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Structural and mechanical property of Si incorporated (Ti,Cr,Al)N coatings deposited by arc ion plating process
Surface & Coatings Technology 200 (2005) 1383 – 1390 www.elsevier.com/locate/surfcoat
Structural and mechanical property of Si incorporated (Ti,Cr,Al)N coatings deposited by arc ion plating process Kenji Yamamoto a,*, Susumu Kujime b, Kazuki Takahara b b
a Materials Research Lab. Kobe Steel Ltd, 1-5-5 Takatsuka-dai Nishi-ku, Kobe, 651-2271 Hyogo, Japan Advanced Products and Technologies Department, Machinery and Engineering Company, Kobe Steel Ltd, 3-1, 2-chome, Shinhama Arai-cho Takasago, Hyogo 676-8670, Japan
Available online 27 September 2005
Abstract (Ti,Cr,Al,Si)N coatings with different Al + Si fractions (0.6 and 0.65) were deposited by an arc ion plating (AIP) apparatus that is equipped with the plasma-enhanced type arc cathode. The (Ti,Cr,Al,Si)N coatings were deposited under different substrate bias voltages and effect of the deposition parameter on the composition, structure and mechanical properties was investigated. X-ray diffraction measurements of the (Ti,Cr,Al,Si)N coating deposited under different substrate bias voltages revealed that formation of the hexagonal phase (Wurzite structure) was only limited to a relatively low bias voltage range of 20 to 30 V. Above this bias voltage, the crystal structure of the coatings was single-phased cubic rock-salt structure (B1 phase) independent of the bias voltage. Grain size of the coating was calculated from the full width of half maximum (FWHM) of the X-ray diffraction peak and it was smaller than the one of conventional (Ti,Al)N or (Ti,Cr,Al)N coating with a comparable Al fraction. The grain size estimated from the cross-sectional TEM observation was less than 10 nm. From the TEM observation, the coating was compositionally homogeneous and there was no evidence that the film had a phase separation such as Si-rich and -poor region. Hardness of the (Ti,Cr,Al,Si)N coating with Al + Si = 0.6 was in the range of 26 to 27 GPa independent of the substrate bias. (Ti,Cr,Al,Si)N coating with Al + Si = 0.65 showed slight increase in hardness from 24 to 27 GPa when the substrate bias was increased to more than 100 V. To evaluate the oxidation resistance, annealing tests in the air at 1000 -C were conducted and surface SEM observations revealed that surface of the conventional (Ti,Al)N and (Ti,Cr,Al)N was covered with coarse oxide grains enriched with TiO2. Whereas only a dense but very thin protective oxide layer was observed in case of (Ti,Cr,Al,Si)N coating after the oxidation. These coatings were applied to the high-speed dry cutting tests against hardened D2 steel (HRC 60) and the result clearly indicated better performance of (Ti,Cr,Al,Si)N compared to the conventional coatings such as (Ti,Al)N and (Ti,Cr,Al)N. D 2005 Elsevier B.V. All rights reserved. Keywords: Si addition; (Ti,Cr,Al,Si)N; Cathodic arc; Al + Si ratio; Oxidation resistance
1. Introduction High hardness and resistance to oxidation at elevated temperatures have always been important criteria for hard coatings particularly in dry cutting applications [1,2]. This is mainly due to the requirement from the various industries to machine harder work-pieces at higher cutting speeds. This was the main reason in 1990s why TiN coating was replaced
* Corresponding author. Tel.: +81 78 992 5505; fax: +81 78 992 5512. E-mail address: [email protected] (K. Yamamoto). 0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.08.025
by (Ti,Al)N coating that had a higher hardness and was more oxidation resistant [3,4]. But of course this was not the end of ever-increasing demands from the industries to improve the oxidation resistance and hardness for better productivity and longer tool life. Recently, much attention has been paid to Si containing coatings such as (Ti,Si)N [5 –12], (Cr,Si)N [13], (Ti,Al,Si)N [14 – 21] and other Si containing coating systems [22,23]. These Si containing coatings have significantly better oxidation resistance compared to the ones without Si [5,14,20,23]. The role of Si in improving the oxidation resistance is not yet fully clarified in many cases. Choi et
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al. reported, however, that preferential oxidation of Si was observed after the oxidation test and this Si-rich oxide film acted as a diffusion barrier for further oxidation [5], much like in case of the preferential oxidation of Al in (Ti,Al)N coating [2]. Also this Si containing (M,Si)N (M: metal) system is known to form so-called ‘‘nano-composite’’ coatings [24,25]. In case of (Ti,Si)N system [6,11,12], authors have reported rather small grain size (less than 10 nm) at the Si content of 10 to 20 at.%. Sometimes small crystalline TiN grains surrounded by amorphous SiNx (aSiNx ) matrix were observed [5,6] and this phase separation was primarily considered as a basic mechanism of the formation of the nano-composite. In the case of (Ti,Al,Si)N system [14,17], addition of Si resulted in refinement of the grain size and formation of a-SiNx phase was observed when coating contained a fairly large amount of Si. (Ti,Al,Si)N coatings [14] containing relatively a small amount of Si (a few atomic percent) consisted of single B1 cubic phase and there was no evidence of formation of second phase such as a-SiNx . Only refinement of the grain was observed in this case. Another interesting aspect of (Ti,Al,Si)N system is change in the crystal structure cubic (B1) to hexagonal (B4) depending on the Al + Si fraction. Tanaka et al. [14] reported crystalline phase composition in (Ti,Al,Si)N system and when Al + Si ratio exceeds 0.61 the crystal structure changed from B1 single phase to B1 + B4 mixed phase. This change in the crystal structure also resulted in loss of the mechanical property where the hardness decreased from at maximum 34 GPa to 25 GPa. In (Ti,Al)N system, there are several reports indicating the phase boundary of B1 and B4 is located around Al ratio of 0.6 to 0.7 [26,27]. Tanaka et al. reported that in case of the coating containing no Si, it had B1 single phase at equivalent Al + Si fraction. This implies that Si has negative effect to maintain the B1 phase, which is usually a harder and a preferred phase for wear resistant applications [28]. In the previous paper [28], we reported properties of (Ti,Cr,Al)N coating system with a high Al fraction of more than 0.65. This (Ti,Cr,Al)N system was characterized by high hardness more than 30 GPa and high oxidation resistance compared to the conventional (Ti,Al)N coating. Motivation of the present work is to improve the properties of (Ti,Cr,Al)N coating further for better tribological performance such as cutting operation. For this purpose, the effect of Si incorporation on the structural and mechanical properties of (Ti,Cr,Al)N coatings was investigated with different Al + Si fractions. Also cutting tests were conducted in comparison with the conventional (Ti,Al)N and (Ti,Cr,Al)N coating.
2. Experimental details (Ti,Cr,Al,Si)N coatings with different Al + Si fractions were deposited by a batch type cathodic arc ion plating
coater equipped with a plasma-enhanced cathode. Details of the deposition equipment and the cathode can be found elsewhere [28]. Two kinds of Ti– Cr –Al – Si targets were prepared by a powder metallurgy process and used for the deposition. Their compositions were Ti0.2Cr0.2Al0.55Si0.05 (target A) and Ti0.15Cr0.2Al0.6Si0.05 (target B). Deposition was conducted in the above-mentioned coater in pure N2 atmosphere at pressure of 2.7 Pa. WC-Co cutting inserts (Mitsubishi Carbide SNGN120408), platinum foils (0.1 mmt) and WC-Co square end-mills (Mitsubishi Carbide 10 mmf, 6 flutes) were used as substrates. Prior to the deposition, these substrates were Ar-ion-etch cleaned for 5 min at Ar pressure of 2.7 Pa. After the cleaning procedure, arc was ignited and deposition was conducted at arc current of 150 A and the substrate temperature was regulated at about 500 -C. The substrate bias was varied from 20 to 150 V to investigate the effect of substrate bias on coating’s properties. The thickness of the coating layer was about 3 Am for all samples unless mentioned. The elemental composition of the coatings was determined using energy dispersive X-ray (EDX) analysis (Horiba EMAX), using ZAF correction and may contain an error approximately 10% at maximum. Crystal structure and preferred orientation of the coating was determined with X-ray diffraction (XRD, Rigaku RINT200-PC) using Cu-ka radiation and the grain size was calculated by the Scherrer’s equation using the full width of half maximum (FWHM) of (111) diffraction peak of the B1 phase [30]. Indentation hardness was measured using a nano-indentation instrument (Elionix ENT-1100) with a Berkovich type diamond indenter. The indenter tip shape correction was conducted using the method proposed by Sawa and Tanaka [31]. No thermal drift correction was used, since this instrument was installed in a thermally regulated chamber in which temperature drift rate is less than 0.1 -C/10 s and one indentation measurement took about 10 s. Microstructure, particularly, to detect the possible existence of separated phase such as a-SiNx , cross-sectional transmission electron microscope observation (TEM: Hitachi) was used. Coatings deposited on WC-Co inserts were cut out and thinned for TEM observation using focused ion beam and micro-lifting technique [32]. Oxidation tests were conducted by annealing the coated platinum foil samples in atmosphere (air) at 1000 -C for 30 min. After the oxidation tests, surface morphology of the samples was investigated by SEM. Depth composition profile of the oxide layer was measured by an Auger electron spectroscopy (AES Perkin Elmer PHI650) using Ar sputter for depth profile measurements. Finally, cutting tests were conducted using hardened AISI D2 cold working die steel (HRC 60) as a work-piece. Cutting parameters were as follows: cutting speed 150 m/min; feed 0.05 mm/flute; axial depth of cut 5 mm; radial depth of cut 0.1 mm; dry cut, air-blow only. Flank wear of the cutting edge was measured after the cutting length of 30 m.
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3. Result and discussion 3.1. Composition and crystal structure Composition of the coatings was slightly different from the targets used for depositions. The coatings deposited from both targets A and B had a few percentage less Al and also a few percentage enriched Cr than the targets, while other elements, Ti and Si, had almost the same composition with the targets. Compositions of (Ti0.2Cr0.23Al0.53Si0.04)N and (Ti0.14Cr0.22Al0.59Si0.05)N were obtained as the resulting compositions of the coatings deposited from targets A and B at the bias voltage of 70 V. This compositional deviation of Al and Cr was more pronounced as the substrate bias was increased. The decrease of Al composition corresponding to the change in the substrate bias can be explained by the preferential sputtering of Al which has a highest sputter yield among the coating’s elements [28,33]. Fig. 1(a) shows X-ray diffraction patterns of (Ti,Cr,Al,Si)N coatings deposited from the targets A and B (hereafter referred as coatings A and B) under various substrate biases. At the substrate bias of 20 V, the diffraction pattern of the coating B contained a weak diffraction peak from the
Fig. 2. Effect of the substrate bias voltage on (a) grain size and (b) position of (111) diffraction peak of (Ti,Cr,Al,Si)N coatings with different Al + Si fractions.
Fig. 1. (a) X-ray diffraction patterns of (Ti,Cr,Al,Si)N coatings with different Al + Si fractions deposited under various substrate biases. Marks (*) denote diffraction peaks from the substrate (WC-Co). (b) Result of deconvolution of overlapping peaks located between the diffraction angle of 32- to 40-.
hexagonal B4 phase. This weak peak, as shown in Fig. 1(b), can be clearly observed by a deconvolution of several overlapping peaks between the diffraction angle from 32- to 40-. This diffraction peak from the B4 phase was only observed for coating B deposited at 20 V, and coatings A and B deposited more than substrate bias of 20 V, only diffraction peaks belonging to the B1 cubic phase were observed. This bias (i.e. ion energy) induced transition in crystal structure was also observed in (Ti,Cr,Al)N system with a high Al fraction [28], yet the nature of this phase transition is not clarified. The preferred orientation of the coating changed from [111] to [100] as the bias was increased for both coatings A and B. Fig. 2 shows the effect of the substrate bias on the (a) grain size and (b) position of (111) peaks. In case of coating A, the grain size was about 12 nm at the bias of 20 V and decreased gradually as the bias was increased. It reached, however, plateau when the bias was increased for more than 50 V and the grain size became a constant value of around 8 nm, whereas the grain size showed nearly constant value of 8 nm for whole bias range in case of coating B. Detailed discussion of the grain size will be given in the Section 3.3. The position of (111) peak was changed by the substrate bias and it shifted toward lower diffraction angle
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Fig. 3. Effect of the substrate bias voltage on (a) indentation hardness and (b) elastic modulus of (Ti,Cr,Al,Si)N coatings.
as the bias was increased. Two factors may influence the position of (111) peak in this case; composition and residual stress. As described at the beginning of this section, the Al fraction decreased and Cr fraction increased as the substrate bias was increased. This change in the composition may induce the change in the lattice parameter thus changing the position of (111) peak. The lattice parameter of CrN and AlN is 0.414 and 0.412 nm and nearly identical, however, and a few atomic percent change of Al and Cr composition can induce the change in diffraction angle of (111) peak smaller than 0.01-. Therefore, the observed change in the (111) peak position was likely due to the change of residual stress that was induced by the substrate bias (ion energy). Usually, PVD-deposited coatings are in compressive stress and an increase in compressive stress results in peak shift to smaller diffraction angle. From this consideration, if we assume the absolute value of the coating’s stress to be compressive, the compressive stress increased as the substrate bias was increased up to 50 V. More than this substrate bias, the stress of the coating showed little change.
Indentation hardness and elastic modulus of coating A showed very little change over the whole bias range and they were about 25 and 400 GPa. In case of coating B, both indentation hardness and elastic modulus increased gradually as the bias was increased. This change in indentation hardness of coating B probably relates to the bias induced crystal structure change observed by the X-ray diffraction measurements. The maximum indentation hardness of coating B was approximately 27 GPa at highest substrate bias of 150 V. These determined indentation hardness and elastic modulus of coatings A and B were both lower than the values of (Ti,Cr,Al)N coatings [28]. In case of (Ti,Cr,Al)N coating, hardness was significantly low, less than 20 GPa, at a lower substrate bias when the coating was composed of mixture of B1 and B4 phase. However, the maximum hardness of approximately 35 GPa was obtained at the bias of 150 V. Incorporation of Si to the (Ti,Cr,Al)N resulted in lower hardness and elastic modulus. However, on the positive side these mechanical properties were more insensitive to the deposition parameter and this is favorable from the viewpoint of robustness of the deposition process. Additionally, Tsui et al. [35] proposed that H 3/E*2 gave information on the resistance of the material to plastic deformation, where H is the indentation hardness and E* is reduced modulus E/(1 m 2). Coatings with high H 3/E*2 values are less likely undergo plastic deformation under external force. Fig. 4 shows E* – indentation hardness relationship for (Ti,Cr,Al,Si)N coatings and (Ti,Cr,Al)N deposited under various substrate biases. Maximum hardnesses of both (Ti,Cr,Al,Si)N coatings were lower than the ones of (Ti,Cr,Al)N coating. When compared at same E*, however, (Ti,Cr,Al,Si)N coatings had higher indentation hardness, thus higher H 3/E*2 value. 3.3. TEM observation Fig. 5 show (a) a high-resolution bright field TEM image of (Ti,Cr,Al,Si)N (coating A), (b) an electron diffraction (ED) pattern with an electron beam diameter
3.2. Mechanical property Fig. 3 shows change in (a) indentation hardness and (b) elastic modulus of coatings deposited under various biases.
Fig. 4. Indentation hardness – reduced modulus (E*) relationship of (Ti,Cr,Al,Si)N and (Ti,Cr,Al)N coatings.
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Fig. 5. (a) High-resolution TEM image of (Ti,Cr,Al,Si)N coating (Al + Si = 0.6), (b) ED pattern with electron beam diameter of approximately 1000 nm and (c) nano-ED pattern of area c.
of about 1000 nm, (c) an ED pattern of region c in the TEM image. Coating A was deposited at the bias of 50 V and substrate temperature of 500 -C. The TEM image showed nano-crystalline nature of the (Ti,Cr,Al,Si)N coating. The grains were sized less than 10 nm and most of them were as small as 5 nm, whereas the grain size of the (Ti0.25Cr0.1Al0.65)N was about 12 nm [34]. The grain sizes of (Ti,Si)N and (Ti,Al,Si)N coatings were reported by several authors using mainly TEM [6,14,17,19]. Tanaka et al. reported that the grain size of (Ti0.41Al0.59)N was approximately 120 to 350 nm, whereas that of (Ti0.42Al0.58Si0.03)N was 50 to 250 nm [14]. ParlinskaWojtan et al. [19] reported the effect of Al + Si content on the grain size of (Ti,Al,Si)N coating with Al + Si content ranging from 10 to about 50 at.% with comparable Si content of about 4 to 6 at.% [14]. The reported grain size showed strong correlation with the Al + Si content and it was less than 10 nm when Al + Si content was more than 40 at.%. These two reports are quite controversial in the absolute grain size, but they agree with each other on the fact that Si has an effect to reduce the grain size. The ED
pattern with fairly large beam diameter (Fig. 5(b)) showed that film only consisted cubic B1 phase and no other phase like hexagonal B4 phase was observed. Additional ED patterns were taken using nano-electron beam with beam diameter was about 1 to 2 nm. Fig. 5(c) shows a typical nano-ED pattern of a single crystal grain. Again only diffraction spots corresponding to the B1 phase were confirmed. 3.4. Oxidation resistance evaluation Fig. 6 shows surface SEM images of different coatings (a) –(d) as deposited and (a-1) –(d-1) after annealed in air at 1000 -C for 30 min. Some macro-particles (MPs) were observed on the surface of the as-deposited coatings. The number of MPs was less for the coatings without Si. After the oxidation tests, surface of coatings without Si (a-1), (b-1) showed a coarse grain-like structure and a fine needle-like structure was observed for the coatings containing Si, (c-1) and (d-1). From the AES depth profiles of sample (b-1) and (c-1) shown in Fig. 7(a) and
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Fig. 6. SEM images of various nitride coatings as deposited, (a) – (d) and after the oxidation test at 1000 -C for 30 min in air (a-1) – (d-1).
(b), we can estimate that grain-like and needle-like structure was corresponding to the formation of Ti- and Al-rich oxide layer. Superior oxidation resistance of
(Ti,Al)N coating is reportedly due to the formation of protective Al2O3 layer by the outward diffusion of Al atom [36]. This protective property is lost at higher temperature when TiO2 layer is preferentially formed. Because this TiO2 layer tends to develop vertical cracks possibly due to the large difference in oxide to metal volume ratio (Pilling –Bedworth ratio [37]). This means that at 1000 -C rapid oxidation was taking place in the coatings without Si. On the other hand, Al-rich protective oxide layer was still formed on the surface of the (Ti,Cr,Al,Si)N coatings as evidenced by the AES depth profile, thus demonstrating the superior oxidation resistance of the (Ti,Cr,Al,Si)N coatings. The oxide layer thickness of the (Ti,Cr,Al,Si)N coating, as compared in Fig. 6(a), was nearly 4 times thinner than (Ti,Cr,Al)N coating. From the AES depth profile of (Ti,Cr,Al,Si)N coating, the surface oxide layer has almost same composition with the un-oxidized part and no specific concentration or preferred oxidation was observed, such as concentration of Si in the oxide layer reported by Choi et al. [5]. The role of Si in improving the oxidation resistance should be clarified to develop further and better coating systems. 3.5. Cutting tests
Fig. 7. AES depth profiles of (a) (Ti 0.25 Cr 0.1 Al 0.65 )N and (b) (Ti0.2Cr0.2Al0.55Si0.05)N after the oxidation test.
High-speed dry cutting tests have been conducted against hardened cold-working die steel (AISI D2, HRC 60) using carbide end-mills. After the cutting length of 30 m, cutting edge of the end-mill was observed using SEM and images are shown in Fig. 8(a) –(c). The flank wear of the conventional (Ti,Al)N coating was about 60 Am after 30 m of cutting and also intensive sticking of the work-piece material was observed. In case of (Ti,Cr,Al)N coating, the flank wear was slightly less than (Ti,Al)N coating, it was 40 Am and no sticking was observed. Finally, the flank wear of
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Fig. 8. SEM images of the worn cutting edge of (a) (Ti0.2Cr0.2Al0.55Si0.05)N, (b) (Ti0.25Cr0.1Al0.65)N and (c) (Ti0.5Al0.5)N coated end-mills after the cutting length of 30 m.
(Ti,Cr,Al,Si)N coating was nearly one third of (Ti,Al)N and half of (Ti,Cr,Al)N coating and also no sticking was observed.
4. Summary In this study, (Ti,Cr,Al,Si)N coatings with different Al + Si fractions were deposited by cathodic arc method and their properties were investigated in relation to the Al + Si fraction and the substrate bias as an influencing deposition parameter. Deposited (Ti,Cr,Al,Si)N coatings had slightly less Al and enriched Cr fractions compared to target compositions and the compositional difference between coatings and target became larger as the substrate bias was increased. (Ti,Cr,Al,Si)N coatings with the Al + Si fraction of 0.6 had cubic B1 structure independent of the substrate bias. But they had hexagonal B4 structure when the Al + Si fraction was 0.65 and the substrate bias was 20 V. When the Al + Si fraction was 0.6, the grain size decreased as the substrate bias was decreased from 14 to 8 nm. The grain size was almost constant value of 8 nm, however, for the coatings with the Al + Si fraction of 0.65 independent of the substrate bias. The grain size was also observed by TEM and it agreed with the results of the XRD. The position change of (111) peak against the substrate bias suggested that the compressive stress linearly increased as the substrate bias was increased up to the substrate bias of 70 V and it stayed constant for further
increase of the substrate bias. The indentation hardness and elastic modulus was lower than the previously reported (Ti,Cr,Al,Si)N coatings. In H –E* relationship, however, (Ti,Cr,Al,Si)N coatings tended to have lower E* compared to (Ti,Cr,Al)N coatings and this suggested that (Ti,Cr,Al,Si)N coatings, having lower H 3/E*2 values, were likely more resistant to plastic deformation. The oxidation resistance of (Ti,Cr,Al,Si)N coating was much higher than (Ti,Cr,Al)N coating that was evidenced by the fact that oxide layer thickness was nearly 4 times thinner than (Ti,Cr,Al)N coating. No concentration of specific element was observed in the oxide layer and this left the mechanism of high oxidation resistance of (Ti,Cr,Al)N coating issue of future investigation. Finally, high-speed dry cutting tests demonstrated that (Ti,Cr,Al,Si)N coating was quite suitable for this purpose.
References [1] Y. Tanaka, T.M. Gu¨r, M. Kelly, S.B. Hagstrom, T. Ikeda, K. Wakihira, H. Satoh, J. Vac. Sci. Technol., A 10 (4) (1992) 1749. [2] H. Ichimura, A. Kawana, J. Mater. Res. 8 (1993) 1093. [3] T. Ikeda, H. Sato, Thin Solid Films 195 (1990) 99. [4] W.-D. Mu¨nz, J. Vac. Sci. Technol., A 4 (6) (1986) 2717. [5] J.B. Choi, K. Cho, M.-H. Lee, K.-H. Kim, Thin Solid Films 447 – 448 (2004) 365. [6] S. Zhang, D. Sun, Y. Fu, H. Du, Thin Solid Films 447 – 448 (2004) 462. [7] F. Vaz, L. Rebouta, P. Goudeau, J. Pacaud, H. Garem, J.P. Riviere, A. Cavaleiro, E. Alves, Surf. Coat. Technol. 133 – 134 (2000) 307.
1390
K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383 – 1390
[8] G. Llauro, F. Gourbilleau, R. Hillel, Thin Solid Films 315 (1998) 336. [9] M. Diserens, J. Patschneider, F. Levy, Surf. Coat. Technol. 120 – 121 (1999) 158. [10] F. Vaz, L. Rebouta, B. Almeida, P. Goudeau, J. Pacaud, J.P. Riviere, J. Bessa e Sousa, Surf. Coat. Technol. 120 – 121 (1999) 166. [11] S.H. Kim, J.K. Kim, K.H. Kim, Thin Solid Films 420 – 421 (2002) 360. [12] P.J. Martin, A. Bendavid, Surf. Coat. Technol. 163 – 164 (2003) 245. [13] E. Martinez, R. Sanjines, O. Banakh, F. Levy, Thin Solid Films 447 – 448 (2004) 332. [14] Y. Tanaka, N. Ichimiya, Y. Onishi, Y. Yamada, Surf. Coat. Technol. 146/147 (2001) 215. [15] I.-W. Park, S.R. Choi, J.H. Suh, C.-G. Park, K.H. Kim, Thin Solid Films 447 – 448 (2004) 443. [16] S. Carvalho, F. Vaz, L. Rebouta, D. Schneider, A. Cavaleiro, E. Alves, Surf. Coat. Technol. 142 – 144 (2001) 110. [17] S. Carvalho, L. Rebouta, E. Ribeiro, F. Vaz, M.F. Denannot, J. Pacaud, J.P. Rivie`re, F. Paumierb, R.J. Gaboriaud, E. Alves, Surf. Coat. Technol. 177 – 178 (2004) 369. [18] S. Carvalho, E. Ribeiro, L. Rebouta, C. Tavares, J.P. Mendonc¸a, A. Caetano Monteiro, N.J.M. Carvalho, J.Th.M. De Hosson, A. Cavaleiro, Surf. Coat. Technol. 177 – 178 (2004) 459. [19] M. Parlinska-Wojtan, A. Karimi, T. Cselle, M. Morstein, Surf. Coat. Technol. 177 – 178 (2004) 376. [20] M. Parlinska-Wojtan, A. Karimi, O. Coddet, T. Cselle, Morstein, Surf. Coat. Technol. 188 – 189 (2004) 344. [21] A. Niederhofer, P. Nelsla´dek, H.-D. Ma¨nnling, K. Moto, S. Vepøek, M. Jilek, Surf. Coat. Technol. 120 – 121 (1999) 173.
[22] A. Bendavid, P.J. Martin, H. Takikawa, Thin Solid Films 420 – 421 (2002) 83. [23] C. Louro, Surf. Coat. Technol. 180 (2004) 544. [24] S. Veprˇek, S. Reiprich, Thin Solid Films 268 (1995) 64. [25] S. Veprˇek, S. Reiprich, L. Shizhi, Appl. Phys. Lett. 66 (1995) 2640. [26] R. Cremer, M. Witthaut, D. Neuschu¨tz, Value Addition Metallurgy, 1998, p. 249. [27] H. Hasegawa, A. Kimura, T. Suzuki, Surf. Coat. Technol. 132 (2000) 76. [28] K. Yamamoto, T. Sato, K. Takahara, K. Hanaguri, Surf. Coat. Technol. 174 – 175 (2003) 620. [30] K. Yamamoto, T. Sato, M. Takeda, Surf. Coat. Technol. 193 (2005) 167. [31] T. Sawa, K. Tanaka, J. Mater. Res. 16 (2001) 3084. [32] United States Patents, 5270552. [33] K. Sato, N. Ichimiya, A. Kondo, Y. Tanaka, Surf. Coat. Technol. 163 – 164 (2003) 135. [34] K. Yamamoto, S. Kujime, K. Takahara, Effect of element addition on structural and mechanical property of TiAlN based coatings deposited by a hybrid coating system, presented ICMCTF (2005) at San Diego B8-1-4. [35] T.Y. Tsui, G.M. Pharr, W.C. Oliver, C.S. Bhatia, Mater. Soc. Symp. Proc. 383 (1995) 447. [36] D. Mclntyre, J.E. Greene, G. Ha˚kansson, J.-E. Sundgren, W.-D. Mu¨nz, J. Appl. Phys. 67 (1990) 1542. [37] N.B. Pilling, R.E. Bedworth, J. Inst. Met. 29 (1923) 529.
Structural and mechanical property of Si incorporated (Ti,Cr,Al)N coatings deposited by arc ion plating process Kenji Yamamoto a,*, Susumu Kujime b, Kazuki Takahara b b
a Materials Research Lab. Kobe Steel Ltd, 1-5-5 Takatsuka-dai Nishi-ku, Kobe, 651-2271 Hyogo, Japan Advanced Products and Technologies Department, Machinery and Engineering Company, Kobe Steel Ltd, 3-1, 2-chome, Shinhama Arai-cho Takasago, Hyogo 676-8670, Japan
Available online 27 September 2005
Abstract (Ti,Cr,Al,Si)N coatings with different Al + Si fractions (0.6 and 0.65) were deposited by an arc ion plating (AIP) apparatus that is equipped with the plasma-enhanced type arc cathode. The (Ti,Cr,Al,Si)N coatings were deposited under different substrate bias voltages and effect of the deposition parameter on the composition, structure and mechanical properties was investigated. X-ray diffraction measurements of the (Ti,Cr,Al,Si)N coating deposited under different substrate bias voltages revealed that formation of the hexagonal phase (Wurzite structure) was only limited to a relatively low bias voltage range of 20 to 30 V. Above this bias voltage, the crystal structure of the coatings was single-phased cubic rock-salt structure (B1 phase) independent of the bias voltage. Grain size of the coating was calculated from the full width of half maximum (FWHM) of the X-ray diffraction peak and it was smaller than the one of conventional (Ti,Al)N or (Ti,Cr,Al)N coating with a comparable Al fraction. The grain size estimated from the cross-sectional TEM observation was less than 10 nm. From the TEM observation, the coating was compositionally homogeneous and there was no evidence that the film had a phase separation such as Si-rich and -poor region. Hardness of the (Ti,Cr,Al,Si)N coating with Al + Si = 0.6 was in the range of 26 to 27 GPa independent of the substrate bias. (Ti,Cr,Al,Si)N coating with Al + Si = 0.65 showed slight increase in hardness from 24 to 27 GPa when the substrate bias was increased to more than 100 V. To evaluate the oxidation resistance, annealing tests in the air at 1000 -C were conducted and surface SEM observations revealed that surface of the conventional (Ti,Al)N and (Ti,Cr,Al)N was covered with coarse oxide grains enriched with TiO2. Whereas only a dense but very thin protective oxide layer was observed in case of (Ti,Cr,Al,Si)N coating after the oxidation. These coatings were applied to the high-speed dry cutting tests against hardened D2 steel (HRC 60) and the result clearly indicated better performance of (Ti,Cr,Al,Si)N compared to the conventional coatings such as (Ti,Al)N and (Ti,Cr,Al)N. D 2005 Elsevier B.V. All rights reserved. Keywords: Si addition; (Ti,Cr,Al,Si)N; Cathodic arc; Al + Si ratio; Oxidation resistance
1. Introduction High hardness and resistance to oxidation at elevated temperatures have always been important criteria for hard coatings particularly in dry cutting applications [1,2]. This is mainly due to the requirement from the various industries to machine harder work-pieces at higher cutting speeds. This was the main reason in 1990s why TiN coating was replaced
* Corresponding author. Tel.: +81 78 992 5505; fax: +81 78 992 5512. E-mail address: [email protected] (K. Yamamoto). 0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2005.08.025
by (Ti,Al)N coating that had a higher hardness and was more oxidation resistant [3,4]. But of course this was not the end of ever-increasing demands from the industries to improve the oxidation resistance and hardness for better productivity and longer tool life. Recently, much attention has been paid to Si containing coatings such as (Ti,Si)N [5 –12], (Cr,Si)N [13], (Ti,Al,Si)N [14 – 21] and other Si containing coating systems [22,23]. These Si containing coatings have significantly better oxidation resistance compared to the ones without Si [5,14,20,23]. The role of Si in improving the oxidation resistance is not yet fully clarified in many cases. Choi et
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al. reported, however, that preferential oxidation of Si was observed after the oxidation test and this Si-rich oxide film acted as a diffusion barrier for further oxidation [5], much like in case of the preferential oxidation of Al in (Ti,Al)N coating [2]. Also this Si containing (M,Si)N (M: metal) system is known to form so-called ‘‘nano-composite’’ coatings [24,25]. In case of (Ti,Si)N system [6,11,12], authors have reported rather small grain size (less than 10 nm) at the Si content of 10 to 20 at.%. Sometimes small crystalline TiN grains surrounded by amorphous SiNx (aSiNx ) matrix were observed [5,6] and this phase separation was primarily considered as a basic mechanism of the formation of the nano-composite. In the case of (Ti,Al,Si)N system [14,17], addition of Si resulted in refinement of the grain size and formation of a-SiNx phase was observed when coating contained a fairly large amount of Si. (Ti,Al,Si)N coatings [14] containing relatively a small amount of Si (a few atomic percent) consisted of single B1 cubic phase and there was no evidence of formation of second phase such as a-SiNx . Only refinement of the grain was observed in this case. Another interesting aspect of (Ti,Al,Si)N system is change in the crystal structure cubic (B1) to hexagonal (B4) depending on the Al + Si fraction. Tanaka et al. [14] reported crystalline phase composition in (Ti,Al,Si)N system and when Al + Si ratio exceeds 0.61 the crystal structure changed from B1 single phase to B1 + B4 mixed phase. This change in the crystal structure also resulted in loss of the mechanical property where the hardness decreased from at maximum 34 GPa to 25 GPa. In (Ti,Al)N system, there are several reports indicating the phase boundary of B1 and B4 is located around Al ratio of 0.6 to 0.7 [26,27]. Tanaka et al. reported that in case of the coating containing no Si, it had B1 single phase at equivalent Al + Si fraction. This implies that Si has negative effect to maintain the B1 phase, which is usually a harder and a preferred phase for wear resistant applications [28]. In the previous paper [28], we reported properties of (Ti,Cr,Al)N coating system with a high Al fraction of more than 0.65. This (Ti,Cr,Al)N system was characterized by high hardness more than 30 GPa and high oxidation resistance compared to the conventional (Ti,Al)N coating. Motivation of the present work is to improve the properties of (Ti,Cr,Al)N coating further for better tribological performance such as cutting operation. For this purpose, the effect of Si incorporation on the structural and mechanical properties of (Ti,Cr,Al)N coatings was investigated with different Al + Si fractions. Also cutting tests were conducted in comparison with the conventional (Ti,Al)N and (Ti,Cr,Al)N coating.
2. Experimental details (Ti,Cr,Al,Si)N coatings with different Al + Si fractions were deposited by a batch type cathodic arc ion plating
coater equipped with a plasma-enhanced cathode. Details of the deposition equipment and the cathode can be found elsewhere [28]. Two kinds of Ti– Cr –Al – Si targets were prepared by a powder metallurgy process and used for the deposition. Their compositions were Ti0.2Cr0.2Al0.55Si0.05 (target A) and Ti0.15Cr0.2Al0.6Si0.05 (target B). Deposition was conducted in the above-mentioned coater in pure N2 atmosphere at pressure of 2.7 Pa. WC-Co cutting inserts (Mitsubishi Carbide SNGN120408), platinum foils (0.1 mmt) and WC-Co square end-mills (Mitsubishi Carbide 10 mmf, 6 flutes) were used as substrates. Prior to the deposition, these substrates were Ar-ion-etch cleaned for 5 min at Ar pressure of 2.7 Pa. After the cleaning procedure, arc was ignited and deposition was conducted at arc current of 150 A and the substrate temperature was regulated at about 500 -C. The substrate bias was varied from 20 to 150 V to investigate the effect of substrate bias on coating’s properties. The thickness of the coating layer was about 3 Am for all samples unless mentioned. The elemental composition of the coatings was determined using energy dispersive X-ray (EDX) analysis (Horiba EMAX), using ZAF correction and may contain an error approximately 10% at maximum. Crystal structure and preferred orientation of the coating was determined with X-ray diffraction (XRD, Rigaku RINT200-PC) using Cu-ka radiation and the grain size was calculated by the Scherrer’s equation using the full width of half maximum (FWHM) of (111) diffraction peak of the B1 phase [30]. Indentation hardness was measured using a nano-indentation instrument (Elionix ENT-1100) with a Berkovich type diamond indenter. The indenter tip shape correction was conducted using the method proposed by Sawa and Tanaka [31]. No thermal drift correction was used, since this instrument was installed in a thermally regulated chamber in which temperature drift rate is less than 0.1 -C/10 s and one indentation measurement took about 10 s. Microstructure, particularly, to detect the possible existence of separated phase such as a-SiNx , cross-sectional transmission electron microscope observation (TEM: Hitachi) was used. Coatings deposited on WC-Co inserts were cut out and thinned for TEM observation using focused ion beam and micro-lifting technique [32]. Oxidation tests were conducted by annealing the coated platinum foil samples in atmosphere (air) at 1000 -C for 30 min. After the oxidation tests, surface morphology of the samples was investigated by SEM. Depth composition profile of the oxide layer was measured by an Auger electron spectroscopy (AES Perkin Elmer PHI650) using Ar sputter for depth profile measurements. Finally, cutting tests were conducted using hardened AISI D2 cold working die steel (HRC 60) as a work-piece. Cutting parameters were as follows: cutting speed 150 m/min; feed 0.05 mm/flute; axial depth of cut 5 mm; radial depth of cut 0.1 mm; dry cut, air-blow only. Flank wear of the cutting edge was measured after the cutting length of 30 m.
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3. Result and discussion 3.1. Composition and crystal structure Composition of the coatings was slightly different from the targets used for depositions. The coatings deposited from both targets A and B had a few percentage less Al and also a few percentage enriched Cr than the targets, while other elements, Ti and Si, had almost the same composition with the targets. Compositions of (Ti0.2Cr0.23Al0.53Si0.04)N and (Ti0.14Cr0.22Al0.59Si0.05)N were obtained as the resulting compositions of the coatings deposited from targets A and B at the bias voltage of 70 V. This compositional deviation of Al and Cr was more pronounced as the substrate bias was increased. The decrease of Al composition corresponding to the change in the substrate bias can be explained by the preferential sputtering of Al which has a highest sputter yield among the coating’s elements [28,33]. Fig. 1(a) shows X-ray diffraction patterns of (Ti,Cr,Al,Si)N coatings deposited from the targets A and B (hereafter referred as coatings A and B) under various substrate biases. At the substrate bias of 20 V, the diffraction pattern of the coating B contained a weak diffraction peak from the
Fig. 2. Effect of the substrate bias voltage on (a) grain size and (b) position of (111) diffraction peak of (Ti,Cr,Al,Si)N coatings with different Al + Si fractions.
Fig. 1. (a) X-ray diffraction patterns of (Ti,Cr,Al,Si)N coatings with different Al + Si fractions deposited under various substrate biases. Marks (*) denote diffraction peaks from the substrate (WC-Co). (b) Result of deconvolution of overlapping peaks located between the diffraction angle of 32- to 40-.
hexagonal B4 phase. This weak peak, as shown in Fig. 1(b), can be clearly observed by a deconvolution of several overlapping peaks between the diffraction angle from 32- to 40-. This diffraction peak from the B4 phase was only observed for coating B deposited at 20 V, and coatings A and B deposited more than substrate bias of 20 V, only diffraction peaks belonging to the B1 cubic phase were observed. This bias (i.e. ion energy) induced transition in crystal structure was also observed in (Ti,Cr,Al)N system with a high Al fraction [28], yet the nature of this phase transition is not clarified. The preferred orientation of the coating changed from [111] to [100] as the bias was increased for both coatings A and B. Fig. 2 shows the effect of the substrate bias on the (a) grain size and (b) position of (111) peaks. In case of coating A, the grain size was about 12 nm at the bias of 20 V and decreased gradually as the bias was increased. It reached, however, plateau when the bias was increased for more than 50 V and the grain size became a constant value of around 8 nm, whereas the grain size showed nearly constant value of 8 nm for whole bias range in case of coating B. Detailed discussion of the grain size will be given in the Section 3.3. The position of (111) peak was changed by the substrate bias and it shifted toward lower diffraction angle
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Fig. 3. Effect of the substrate bias voltage on (a) indentation hardness and (b) elastic modulus of (Ti,Cr,Al,Si)N coatings.
as the bias was increased. Two factors may influence the position of (111) peak in this case; composition and residual stress. As described at the beginning of this section, the Al fraction decreased and Cr fraction increased as the substrate bias was increased. This change in the composition may induce the change in the lattice parameter thus changing the position of (111) peak. The lattice parameter of CrN and AlN is 0.414 and 0.412 nm and nearly identical, however, and a few atomic percent change of Al and Cr composition can induce the change in diffraction angle of (111) peak smaller than 0.01-. Therefore, the observed change in the (111) peak position was likely due to the change of residual stress that was induced by the substrate bias (ion energy). Usually, PVD-deposited coatings are in compressive stress and an increase in compressive stress results in peak shift to smaller diffraction angle. From this consideration, if we assume the absolute value of the coating’s stress to be compressive, the compressive stress increased as the substrate bias was increased up to 50 V. More than this substrate bias, the stress of the coating showed little change.
Indentation hardness and elastic modulus of coating A showed very little change over the whole bias range and they were about 25 and 400 GPa. In case of coating B, both indentation hardness and elastic modulus increased gradually as the bias was increased. This change in indentation hardness of coating B probably relates to the bias induced crystal structure change observed by the X-ray diffraction measurements. The maximum indentation hardness of coating B was approximately 27 GPa at highest substrate bias of 150 V. These determined indentation hardness and elastic modulus of coatings A and B were both lower than the values of (Ti,Cr,Al)N coatings [28]. In case of (Ti,Cr,Al)N coating, hardness was significantly low, less than 20 GPa, at a lower substrate bias when the coating was composed of mixture of B1 and B4 phase. However, the maximum hardness of approximately 35 GPa was obtained at the bias of 150 V. Incorporation of Si to the (Ti,Cr,Al)N resulted in lower hardness and elastic modulus. However, on the positive side these mechanical properties were more insensitive to the deposition parameter and this is favorable from the viewpoint of robustness of the deposition process. Additionally, Tsui et al. [35] proposed that H 3/E*2 gave information on the resistance of the material to plastic deformation, where H is the indentation hardness and E* is reduced modulus E/(1 m 2). Coatings with high H 3/E*2 values are less likely undergo plastic deformation under external force. Fig. 4 shows E* – indentation hardness relationship for (Ti,Cr,Al,Si)N coatings and (Ti,Cr,Al)N deposited under various substrate biases. Maximum hardnesses of both (Ti,Cr,Al,Si)N coatings were lower than the ones of (Ti,Cr,Al)N coating. When compared at same E*, however, (Ti,Cr,Al,Si)N coatings had higher indentation hardness, thus higher H 3/E*2 value. 3.3. TEM observation Fig. 5 show (a) a high-resolution bright field TEM image of (Ti,Cr,Al,Si)N (coating A), (b) an electron diffraction (ED) pattern with an electron beam diameter
3.2. Mechanical property Fig. 3 shows change in (a) indentation hardness and (b) elastic modulus of coatings deposited under various biases.
Fig. 4. Indentation hardness – reduced modulus (E*) relationship of (Ti,Cr,Al,Si)N and (Ti,Cr,Al)N coatings.
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Fig. 5. (a) High-resolution TEM image of (Ti,Cr,Al,Si)N coating (Al + Si = 0.6), (b) ED pattern with electron beam diameter of approximately 1000 nm and (c) nano-ED pattern of area c.
of about 1000 nm, (c) an ED pattern of region c in the TEM image. Coating A was deposited at the bias of 50 V and substrate temperature of 500 -C. The TEM image showed nano-crystalline nature of the (Ti,Cr,Al,Si)N coating. The grains were sized less than 10 nm and most of them were as small as 5 nm, whereas the grain size of the (Ti0.25Cr0.1Al0.65)N was about 12 nm [34]. The grain sizes of (Ti,Si)N and (Ti,Al,Si)N coatings were reported by several authors using mainly TEM [6,14,17,19]. Tanaka et al. reported that the grain size of (Ti0.41Al0.59)N was approximately 120 to 350 nm, whereas that of (Ti0.42Al0.58Si0.03)N was 50 to 250 nm [14]. ParlinskaWojtan et al. [19] reported the effect of Al + Si content on the grain size of (Ti,Al,Si)N coating with Al + Si content ranging from 10 to about 50 at.% with comparable Si content of about 4 to 6 at.% [14]. The reported grain size showed strong correlation with the Al + Si content and it was less than 10 nm when Al + Si content was more than 40 at.%. These two reports are quite controversial in the absolute grain size, but they agree with each other on the fact that Si has an effect to reduce the grain size. The ED
pattern with fairly large beam diameter (Fig. 5(b)) showed that film only consisted cubic B1 phase and no other phase like hexagonal B4 phase was observed. Additional ED patterns were taken using nano-electron beam with beam diameter was about 1 to 2 nm. Fig. 5(c) shows a typical nano-ED pattern of a single crystal grain. Again only diffraction spots corresponding to the B1 phase were confirmed. 3.4. Oxidation resistance evaluation Fig. 6 shows surface SEM images of different coatings (a) –(d) as deposited and (a-1) –(d-1) after annealed in air at 1000 -C for 30 min. Some macro-particles (MPs) were observed on the surface of the as-deposited coatings. The number of MPs was less for the coatings without Si. After the oxidation tests, surface of coatings without Si (a-1), (b-1) showed a coarse grain-like structure and a fine needle-like structure was observed for the coatings containing Si, (c-1) and (d-1). From the AES depth profiles of sample (b-1) and (c-1) shown in Fig. 7(a) and
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Fig. 6. SEM images of various nitride coatings as deposited, (a) – (d) and after the oxidation test at 1000 -C for 30 min in air (a-1) – (d-1).
(b), we can estimate that grain-like and needle-like structure was corresponding to the formation of Ti- and Al-rich oxide layer. Superior oxidation resistance of
(Ti,Al)N coating is reportedly due to the formation of protective Al2O3 layer by the outward diffusion of Al atom [36]. This protective property is lost at higher temperature when TiO2 layer is preferentially formed. Because this TiO2 layer tends to develop vertical cracks possibly due to the large difference in oxide to metal volume ratio (Pilling –Bedworth ratio [37]). This means that at 1000 -C rapid oxidation was taking place in the coatings without Si. On the other hand, Al-rich protective oxide layer was still formed on the surface of the (Ti,Cr,Al,Si)N coatings as evidenced by the AES depth profile, thus demonstrating the superior oxidation resistance of the (Ti,Cr,Al,Si)N coatings. The oxide layer thickness of the (Ti,Cr,Al,Si)N coating, as compared in Fig. 6(a), was nearly 4 times thinner than (Ti,Cr,Al)N coating. From the AES depth profile of (Ti,Cr,Al,Si)N coating, the surface oxide layer has almost same composition with the un-oxidized part and no specific concentration or preferred oxidation was observed, such as concentration of Si in the oxide layer reported by Choi et al. [5]. The role of Si in improving the oxidation resistance should be clarified to develop further and better coating systems. 3.5. Cutting tests
Fig. 7. AES depth profiles of (a) (Ti 0.25 Cr 0.1 Al 0.65 )N and (b) (Ti0.2Cr0.2Al0.55Si0.05)N after the oxidation test.
High-speed dry cutting tests have been conducted against hardened cold-working die steel (AISI D2, HRC 60) using carbide end-mills. After the cutting length of 30 m, cutting edge of the end-mill was observed using SEM and images are shown in Fig. 8(a) –(c). The flank wear of the conventional (Ti,Al)N coating was about 60 Am after 30 m of cutting and also intensive sticking of the work-piece material was observed. In case of (Ti,Cr,Al)N coating, the flank wear was slightly less than (Ti,Al)N coating, it was 40 Am and no sticking was observed. Finally, the flank wear of
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Fig. 8. SEM images of the worn cutting edge of (a) (Ti0.2Cr0.2Al0.55Si0.05)N, (b) (Ti0.25Cr0.1Al0.65)N and (c) (Ti0.5Al0.5)N coated end-mills after the cutting length of 30 m.
(Ti,Cr,Al,Si)N coating was nearly one third of (Ti,Al)N and half of (Ti,Cr,Al)N coating and also no sticking was observed.
4. Summary In this study, (Ti,Cr,Al,Si)N coatings with different Al + Si fractions were deposited by cathodic arc method and their properties were investigated in relation to the Al + Si fraction and the substrate bias as an influencing deposition parameter. Deposited (Ti,Cr,Al,Si)N coatings had slightly less Al and enriched Cr fractions compared to target compositions and the compositional difference between coatings and target became larger as the substrate bias was increased. (Ti,Cr,Al,Si)N coatings with the Al + Si fraction of 0.6 had cubic B1 structure independent of the substrate bias. But they had hexagonal B4 structure when the Al + Si fraction was 0.65 and the substrate bias was 20 V. When the Al + Si fraction was 0.6, the grain size decreased as the substrate bias was decreased from 14 to 8 nm. The grain size was almost constant value of 8 nm, however, for the coatings with the Al + Si fraction of 0.65 independent of the substrate bias. The grain size was also observed by TEM and it agreed with the results of the XRD. The position change of (111) peak against the substrate bias suggested that the compressive stress linearly increased as the substrate bias was increased up to the substrate bias of 70 V and it stayed constant for further
increase of the substrate bias. The indentation hardness and elastic modulus was lower than the previously reported (Ti,Cr,Al,Si)N coatings. In H –E* relationship, however, (Ti,Cr,Al,Si)N coatings tended to have lower E* compared to (Ti,Cr,Al)N coatings and this suggested that (Ti,Cr,Al,Si)N coatings, having lower H 3/E*2 values, were likely more resistant to plastic deformation. The oxidation resistance of (Ti,Cr,Al,Si)N coating was much higher than (Ti,Cr,Al)N coating that was evidenced by the fact that oxide layer thickness was nearly 4 times thinner than (Ti,Cr,Al)N coating. No concentration of specific element was observed in the oxide layer and this left the mechanism of high oxidation resistance of (Ti,Cr,Al)N coating issue of future investigation. Finally, high-speed dry cutting tests demonstrated that (Ti,Cr,Al,Si)N coating was quite suitable for this purpose.
References [1] Y. Tanaka, T.M. Gu¨r, M. Kelly, S.B. Hagstrom, T. Ikeda, K. Wakihira, H. Satoh, J. Vac. Sci. Technol., A 10 (4) (1992) 1749. [2] H. Ichimura, A. Kawana, J. Mater. Res. 8 (1993) 1093. [3] T. Ikeda, H. Sato, Thin Solid Films 195 (1990) 99. [4] W.-D. Mu¨nz, J. Vac. Sci. Technol., A 4 (6) (1986) 2717. [5] J.B. Choi, K. Cho, M.-H. Lee, K.-H. Kim, Thin Solid Films 447 – 448 (2004) 365. [6] S. Zhang, D. Sun, Y. Fu, H. Du, Thin Solid Films 447 – 448 (2004) 462. [7] F. Vaz, L. Rebouta, P. Goudeau, J. Pacaud, H. Garem, J.P. Riviere, A. Cavaleiro, E. Alves, Surf. Coat. Technol. 133 – 134 (2000) 307.
1390
K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383 – 1390
[8] G. Llauro, F. Gourbilleau, R. Hillel, Thin Solid Films 315 (1998) 336. [9] M. Diserens, J. Patschneider, F. Levy, Surf. Coat. Technol. 120 – 121 (1999) 158. [10] F. Vaz, L. Rebouta, B. Almeida, P. Goudeau, J. Pacaud, J.P. Riviere, J. Bessa e Sousa, Surf. Coat. Technol. 120 – 121 (1999) 166. [11] S.H. Kim, J.K. Kim, K.H. Kim, Thin Solid Films 420 – 421 (2002) 360. [12] P.J. Martin, A. Bendavid, Surf. Coat. Technol. 163 – 164 (2003) 245. [13] E. Martinez, R. Sanjines, O. Banakh, F. Levy, Thin Solid Films 447 – 448 (2004) 332. [14] Y. Tanaka, N. Ichimiya, Y. Onishi, Y. Yamada, Surf. Coat. Technol. 146/147 (2001) 215. [15] I.-W. Park, S.R. Choi, J.H. Suh, C.-G. Park, K.H. Kim, Thin Solid Films 447 – 448 (2004) 443. [16] S. Carvalho, F. Vaz, L. Rebouta, D. Schneider, A. Cavaleiro, E. Alves, Surf. Coat. Technol. 142 – 144 (2001) 110. [17] S. Carvalho, L. Rebouta, E. Ribeiro, F. Vaz, M.F. Denannot, J. Pacaud, J.P. Rivie`re, F. Paumierb, R.J. Gaboriaud, E. Alves, Surf. Coat. Technol. 177 – 178 (2004) 369. [18] S. Carvalho, E. Ribeiro, L. Rebouta, C. Tavares, J.P. Mendonc¸a, A. Caetano Monteiro, N.J.M. Carvalho, J.Th.M. De Hosson, A. Cavaleiro, Surf. Coat. Technol. 177 – 178 (2004) 459. [19] M. Parlinska-Wojtan, A. Karimi, T. Cselle, M. Morstein, Surf. Coat. Technol. 177 – 178 (2004) 376. [20] M. Parlinska-Wojtan, A. Karimi, O. Coddet, T. Cselle, Morstein, Surf. Coat. Technol. 188 – 189 (2004) 344. [21] A. Niederhofer, P. Nelsla´dek, H.-D. Ma¨nnling, K. Moto, S. Vepøek, M. Jilek, Surf. Coat. Technol. 120 – 121 (1999) 173.
[22] A. Bendavid, P.J. Martin, H. Takikawa, Thin Solid Films 420 – 421 (2002) 83. [23] C. Louro, Surf. Coat. Technol. 180 (2004) 544. [24] S. Veprˇek, S. Reiprich, Thin Solid Films 268 (1995) 64. [25] S. Veprˇek, S. Reiprich, L. Shizhi, Appl. Phys. Lett. 66 (1995) 2640. [26] R. Cremer, M. Witthaut, D. Neuschu¨tz, Value Addition Metallurgy, 1998, p. 249. [27] H. Hasegawa, A. Kimura, T. Suzuki, Surf. Coat. Technol. 132 (2000) 76. [28] K. Yamamoto, T. Sato, K. Takahara, K. Hanaguri, Surf. Coat. Technol. 174 – 175 (2003) 620. [30] K. Yamamoto, T. Sato, M. Takeda, Surf. Coat. Technol. 193 (2005) 167. [31] T. Sawa, K. Tanaka, J. Mater. Res. 16 (2001) 3084. [32] United States Patents, 5270552. [33] K. Sato, N. Ichimiya, A. Kondo, Y. Tanaka, Surf. Coat. Technol. 163 – 164 (2003) 135. [34] K. Yamamoto, S. Kujime, K. Takahara, Effect of element addition on structural and mechanical property of TiAlN based coatings deposited by a hybrid coating system, presented ICMCTF (2005) at San Diego B8-1-4. [35] T.Y. Tsui, G.M. Pharr, W.C. Oliver, C.S. Bhatia, Mater. Soc. Symp. Proc. 383 (1995) 447. [36] D. Mclntyre, J.E. Greene, G. Ha˚kansson, J.-E. Sundgren, W.-D. Mu¨nz, J. Appl. Phys. 67 (1990) 1542. [37] N.B. Pilling, R.E. Bedworth, J. Inst. Met. 29 (1923) 529.