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CrN coatings on cemented carbide
Surface and Coatings Technology 116–119 (1999) 108–115 www.elsevier.nl/locate/surfcoat
Deposition and characterisation of multilayered PVD TiN/CrN coatings on cemented carbide Maria Nordin *, Mats Larsson Uppsala University, Department and Division of Materials Science, Box 534, S-751 21 Uppsala, Sweden
Abstract In laboratory as well as in application tests, multilayered PVD coatings have shown enhanced mechanical and tribological properties as compared to today’s single layered PVD coatings. A coating which has shown very interesting properties, such as high temperature stability and high fracture toughness, is multilayered PVD TiN/CrN. Our knowledge about the connection between the growth dynamics and the properties of this coating is, however, rather poor. Therefore, to further develop and optimise this coating, it is necessary to study the correlations between on the one hand the growth process, and on the other the microstructure, composition and mechanical properties of the coating. In this work growth rate, morphology, microstructure, chemical and phase composition, together with coating hardness, Young’s modulus and fracture toughness of different multilayered PVD TiN/CrN coatings have been evaluated. All coatings have been deposited on cemented carbide substrates using a combination of reactive electron beam evaporation ( Ti) and reactive d.c. magnetron sputtering (Cr). The influence of lamella thickness and different deposition parameters: substrate bias, magnetron sputtering power and nitrogen flow, on the above mentioned characteristics has been examined. In addition to the multilayered TiN/CrN coatings, homogeneous TiN and CrN have been included and compared. The investigation showed that a dense, fully cubic NaCl phase, multilayered PVD TiN/CrN coating with high fracture toughness can be deposited provided that the lamella thickness is kept less than 14 nm and 5 nm of TiN and CrN respectively. Thin lamellae seem to inhibit transformation from growth of the cubic NaCl phase to new phases, e.g. hexagonal b-Cr N and 2 metallic Cr. Furthermore, thin lamellae yielded a (200) preferred growth orientation while thicker lamellae generated a mixture of growth orientations. In addition, very thin lamellae must be deposited to obtain good fracture toughness. It was also found that it was necessary to use a negative substrate bias in order to obtain a high quality coating. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cemented carbide; Coating; Multilayer; PVD; TiN/CrN
1. Introduction For metal cutting tools two of the most common commercial PVD coatings are TiN and CrN. Two other frequently used coatings, not considered in this work, are Ti(C,N ) and ( Ti,Al )N. Of these TiN is by far the most extensively investigated and used material. This is partly due to the ease with which TiN can be deposited to its very hard and wear resistant cubic NaCl phase (it has a large nitrogen solubility) and partly due to its beautiful golden colour. CrN, with its excellent high temperature stability [1–3], is often utilised in applications where a hard, oxidation- and corrosion-resistant * Corresponding author. Fax: +46-18-471-35-72. E-mail address: [email protected] (M. Nordin)
coating is required, e.g. on moulds for Al alloy die casting [2]. CrN is also known for its high toughness, which facilitates growth of thicker coatings (5–20 mm). A characteristic of CrN is that it is more difficult to deposit to its cubic NaCl phase, due to its lower nitrogen solubility and lower stability as compared to TiN. Nevertheless, some papers have shown that it is possible to grow stoichiometric cubic CrN coatings, e.g. Refs. [2,4]. There have been several scientific works on multilayered PVD coatings, i.e. coatings obtained by alternately depositing two chemically and/or mechanically different materials to form a layered structure, during the last years. In many cases multilayered coatings have shown enhanced mechanical and tribological properties as compared to today’s single layered coatings [5,6 ]. It has also been shown that by layering thin layers of a metal
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nitride (preferably less than a few nm of coatings such as VN, NbN and CrN ) with layers of TiN, fully NaCl cubic coatings with a high hardness can be grown [7– 9]. Up to today, most of the research on multilayered or superlattice coatings has been on the systems TiN/VN and TiN/NbN [7,8]. Another multilayered coating which has shown very interesting properties, such as high temperature stability and high fracture toughness, is made of the two nitrides TiN and CrN [10,11]. Our knowledge about the correlations between on the one hand the growth dynamics, chemical composition, phase composition and microstructure, and on the other the mechanical/tribological properties of multilayered TiN/CrN coatings is, however, scarce. Only a few works concerning this coating can be found. Yashar et al. have investigated the growth process and phase composition of multilayered TiN/CrN as a function of lamella 0.6 thickness [9]. They showed that CrN , which normally 0.6 is hexagonal, can be stabilised into a cubic NaCl phase provided that the lamellae are kept less than 10 nm. Furthermore, Panjan et al. showed that it is possible to increase the activation energy for oxidation by layering TiN with CrN (modulation periods in the range 83– 425 nm) as compared to the single layered coatings made of the constituents [12]. The aim of this work is to increase our knowledge of the growth dynamics and properties of multilayered PVD TiN/CrN coatings. A number of multilayered TiN/CrN coatings have been characterised with respect to their growth rate, morphology, microstructure, chemical composition and phase composition as well as their composite hardness and fracture toughness. The influence of the lamella thickness and the deposition parameters: substrate bias, magnetron sputtering power and nitrogen flow, on those characteristics has been evaluated. In addition to the multilayered coatings, homogeneous TiN and CrN have been included as reference coating materials.
2. Experimental 2.1. Substrate material Cemented carbide (10 wt% Co and 90 wt% WC ) was used as substrate material. The substrate hardness was 1350 HV . All substrates were polished to a surface 30 kg roughness of approximately 5 nm and then ultrasonically cleaned in a heated alkaline solution followed by ethanol (5 min each). 2.2. Coating deposition All coatings were deposited in a commercial Balzers BAI 640R coating unit. Prior to coating deposition, the samples were resistively heated to 450°C for 60 min and
thereafter Ar ion etched for 15 min using a substrate bias of −200 V. The coatings were reactively deposited using a combination of electron beam gun (e-gun) evaporation of Ti and d.c. magnetron sputtering of Cr. For all the multilayered coatings the deposition sequence started with the deposition of a 30 nm Ti layer followed by a 200 nm TiN layer to promote coating adhesion. This was obtained while the substrates were held stationary above the Ti source. Simultaneously the Cr target was Ar ion etched. By keeping both sources active and by alternately expose the substrates to the Ti source and the Cr source using substrate rotation, a multilayered coating was grown. The total pressure was kept constant at 2×10−3 mbar. Before the deposition of single layered CrN and TiN commenced, a 30 nm layer (Cr and Ti, respectively) was deposited. Single layered TiN and CrN were deposited by keeping only one of the material sources active and admitting nitrogen gas into the chamber. The total pressure was kept at 1.7×10−3 mbar during the TiN deposition and at 3×10−3 mbar during the CrN deposition. The total deposition time was 40 min for all the multilayered coatings, 60 min for TiN and 160 min for CrN. In Table 1 an overview of the evaluated process parameters are presented. 2.3. Coating characterisation From studies of fractured coating cross-sections, information about the coating thickness, multilayer period (L; the thickness of one lamella of TiN together with one lamella of CrN ) and coating morphology could be obtained using scanning electron microscopy fitted with a field emission gun (FEG-SEM ), see e.g. Fig. 1. An atomic force microscope (AFM ) was used to reveal the surface topography. Cross-sectional transmission electron microscopy ( XTEM ) in a JEOL 2000 FXII Table 1 An overview of the deposition process parameters used. M(x) indicates a multilayered TiN/CrN coating with the varied parameter within brackets [in comparison to the reference coating designated M(ref )] Coating
Substrate Magnetron N Substrate rotary 2 bias ( V ) power (kW ) flow (sccm) speed (rpm)
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
−110 −110 −110 −110 −110 0 −200 −110 −110 −110 −110
– 2 2 2 2 2 2 4 6 2 2
140 40 155 155 155 155 155 155 155 115 195
10 10 10 5 1 10 10 10 10 10 10
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developed by Oliver and Pharr [14]. A Berkovich diamond tip and an indentation depth of 150 nm were used. By performing Rockwell C indentation tests (in a commercial scratch tester with the specimen translation disconnected ), information about the coating toughness could be gained. The load was continuously increased up to 100 N at a loading rate of 10 N/min and subsequently decreased at the same rate. As a measure of the coating fracture toughness, the fracture load, L , i.e. the F load at the first coating fracture as detected using acoustic emission (AE) was utilised [Fig. 2(a)]. In all cases, this load corresponded to a circular crack at the rim of the indent, see Fig. 2(b). All indentations were post-experimentally studied in a light optical microscope (LOM ). Fig. 1. Representative example of a SEM micrograph [M(6 kW )] showing the layered structure. In the micrograph the column diameter is indicated as well as 10 multilayer periods.
equipped with a LaB filament and a 200 kV voltage 6 was employed to evaluate the microstructure and selected area electron diffraction (SAED) to determine the phase composition of one multilayered coating which in the following is denoted M(ref ). The chemical composition of the coatings was determined using Auger electron spectroscopy (AES). To subdivide the intensity from the overlapping transitions of Ti and N at approximately 382 eV into separate components the method obtained by Dawson and Tzatzov [13] was used. The electron beam energy and emission current were 10 keV and 5 mA, respectively. All the sensitivity factors ( Ti, Cr and N ) were obtained from nitride reference powders. The phase composition and preferred growth orientation were determined using a Siemens D5000 X-ray diffractometer. Cu Ka radiation over an area of approximately 1 mm2 was employed and the measurement range was 2h=30–100°. No sample rotation was assessed during measurement. The composite hardness and Young’s modulus were determined using nanoindentation and the method
3. Results 3.1. Coating thickness and multilayer period For the multilayered coatings it was found that the coating thickness (and the multilayer period), which is directly related to the coating growth rate since the deposition time was the same for all multilayered coatings, increased with magnetron power and nitrogen flow and slightly decreased with substrate bias, see Table 2. The coating thickness was, however, not affected by the substrate rotary speed, while, naturally, L decreased with increasing rotary speed. 3.2. Morphology and microstructure All coatings deposited (both single- and multilayered) displayed a dense and homogeneous surface morphology, see Fig. 3(a). The only exception was M(0 V ), which appeared to be more porous, see Fig. 3(b). No significant influence of the magnetron power, nitrogen flow or rotary speed on the surface morphology was observed.
Fig. 2. The fracture load (L ) was defined as the load corresponding to the first significant increase in the AE signal. (a) A typical example of the F AE signal recorded during indentation. (b) Representative appearance of an indented coating, note the circular cracks each corresponding to one of the AE peaks (SEM ).
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Table 2 Coating thickness, t, multilayer period, L, chemical composition (approximate measurement scatter was 5%), Ti to Cr ratio (Ti:Cr), Me to N ratio (Me:N ) and fracture load, L F Coating
t (nm)
L (nm)
at% Ti
at% Cr
at% N
Ti:Cr
Me:N
L (N ) F
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
3800±100 4400±150 3700±100 3600±100 3900±100 3900±100 3400±100 4200±150 5000±200 2500±100 4800±200
– –
51 – 45 46 43 44 48 38 32 37 48
– 53 14 11 12 14 12 24 35 21 9
49 47 41 43 45 42 40 38 33 42 43
– – 3.2 4.2 3.6 3.1 4.0 1.6 0.9 1.8 5.3
1.0 1.1 1.4 1.3 1.2 1.4 1.5 1.6 2.0 1.4 1.3
27±2 32±1 27±1 26±1 23±1 11±3 27±9 23±2 24±3 24±2 26±1
9.7±0.5 19.1±1 110.0±5 9.8±0.5 9.5±0.5 12.0±0.6 14.5±0.7 6.1a±0.5 12.4±0.6
a Period is calculated from the coating growth rate since it was too thin to measure by SEM.
All coatings were found to have a more or less columnar structure when studied in SEM. When comparing the multilayered coatings, the largest column diameter was found for the coating deposited at 0 V substrate bias. Furthermore, it was found that the
diameter increased with magnetron sputtering power. For the other multilayered coatings no clear difference in column size could be seen. The TiN coating displayed a more pronounced columnar microstructure than both CrN and the multilayered coatings. XTEM displayed that M(ref ) was columnar, dense and fine grained (see Fig. 4), which is in agreement with the results obtained using SEM. XTEM also showed that coating growth, and hence coating surface profile, well followed the surface profile of the substrate. The substrate roughness developed during Ar ion etching of the substrates prior to coating deposition. The lamella thickness of CrN and TiN was approximately 3 and 7 nm, respectively. The SAED analysis revealed a 100% cubic NaCl phase composition, see diffraction pattern insert in Fig. 4. It was also observed that the diffraction
Fig. 3. Typical coating surface topography (AFM ) of (a) the coatings deposited using substrate bias [M(ref )] and (b) without substrate bias [M(0 V )].
Fig. 4. XTEM micrograph and SAED pattern of M(ref ). The broadening of the diffraction peaks was due to coating growth on a not perfectly smooth substrate.
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Fig. 5. XRD patterns of the multilayered coatings with different L. S refers to peaks originating from the substrate.
spots were somewhat broadened as a result of coating growth on a relatively rough substrate material. 3.3. Chemical and phase composition The single layered coatings had an elemental composition close to that of stoichiometric TiN and CrN ( Table 2). All multilayered coatings were found to be understoichiometric (with respect to the cubic TiN and CrN phase), with a metal to nitrogen (Me:N ) ratio ranging from 1.3 to 2.0, see Table 2. It was also found that the substrate bias had no influence on the composition of the coatings. As expected, an increased magnetron power yielded a decrease in the Ti to Cr ratio as well as the nitrogen content in the coatings (see Table 2). An increased nitrogen flow was found to result in an increased Ti to Cr ratio. The Me to N ratio was not affected by the nitrogen flow, see Table 2. Single layered TiN displayed a cubic NaCl phase (PDF file 31-1403) with a (111) preferred growth orientation. Single layered CrN showed a peak corresponding ˚ . Due to the large to a lattice plane spacing of 2.03 A width of this peak it could not be concluded whether it corresponded to the cubic NaCl phase (PDF file 11-0065) with a (200) preferred growth orientation or the hexagonal b-Cr N (PDF file 35-0803) with a (111) 2 or (200) preferred growth orientation. However, the chemical composition suggests that it was the cubic NaCl phase of CrN. All multilayered coatings possessed a peak originating ˚ , see Fig. 5. from a plane spacing of approximately 2.13 A This corresponds well to the (200) peak of the cubic TiN phase. Unfortunately, due to a slight peak overlap of the (200) peak of cubic CrN, the (111) and (200)
peaks of b-Cr N, together with the fact that these 2 coatings are understoichiometric (with respect to the cubic TiN and CrN ), it was not possible to determine the origin of this peak. For simplicity, this peak will be referred to as the (200) peak of TiN for the remainder of this paper. M(1 rpm) showed, apart from the (200) peak, a peak resulting from (111) oriented TiN crystals as well as from (302) oriented b-Cr N. No influence of the nitro2 gen flow on the phase composition could be found. M(4 kW ) displayed a broad peak of extremely low intensity revealing small amounts of metallic Cr (PDF file 06-0694) with a (200) preferred growth orientation. The intensity of the (200) Cr peak was found to increase with magnetron power. 3.4. Coating hardness, Young’s modulus and fracture toughness The coating hardness and Young’s modulus of the coatings are presented in Fig. 6. The hardness of the TiN and CrN was 26.5 and 23.3 GPa, respectively. The hardness of the multilayered coatings ranged from 29.1 GPa to 31.9 GPa. In addition, the coating hardness decreased with increasing magnetron power, i.e. increasing Cr:Ti and Me:N ratio. It was also found that the hardness decreased with the rotary speed of the substrates, i.e. with increasing lamella thickness. Young’s modulus of the multilayered coatings was found to range between the value for CrN (340 GPa) and the value for TiN (500 GPa). The modulus was found to be relatively insensitive to changes in the process parameters. The only exception was found for the magnetron power, for which the modulus decreased
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Fig. 6. Coating hardness and Young’s modulus of the coatings.
with increasing magnetron power, i.e. with increasing Cr:Ti and Me:N ratio, see Fig. 6. The fracture load, L , ranged between 11 and 32 N, F see Table 2. The lowest load was found for M(0 V ), whereas the highest was found for the single layered CrN coating. For TiN and all multilayered coatings the level of the fracture load was comparable, in the range of 23 N to 27 N. The intensity of the AE signal, detected at the fracture load L , was in all cases 0.5–6 dB [cf. F Fig. 2(a)], and corresponded to fracture at the rim of the indent, cf. Fig. 2(b). At higher load some of the coatings displayed a high intensity signal of 30–150 dB [at 33 N for CrN, 76 N for M(1 rpm), 59 N for M(0 V ), 57 N for M(6 kW ) and 88 N for M(115 sccm)], see Fig. 7(a). This load was found to correspond to spalling of the coating, see Fig. 7(b).
4. Discussion 4.1. Chemical and phase composition By increasing the magnetron power from 2 to 6 kW it was found that the Cr to Ti ratio increased from 0.3
to 1.1 and that the nitrogen content decreased from 41 at% to 33 at%. In addition, the coatings deposited using a high magnetron power, i.e. M(4 kW ) and M(6 kW ), contained small amounts of metallic (200) oriented Cr. Evidently, for these two coatings, the nitrogen partial pressure must be increased in order to retain a stoichiometric multilayered TiN/CrN coating. However, this was not possible as an increase in nitrogen partial pressure will act detrimentally on the quality of the TiN phase. The reason for this is that the electron beam evaporation process of TiN requires a relatively low total pressure as compared to that of the d.c. magnetron sputtering process. By increasing the nitrogen flow from 115 to 195 sccm an increase in the Ti to Cr ratio from 1.8 to 5.3 was obtained, whereas the nitrogen content in the coating was unaffected. This was believed to be a result of the fact that the evaporation rate of Ti increases with the nitrogen flow, since the Ti emission was controlled to obtain a constant total pressure (and therefore also a constant partial pressure of N ) in the chamber. As a 2 result, the nitrogen content in the coating will be unchanged. Another explanation for the increase of Ti as compared to Cr could be that the nitrogen reactivity
Fig. 7. Representative examples of (a) high intensity AE signal and (b) the corresponding coating failure.
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Table 3 Estimated TiN and CrN lamella thickness Coating
t (nm) TiN
t (nm) CrN
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
– –
– –
7 14 80 7 7 7 7 3 10
3 5 30 3 3 5 8 3 3
is higher for Ti than for Cr. This was found by Jehn et al. [3] who showed that the Ti to Cr ratio increased from 2.5 to 4.0 as the nitrogen pressure was increased from 8×10−5 to 3×10−3 mbar, although no pressure control was utilised. It is clear that when discussing the influence of the deposition parameters on the phase composition of the coatings it would be more convenient to use the lamella thickness, i.e. the thickness of the individual layers of TiN and CrN, instead of the multilayer period, since the two individual components have different properties. Unfortunately, the individual lamella thickness cannot be determined accurately for all the multilayered coatings using FEG-SEM. Therefore an estimation of the lamella thicknesses was done and used when discussing the phase composition and mechanical properties of the coatings, see Table 3. An interesting finding was that M(1 rpm) showed diffraction peaks originating from cubic NaCl TiN (111) as well as of TiN (200), cf. Fig. 5. Homogeneous TiN displayed only a (111) peak, whereas thin TiN lamellae in e.g. M(ref ) and M(5 rpm) only displayed a (200) peak. This ought to mean that the thickness of the TiN lamellae in M(1 rpm) was higher than the critical thickness for transformation of the preferred growth orientation from (200) to (111). That is, the critical lamella thickness for the transformation is in the range 14 to 80 nm. A similar finding was done for the CrN. For the M(1 rpm) coating four peaks originating from (200) CrN, (200) b-Cr N, (111) b-Cr N as well as (302) 2 2 b-Cr N were observed ( Fig. 5). For the coatings with 2 thinner CrN lamellae, i.e. M(ref ) and M(5 rpm), only the (200) CrN peak and/or (200) and (111) b-Cr N 2 peak were observed. The SAED analysis showed that the M(ref ) consisted of 100% cubic phase, i.e. no traces of the b-Cr N were found. This indicated that the 2 thickness of the CrN lamella in the M(1 rpm) was above, and in the M(ref ) [and probably also M(5 rpm)] below, the critical thickness for transformation from the NaCl phase with a (200) preferred growth orientation
to a mixed phase and texture composition of (200) and/or (111) and (302) b-Cr N. Hence, the critical 2 thickness for transformation must be somewhere between 5 and 31 nm. It is clear that the lamella thickness, i.e. the thickness of deposited material per revolution, has to be kept low in order to stabilise the cubic NaCl phase as well as the (200) preferred growth orientation. The reason why the single layered CrN most likely exhibits a NaCl phase with 100% (200) preferred growth orientation is probably the high rotary speed of the substrates used during CrN deposition. This will result in a short exposure time above the sputtering target. As a result, only thin (approximately 3 nm) CrN layers will be deposited at each revolution. This also ensured a thorough argon ion etching of each CrN layer deposited, since the etching takes place almost solely away from the sputtering target. To confirm this theory an additional CrN coating with the substrates held stationary over the sputtering target was deposited. It was observed that the coating consisted of a mixture of cubic CrN and hexagonal b-Cr N, or more likely only hexagonal b-Cr N. 2 2 Apart from a peak representing (200) CrN and/or (111) b-Cr N, which was still the dominating diffraction peak, 2 several peaks were present such as (200) Cr, (111) b-Cr N, (112) b-Cr N and (302) b-Cr N. 2 2 2 4.2. Fracture toughness The coatings that displayed cohesive/adhesive failures at high indentation loads have one feature in common, namely thick CrN lamellae [approximately 30 nm in M(1 rpm)] or a high CrN to TiN ratio [8:7 for M(6 kW ) and 3:3 for M(115 sccm)]. From this it can be concluded that CrN lamellae should be kept thin if a coating with high cohesive strength is to be made. Finally it can be mentioned that the results of the indentation tests were in agreement with the results obtained in an earlier investigation of the coating cohesion/adhesion using the well-known scratch test [15].
5. Summary The investigation was performed in order to increase the understanding of multilayered PVD coatings and especially of the TiN/CrN coating system. It was found that a dense, fully NaCl cubic, multilayered PVD TiN/CrN coating with a high cohesive strength can be deposited using the hybrid technique combining reactive electron beam evaporation of Ti and d.c. magnetron sputtering of Cr. To obtain a high quality coating it was necessary to use not only a negative substrate bias, but also to deposit very thin TiN lamellae, ≤14 nm, and thin CrN lamellae, ≤5 nm. Thin lamellae seem to inhibit transformation from growth of the cubic NaCl phase to new phases, e.g. the hexagonal b-Cr N and the metallic 2
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Cr. Furthermore, thin lamellae coatings grew in a 100% (200) preferred orientation, while thicker lamellae coatings displayed a mixed texture.
Acknowledgements The financial support of AB Sandvik Coromant and the Swedish Research Council for Engineering Sciences ( TFR) is gratefully acknowledged. Torbjo¨rn Selinder at AB Sandvik Coromant is recognised for providing the cemented carbide substrates.
References [1] O. Knotek, F. Lo¨ffler, H.-J. Scholl, Surf. Coat. Technol. 45 (1991) 53.
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[2] B. Navinsek, P. Panjan, Surf. Coat. Technol. 74/75 (1995) 919. [3] H.A. Jehn, F. Thiergarten, E. Ebersbach, D. Fabian, Surf. Coat. Technol. 50 (1991) 45. [4] S.B. Sant, K.S. Gill, Surf. Coat. Technol. 68/69 (1994) 152. [5] C. Subramanian, K.N. Strafford, Wear 165 (1993) 85. [6 ] H. Holleck, V. Schier, Surf. Coat. Technol. 76/77 (1–3) (1995) 328. [7] U. Helmersson, S. Todorova, S.A. Barnett, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Appl. Phys. 62 (2) (1987) 481. [8] X. Chu, M.S. Wong, W.D. Sproul, S.L. Rohde, S.A. Barnett, J. Vac. Sci. Technol. A 10 (4) (1992) 1604. [9] P. Yashar, X. Chu, S.A. Barnett, J. Rechner, Y.Y. Wang, M.S. Wong, W.D. Sproul, Appl. Phys. Lett. 72 (8) (1998) 987. [10] U. Wiklund, O. Wa¨nstrand, M. Larsson, S. Hogmark, Wear (1998) in press. [11] M. Nordin, M. Larsson, S. Hogmark, Surf. Coat. Technol. 106 (2/3) (1998) 234. [12] P. Panjan, B. Navinsek, A. Cvelbar, A. Zalar, J. Vlcek, Surf. Coat. Technol. 98 (1998) 1497. [13] P.T. Dawson, K.K. Tzatzov, Surf. Sci. 149 (1985) 105. [14] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (6) (1992) 1564. [15] M. Nordin, M. Larsson, S. Hogmark, Wear (1998) in press.
Deposition and characterisation of multilayered PVD TiN/CrN coatings on cemented carbide Maria Nordin *, Mats Larsson Uppsala University, Department and Division of Materials Science, Box 534, S-751 21 Uppsala, Sweden
Abstract In laboratory as well as in application tests, multilayered PVD coatings have shown enhanced mechanical and tribological properties as compared to today’s single layered PVD coatings. A coating which has shown very interesting properties, such as high temperature stability and high fracture toughness, is multilayered PVD TiN/CrN. Our knowledge about the connection between the growth dynamics and the properties of this coating is, however, rather poor. Therefore, to further develop and optimise this coating, it is necessary to study the correlations between on the one hand the growth process, and on the other the microstructure, composition and mechanical properties of the coating. In this work growth rate, morphology, microstructure, chemical and phase composition, together with coating hardness, Young’s modulus and fracture toughness of different multilayered PVD TiN/CrN coatings have been evaluated. All coatings have been deposited on cemented carbide substrates using a combination of reactive electron beam evaporation ( Ti) and reactive d.c. magnetron sputtering (Cr). The influence of lamella thickness and different deposition parameters: substrate bias, magnetron sputtering power and nitrogen flow, on the above mentioned characteristics has been examined. In addition to the multilayered TiN/CrN coatings, homogeneous TiN and CrN have been included and compared. The investigation showed that a dense, fully cubic NaCl phase, multilayered PVD TiN/CrN coating with high fracture toughness can be deposited provided that the lamella thickness is kept less than 14 nm and 5 nm of TiN and CrN respectively. Thin lamellae seem to inhibit transformation from growth of the cubic NaCl phase to new phases, e.g. hexagonal b-Cr N and 2 metallic Cr. Furthermore, thin lamellae yielded a (200) preferred growth orientation while thicker lamellae generated a mixture of growth orientations. In addition, very thin lamellae must be deposited to obtain good fracture toughness. It was also found that it was necessary to use a negative substrate bias in order to obtain a high quality coating. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cemented carbide; Coating; Multilayer; PVD; TiN/CrN
1. Introduction For metal cutting tools two of the most common commercial PVD coatings are TiN and CrN. Two other frequently used coatings, not considered in this work, are Ti(C,N ) and ( Ti,Al )N. Of these TiN is by far the most extensively investigated and used material. This is partly due to the ease with which TiN can be deposited to its very hard and wear resistant cubic NaCl phase (it has a large nitrogen solubility) and partly due to its beautiful golden colour. CrN, with its excellent high temperature stability [1–3], is often utilised in applications where a hard, oxidation- and corrosion-resistant * Corresponding author. Fax: +46-18-471-35-72. E-mail address: [email protected] (M. Nordin)
coating is required, e.g. on moulds for Al alloy die casting [2]. CrN is also known for its high toughness, which facilitates growth of thicker coatings (5–20 mm). A characteristic of CrN is that it is more difficult to deposit to its cubic NaCl phase, due to its lower nitrogen solubility and lower stability as compared to TiN. Nevertheless, some papers have shown that it is possible to grow stoichiometric cubic CrN coatings, e.g. Refs. [2,4]. There have been several scientific works on multilayered PVD coatings, i.e. coatings obtained by alternately depositing two chemically and/or mechanically different materials to form a layered structure, during the last years. In many cases multilayered coatings have shown enhanced mechanical and tribological properties as compared to today’s single layered coatings [5,6 ]. It has also been shown that by layering thin layers of a metal
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nitride (preferably less than a few nm of coatings such as VN, NbN and CrN ) with layers of TiN, fully NaCl cubic coatings with a high hardness can be grown [7– 9]. Up to today, most of the research on multilayered or superlattice coatings has been on the systems TiN/VN and TiN/NbN [7,8]. Another multilayered coating which has shown very interesting properties, such as high temperature stability and high fracture toughness, is made of the two nitrides TiN and CrN [10,11]. Our knowledge about the correlations between on the one hand the growth dynamics, chemical composition, phase composition and microstructure, and on the other the mechanical/tribological properties of multilayered TiN/CrN coatings is, however, scarce. Only a few works concerning this coating can be found. Yashar et al. have investigated the growth process and phase composition of multilayered TiN/CrN as a function of lamella 0.6 thickness [9]. They showed that CrN , which normally 0.6 is hexagonal, can be stabilised into a cubic NaCl phase provided that the lamellae are kept less than 10 nm. Furthermore, Panjan et al. showed that it is possible to increase the activation energy for oxidation by layering TiN with CrN (modulation periods in the range 83– 425 nm) as compared to the single layered coatings made of the constituents [12]. The aim of this work is to increase our knowledge of the growth dynamics and properties of multilayered PVD TiN/CrN coatings. A number of multilayered TiN/CrN coatings have been characterised with respect to their growth rate, morphology, microstructure, chemical composition and phase composition as well as their composite hardness and fracture toughness. The influence of the lamella thickness and the deposition parameters: substrate bias, magnetron sputtering power and nitrogen flow, on those characteristics has been evaluated. In addition to the multilayered coatings, homogeneous TiN and CrN have been included as reference coating materials.
2. Experimental 2.1. Substrate material Cemented carbide (10 wt% Co and 90 wt% WC ) was used as substrate material. The substrate hardness was 1350 HV . All substrates were polished to a surface 30 kg roughness of approximately 5 nm and then ultrasonically cleaned in a heated alkaline solution followed by ethanol (5 min each). 2.2. Coating deposition All coatings were deposited in a commercial Balzers BAI 640R coating unit. Prior to coating deposition, the samples were resistively heated to 450°C for 60 min and
thereafter Ar ion etched for 15 min using a substrate bias of −200 V. The coatings were reactively deposited using a combination of electron beam gun (e-gun) evaporation of Ti and d.c. magnetron sputtering of Cr. For all the multilayered coatings the deposition sequence started with the deposition of a 30 nm Ti layer followed by a 200 nm TiN layer to promote coating adhesion. This was obtained while the substrates were held stationary above the Ti source. Simultaneously the Cr target was Ar ion etched. By keeping both sources active and by alternately expose the substrates to the Ti source and the Cr source using substrate rotation, a multilayered coating was grown. The total pressure was kept constant at 2×10−3 mbar. Before the deposition of single layered CrN and TiN commenced, a 30 nm layer (Cr and Ti, respectively) was deposited. Single layered TiN and CrN were deposited by keeping only one of the material sources active and admitting nitrogen gas into the chamber. The total pressure was kept at 1.7×10−3 mbar during the TiN deposition and at 3×10−3 mbar during the CrN deposition. The total deposition time was 40 min for all the multilayered coatings, 60 min for TiN and 160 min for CrN. In Table 1 an overview of the evaluated process parameters are presented. 2.3. Coating characterisation From studies of fractured coating cross-sections, information about the coating thickness, multilayer period (L; the thickness of one lamella of TiN together with one lamella of CrN ) and coating morphology could be obtained using scanning electron microscopy fitted with a field emission gun (FEG-SEM ), see e.g. Fig. 1. An atomic force microscope (AFM ) was used to reveal the surface topography. Cross-sectional transmission electron microscopy ( XTEM ) in a JEOL 2000 FXII Table 1 An overview of the deposition process parameters used. M(x) indicates a multilayered TiN/CrN coating with the varied parameter within brackets [in comparison to the reference coating designated M(ref )] Coating
Substrate Magnetron N Substrate rotary 2 bias ( V ) power (kW ) flow (sccm) speed (rpm)
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
−110 −110 −110 −110 −110 0 −200 −110 −110 −110 −110
– 2 2 2 2 2 2 4 6 2 2
140 40 155 155 155 155 155 155 155 115 195
10 10 10 5 1 10 10 10 10 10 10
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developed by Oliver and Pharr [14]. A Berkovich diamond tip and an indentation depth of 150 nm were used. By performing Rockwell C indentation tests (in a commercial scratch tester with the specimen translation disconnected ), information about the coating toughness could be gained. The load was continuously increased up to 100 N at a loading rate of 10 N/min and subsequently decreased at the same rate. As a measure of the coating fracture toughness, the fracture load, L , i.e. the F load at the first coating fracture as detected using acoustic emission (AE) was utilised [Fig. 2(a)]. In all cases, this load corresponded to a circular crack at the rim of the indent, see Fig. 2(b). All indentations were post-experimentally studied in a light optical microscope (LOM ). Fig. 1. Representative example of a SEM micrograph [M(6 kW )] showing the layered structure. In the micrograph the column diameter is indicated as well as 10 multilayer periods.
equipped with a LaB filament and a 200 kV voltage 6 was employed to evaluate the microstructure and selected area electron diffraction (SAED) to determine the phase composition of one multilayered coating which in the following is denoted M(ref ). The chemical composition of the coatings was determined using Auger electron spectroscopy (AES). To subdivide the intensity from the overlapping transitions of Ti and N at approximately 382 eV into separate components the method obtained by Dawson and Tzatzov [13] was used. The electron beam energy and emission current were 10 keV and 5 mA, respectively. All the sensitivity factors ( Ti, Cr and N ) were obtained from nitride reference powders. The phase composition and preferred growth orientation were determined using a Siemens D5000 X-ray diffractometer. Cu Ka radiation over an area of approximately 1 mm2 was employed and the measurement range was 2h=30–100°. No sample rotation was assessed during measurement. The composite hardness and Young’s modulus were determined using nanoindentation and the method
3. Results 3.1. Coating thickness and multilayer period For the multilayered coatings it was found that the coating thickness (and the multilayer period), which is directly related to the coating growth rate since the deposition time was the same for all multilayered coatings, increased with magnetron power and nitrogen flow and slightly decreased with substrate bias, see Table 2. The coating thickness was, however, not affected by the substrate rotary speed, while, naturally, L decreased with increasing rotary speed. 3.2. Morphology and microstructure All coatings deposited (both single- and multilayered) displayed a dense and homogeneous surface morphology, see Fig. 3(a). The only exception was M(0 V ), which appeared to be more porous, see Fig. 3(b). No significant influence of the magnetron power, nitrogen flow or rotary speed on the surface morphology was observed.
Fig. 2. The fracture load (L ) was defined as the load corresponding to the first significant increase in the AE signal. (a) A typical example of the F AE signal recorded during indentation. (b) Representative appearance of an indented coating, note the circular cracks each corresponding to one of the AE peaks (SEM ).
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Table 2 Coating thickness, t, multilayer period, L, chemical composition (approximate measurement scatter was 5%), Ti to Cr ratio (Ti:Cr), Me to N ratio (Me:N ) and fracture load, L F Coating
t (nm)
L (nm)
at% Ti
at% Cr
at% N
Ti:Cr
Me:N
L (N ) F
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
3800±100 4400±150 3700±100 3600±100 3900±100 3900±100 3400±100 4200±150 5000±200 2500±100 4800±200
– –
51 – 45 46 43 44 48 38 32 37 48
– 53 14 11 12 14 12 24 35 21 9
49 47 41 43 45 42 40 38 33 42 43
– – 3.2 4.2 3.6 3.1 4.0 1.6 0.9 1.8 5.3
1.0 1.1 1.4 1.3 1.2 1.4 1.5 1.6 2.0 1.4 1.3
27±2 32±1 27±1 26±1 23±1 11±3 27±9 23±2 24±3 24±2 26±1
9.7±0.5 19.1±1 110.0±5 9.8±0.5 9.5±0.5 12.0±0.6 14.5±0.7 6.1a±0.5 12.4±0.6
a Period is calculated from the coating growth rate since it was too thin to measure by SEM.
All coatings were found to have a more or less columnar structure when studied in SEM. When comparing the multilayered coatings, the largest column diameter was found for the coating deposited at 0 V substrate bias. Furthermore, it was found that the
diameter increased with magnetron sputtering power. For the other multilayered coatings no clear difference in column size could be seen. The TiN coating displayed a more pronounced columnar microstructure than both CrN and the multilayered coatings. XTEM displayed that M(ref ) was columnar, dense and fine grained (see Fig. 4), which is in agreement with the results obtained using SEM. XTEM also showed that coating growth, and hence coating surface profile, well followed the surface profile of the substrate. The substrate roughness developed during Ar ion etching of the substrates prior to coating deposition. The lamella thickness of CrN and TiN was approximately 3 and 7 nm, respectively. The SAED analysis revealed a 100% cubic NaCl phase composition, see diffraction pattern insert in Fig. 4. It was also observed that the diffraction
Fig. 3. Typical coating surface topography (AFM ) of (a) the coatings deposited using substrate bias [M(ref )] and (b) without substrate bias [M(0 V )].
Fig. 4. XTEM micrograph and SAED pattern of M(ref ). The broadening of the diffraction peaks was due to coating growth on a not perfectly smooth substrate.
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Fig. 5. XRD patterns of the multilayered coatings with different L. S refers to peaks originating from the substrate.
spots were somewhat broadened as a result of coating growth on a relatively rough substrate material. 3.3. Chemical and phase composition The single layered coatings had an elemental composition close to that of stoichiometric TiN and CrN ( Table 2). All multilayered coatings were found to be understoichiometric (with respect to the cubic TiN and CrN phase), with a metal to nitrogen (Me:N ) ratio ranging from 1.3 to 2.0, see Table 2. It was also found that the substrate bias had no influence on the composition of the coatings. As expected, an increased magnetron power yielded a decrease in the Ti to Cr ratio as well as the nitrogen content in the coatings (see Table 2). An increased nitrogen flow was found to result in an increased Ti to Cr ratio. The Me to N ratio was not affected by the nitrogen flow, see Table 2. Single layered TiN displayed a cubic NaCl phase (PDF file 31-1403) with a (111) preferred growth orientation. Single layered CrN showed a peak corresponding ˚ . Due to the large to a lattice plane spacing of 2.03 A width of this peak it could not be concluded whether it corresponded to the cubic NaCl phase (PDF file 11-0065) with a (200) preferred growth orientation or the hexagonal b-Cr N (PDF file 35-0803) with a (111) 2 or (200) preferred growth orientation. However, the chemical composition suggests that it was the cubic NaCl phase of CrN. All multilayered coatings possessed a peak originating ˚ , see Fig. 5. from a plane spacing of approximately 2.13 A This corresponds well to the (200) peak of the cubic TiN phase. Unfortunately, due to a slight peak overlap of the (200) peak of cubic CrN, the (111) and (200)
peaks of b-Cr N, together with the fact that these 2 coatings are understoichiometric (with respect to the cubic TiN and CrN ), it was not possible to determine the origin of this peak. For simplicity, this peak will be referred to as the (200) peak of TiN for the remainder of this paper. M(1 rpm) showed, apart from the (200) peak, a peak resulting from (111) oriented TiN crystals as well as from (302) oriented b-Cr N. No influence of the nitro2 gen flow on the phase composition could be found. M(4 kW ) displayed a broad peak of extremely low intensity revealing small amounts of metallic Cr (PDF file 06-0694) with a (200) preferred growth orientation. The intensity of the (200) Cr peak was found to increase with magnetron power. 3.4. Coating hardness, Young’s modulus and fracture toughness The coating hardness and Young’s modulus of the coatings are presented in Fig. 6. The hardness of the TiN and CrN was 26.5 and 23.3 GPa, respectively. The hardness of the multilayered coatings ranged from 29.1 GPa to 31.9 GPa. In addition, the coating hardness decreased with increasing magnetron power, i.e. increasing Cr:Ti and Me:N ratio. It was also found that the hardness decreased with the rotary speed of the substrates, i.e. with increasing lamella thickness. Young’s modulus of the multilayered coatings was found to range between the value for CrN (340 GPa) and the value for TiN (500 GPa). The modulus was found to be relatively insensitive to changes in the process parameters. The only exception was found for the magnetron power, for which the modulus decreased
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Fig. 6. Coating hardness and Young’s modulus of the coatings.
with increasing magnetron power, i.e. with increasing Cr:Ti and Me:N ratio, see Fig. 6. The fracture load, L , ranged between 11 and 32 N, F see Table 2. The lowest load was found for M(0 V ), whereas the highest was found for the single layered CrN coating. For TiN and all multilayered coatings the level of the fracture load was comparable, in the range of 23 N to 27 N. The intensity of the AE signal, detected at the fracture load L , was in all cases 0.5–6 dB [cf. F Fig. 2(a)], and corresponded to fracture at the rim of the indent, cf. Fig. 2(b). At higher load some of the coatings displayed a high intensity signal of 30–150 dB [at 33 N for CrN, 76 N for M(1 rpm), 59 N for M(0 V ), 57 N for M(6 kW ) and 88 N for M(115 sccm)], see Fig. 7(a). This load was found to correspond to spalling of the coating, see Fig. 7(b).
4. Discussion 4.1. Chemical and phase composition By increasing the magnetron power from 2 to 6 kW it was found that the Cr to Ti ratio increased from 0.3
to 1.1 and that the nitrogen content decreased from 41 at% to 33 at%. In addition, the coatings deposited using a high magnetron power, i.e. M(4 kW ) and M(6 kW ), contained small amounts of metallic (200) oriented Cr. Evidently, for these two coatings, the nitrogen partial pressure must be increased in order to retain a stoichiometric multilayered TiN/CrN coating. However, this was not possible as an increase in nitrogen partial pressure will act detrimentally on the quality of the TiN phase. The reason for this is that the electron beam evaporation process of TiN requires a relatively low total pressure as compared to that of the d.c. magnetron sputtering process. By increasing the nitrogen flow from 115 to 195 sccm an increase in the Ti to Cr ratio from 1.8 to 5.3 was obtained, whereas the nitrogen content in the coating was unaffected. This was believed to be a result of the fact that the evaporation rate of Ti increases with the nitrogen flow, since the Ti emission was controlled to obtain a constant total pressure (and therefore also a constant partial pressure of N ) in the chamber. As a 2 result, the nitrogen content in the coating will be unchanged. Another explanation for the increase of Ti as compared to Cr could be that the nitrogen reactivity
Fig. 7. Representative examples of (a) high intensity AE signal and (b) the corresponding coating failure.
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Table 3 Estimated TiN and CrN lamella thickness Coating
t (nm) TiN
t (nm) CrN
TiN CrN M(ref ) M(5 rpm) M(1 rpm) M(0 V ) M(−200 V ) M(4 kW ) M(6 kW ) M(115 sccm) M(195 sccm)
– –
– –
7 14 80 7 7 7 7 3 10
3 5 30 3 3 5 8 3 3
is higher for Ti than for Cr. This was found by Jehn et al. [3] who showed that the Ti to Cr ratio increased from 2.5 to 4.0 as the nitrogen pressure was increased from 8×10−5 to 3×10−3 mbar, although no pressure control was utilised. It is clear that when discussing the influence of the deposition parameters on the phase composition of the coatings it would be more convenient to use the lamella thickness, i.e. the thickness of the individual layers of TiN and CrN, instead of the multilayer period, since the two individual components have different properties. Unfortunately, the individual lamella thickness cannot be determined accurately for all the multilayered coatings using FEG-SEM. Therefore an estimation of the lamella thicknesses was done and used when discussing the phase composition and mechanical properties of the coatings, see Table 3. An interesting finding was that M(1 rpm) showed diffraction peaks originating from cubic NaCl TiN (111) as well as of TiN (200), cf. Fig. 5. Homogeneous TiN displayed only a (111) peak, whereas thin TiN lamellae in e.g. M(ref ) and M(5 rpm) only displayed a (200) peak. This ought to mean that the thickness of the TiN lamellae in M(1 rpm) was higher than the critical thickness for transformation of the preferred growth orientation from (200) to (111). That is, the critical lamella thickness for the transformation is in the range 14 to 80 nm. A similar finding was done for the CrN. For the M(1 rpm) coating four peaks originating from (200) CrN, (200) b-Cr N, (111) b-Cr N as well as (302) 2 2 b-Cr N were observed ( Fig. 5). For the coatings with 2 thinner CrN lamellae, i.e. M(ref ) and M(5 rpm), only the (200) CrN peak and/or (200) and (111) b-Cr N 2 peak were observed. The SAED analysis showed that the M(ref ) consisted of 100% cubic phase, i.e. no traces of the b-Cr N were found. This indicated that the 2 thickness of the CrN lamella in the M(1 rpm) was above, and in the M(ref ) [and probably also M(5 rpm)] below, the critical thickness for transformation from the NaCl phase with a (200) preferred growth orientation
to a mixed phase and texture composition of (200) and/or (111) and (302) b-Cr N. Hence, the critical 2 thickness for transformation must be somewhere between 5 and 31 nm. It is clear that the lamella thickness, i.e. the thickness of deposited material per revolution, has to be kept low in order to stabilise the cubic NaCl phase as well as the (200) preferred growth orientation. The reason why the single layered CrN most likely exhibits a NaCl phase with 100% (200) preferred growth orientation is probably the high rotary speed of the substrates used during CrN deposition. This will result in a short exposure time above the sputtering target. As a result, only thin (approximately 3 nm) CrN layers will be deposited at each revolution. This also ensured a thorough argon ion etching of each CrN layer deposited, since the etching takes place almost solely away from the sputtering target. To confirm this theory an additional CrN coating with the substrates held stationary over the sputtering target was deposited. It was observed that the coating consisted of a mixture of cubic CrN and hexagonal b-Cr N, or more likely only hexagonal b-Cr N. 2 2 Apart from a peak representing (200) CrN and/or (111) b-Cr N, which was still the dominating diffraction peak, 2 several peaks were present such as (200) Cr, (111) b-Cr N, (112) b-Cr N and (302) b-Cr N. 2 2 2 4.2. Fracture toughness The coatings that displayed cohesive/adhesive failures at high indentation loads have one feature in common, namely thick CrN lamellae [approximately 30 nm in M(1 rpm)] or a high CrN to TiN ratio [8:7 for M(6 kW ) and 3:3 for M(115 sccm)]. From this it can be concluded that CrN lamellae should be kept thin if a coating with high cohesive strength is to be made. Finally it can be mentioned that the results of the indentation tests were in agreement with the results obtained in an earlier investigation of the coating cohesion/adhesion using the well-known scratch test [15].
5. Summary The investigation was performed in order to increase the understanding of multilayered PVD coatings and especially of the TiN/CrN coating system. It was found that a dense, fully NaCl cubic, multilayered PVD TiN/CrN coating with a high cohesive strength can be deposited using the hybrid technique combining reactive electron beam evaporation of Ti and d.c. magnetron sputtering of Cr. To obtain a high quality coating it was necessary to use not only a negative substrate bias, but also to deposit very thin TiN lamellae, ≤14 nm, and thin CrN lamellae, ≤5 nm. Thin lamellae seem to inhibit transformation from growth of the cubic NaCl phase to new phases, e.g. the hexagonal b-Cr N and the metallic 2
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Cr. Furthermore, thin lamellae coatings grew in a 100% (200) preferred orientation, while thicker lamellae coatings displayed a mixed texture.
Acknowledgements The financial support of AB Sandvik Coromant and the Swedish Research Council for Engineering Sciences ( TFR) is gratefully acknowledged. Torbjo¨rn Selinder at AB Sandvik Coromant is recognised for providing the cemented carbide substrates.
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