www.elsevier.com/locate/surfcoat

Surface & Coatings Technolog

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

a Materials Research Lab. Kobe Steel Ltd, 1-5-5 Takatsuka-dai Nishi-ku, Kobe, 651-2271 Hyogo, Japanb 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

0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.surfcoat.2005.08.025

* Corresponding author. Tel.: +81 78 992 5505; fax: +81 78 992 5512.

E-mail address: kenji-y@kobelco.jp (K. Yamamoto).

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

y 200 (2005) 1383 – 1390

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–13901384

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 (a-

SiNx) 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 inves-

tigated 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 RINT-

200-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.

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–1390 1385

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. 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 de-

convolution of overlapping peaks located between the diffraction angle of

32- to 40-.

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.

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- to40-. 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 Vand 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

Fig. 3. Effect of the substrate bias voltage on (a) indentation hardness and

(b) elastic modulus of (Ti,Cr,Al,Si)N coatings.

Fig. 4. Indentation hardness – reduced modulus (E*) relationship of

(Ti,Cr,Al,Si)N and (Ti,Cr,Al)N coatings.

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–13901386

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-. There-fore, 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 compres-

sive, 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.

3.2. Mechanical property

Fig. 3 shows change in (a) indentation hardness and (b)

elastic modulus of coatings deposited under various biases.

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 gradu-

ally 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 H3/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�m2). Coatings with high H3/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 hard-

nesses 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 H3/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

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.

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–1390 1387

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 coat-

ing. 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]. Parlinska-

Wojtan 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

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).

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–13901388

(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

Fig. 7. AES depth profiles of (a) (Ti0.25Cr0.1Al0.65)N and (b)

(Ti0.2Cr0.2Al0.55Si0.05)N after the oxidation test.

(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 resist-

ance 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

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 conven-

tional (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

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.

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–1390 1389

(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 H3/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. Gur, 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. Munz, 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.

K. Yamamoto et al. / Surface & Coatings Technology 200 (2005) 1383–13901390

[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. Riviere, 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. Mendonca, A.

Caetano Monteiro, N.J.M. Carvalho, J.Th.M. De Hosson, A. Cav-

aleiro, 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. Nelsladek, H.-D. Mannling, 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. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64.

[25] S. Veprek, S. Reiprich, L. Shizhi, Appl. Phys. Lett. 66 (1995) 2640.

[26] R. Cremer, M. Witthaut, D. Neuschutz, 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. Hakansson, J.-E. Sundgren, W.-D. Munz,

J. Appl. Phys. 67 (1990) 1542.

[37] N.B. Pilling, R.E. Bedworth, J. Inst. Met. 29 (1923) 529.