Microstructure modifications of CrN coatings by pulsed bias sputtering

Surface & Coatings Technology 206 (2012) 4666–4671

Contents lists available at SciVerse ScienceDirect

Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat

Microstructure modifications of CrN coatings by pulsed bias sputtering S. Grasser ⁎, R. Daniel, C. Mitterer Department of Physical Metallurgy and Materials Testing, Montanuniversität Leoben, Franz-Josef-Strasse 18, 8700 Leoben, Austria

a r t i c l e

i n f o

Article history: Received 1 February 2012 Accepted in revised form 12 May 2012 Available online 22 May 2012 Keywords: Reactive magnetron sputtering Chromium nitride Pulsed bias voltage Ion bombardment Electron bombardment

a b s t r a c t This paper investigates the influence of a combined electron and ion bombardment via asymmetric bipolar pulsed substrate bias on the microstructure and mechanical properties of reactively magnetron-sputtered CrN hard coatings for two different ion-to-atom flux ratios. A short ratio of electron-to-ion bombardment duration leads to a preferred (110) orientation of the CrN phase, high residual stress up to − 4.1 GPa and hardness values up to 17.9 GPa. An increased electron-toion bombardment ratio promotes a change in preferred orientation from (110) towards (100). Concurrently, residual stress and hardness values are reduced to a few ± 100 MPa and 15.1 GPa, respectively. The texture cross-over from (110) to (100) is shifted to longer electron bombardment duration for higher ion-to-atom flux ratios. © 2012 Elsevier B.V. All rights reserved.

1. Introduction

2. Experimental details

The outstanding properties of transition metal nitride hard coatings (e.g. CrN, TiN) deposited by magnetron sputtering are closely related to their microstructure, which results from the growth condition controlled by kinetic restrictions and thermodynamics. They both depend on the substrate temperature as well as the energy and flux of incident particles incoming at the substrate surface. However, for deposition, mostly macroscopic deposition parameters (e.g. magnetron power, gas flow rates, heating power, substrate bias voltage, current of Helmholtz coil system, etc.) are used to describe the process. While the effects of a negative substrate bias voltage (such as enhanced adatom mobility, formation of defects, implantation and re-sputtering) on the evolution of the microstructure are well investigated [1–8], little is known about the influence of a pulsed bias voltage with a positive fraction. Here, an additional electron bombardment facilitates migration of adatoms on the growing surface and recovery of defects during deposition [9–11]. Thus, kinetic restrictions and thermodynamics during growth can be effectively balanced, promoting the evolution of a favorable microstructure of transition metal nitride hard coatings [12,13]. In this paper, we report on the influence of an asymmetric bipolar pulsed substrate bias voltage on the growth of CrN coatings deposited by reactive unbalanced magnetron sputtering. We show how an alternating electron and ion bombardment effectively controls the evolution of microstructure and thus affects the mechanical properties. Special emphasis will also be laid on the effect of different ion-toatom flux ratios during deposition.

CrN coatings were grown onto polished Si (100) substrates of dimension 7 × 20 × 0.525 mm 3 in a lab-scale dc reactive unbalanced magnetron sputtering system, described in detail elsewhere [14]. A Cr disc (∅ 150 mm, PLANSEE Composite Materials) mounted onto a Gencoa PP150 magnetron was used as target at a target-to-substrate distance of 75 mm. A pivot-mounted shutter is situated midway between the target and substrate holder. The deposition chamber (10ℓ) was evacuated by a turbomolecular pumping unit down to a base pressure of 2 × 10 −3 Pa. Prior deposition, target precleaning was performed for 5 min followed by ion-etching (−1250 V, 20 mA) of the substrates for 10 min, both by argon glow discharge. All CrN coatings were synthesized with a constant target current of 1.5 A resulting in a power density of 3 W/cm 2. The substrate temperature was maintained at 350 °C during the whole deposition process. The working gas total pressure (pT = pAr + pN2) was controlled by a mass flow controller and adjusted to 0.4 Pa with a partial pressure ratio pN2/pT of 0.31. The substrate holder was biased using an MKS ENI RPG-50 pulsed dc plasma generator, which can be operated in dc as well as in asymmetric bipolar pulsed mode (Fig. 1). The latter can be characterized by the period T or frequency f =1/T and a duty cycle D= f × τ, where τ is the positive pulse duration of the electrical signal. The bias power supply operated in pulsed mode can provide a maximum power output of 5 kW with frequencies ranging from 50 to 250 kHz. The positive pulse duration can be set between 0.5 μs and 40% of the period T. In our experiments a constant negative base voltage of − 70V and a pre-set positive reverse voltage of +37 V was used. In order to investigate the influence of a pulsed bias on the growth of CrN coatings, the duty cycle was varied (2.5, 15, 20, 30 and 40%) at a constant pulse frequency of 50 kHz.

⁎ Corresponding author. Tel.: + 43 3842 402 4232; fax: + 43 3842 402 4202. E-mail address: [email protected] (S. Grasser). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.05.043

S. Grasser et al. / Surface & Coatings Technology 206 (2012) 4666–4671

b

+

Table 1 Ion and atom fluxes and ion-to-atom flux ratio in the regions A and B.

Time

A B

  1 J i 1019 s⋅m 2

  1 J a 1019 s⋅m 2

Ji/Ja

2.4 9.1

6.0 5.6

0.4 1.6

-

Time

Pulse duration (τ)

Voltage

Reverse voltage

-

Voltage

+

a

4667

Base voltage

Period (T)

Fig. 1. Schematics of an asymmetric bipolar pulse with negative base voltage and positive reverse voltage with duty cycles D = 1/T × τ of (a) 2.5% and (b) 40%.

The voltage characteristics of the applied substrate bias was recorded with a Tektronix TPS 2024 digital oscilloscope using a PMK PHV 1000 high voltage probe with a probe bandwidth of 400 MHz, a rise time of 0.9 ns and a maximum rated input voltage of 1000 V. The deposition rates were calculated from the coating thickness measured by the ball crater technique and the corresponding deposition time. Plasma potential Vp and ion flux Ji were determined from Langmuir wire probe measurements (Hiden ESP). The arrangement of the electrically grounded adjustable shielding between target and substrate within the deposition chamber creates two regions (A and B) which differ in the ion flux incoming to the substrate surface (Fig. 2). The tungsten tip of the Langmuir probe was situated directly in front of the silicon substrates for both regions A and B. The atom flux Ja was derived from deposition rate measurements. In Table 1 typical average values of ion and atom fluxes and ion-to-atom flux ratio (Ji/Ja) are given corresponding to the positions directly in front of the substrate holder above the center of the target erosion track. The structural investigation of the as-deposited samples was conducted by X-ray diffraction (XRD) in a Siemens D500 diffractometer in θ − 2θ Bragg–Brentano configuration operated at 40 kV and 25 mA using Cu-Kα radiation. The diffraction patterns obtained were compared with the PDF 01-077-0047 (ICDD, 1977). The grain size (size of coherently diffracting domains) of the coatings was determined with a single-line method by evaluation of the (200)-reflection using a pseudo-Voigt function [15]. The texture coefficient Thkl was gained from the normalized hkl XRD intensities as T hkl ¼ Ihkl =½I111 þ I200 þ I 220 ;

ð1Þ

where hkl represents the (111), (200) and (220)-reflections. Morphology and elemental composition of the CrN coatings were investigated using a scanning electron microscope (SEM, Zeiss EVO 50) equipped with an electron-dispersive (EDX, Oxford Instruments

INCA) and wavelength-dispersive (WDX, Oxford Instruments INCAWave) X-ray analyzer. The coating hardness was studied by means of nanoindentation (UMIS, Fischer-Cripps Laboratories) using a Berkovich diamond tip. 20 indents per sample were performed at a constant maximum load of 35 mN in order to avoid the effect of the underlying substrate and to obtain statistically reliable values. The Oliver and Pharr method [16] was used for evaluation of hardness from the load-displacement data. Residual stress of the CrN coatings in their as-deposited state was determined by the substrate-curvature method. The average total residual biaxial coating stress σtot was derived from the deflection of two parallel laser beams caused by the substrate-curvature radius r using the simplified Stoney equation [17,18]: σ tot ¼ M

t 2s 1 : 6t c r

ð2Þ

Here, M is the biaxial elastic modulus of the Si (100) substrate (180 GPa [19]) and ts is the thickness of the substrate. 3. Results and discussion 3.1. Deposition characteristics The MKS ENI RPG-50 power supply used for substrate biasing allows the application of an asymmetric bipolar pulsed voltage. The applied signal features an intensive negative voltage peak supplemental to the negative base voltage and the positive reverse voltage (see Fig. 3). This allows in general to study the influence of the combination of an intensive ion bombardment and mild electron bombardment on the evolution of the microstructure of CrN coatings. The high negative voltage peak is due to the combination of two different power units coupled within the MKS ENI RPG-50 power supply: (i) a primary power unit responsible for the negative base voltage and (ii) a secondary power unit providing the positive reverse voltage. The latter is switched on during the positive fraction of the cycle only. In contrast, the primary power unit is running continuously and the flow of energy is either directed to the output of the power unit or temporarily stored in internal inductors. At the beginning of each negative fraction of the cycle the flow of energy is redirected from the inductors towards the power supply output. Thus, the surplus energy stored within the

A

75mm

Shielding

a B

150m

m

Target TiB2 Fig. 2. Schematic drawing of the target-substrate arrangement and the shielding in between. The two regions (A and B) with different ion-to-atom flux ratio are indicated.

Voltage [V]

Substrate holder

b

200 0 -200 -400 -600 -800 0

10

20

30

40

0

10

20

30

40

Fig. 3. Measured evolution of the asymmetric bipolar pulsed substrate bias voltage over time for a frequency of 50 kHz and a duty cycle of (a) 2.5% and (b) 40%, respectively.

4668

S. Grasser et al. / Surface & Coatings Technology 206 (2012) 4666–4671

inductors is released causing a negative voltage peak as can be seen in Fig. 3. The negative voltage peak is proportional to the energy stored within the inductors during the positive pulse duration τ and can thus be controllably varied. The relation between τ resulting in the negative voltage peak and the duty cycle is shown in Fig. 4. A substrate bias voltage Vb negative with respect to the floating potential of the substrate holder attracts positive ions during deposition, causing an energetic particle bombardment of the coating during its growth. The energy of the incident ions Ei can be estimated from the difference between substrate bias potential and the plasma potential Vp with Ei ≃ e|Vb − Vp| [20]. Due to the voltage peak (see Fig. 3), the ion energy may be as high as 600 eV for several hundreds of nanoseconds, in particular for high duty cycles (see Fig. 4). Incident particles are known to exhibit extensive influence of the formation of the coating microstructure depending on their Ei. Low energy ions activate surface processes and thus enhance nucleation rates and coating density, and subsequently decrease the average grain size. High energetic incident ions also affect the subsurface of the coating mainly by linear collision cascade effects. This results in formation of structural defects and implementation and after exceeding a certain critical limit of Ei, suppression of columnar growth through re-nucleation processes may occur. Further increase in Ei is typically related with re-sputtering, which is, however, strongly texture dependent—an open in-plane geometry promotes channeling effects and thus higher damage resistance [3,21]. A positive substrate bias voltage attracts electrons from the plasma discharge causing a plasma based electron bombardment of the growing coating. The mass of electrons me is much lower than that of the argon ions (me ≃ 10 − 5 × mAr) resulting predominantly in inelastic collisions with the coating surface. Here, the kinetic energy of the electrons is turned into heat, which subsequently enhances the adatom mobility and may result also in recovery of defects, while particle bombardment-induced damage or re-sputtering of the growing coating is reduced [10,12,22]. 3.2. Chemical composition, morphology and coating thickness The CrN atomic ratio of all CrN coatings in the as-deposited state was found to be 1.0 as revealed by EDX measurements. WDX measurements confirmed the presence of argon in coatings synthesized with duty cycles of 20% or higher irrespective of the Ji/Ja ratio. The incorporation of argon is related to the negative voltage peak following each positive pulse (up to − 600 V for a duty cycle of 40%, Fig. 4). Here, Ar + ions from the near-substrate region are attracted and gain energies in the range of several hundreds eV. Thus, their deeper penetration and subsequent trapping in the lattice is very likely [2,23]. The oxygen concentration in the coatings originating from the residual atmosphere was found to be below the detection limit. 0

10

1000

Pulse duration Voltage peak

-200

8

800

-400

6

600

-600

4

400

-800

2

200

-1000

0

Fig. 5. SEM micrographs of CrN coatings in cross-section grown onto Si (100) substrates with Ji/Ja = 1.6 and a duty cycle of (a) 2.5% and (b) 40%, respectively.

The SEM cross-sectional images shown in Fig. 5 illustrate the microstructural evolution of the CrN layers grown with different duty cycles under intense ion bombardment (Ji/Ja = 1.6). All coatings investigated exhibit a compact and fully dense microstructure. Close to the substrate/coating interface very fine randomly orientated grains are dominating. With increasing distance from the interface, the evolution of larger V-shaped columns can be observed for the CrN coatings grown with low duty cycle (2.5%). A high duty cycle of 40% associated with a shorter but more intensive ion bombardment impedes the formation of pronounced columns. This is due to the ion-irradiation induced re-nucleation, which gives rise to the development of a fine crystalline, dense and almost structureless morphology. The development of the coating thickness with duty cycle and Ji/Ja is depicted in Fig. 6. The thickness increases with increasing duty cycle from 3.2 to 3.3 μm and from 2.9 to 3.1 μm for Ji/Ja of 0.4 and 1.6, respectively. This increase is related to the shorter ion bombardment duration at higher duty cycles and therefore reduces re-sputtering during the coating growth. The lower overall thickness at higher Ji/Ja = 1.6 stems from the higher ion flux, leading to enhanced re-sputtering [3,24]. 3.3. Coating structure Figs. 7 and 8 show the XRD-patterns from as-deposited CrN coatings synthesized on Si (100) substrates with different Ji/Ja of 0.4 and 1.6, respectively. All coatings are well crystalline, exhibiting a typical mono-phase B1-NaCl structure [25]. 3.6 3.4 3.2 3.0

Ion energy 0

10

20

30

0 40

Duty cycle [%] Fig. 4. Relationship between duty cycle and pulse duration τ for a given frequency of 50 kHz and implication on the negative voltage peak and corresponding maximum ion energy Ei.

2.8

Ji/Ja = 0.4 Ji/Ja = 1.6

2.6 0

10

20

30

40

Duty cycle [%] Fig. 6. Development of coating thickness versus duty cycle for both low and high Ji/Ja. The symbols and the error bars represent the mean value and standard deviation of 3 individual calottes.

S. Grasser et al. / Surface & Coatings Technology 206 (2012) 4666–4671

(111)

(200)

4669

35

(220) (311) (222)

Ji/Ja = 0.4 Ji/Ja = 1.6

Intensity [a.u.]

Domain size [nm]

30 25 20 15 10 5 0 0 35

40

45

50

55

60 θ

75

Fig. 7. XRD patterns of CrN coatings grown with Ji/Ja = 0.4 and different duty cycles. The x-axis is interrupted from 65 to 74° in order to mask the intensive (004)-reflection of the Si (100) substrate.

CrN grown at Ji/Ja = 0.4 and a duty cycle of 2.5% exhibit a preferred (110) crystallographic orientation with T(220) = 0.69 and a small volume fraction of 111, 200 and 311 grains. The diffraction peaks are shifted to lower 2θ values with respect to the relaxed structure of randomly orientated CrN powder indicating compressive stress in the coating. Increasing duty cycle from 15 to 40% gives rise to pronounced changes in the coating texture, where T(220) decreases from 0.19 to 0.12 and T(200) increases from 0.63 to 0.78. This demonstrates the effective control of the coating texture by variation of the ratio of the electron-to-ion bombardment duration. The less pronounced peak shift at higher duty cycles reveals, in addition, the relation of electron-to-ion bombardment duration ratio and the stress state in the deposited coatings. The domain size calculated from the peak broadening decreases from initially 30 nm at a duty cycle of 2.5% to 12 nm at a duty cycle of 40% (see Fig. 9). Increasing Ji/Ja to 1.6 results again in preferred (110) orientation (T(220) = 0.67) for a duty cycle of 2.5% (Fig. 8). With increasing duty cycle the (220) reflection becomes even more dominating with T(220) = 0.92 at a duty cycle of 20%. A further increase of duty cycle leads to a complete cross-over in texture from (110) to (100) orientation as indicated by T(220) = 0.11 and T(200) = 0.79 for a duty cycle of 40%. Concurrently, all reflections are shifted to lower 2θ for low duty cycles with a maximum at a duty cycle of 20% and diminishing distortions for higher duty cycles. The domain size constantly decreases from initially 20 nm at duty cycle of 2.5% to 10 nm with increasing duty cycle (Fig. 9).

(111)

(200)

10

80

20 30 Duty cycle [%]

40

Fig. 9. Development of the domain size in CrN coatings over duty cycle determined from the (200) diffraction peak of the CrN phase for different Ji/Ja.

The development of a coating texture is typically associated with competitive growth controlled by thermodynamics and kinetic restrictions during deposition. Under thermodynamical conditions, CrN develops with the (100)-texture, which corresponds to the lowest surface energy atomic in-plane arrangement (γ(111) > γ(110) > γ(100)) [2,26]. If kinetic mechanisms are dominating, the (111) orientation develops more likely [27–30]. Under high-energy high-flux irradiation conditions, re-sputtering effects may significantly affect the texture development of the coatings. In such case, an open atomic arrangement of (110) in-plane coating texture is preferred, which is less susceptible for re-sputtering. These conditions are met at low duty cycles of 2.5%, where ion bombardment of high-energetic incident ions reaching energies up to 105 eV dominates. An increase in duty cycle shifts the ratio of electron-to-ion bombardment duration towards longer electron bombardment, thus promoting electron induced annealing effects. Concurrently, the negative peak voltage and thus maximum ion energy is increasing (Ei = 600 eV at duty cycle of 40%, Fig. 4). As a consequence, the conditions controlling competitive growth change and give rise to a thermodynamically favored (100)texture. Additionally, defect density reduces as indicated by less promoted peak shift to lower 2θ angle corresponding to only slight lattice deformation of the deposited coatings. At a higher ion-to-atom flux ratio (Ji/Ja = 1.6) the cross-over in texture from (110) towards (100) orientation is retarded to higher duty cycles corresponding to longer electron bombardment (see Figs. 7 and 8). This is attributed to the more intensive ion flux inducing more defects. Hence, a higher electron-to-ion flux ratio is required to prevail the effects of intensive high flux ion bombardment. Furthermore, crystal growth is impeded more effectively than in the case of Ji/Ja = 0.4, leading to an overall reduction in grain size [21,24].

(220) (311) (222)

Intensity [a.u.]

3.4. Mechanical properties

35

40

45

50

55

60 θ

75

80

Fig. 8. XRD patterns of CrN coatings grown with Ji/Ja = 1.6 and different duty cycles. The x-axis is interrupted from 65 to 74° in order to mask the intensive (004)-reflection of the Si (100) substrate.

The development of the average total residual biaxial stress σtot in the CrN coatings grown on Si (100) substrates with Ji/Ja = 0.4 and 1.6 and various duty cycles is shown in Fig. 10. The results demonstrate the strong relationship between energy and flux of charged particles and neutrals and the residual stress in CrN coatings. Depending on the duty cycle, two distinct groups can be distinguished. At low duty cycles, corresponding to a low ratio of electron-to-ion bombardment duration, coatings show high compressive stress ranging from − 1.7 to − 4.1 GPa. These coatings exhibit a preferred (110) crystallographic orientation. In contrast, at high duty cycles, residual stress is reduced to a few ± 100 MPa, and the (100)-orientated grains are dominating. The results gained from nanoindentation measurements of CrN coatings deposited with Ji/Ja = 0.4 and 1.6 at duty cycles ranging from 2.5 to 40% are shown in Fig. 11. CrN coatings grown with a

4670

S. Grasser et al. / Surface & Coatings Technology 206 (2012) 4666–4671

1

19 18

(100)

Hardness [GPa]

Residual stress [GPa]

tension compression tension compression

0 -1 -2 -3

(110) -4

Ji/Ja = 0.4 Ji/Ja = 1.6

-5

17 16 15

J i/J a = 0.4 J i/J a = 1.6

14 0

10

20

30

40

1

Fig. 10. Evolution of residual stress of CrN coatings vs duty cycle for different Ji/Ja.

duty cycle of 2.5% show hardness values between 16.6 and 17.9 GPa for Ji/Ja of 0.4 and 1.6, respectively. Increasing the duty cycle to 40% decreases the hardness values to 15.8 and 15.1 GPa for Ji/Ja of 0.4 and 1.6, respectively. The relationship between hardness and residual stress for CrN coatings grown with Ji/Ja of 0.4 and 1.6 is depicted in Fig. 12. Coatings with preferred (100) orientation (T(200) > 0.51) show low hardness values (15.1 to 15.9 GPa) and compressive or tensile stress levels in the range of a few 100 MPa. In contrast, (110)-orientated grains are dominating (T(220) > 0.76) in CrN coatings showing high hardness (16.6 to 17.9 GPa) in connection with high residual compressive stress (− 1.7 to − 4.1 GPa). The resulting residual stress state as well as the hardness of the CrN coatings are caused by the complex interplay of ion and electron bombardment during growth. High energetic ion bombardment promotes the formation of surface and sub-surface defects (e. g. implantation of ions, point defects). Consequently, the CrN lattice becomes distorted, causing compressive in-plane stress. Additionally, the movement of dislocations is hindered, increasing the hardness of the CrN coatings [3,6,21,30]. In contrast, the electron bombardment during growth stimulates the mobility of adsorbed particles. As a result, defect density and lattice distortions are reduced by recovery effects and thereby hardness as well as residual compressive stress decreases [10]. Depending on the ratio of electron-to-ion bombardment duration, reflected by the duty cycle, either formation or annihilation of defects prevails. This is also illustrated by the influence of the ion-to-atom flux ratio (0.4 and 1.6) on the development of residual stress and hardness. At Ji/Ja = 1.6, a higher duty cycle, corresponding to a longer electron bombardment duration, is required in order to compensate for the higher defect density generated by the more intense high energetic ion flux. The effect of recovery is even more pronounced due to the higher defect density

-1

-2

-3

-4

-5

Fig. 12. Interrelationship between residual stress and hardness of CrN coatings grown with different Ji/Ja.

generated by the higher ion flux. It is remarkable, that the electron bombardment is sufficient to compensate even for the damage caused by the intense ion bombardment (Ei up to 600 eV, see Fig. 4) at high duty cycles. 4. Conclusions Reactively magnetron-sputtered CrN hard coatings were synthesized using asymmetric bipolar pulsed substrate bias with duty cycle from 2.5 to 40%. We established the interrelationship between the ratio of electron-to-ion bombardment duration during growth and microstructure and mechanical properties for different ion-toatom flux ratios (Ji/Ja = 0.4 and 1.6). At low duty cycles (corresponding to a low ratio of electron-to-ion bombardment duration) the influence of the intense ion bombardment dominates. The CrN coatings show a preferred (110) crystallographic orientation with high residual stress and hardness of up to − 4.1 GPa and 17.9 GPa, respectively. A higher duty cycle stimulates adatom mobility through enhanced electron bombardment of the coating surface-near region, thus, altering growth kinetics. The texture changes from preferential (110) to (100) orientation. Concurrently, residual stress and hardness are reduced to a few ± 100 MPa and 15.1 GPa, respectively. Increasing Ji/Ja from 0.4 to 1.6 defers the shift in texture, residual stress and hardness to higher duty cycles, corresponding to longer electron bombardment duration. In summary, asymmetric bipolar pulsed substrate biasing provides an effective tool to control the formation and annihilation of defects in B1-NaCl structured transition metal nitrides. Consequently, it allows the design of a desired microstructure with applicationtailored mechanical properties. Acknowledgments

19

Ji/Ja = 0.4 Ji/Ja = 1.6

18

Hardness [GPa]

0

Residual stress [GPa]

Duty cycle [%]

This work was supported by the Austrian NANO Initiative via a grant from the Austrian Research Promotion Agency FFG within the project “HIPIMS coating”.

17

References 16 15 14 0

10

20

30

40

Duty cycle [%] Fig. 11. Evolution of hardness of CrN coatings vs duty cycle for different Ji/Ja.

[1] R. Messier, A.P. Giri, R.A. Roy, J. Vac. Sci. Technol. A 2 (2) (1984) 500. [2] I. Petrov, L. Hultman, J.-E. Sundgren, J.E. Greene, J. Vac. Sci. Technol. A 10 (2) (1992) 265. [3] W. Ensinger, Nucl. Instrum. Methods Phys. Res. B 127-128 (1997) 796. [4] J. Musil, Surf. Coat. Technol. 100–101 (1998) 280. [5] T. Hurkmans, D.B. Lewis, H. Paritong, J.S. Brooks, W.D. Münz, Surf. Coat. Technol. 114 (1) (1999) 52. [6] P.H. Mayrhofer, G. Tischler, C. Mitterer, Surf. Coat. Technol. 142–144 (2001) 78. [7] J.J. Olaya, G. Wei, S.E. Rodil, S. Muhl, B. Bhushan, Vacuum 81 (5) (2007) 610. [8] R. Daniel, D. Holec, M. Bartosik, J. Keckes, C. Mitterer, Acta Mater. 59 (7) (2011) 6631. [9] M. Andritschky, F. Guimarães, V. Teixeira, Vacuum 44 (8) (1993) 809.

S. Grasser et al. / Surface & Coatings Technology 206 (2012) 4666–4671 [10] A.G. Fedorus, E.V. Klimenko, A.G. Naumovets, E.M. Zasimovich, I.N. Zasimovich, Nucl. Instrum. Methods Phys. Res. B 101 (1-2) (1995) 207. [11] J.M. White, J. Mol. Catal. A: Chem. 131 (1–3) (1998) 71. [12] P.J. Kelly, R. Hall, J. O'Brien, J.W. Bradley, G. Roche, R.D. Arnell, Surf. Coat. Technol. 142–144 (2001) 635. [13] J.-W. Lee, S.-K. Tien, Y.-C. Kuo, C.-M. Chen, Surf. Coat. Technol. 200 (10) (2006) 3330. [14] P. Losbichler, C. Mitterer, Surf. Coat. Technol. 97 (1–3) (1997) 567. [15] T.H. de Keijser, J.I. Langford, E.J. Mittemeijer, A.B.P. Vogels, J. Appl. Crystallogr. 15 (3) (1982) 308. [16] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (6) (1992) 1564. [17] G.G. Stoney, Proc. R. Soc. London A 82 (553) (1909) 172. [18] W.D. Nix, Metall. Mater. Trans. A 20A (11) (1989) 2217. [19] J.J. Wortman, R.A. Evans, J. Appl. Phys. 36 (1) (1965) 153. [20] B. Chapman, Glow Discharge Processes, Wiley Interscience, New York, 1980. [21] I. Petrov, P.B. Barna, L. Hultman, J.E. Greene, J. Vac. Sci. Technol. A 21 (5) (2003) 117.

4671

[22] M. Berger, L. Karlsson, M. Larsson, S. Hogmark, Thin Solid Films 401 (1–2) (2001) 179. [23] L. Hultman, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Vac. Sci. Technol. A 7 (3) (1989) 1187. [24] W. Olbrich, G. Kampschulte, Surf. Coat. Technol. 61 (1–3) (1993) 262. [25] M. Venkatraman, J.P. Neumann, in: T. Massalski (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, 1990. [26] I. Petrov, L. Hultman, U. Helmersson, J.-E. Sundgren, J. Greene, Thin Solid Films 169 (2) (1989) 299. [27] L. Hultman, J.-E. Sundgren, J.E. Greene, D.B. Bergstrom, I. Petrov, J. Appl. Phys. 78 (9) (1995) 5395. [28] D. Gall, S. Kodambaka, M.A. Wall, I. Petrov, J.E. Greene, J. Appl. Phys. 93 (11) (2003) 9086. [29] R. Daniel, K.J. Martinschitz, J. Keckes, C. Mitterer, J. Phys. D: Appl. Phys. 42 (7) (2009) 75401. [30] R. Daniel, K.J. Martinschitz, J. Keckes, C. Mitterer, Acta Mater. 58 (7) (2010) 2621.