Background of the invention

The present invention relates to a cutting tool for machining by chip removal which cutting tool comprises a substrate of cemented carbide, cermet, ceramics, cubic boron nitride based material, or high speed steel, and a hard and wear resistant refractory coating of which at least one layer of the coating comprises a cubic c- (Me, Si)N-phase without co-existing amorphous phase.

TiN has been widely used as hard layer on cutting tools, but the poor oxidation resistance at temperatures above 500 ⁰ C has created an interest in ternary or quar- ternary compounds, e.g. Ti-Al-N and Ti-Cr-Al-N. More complex quaternaries, e.g., Ti-Al-Si-N have been reported and described as super hard H>45 GPa due to a two phase structure consisting of crystalline phase of NaCl-type in combination with amorphous Si ₃ N ₄ or SiN _x . These coating materials show an improved oxidation re- , sistance and improved performance in machining of hardened steels.

However, material based on an amorphous phase tends to plastically deform by grain boundary sliding at elevated temperature which is why optimum performance is not obtained by that solution.

It is an object of the present invention to provide a coating comprising at least one layer containing a c- (Me, Si)N phase without any amorphous phase to be applied to a cutting tool for machining by chip removal . It is a further object of the present invention to provide a method for depositing layers containing a c- (Me, Si)N phase without any amorphous phase by PVD technology.

It has been found that, by balancing the chemical composition, the amount of thermal energy and the degree

of ion induced surface activation during growth, growth rate, and pressure, layers containing C-(Me, Si)N phase without any amorphous phase can be obtained which compared to prior art display improved performance in metal cutting.

Brief description of the drawings

Fig 1. X-ray diffractograms from c- (Tii- _x Si _x )N layers in a) as-deposited, b) annealed at 900 ⁰ C, and c) annealed at 1100 ⁰ C states. The Si-content of each layer is indicated.

Fig 2. Cross-section transmission electron micrographs of a) TiN and b) Tio. ₉ 2Sio.osN layers on WC-Co substrates . Fig 3. Cross-section TEM micrographs from an as- deposited Tio. ₈₆ Sio. ₁₄ N thin layer on WC-Co substrate in, a) overview with selected area electron diffraction pattern and b) HREM image. The columnar microstructure with internal branching of subgrains of (002) crystallo- graphic orientation is indicated in a) . The trace of

(200) and (002) planes in neighbouring fibres I-III with zone axes [010], [hko] , and [110], respectively, after mutual rotation around [001] is shown in b) .

Fig 4. a) Hardness and b) Young's modulus of as- deposited and annealed at 900 ⁰ C and 1100 ⁰ C for 2h Ti ₁ - _X Si _x N layers .

Detailed description of the invention

According to the present invention, a cutting tool for machining by chip removal is provided comprising a body of a hard alloy of cemented carbide, cermet, ceramics, cubic boron nitride based material, or high speed steel, onto which a wear resistant coating is deposited composed of one or more layers of refractory compounds comprising at least one layer of C-(Me, Si)N phase with-

out any amorphous phase. Additional layers are composed of nitrides and/or carbides and/or oxides with the elements chosen from Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W, Si and Al, grown using Physical vapour deposition (PVD) , or other deposition technologies than PVD such as plasma enhanced chemical vapour deposition (PACVD) and/or chemical vapour deposition (CVD) . Tools according to the present invention are particularly useful in metal cutting applications where the chip thickness is small and the work material is hard such as copy milling using solid end mills, insert milling cutters or drilling of hardened steels.

The cutting tool insert, solid end mill, or drill, comprising a substrate and a coating composed of one or more layers of refractory compounds of which at least one layer comprises crystalline cubic phase, c- (Mei- _x Si _x ) _c (Ni-bXb)d/ where Me is one or several of the elements Ti, Zr, V, Cr, and Al and X is one or several of the elements C, 0 and B, the ratio R = d/c of c- (Mei- _x Si _x ) _c (Ni- _b Xb)d phase is between 0.85 and 1.05, and that b is less than 0.1 and x is between 0.05 and 0.20 and that it contains no amorphous phase detected by XRD or/and TEM (transmission electron microscope) or/and XPS (X-ray photoelectron spectroscopy) . The c- (Me, Si)N layer is characterised by:

The existence of a crystalline C-(Me, Si)N, as detected by X-ray diffraction (XRD) in Θ-2Θ and/or gracing incidence geometry showing one or several of the following features : - a c- (Me, Si)N (111) peak, using CuKa radiation, at about 37°2Θ,

- a c- (Me, Si)N (200) peak, using CuKa radiation, at about 43°2Θ,

- a c- (Me, Si)N (102) peak, using CuKa radiation, at about 62°2Θ,

- When Me contains other elements than Ti the peak positions could be shifted.

- The structure of the C-(Me, Si)N is preferably of sodium chloride, NaCl-type, and/or of zinc sulphide, ZnS- type.

- The texture defined as the ratio, K, between the area of the C-(Me, Si)N (200) peak, using CuKa radiation in Θ-2Θ geometry, and the c- (Me, Si)N (111) is above 5 and preferably above 10.

- The FWHM (Full Width Half Maximum) value, using CuKα radiation in Θ-2Θ geometry, caused by small grains and/or inhomogeneous stresses, of the c- (Me, Si)N (200) is above 0.5°2θ when x<0.1 and above 0.9 when x^0.1.

The layer comprising c- (Me ₇ Si)N has a considerably increased hardness compared to a cubic single phase layer of a NaCl type c-TiN structure, see Example 1 and improved high temperature properties since it does not contain any amorphous phase.

The total coating thickness, if the C-(Me, Si)N containing layer (s) according to the present invention if combined with other layer (s), is between 0.1 and 10 μm, preferably between 0.5 and 5 μm with the total thickness of the non C-(Me, Si)N containing layer (s) varying between 0.5 and 8 μm.

In an alternative embodiment the C-(Me, Si)N containing layer (s) of 0.5 and 10 μm thickness, with or without other layer (s) according to the above described, an outer 0.5 to 5 μm thick layer consisting of a solid low friction material based on M0S ₂ , DLC (Diamond Like Carbon) or a MeC/C, where Me is Cr, W, Ti or Ta, can be deposited on top of the coating.

In yet an alternative embodiment the C-(Me, Si)N layer(s), between 0.1 and 2 μm thickness, are one of one

to five different materials in a 1.0 to 15 μm thick multilayer coating consisting of individually 2-500, preferably 5-200, layers.

In yet an alternative embodiment at least one c- (Me, Si)N containing layer of 0.1 and 1.0 μm thickness is used for metal cutting application where the chip thickness is very small. In this embodiment a post-treatment resulting in a surface roughness of Ra on the edge of the cutting tool better than 0.15 μm is preferable. The method used to grow the layers of the present invention, comprising C-(Me, Si)N phase without any amorphous phase, here exemplified by the system Ti-Si-N, is based on arc evaporation of an alloyed, or composite cathode, under the following conditions: The Ti-Si cathode composition is 10-20 at-% Si, preferably 15-20 at-% Si and balance of Ti.

The evaporation current is between 50 A and 200 A depending on cathode size and cathode material. When using cathodes of 63 mm in diameter the evaporation cur- rent is preferably between 70 A and 100 A.

The substrate bias is between -10 V and -80 V, preferably between -40 V and -60 V.

The deposition temperature is between 400°C and 700°C, preferably above 50O ⁰ C and most preferably above 520-600°C.

When growing layer (s) containing C-(Me, Si)N an Ar+N ₂ atmosphere consisting of 0-50 vol-% Ar, preferably 0-20 vol-%, at a total pressure of 0.5 Pa to 9.0 Pa, preferably 1.5 Pa to 7.0 Pa, is used. For the growth of C-(Me ₇ Si) (N, X) layer (s) where X includes C and O, carbon and/or oxygen containing gases have to be added to the N ₂ and/or Ar+N ₂ atmosphere (e.g. C ₂ H ₂ , CH ₄ , CO, CO ₂ , O ₂ ) . If X also includes B it could be added either by alloying the target with B or by adding a boron containing gas to the atmosphere.

In order to obtain the preferred structure according to this invention, i.e. a layer containing c- (Me, Si)N here exemplified by C-(Ti, Si)N, we have found that several deposition parameters have to be fine- tuned. One important factor is to keep the deposition temperature optimized in combination with the degree of surface activation of the adatoms by a proper choice of energy of impinging ions by controlling pressure and negative substrate bias. The best parameters depend probably on a combination of the distance between the cathode and the substrate and pressure which determine the average charge state of the ions reaching the border of dark space surrounding the substrates.

Improved oxidation resistance could be achieved by replacing some of the Ti with one or several of the Zr, Al and Cr.

When growing layer (s) containing C-(Me, Si)N phase there is a risk that the compressive residual stress becomes very high, which will influence the performance negatively in cutting applications when sharp cutting edges are used and/or when the demand on good adhesion is of outermost importance. One possibility to decrease the compressive residual stresses is to apply a post annealing process, or in-situ annealing, preferably in an atmosphere of Ar and/or N ₂ at temperatures between 600°C and 1000°C for a period of 20 to 600 min which due to materials high temperature stability will not negatively affect the layer, as demonstrated in Example 1. However, an annealing temperature of 1100 °C is too high. The present invention has been described with reference to layer (s) containing C-(Me, Si)N phase deposited using arc evaporation. It is obvious that C-(Me, Si)N phase layer (s) also could be produced using other PVD- technologies such as magnetron sputtering, electron beam

evaporation, ion plating, or laser ablation if the correct deposition parameters could be fine tuned.

Example 1 12x12x4 mm plates of cemented carbide WC-Co (6 wt-%) were used as substrates. Before the deposition the substrates were ground and polished to a mirror like finish and cleaned in an ultrasonic alkaline degreasing agent. The layers were deposited by an arc evaporation system. Three cathodes of composition Ti, Tig _O Siio, and Ti ₈₀ Si ₂ O/ respectively, mounted on top of each other were used to produce Tii- _x Si _x N layers of varying composition from one batch.

The shortest cathode-to-substrate distance was 160 mm. The system was evacuated to a pressure of less than 2.0XlO ^"3 Pa, after which the substrates were sputter cleaned with Ar ions .

The deposition was carried out in a 99.995% pure N ₂ atmosphere at a total pressure of 5.0 Pa, using a sub- strate bias of -45 V. The layer thickness was about 2 μm. The deposition temperature was about 520°C. Immediately after deposition the chamber was vented with dry N ₂ .

The variants are hereafter referred to by their Si- content .

Chemical analysis of the layer compositions performed using EDX analysis of as-deposited layers from the different positions within the deposition chamber showed a continuous composition range between 0≤x≤0.14 Si for samples positioned between Ti and Tio.gSio.i or

Ti _o .8Sio.2 targets. The layer closest to the Tio.8Sio.2 target was expected, due to setup geometry, to have higher Si content than the actual x=0.14. However, the apparent loss of Si during arc evaporation can be explained by

8 I H J i_ va-> -«• ^■ «»« β w w w ¹ ^" y

the stronger impact from Ti atoms than Si atoms on the sample surface.

The chemical bonding structure in the near-surface region was analysed by X-ray photoelectron spectroscopy (XPS) . The XPS was equipped with a non-monochromated Al K _α at 1486.6 eV x-ray source and a hemispherical electron energy analyzer. The samples were then Ar-etched and survey scans of the binding energy 0-1100 eV was recorded for each sample. For accurate determination of the exact peak positions of the Ti2p, Nls, Si2s, Si2p and CIs peaks, local region scans were recorded.

Nanoindentation analysis of layers after mechanical polishing of the surface was performed using a Nano Instruments Nanolndenter II with a Berkovich diamond tip. The maximum tip load was 25 inN.

The stoichiometry i.e. the ratio R = d/c of c- (Mei_ xSiχ) _c (Ni-bXb)d was found to be between 0.9 and 1.0 and index b <0.05 for all variants.

Fig. Ia) shows XRD diffractograms from the as-de- posited Tii_ _x Si _x N, layers. The diffractograms for all compositions revealed a NaCl-structure compound with a lattice parameter very close to TiN, 0.424 nm. The preferential layer growth orientation changed from mixed <111> to exclusive <200> with increasing Si content. The tex- ture, K, was above 10 for all variants with χ:>0.03. The FWHM-value was 0.5° for x=0.05, 1.0° for x=0.11 and 1.6° for x=0.14.

Cross-section TEM micrographs from the as-deposited samples are presented in Figs. 2a) TiN, 2b) Tio. ₉₂ Sio.osN, and 3) Tio. ₈₆ Sio. ₁₄ N, respectively. The layers exhibited a dense columnar structure where the column width decreased from 200 nm to 100 nm with increasing Si content. Interestingly, for all layers, the top surface was retaining the roughness and features of the sub- strate surface. This implies that 3D-island growth and

possible faceting was effectively suppressed during the deposition. From Figs. 2 and 3 it is evident that the Si content also affected the structure and increased the defect density. The as-deposited Tio. ₈₆ Sio.14N, see Fig. 3, exhibited within each column a fine-fibrous structure. Higher magnification imaging revealed nm-structure of fibers (elongated crystalline grains) with large strain contrast and moire fringes from overlapping features . The HREM image in Fig. 3b) shows a typical appearance of three fibers of the cubic (Ti, Si)N phase with high defect density from dislocations. The observations in Fig. 3 show an interesting growth mode with a rotating lattice by branching of fibers over each column to solid angles up to 20 ⁹ (the width of the (002) arch in Fig. 3a) . Branching begins at the column boundaries and the subgrains merge at the apparent stem of the columns. This takes place to maintain the (002) growth surface. No volumes of any amorphous phase were found by the TEM analysis . Fractured cross-sections from as-deposited

Ti _! _ _x Si _x N, x=0, x=0.05, x=0.08, x=0.14 were investigated by scanning electron microscopy. The micrographs showed dense columnar microstructure with macro particles. The thickness of the (Ti, Si)N layer ranged between 1.9 and 2.4 μm. As-deposited Tio.86Sio.14N behaved as a fine-fi- bered structure when fractured in agreement with the nanostructure seen by XTEM.

Results from XRD of Tii- _x Si _x N layers isothermally annealed at 900 ⁰ C and 1100 ⁰ C are in Ar-atmosphere are presented in Fig. 1 b) and c) . The texture of the as-deposited samples was maintained. However, the peak broadening decreases for increasing annealing temperatures.

Tii_ _x Si _x N solid solution layers undergo recovery from a growth-induced compressive stress state during annealing above the deposition temperature. From Fig. 1,

the recovery is evident from a decreasing peak broadening.

Fig. 4 shows how the Si content influences the hardness and Young's modulus for the as-deposited layers and those annealed at 900 ⁰ C and 1100 ⁰ C. The addition of Si increased the hardness throughout the whole gradient series for the as-deposited layers from 31.3+1.3 GPa up to 44.7+1.9 GPa for Tio. ₈₆ Sio. ₁₄ N. From XRD and TEM it is evident that the Tio. ₈₆ Sio. ₁₄ N layer contains the high- est defect density. Thus, both solid solution and defect hardening may be active. The 900 ⁰ C annealed samples retained their hardness with a similar trend for Si content coupled to hardness. At 1100 ⁰ C, however, the hardness has decreased to below 30 GPa for all compositions. The Young's modulus of the layers increases slightly with increasing Si content for the as-deposited layers from 670 GPa, for x=0, to 700 GPa, for x=0.05, Fig 4b. For higher amount of Si, however, the Young's modulus decreased to -600 GPa. X-ray photoelectron spectroscopy from as-deposited Tio. ₈₆ Sio. ₁₄ N showed the presence of the elements Ti, Si, N, as well as a very small amount of O after Ar-etching. Analysis of the Si2p peak showed a binding energy 100.9 eV for the Tio.86Sio. ₁₄ N layer, which suggests Si-N bond- ing, close to the reported value of 0-Si ₃ N ₄ at 100.6 eV. Amorphous phase silicon nitride, a-Si ₃ N ₄ , usually positioned at 101.8 eV was not detected.

Tii_ _x Si _x N solid solution layers of NaCl-structure containing as much as 14 at-% Si was synthesized by arc- evaporation at a substrate temperature of 520 ⁰ C. For the highest Si contents, however, XRD and TEM reveal a defect-rich lattice whereas XPS do not imply the presence of amorphous Si ₃ N ₄ . Since the lattice parameters for all solid solutions including c-NaCl structured SiN are found to be very close to TiN, the identification by

electron or X-ray diffraction techniques of any compositional inhomogeneities from synthesis or secondary phase transformation will be difficult. However, XRD analysis clearly shows that no hump from amorphous phase, as ex- pected typically FWHM of 4-6°2θ centred at about 37°2θ, could be detected.

Example 2

A copy milling test was performed using RDHW10T3M0T-MD06 inserts coated similarly as in Example 1 (variants x= 0, 0.06 and 0.14). Inserts with a (Ti, Si)N-layer with a thickness of 3.5 μm containing amorphous phase were used as reference. The tool life was measured when the inserts were worn out as defined as when sparkles were created and the material got an uneven surface. Tool life is reported in table 1.

Material: DIN X155 CrMoV 12 1, hardened to 58HRC Dry machining v _c = 250 m/min f _z = 0,2 mm/tooth a _p = 1 mm, a _e 2 mm

Table 1.

Variant Tool life Tool life

Edge 1 (min) Edge 2 (min)

X=O 3.0 2.6

X=O. .05 4.1 3.4

X=O. .14 5.5 4.0

Ref 3.3 3.0

This test shows that the variants with high amount of Si in C-(Ti, Si)N have the longest tool life.