitrided to form in situ extremely finegrained hard constituent particles

with a diameter <0.1 μm within the binder phase metals. In this way, a ready to press powder is obtained in which no additives of powder other than very highmelting and consequently eltmetallurgically difficult to handle phases are present. Powder manufactured in this way is, thus, characterized by its fine-grained particles, .<0.1 μm. In the second step the additional hard constituents according to the invention are added which gives the material special properties due to a unique structure in sintered state. Said additional hard constituents comprise carbides, nitrides and/or carbonitrides of metals from the groups IV, V and VI of the Periodic System of the elements, preferably Ti(C,N), (Ti,Ta)C, (Ti,Ta) (C,N) , (Ti,W) (C,N), (Ti,Ta,V) (C,N) and/or (Ti,Ta, ) (C,N) etc depending on the desired property profile of the sintered alloy.

The additional hard constituent powders preferably have essentially equiaxial grains with a narrow grain size distribution and high chemical purity, preferably being produced directly by carbonitriding of the oxides of the metals or of the metals themselves. The mixing to powder obtained in the first step above shall in such case take place in a very careful manner in order to avoid excessive milling. One way of obtaining a good mixture is first to make sure that the additional hard constituent powder is deagglomerated in a separate step and after that mix in a conventional dry mixer. Another way is to add said powder immediately before the end of the milling time. In an alternative more wear resisting embodiment the additional hard constituents are predominantly added as agglomerates with a diameter of 20-200, preferably 40-80 μ and care is being taken not to crush the agglomerates and yet obtain a good mixing. In addition, it is very important to choose the grain size of the additional hard constituents such that

after the sintering they have a mean grain size of 1.5- 15, preferably 1.5-5 times greater than the mean grain size of the prealloy matrix which is <1 μm. Suitable grain size of the additional hard constituent powders is consequently 0.8-5 μ , preferably 0.8-3 μm. Suitable contents of the additional hard constituents are 10-50 weight-% of the prealloyed powder, preferably 20-40 wt-%.

The milled mixture of prealloyed powder and additional hard constituents is used for the manufacture of cutting inserts with known methods such as spraydrying, pressing and sintering.

A sintered structure of a material according to the invention is shown in fig 1 and is characterized by a very finegrained prealloyed matrix in which the hard constituent particles have a core-rim structure. The additional hard constituent grains (black in the image) more or less welldispersed in this matrix have preserved their virgin character i.e. they essentially lack rims and remain with their original morphology, chemical composition and grain size.

Figure 2 shows the microstructure of the alternative embodiment. A is an agglomerate filled with binder phase in the fine grained matrix, B. As in figure 1 the hard constituent grains of the agglomerates essentially lack rims and remain with their original morphology, chemical composition and grain size. The agglomerates are surrounded by a few μm thick zone depleted in binder phase, C. Cutting inserts according to the invention show superior wear resistance but also toughness properties. There are several reasons to this: Adding additional hard constituents to a given hard material composition means that it is possible to increase the share of hard constituents with a corresponding decrease in the total binder phase content which, of course, increases the

wear resistence. That the toughness as in this case also increases might be due to the resulting structure. As mentioned the mean grain size of the additional hard constituents is greater than the fine grained prealloyed fraction. This means, however, that the binder phase volume relative to the hard constituent volume increases and, thus, counteracts a decrease in toughness. In addition, the additional hard constituents are present not as usual in conventional titanium based carbonitride alloys in the sense that said hard constituents almost completely lack rim which is known to be the brittle phase in such alloys. Yet another possible explanation to the improved toughness behaviour is the fact that additions of e.g. TiCo.sNg.s or TiCQ.3No.7 to a carbide rich starting composition lead to a more nitride rich total composition which of course is favourable from toughness point of view. It can not be excluded that the structure obtained gives rise to favourable inner stresses which may have crack stopping effects which leads to the good toughness behaviour.

Example 1

A prealloy of the metals Ti, Ta, V, Co, Ni was made in a vacuum induction furnace at 1450°C in Ar protecting gas (400 mbar) . The composition of the ingot after casting in the ladle was in % by weight: Ti 66, Ta 8, V 6, Ni 8 and Co 12. After cooling, the ingot was crushed to a grain size < 1 mm. The crushed powder was milled together with necessary carbon addition in a ball mill with paraffin as milling liquid to a grain size <_ 50 μm. The mixture was poured on a stainless plate and placed in a furnace with a tight muffle. The removal of the milling liquid was done in flowing hydrogen gas at the temperature 100-300°C. After that, the powder was carbonitrided in solid phase by addition of nitrogen gas. The total cycle time was 7 h including three

evacuations in order to retard the procedure. The carburizing occurs essentially at the temperature 550- 900°C. Then the final carbonitride charge was cooled in nitrogen gas. The finishing powder manufacture was done in conventional ways, i.e., additional powders (WC and M02C) were added and milled together with the carbonitride charge to final powder which was spray- dried in usual ways. This powder was wet mixed with 20 % by weight TiCN 50/50 with spherical morphology and with mean grain size 1.4 μ (FSSS) . This powder was not yet deagglomerated. The mixture was stirred with reduced speed in a ball mill for 30 min and spray dried to ready to press powder.

Example 2

Inserts of type TNMG 160408-QF were manufactured of prealloyed powder without addition of additional hard constituents according to example 1 (alloy A) with the following metal composition in mole-%: Ti 62.4, Ta 2.3, V 4.7, W 6.2, Mo 7, Co 10 and Ni 7.4. An alloy according to the invention of example 1 i.e. alloy A with an addition of 20 % by weight TiCN 50/50 was used for the manufacture of inserts with the same geometry (alloy B) . The inserts of both alloys were sintered at the same time and had the same edge radius and -rounding. They were tested in an interrupted cutting operation until fracture. 20 inserts of each alloy were tested. Cutting data at the initial engagement was:

V = 110 m/min f ₀ = 0.11 mm/rev a = 1.5 mm

Work piece material = SS 2244 The feed was increased linearly until all of the inserts had fractured. Thereafter, the accumulated

fracture frequency was determined as a function of the time to fracture. The 50 % value fracture frequency for a certain feed in mm/rev is used as comparison number for the toughness behaviour. Result:

Alloy A 0.145

Alloy B, according to the invention 0.22 Looking at the individual results alloy B wins in 19 cases out of 20. The result is, thus, statistically very convincing.

With the same alloys also a wear resistance test was performed. Test 1 was longitudinal turning in SS 2541 with the following cutting data: V = 400 m/min f ₀ = 0.15 mm/rev a = 1.0 mm

The flank wear was measured continuously. Three tests were run. VB after 15 min engagement time is given in the table below (VB in 1/10 mm) 1 2 3

Average Alloy A 0.35 0.36 0.29 0.33

Alloy B, invention 0.21 0.24 0.24 0.23 The flank wear is as apparent significantly better for alloy B according to the invention.

In the same work piece also the tool life of the alloys was determined. The tool life criterion was plastic deformation. Cutting data were as above. Result Tool life, min

1 2 3 Average Alloy A 20 16 17.5 17.8

Alloy B, invention 63.5 35 27.5 42 The alloy according to the invention, thus, had a much better- tool life.

Test 2 was a wear resistance test made as plane turning of tubes in SS 2234 with the following cutting data:

V = 400 m/min f = 0.15 mm/rev a = 1.0 mm

Engagement time/section: 0.11 min. As predetermined tool life VB=0.3 mm was chosen. This value was reached for alloy A after 130 sections (corresponding to 14.3 min) whereas alloy B according to the invention got the same VB after 185 sections (corresponding to 20.4 min). Again, a pronounced difference, thus, in favour of alloy B is found. Complete tool life was not tested in this operation. None of the alloys had fracture after this predetermined VB-value.