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Thin Solid Films 493

Constitution and microstructure of magnetron sputtered nanocomposite

coatings in the system Ti–Al–N–C

M. Stueber a,*, P.B. Barna b, M.C. Simmonds c, U. Albers a, H. Leiste a, C. Ziebert a,

H. Holleck a, A. Kovacs b, P. Hovsepian c, I. Gee c

a Forschungszentrum Karlsruhe, Institute of Materials Research I, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germanyb Research Institute for Technical Physics and Materials Science, Hungarian Academy of Sciences, Konkoly Thege St. 29-33, H-1121 Budapest, Hungary

c Sheffield Hallam University, Materials Research Institute, City Campus, Howard Street, S1 1WB Sheffield, United Kingdom

Received 18 August 2004; received in revised form 30 June 2005; accepted 25 July 2005

Available online 26 August 2005

Abstract

An ambitious objective in the development of self-lubricating wear-resistant coatings is to make use of lubricious phases such as graphite,

amorphous carbon or MoS2 incorporated into coatings. A series of (Ti,Al)(N,C) coatings with different carbon contents (0–28 at.%) have

been deposited by reactive magnetron sputtering of TiAl in a mixture of Ar, N2 and CH4 gases. The microstructure and constitution of these

coatings have been investigated using electron microprobe analysis, atomic force microscopy, X-Ray photoelectron spectroscopy, cross

sectional transmission electron microscopy, Raman spectroscopy, X-ray diffraction and pole figure analyses. Starting from a pure TiAlN

coating significant changes in the microstructure of the coatings were observed dependent on the carbon concentration. Under optimum

conditions nanocomposite coatings with a structure of a coexisting metastable hard, nanocrystalline fcc TiAlNC phase and an amorphous

carbon phase (a-C) have been deposited. The localization of an amorphous carbon phase has been shown by high-resolution transmission

electron microscopy.

D 2005 Elsevier B.V. All rights reserved.

Keywords: Nanocomposite; Sputtering; Transmission electron microscopy; X-Ray photoelectron spectroscopy

1. Introduction

High-speed manufacturing processes and tribologically

stressed components in many applications have a demand

for innovative materials and surfaces providing complex

multifunctional property profiles covering wear resistance,

oxidation and corrosion resistance and low friction. Nano-

structured and nanocomposite materials and thin films are

promising concepts for the engineering design and realisa-

tion of such property profiles [1–7]. In the field of surface

science and technology the combination of a nanocrystal-

line, hard and wear resistant phase and a lubricious phase

such as amorphous carbon (a-C) or MoS2, coexisting in a

single-layer coating has been reported many times in the

literature [8–17]. The application of crystalline binary hard

0040-6090/$ - see front matter D 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.tsf.2005.07.290

* Corresponding author. Tel.: +49 7247 82 3889; fax: +49 7247 82 4567.

E-mail address: michael.stueber@imf.fzk.de (M. Stueber).

phases, either thermodynamically stable or metastable such

as WC1�x [17–22], TiC1�x [11,23–28], TaC1�x [29], or

TiB2�x [30] in nanocomposites together with a-C are mostly

state-of-the-art and some coatings have been commercially

available for many years. Nanocomposite coatings combin-

ing ternary nanocrystalline metastable hard phases and a-C

have not yet been reported in detail.

The material system Ti–Al–N–C is a good candidate for

examining the synthesis of such coatings as it offers well-

known metastable hard phases such as TiAlN and TiAlNC

which are the basis for the presently most frequently used

coatings in the tool industry. With (Ti,Al)N coatings for

example, a higher oxidation resistance up to 750 -C, anincreased hardness and a better cutting performance

compared to pure TiN films was reported first by Munz

and Knotek et al. [31,32]. The development and optimisa-

tion of (Ti,Al)N based thin films since then is described in

the excellent review article of PalDey and Deevi [33].

(2005) 104 – 112

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112 105

Detailed information on the constitution, microstructure and

thermodynamic stability of phases in the Ti–Al–N system

as well as the properties and performance of Ti–Al–N thin

films prepared by various physical vapour deposition (PVD)

methods such as magnetron sputtering, ion beam assisted

deposition, arc evaporation, arc ion plating or chemical

vapour deposition (CVD) methods such as plasma enhanced

CVD are given in the literature [34–46].

Zhang et al. prepared (Ti,Cr)CN nanocrystalline films by

magnetron sputtering [47]. Shieh and Hon [48] reported on

chemical vapour deposited nanocomposite coatings of

(Ti,Al)(C,N), while Lackner et al. investigated the chemical,

microstructural and tribological properties of (Ti,Al)(C,N)

coatings deposited by pulsed laser deposition [49].

In the present work, a series of (Ti,Al)(N,C) coatings

with different carbon contents (0–28 at.%) have been

deposited by reactive magnetron sputtering of TiAl in a

mixture of Ar, N2 and CH4 gases. Starting from a pure

(Ti,Al)N coating significant changes in the microstructure of

the coatings were observed dependent on the carbon

concentration. Under optimum conditions nanocomposite

coatings with a structure of a coexisting metastable hard,

nanocrystalline fcc TiAlNC phase and an a-C phase have

been deposited. This paper presents and discusses the

preparation, growth, constitution and microstructure of

these coatings.

2. Experimental details

2.1. Deposition method

(Ti,Al)(N,C) single layer coatings of various micro-

structure and constitution have been prepared by reactive

magnetron sputtering of a TiAl target in a mixture of Ar, N2

and CH4 gases. The equipment used was a Hauzer HTC 625

machine. The target applied was made of a commercially

available TiAl alloy composed of 50 mol% Ti and 50 mol%

Al. The target size was 400�125 mm. All deposition

experiments have been carried out at a target power of 6 kW.

Commercially available hard metal tool inserts have been

used as substrates for all samples apart from those for the

pole figure measurements, where thermally oxidized Si

single crystals have been used to avoid substrate interfe-

rence effects. Prior to deposition the substrate materials have

been ultrasonically cleaned in acetone and isopropanol for

15 min and have been plasma-etched by a dc glowing

discharge in argon for 90 min. The SiO2 surface layer of the

Si single crystals has been completely etched away during

the plasma etching. During deposition the substrate temper-

ature was kept constant at 400 -C for all films and a constant

bias voltage of �80 V dc was applied to the substrates. The

argon gas flow was kept constant at 200 sccm while the

nitrogen gas flow was kept constant at 32 sccm, giving a

total pressure of 0.34 Pa. The CH4 gas flow was varied

systematically between 0 and 30 sccm.

2.2. Characterisation of microstructure, constitution and

properties

The elemental composition of the (Ti,Al)(N,C) coatings

was characterised by electron microprobe analysis (EMPA)

using a Cameca microbeam system.

The microstructure of the coatings was investigated by

cross sectional transmission electron microscopy (XTEM,

Philips CM20 and JEOL 3010 UHR) as well as by selected

area electron diffraction (SAED). Cross sectional specimens

were prepared in a Technoorg Linda IV/H/L ion beam

thinning apparatus. Ion beam thinning was completed with

low angle and low energy ion beam (<200 eV) according to

the method reported by Barna and Barna and Pecz [50,51].

The diffraction patterns were evaluated with the Process

Diffraction software tool developed by Labar [52,53].

X-ray photoelectron spectroscopy (XPS) analysis was

performed on a Kratos Axis 165 XPS spectrometer using

monochromatic Al Ka (1486.6 eV) radiation in a vacuum of

around 5�10�9 mbar. Survey scans were performed at a

pass energy of 80 eV, high resolution scans of 20 eV; in this

configuration the FWHM peak width of the Ag 3d 5/2 peak

is around 0.6 eV. The energy analyser was calibrated using

the Ag 3d and Auger lines prior to use and the Ag 3d peak

was monitored during the XPS measurements to control

energy shifts due to electrostatic charging [54].

Raman spectroscopy was performed using a Renishaw

Raman System 2000 spectrometer. The holographic notch

filters allowed a lower spectral limit of ¨150 cm�1. The

excitation wavelength used was 632.8 nm (25 mW power)

from a HeNe laser. A maximum throughput of 4–5 mW was

possible at the sample in a 1–2 Am spot through at standard

50� microscope objective. The penetration depth into these

PVD coatings is low, approximately 100–200 nm, the main

controlling factors are: reflectivity of the surface, density,

colouration, and topography.

Glancing angle X-ray diffraction (GAXRD) was per-

formed on a Phillips PW1710 X-ray diffractometer. Typi-

cally, an incidence angle of 5 degrees was used with Cu Ka

radiation (k=0.15418 nm). X-ray pole figures were

obtained on a Phillips X’Pert X-ray diffractometer again

using Cu Ka radiation.

Mechanical properties and the performance of the coa-

tings on tool steel substrates have been characterised by

measuring the microhardness applying the Vickers method

(0.5 N load) and the critical load of failure in the scratch test

(CSEM Revetest). The coating thickness was determined by

surface profilometry.

3. Results and discussion

The chemical composition and the constitution of the

deposited (Ti,Al)(N,C) films are given in Table 1. The argon

contamination of the coatings is not specified in this table as

it is below 0.1 at.% for all films, as measured by EMPA. The

Table 1

Chemical composition and constitution of (Ti,Al)(N,C) coatings deposited

by reactive dc magnetron sputtering from a TiAl alloy target under variation

of the methane gas flow

Sample

no.

Coating type and

microstructure

Elemental composition (at.%)

Ti Al N C O

1 nc fcc TiAlN 26.7 26.5 45.0 0.7 1.0

2 nc fcc TiAlNC 21.7 26.5 45.3 5.5 0.9

3 nc fcc TiAlNC+a-C 20.6 24.1 46.4 8.1 0.8

4 nc fcc TiAlNC+a-C 21.2 23.0 45.5 9.5 0.7

5 nc fcc TiAlNC+a-C 20.4 22.6 42.7 13.5 0.7

6 nc fcc TiAlNC+a-C 19.5 21.8 41.4 16.3 1.1

7 nc fcc TiAlNC+a-C 17.0 19.0 34.9 28.0 0.9

All other deposition parameters were kept constant.

36 37 38 39 40

(Rel

ativ

e in

tens

ity (

arb.

uni

ts)

Diffraction angle, 2θ (degrees)

0 at.% C

5.5 at.% C8.1 at.% C

9.5 at.% C

28.0 at.% C

Fig. 1. GAXRD diffractograms of reactively sputtered (Ti,Al)(N,C) films

showing the shift of the (111) peaks of TiAlN like structure with increasing

carbon content. The dashed lines are an aid to the eye to show the shift in

peak position.

8.1 at.% C

nten

sity

(ar

b. u

nits

)

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112106

Ti / (Ti+Al) and the non-metal /metal ratio are given in

Table 2. It is clear that the Ti / (Ti+Al) ratio is nearly 0.5 and

remains constant for all coatings independent on the non-

metal /metal ratio. In the case of the nearly carbon-free

(Ti,Al)N coatings the non-metal /metal ratio is 0.85 while

with increasing carbon concentration the overall non-metal /

metal (N+C) / (Ti+Al) ratio increases significantly. This

behaviour is clearly related to the increasing carbon

concentration caused by the increasing methane gas flow.

The N/ (Ti+Al) ratio is between 0.85 and 1.04.

Glancing angle X-ray diffractograms are shown in Figs.

1 and 2. GAXRD differs from conventional Bragg–

Brentano XRD in that in the Bragg–Brentano geometry

the diffraction vector is normal to the surface. For a glancing

angle incidence of 5-, as used in this study, the diffraction

vector is approximately 13- and 15- off the surface normal

for the (111) and (200) orientations, respectively. GAXRD

was used in preference to Bragg–Brentano XRD in this

study to avoid strong substrate interference effects. Figs. 1

and 2 and analysis of diffractograms taken over a large

range of diffraction angles indicates the presence of the

cubic B1-NaCl structure only in all coatings. From Fig. 1 it

can be seen that the (111) peak position moves to a lower

diffraction angle indicating an expansion in the d-spacing

with increasing carbon content. The change in the d-spacing

with increasing carbon content is shown in Fig. 3a. For pure

(Ti,Al)N coatings the lattice spacing corresponding to the

(111) reflection was calculated to be d111=0.2402 nm which

shows reasonable agreement with values for TiAlN in the

literature [55]. With increasing carbon content up to around

Table 2

Ti / (Ti+Al) and non-metal /metal atomic concentration ratios of magnetron-

sputtered (Ti,Al)(N,C) coatings

Sample no. Ti / (Ti+Al) N/ (Ti+Al) C/ (Ti+Al) (N+C) / (Ti+Al)

1 0.50 0.85 0.01 0.86

2 0.45 0.94 0.11 1.05

3 0.46 1.04 0.18 1.22

4 0.48 1.03 0.22 1.25

5 0.47 0.99 0.32 1.31

6 0.47 1.00 0.39 1.40

7 0.47 0.97 0.78 1.75

9.5 at.% C the lattice spacing increases approximately

linearly to 0.2434 nm. Above 9.5% the d-spacing is nearly

constant and only increases slightly. The film with the

highest carbon concentration of 28 at.% has a value of

d111=0.2437 nm. A similar relationship between d-spacing

and the (200) reflection was also observed (Fig. 3b). The

lattice parameters of fcc-TiC and fcc-TiN are 0.4328 and

0.4238 nm, respectively. The expansion of the lattice that is

observed here may be due to the direct substitution of

carbon into the nitrogen lattice positions, because of the

larger atomic radius of carbon in comparison to nitrogen.

42 43 44 45

5.5 at.% C

9.5 at.% C

Diffraction angle, 2θ (degrees)

Rel

ativ

e i

0 at.% C

28.0 at.% C

Fig. 2. GAXRD diffractograms of reactively sputtered (Ti,Al)(N,C) films

showing the shift of the (200) peaks of the TiAlN like structure with

increasing carbon content. The dashed lines are an aid to the eye to show

the shift in peak position.

0 5 10 15 20 25 30

2.40

2.41

2.42

2.43

2.44

d-sp

acin

g (A

ngst

rom

)

Carbon content (at.%)

0 5 10 15 20 25 30

2.08

2.09

2.10

2.11

d-sp

acin

g (A

ngst

rom

)

Carbon content (at.%)

(b)

(a)

Fig. 3. Graph showing the variation in the d-spacing of (a) the (111) planes

and (b) the (200) plane versus the carbon content of the films. The error

bars represent the error estimate on the determination of the d-spacing.

(111) peak (200) peak

5.5

at.%

C

8.1

at.%

C16

.3 a

t.%

C28

.0 a

t.%

C

Fig. 4. X-Ray Diffraction pole figures of the (111) and (200) peaks for the

(Ti,Al)(N,C) coating with 5.5, 8.1, 16.3, and 28.0 at.�% carbon.

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112 107

The threshold at around 9.5 at.% suggest that this

substitution occurs readily at low carbon contents but

saturates at values above this value. Where the excess

carbon is situated above this threshold is not indicated from

the XRD.

Fig. 1 shows a peak broadening of the (111) reflection

with increasing C content. This indicates a decreasing grain

size with increasing carbon content. Further insight into the

growth of these films is shown in the X-ray pole figures of

films prepared on Si substrates (Fig. 4). The pole figures

indicate that the films of C content up to 13.5 at.% C have

well pronounced two-axis texture [56] with (111) texture

axis off to the substrate normal. The (200) orientation pole

figures of this range of C content show three well defined

peaks with a tilt angle of approximately 55 - with respect to

the (111) peak maximum. These peaks are related to the

(100) planes of the (111) oriented crystals and clearly prove

the crystals are aligned with respect to the surface. This

alignment can be related to the orientation effect of the ion-

sputtered Si single crystal substrates. In a recent study

Kovacs et al. have shown that the (Ti,Al)(N,C) coatings

with 5.5 at.% grow epitaxially on hard metal substrate [57].

However the substrate interference effects did not allow to

record pole figures on these samples grown on hard metal.

The pole figures of the coatings with 16.3 and 28 at.% C

indicate the decreasing texturing with increasing C content.

No (200) texture could be detected in any of the coatings

investigated. Evidence for the existence of the two-axis

texture can also be seen in the AFM image of a coating with

5.5 at.% C (Fig. 5). This figure is representative for the

whole surface of this sample and clearly illustrates the three-

fold symmetry of (100) planes.

In TEM characterisation the strong correlation between

the carbon content, grain size and morphology can be seen

in more detail (Figs. 6–8). All coatings show a columnar

morphology of structure zone T according to the structure

zone models of Thornton and of Barna and Adamik [58,59].

The columnar morphology of carbon-free (Ti,Al)N thin

coatings agrees well to similar findings documented in the

literature [37,60]. The columns are single crystals and are

separated by tight grain boundaries. Doping the (Ti,Al)N

coating with C very pronounced grain boundaries show up

Fig. 7. Cross sectional TEM images of a (Ti,Al)(N,C) coating with a C

content of 13.5 at.%: a) bright field image with corresponding SAED

pattern, b) dark field image taken by (200) reflection of fcc TiAlNC phase,

c) HRTEM image: the arrows indicate the amorphous carbon cluster

inclusions at the grain boundaries.

Fig. 5. AFM image of a (Ti,Al)(N,C) coating with a C content of 5.5 at.%.

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112108

on the XTEM images at lower C concentrations <9.5 at.%

(Fig. 6). The columns are single crystals and the surface

topography corresponds to the topography of the points of

individual columns. In the range of 5.5 to 13.5 at.% C the

diameter of columns is decreasing with increasing carbon

concentration (for example, the average column diameter

was 440 nm in the case of 5.5 at.% C and 300 nm,

respectively in the case of 8.1 at.% C). The columns are

separated by pronounced grain boundaries containing an

amorphous phase. In the range of C concentration of 9.5–

16.3 at.% the columnar structure transforms into a dense

fibre-like structure (Fig. 7a). Dark field images (Fig. 7b)

indicate that the fibres are mostly single crystals and no

pronounced grain boundaries can be identified at this

resolution. However, in the high-resolution TEM image

(Fig. 7c) amorphous carbon cluster inclusions of 1–2 nm

size at grain boundaries can be seen. The average column

diameter is further reduced to a minimum value of 52 nm for

coatings with 13.5 at.% C and then increases with increasing

carbon concentration. The structure of coatings with high C

concentration (C>20 at.%) is homogenous with nano-

crystalline structure in the near substrate region and trans-

Fig. 6. XTEM images of a (Ti,Al)(N,C) coating with a C content of 5.5

at.%: a) bright field image, b) dark field image, c) selected area diffraction

pattern.

forms into a columnar structure with increasing film

thickness. The columns have large diameter (about 200–

250 nm for films with 28 at.% C) and well defined column

boundaries (Fig. 8). The phase contrast images clearly show

that columns are composed of frond-like lamellae (Fig. 8b).

According to the high-resolution TEM images (Fig. 8c) the

frond-like lamellae are separated by an amorphous phase.

Fig. 8. Cross sectional TEM images of a (Ti,Al)(N,C) coating with a C

content of 28.0 at.%: a) bright field image with corresponding SAED

pattern, b) phase contrast image, c) HRTEM image: the arrows indicate the

amorphous phase between the frond-like lamellae.

0 5 10 15 20 25 30454.0

454.2

454.4

454.6

454.8

455.0

455.2

455.4

455.6

396.2

396.4

396.6

396.8

397.0

397.2

397.4

397.6

397.8

Ti

N p

eak

pos

ition

(eV

)

Ti p

eak

pos

ition

(eV

)

Atomic fraction of carbon (at.%)

N

Fig. 10. Change in the binding energy of the Ti 2p3 / 2 and N 1s peaks of

(Ti,Al)(N,C) films with increasing carbon content after etching with 5 keV

Ar ions. The error bars represent the error estimate on the determination of

the peak positions.

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112 109

The grain size of the TiAlNC phase in this composite

structure is 3–6 nm while the thickness of the frond–

lamellae is about 2–3.5 nm and their length is 40–60 nm.

The columns are tilted to the substrate plane in all coatings

in agreement with the pole figures. The average column

diameter versus the carbon concentration of (Ti,Al)(N,C)

coatings is shown in Fig. 9. In agreement with pole figure

analysis coatings with 28 at.% C show a nearly random

crystal orientation.

The bonding states of the (Ti,Al)(N,C) coatings were

characterised by XPS. A significant change in the binding

energy in terms of a shift of the Ti 2p3 / 2 and N 1s peak

positions to higher energy with increasing carbon concen-

tration was observed (Fig. 10). The Ti 2p3 / 2 peak position

for example shifts from 454.7 eV in the case of pure

(Ti,Al)N films, which is in good agreement with a published

value of 454.8 eV for Ti0.59Al0.41N films [61] to 455.3 eV

for (Ti,Al)(N,C) films with 13.5 at.% carbon and to 455.4

eV for films with 28 at.% carbon. The same position has

been found by Phase and Sathe in TiC0.76N0.43 films

synthesized by nitrogen bombarding of a C/Ti bilayer and

was assigned to Ti–C–N bonding [54]. According to their

results a complex structure, which was found between the Ti

2p3 / 2 Ti 2p1 / 2 doublet (456–459 eV), can be attributed to

the presence of several reduced oxide phases such as

TiNxOy, Ti2O3 and TiO2. Further evidence gave the

increasing intensity of these structures after annealing of

the samples at 630 -C for 120 min. The most significant

shift in the binding energy is seen for carbon concentrations

below 10 at.%. A similar shift to higher binding energies has

recently been reported in the case of tantalum in cubic

tantalum carbonitride [62] and was ascribed to the positive

charge on the tantalum atoms increasing with the sub-

stitution of nitrogen atoms with carbon atoms. The shift for

Ta was much smaller (0.35 eV) over a wider composition

range than observed here (0.7 eV) over a smaller compo-

sition range. This may be because in this study the TiAlN

0 10 20 300

125

250

375

500

Col

umn

diam

eter

(nm

)

Atomic fraction of carbon (at.%)

Fig. 9. Average column diameter of (Ti,Al)(N,C) coatings dependent on the

carbon content. Note that columns are single crystals for carbon

concentrations below 20 at.%. At C>20 at.% the columns are composite

structures.

films are sub-stoichiometric and the carbon atoms can

substitute into nitrogen vacancies rather than substituting

nitrogen atoms, resulting in a more pronounced shift in the

observed binding energy. It is also interesting to note that a

smaller shift is also seen in the nitrogen peak.

The N 1s peak shifts from 396.35 eV in the case of pure

(Ti,Al)N films to 396.7 eV for (Ti,Al)(N,C) films with 13.5

at.% carbon or more. The Al 2p3 / 2 peak showed that the

Al2O3 surface layer, which was present after deposition, was

almost completely removed by sputter etching. Its position

of 74.2 eV after etching, together with the N 1s position

might indicate Al–N bonding as described in Wang and

Chiu [63].

The C 1s peak feature of the samples is difficult to

evaluate as there is adventitious carbon and oxygen

contamination that is difficult to remove by ion etching. A

high resolution scan of the C 1s peak of a (Ti,Al)(N,C)

coating with 28 at.% carbon content after 50 min of ion

etching is shown in Fig. 11. The optimised fit of the peak

feature suggests at least five bands at 281.8, 283.7, 285.5,

278 280 282 284 286 288 290 292 294 2968.0x103

1.0x103

1.2x103

1.4x103

1.6x103

1.8x103

2.0x103

2.2x103

Inte

nsity

(co

unts

)

Binding Energy (eV)

Fig. 11. XPS spectrum showing the C 1s peak for the (Ti,Al)(N,C) coating

with 28.0 at.% carbon after etching with 5 keV Ar ions.

0 5 10 15 20 25 300.0

5.0x104

1.0x105

1.5x105

2.0x105

2.5x105

3.0x105

3.5x105

4.0x105

Inte

nsity

(co

unts

)

Atomic fraction of Carbon (%)

Fig. 13. Integrated peak intensity of the D–G band versus the carbon

content of the films.

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112110

286.5 and 287.9 eV. The signal at 281.8 eV can be attributed

to Ti–C bonding, whereas the signals at 285.5 and 286.5 eV

are typical for C–OH bonding and aliphatic/h-carbons andother adventitious hydrocarbons as sputter etching residues,

respectively [64]. The nature of the other broader bands at

283.7 and 287.9 eV is not absolutely clear but might be

attributable to interfacial carbon at the grain boundaries [64]

and CO species due to recontamination during the XPS

analysis, respectively. Moreover XPS indicates that a two-

phase microstructure dominates in the coatings with high

carbon content.

Additional information on the coating constitution can be

obtained from Raman spectroscopy which is a sensitive

measure for various forms of carbon. The Raman spectra of

the (Ti,Al)(N,C) films are given in Fig. 12. The onset of a

broad band between 1350 and 1600 cm�1 characteristic for

carbon–carbon bonding of a highly disordered amorphous

carbon phase of mixed sp2 and sp3 bonding is visible at

films with higher carbon concentrations. The integrated

peak intensity shows that the D–G bands are seen at carbon

contents of 8.1 at.% and higher and that both bands increase

strongly with increasing carbon content (Fig. 13). This

confirms the existence of a free a-C phase in the

(Ti,Al)(N,C) films at higher C concentrations as indicated

by XTEM and XPS characterisation.

Thickness, critical load of failure in the scratch test and

the Vickers microhardness of the (Ti,Al)(N,C) coatings are

summarised in Table 3.

From the chemical composition of the carbon-free

metastable nanocrystalline fcc TiAlN films a N-deficit is

obvious from the (Ti+Al) /N ratio. Thus the microstructure

evolution of these coatings dependent on the carbon content

might be explained by a five-step growth model: 1) The

growth starts from a sub-stoichiometric (Ti,Al)N1�x phase

with a significant N-deficit. 2) By adding carbon to the

coating these lattice places are filled and a single-phase

nanocrystalline fcc TiAlNC structure is built. 3) Once the N-

vacancies are filled, carbon atoms can substitute regular N

0 1000 2000 3000 4000

28.0 at. % C

13.5 at. % C

5.5 at.% C

Inte

nsity

(ar

b. u

nits

)

Raman Shift (cm-1)

8.1 at.% C

9.5 at.% C

16.3 at. % C

Fig. 12. Raman spectra showing the onset of the D–G band (the dotted

box) characteristic for amorphous carbon with increasing carbon content in

the (Ti,Al)(N,C) coatings.

positions. 4) Exceeding the solubility range of carbon within

the TiAlN lattice, which seems to occur above 8 at.% of

carbon, the carbon starts to build nanoclusters or agglome-

rates at grain boundaries of the (Ti,Al)(N,C) lattice. 5) When

increasing the C content further, the size of carbon clusters

increases and larger domains of amorphous carbon phase or

even an amorphous carbon matrix can be formed on the

growth surface of TiAlNC crystals. The growth of TiAlNC

crystals is interrupted by these C domains decreasing the

grain size of TiAlNC further. This process leads to the

formation of a frond-like structure characterizing the films of

28 at.% C. Thus if the maximum solubility for carbon in the

nanoscale fcc lattice is achieved, the growth of phase-

separated nanocomposite coatings starts. This phase separa-

tion seems to take place during the structure formation and to

follow the thermodynamic rules and kinetics of PVD phase

transitions. The occurrence of coexisting nanocrystalline/

amorphous phases in ternary and quaternary systems was

predicted by Holleck already in 1986 on the basis of

theoretical considerations on the thermodynamics and

kinetics of PVD processes [65] and again by Veprek on the

basis of his extensive experimental and theoretical work on

the Ti–Si–N system [66]. Furthermore our results are in

excellent agreement with a recently published paper by

Lackner et al., who identified very similar growth mecha-

Table 3

Thickness, critical load of failure in the scratch test and Vickers

microhardness of magnetron-sputtered (Ti,Al)(N,C) coatings

Sample

no.

(Ti,Al)(N,C)

coating, carbon

content (at.%)

Thickness

(Am)

Critical load

of failure (N)

Microhardness

(HV0.05)

1 0 4.7 15 1852

2 5.5 5.8 120 1401

3 8.1 4.2 25 1528

4 9.5 3.5 45 1622

5 13.5 4.7 45 1635

6 16.3 5.4 35 1554

7 28.0 5.1 30 1141

M. Stueber et al. / Thin Solid Films 493 (2005) 104–112 111

nisms and microstructures for (Ti,Al)(C,N) coatings depo-

sited by pulsed laser deposition [49].

4. Conclusions

It has been shown that it is possible to create by PVD

methods a nanocomposite coating coexisting of two

metastable phases, one of them being nanocrystalline and

the other being amorphous. The microstructural arrange-

ment of both these phases, i.e., in the form of isolated a-C

clusters within a nanocrystalline matrix or as a continuous

amorphous network or matrix percolated with the nano-

crystalline phase, clearly depends on the process parameters

selected for deposition and can be explained by the

proposed five-step growth model. From this point of view

and as well from the literature it can be concluded that these

different microstructures will result in completely different

property profiles of such coatings. The properties of

nanocomposite coatings are determined by phase boundaries

and by the materials selected to build them. The ability to

synthesise metastable nanocrystalline phases within a nano-

composite coating and the combination of various types of

nanocomposite layers within a nanocomposite multilayer

thin film will offer to the material scientist a powerful tool

for an engineering design of the coatings properties and

performance. For example, the idea of superlattice coatings

built of nanocomposite layers of similar microstructure

seems very promising for future applications.

Taking into account the similarities and differences

between the reported type of nanocomposite coatings and

the literature it must be concluded that the modelling and

simulation of nanocomposite growth processes should focus

on the fundamentals of materials science and thermody-

namics. Exciting properties such as super-toughness and

super-hardness are desired for many applications but are

always only a result of the design of the constitution and

microstructure of a coating. To achieve success in future

engineering of multifunctional PVD coating properties for

specific applications, a deep knowledge and understanding

of phase relations in the phase diagrams, of the thermody-

namics and kinetics of the deposition process and of the

corresponding PVD phase diagrams is desired.

Acknowledgements

The authors would like to appreciate the funding of this

work by the European Commission within Framework

Programme 5 under the Growth Programme, Grant G5RD-

CT2001-00578. Parts of the TEM characterisation is grate-

fully acknowledged also by the Hungarian National Research

Foundation, contract no. OTKAT033075 and T043437. The

technical assistance of D. Zahoor and S. Creasey as well as

discussions with Dr. B. Lewis, Dr. S. Ulrich and Dr. J.

Rodenburg are also gratefully acknowledged.

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