Growth, stress and hardness of reactively sputtered tungsten nitride thin films

Surface & Coatings Technology 205 (2010) 1953–1961

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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t

Growth, stress and hardness of reactively sputtered tungsten nitride thin films M. Wen a, Q.N. Meng a, W.X. Yu a, W.T. Zheng a,⁎, S.X. Mao b, M.J. Hua b a b

Department of Materials Science, Key Laboratory of Automobile Materials, MOE, and State Key Laboratory of Superhard Materials, Jilin University, China Department of Mechanical Engineering and Materials Science, University of Pittsburgh, Pittsburgh, PA 15261, USA

a r t i c l e

i n f o

Article history: Received 6 May 2010 Accepted in revised form 19 August 2010 Available online 25 August 2010 Keywords: WNx thin films Magnetron sputtering Phase transition Stress Hardness

a b s t r a c t Tungsten nitride (WNx) thin films were deposited on Si(100) substrates using direct current reactive magnetron sputtering in discharging a mixture of N2 and Ar gas. The effects of nitrogen flow rate (FN2) and substrate bias voltage (Vb) on the composition, phase structure, and mechanical properties for the obtained films were evaluated by means of X-ray photoelectron spectroscopy, X-ray diffraction, high-resolution transmission electron microscopy and nanoindentation. The evolution of phase structure is found closely correlated to N concentration in the films. When Vb = −40 V, with increasing FN2, the N/W atomic ratio gradually increases in the film, accompanied by a phase transition from cubic β-W to hexagonal WN through face centered-cubic (fcc)-W2N. At FN2 = 15 sccm, the N/W atomic ratio gradually decreases with increasing the absolute value of Vb, resulting in a transition from fcc-W2N to cubic β-W(N) through a mixture of fccW2N + β-W(N). In addition, the increase in implanted nitrogen causes the increase in the compressive stress with increasing FN2. In contrast, although with increasing the absolute value of Vb from 80 to 160 V the N/W atomic ratio decreases, the increase of the defects caused by increasing ion bombarding energy, dominates the increase of the compressive stress. Furthermore, the maximum hardness value for the films arrives at 38.9 GPa, which is obtained at Vb = −120 V when fcc-W2N + β-W(N) mixed structure is formed. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Tungsten nitride (WNx) belongs to a class of refractory metal nitrides that have the unique properties of excellent hardness, chemical inertness, high melting point, good chemical stability and high conductivity [1]. The combination of these properties makes tungsten nitride a suitable material for diffusion barriers in microelectronic devices [2–8], Schottky contacts to GaAs [9,10], gate electrodes in metaloxide-semiconductor field effect transistors [11–14], and hard wearresistant protective coatings for cutting tools [15–17]. Furthermore, many investigations [8,14,15,18] have revealed that the phase structure and composition in WNx films have a significant effect on the performance of the layer in certain applications, such as diffusion in microelectronic devices, electrical resistivity, work function, and hardness. However, due to the complex relationship between processing parameters and physical properties of the layers, an understanding of these relationships is crucial for designing layers with the desired properties and exceptional performances. So far, many techniques have been used to achieve tungsten nitride films, such as reactive magnetron sputtering [14–17,19–28], pulsed laser deposition [29], ion-beam sputter deposition [30], atomic layer deposition [8], metal organic chemical vapor deposition [6], and plasma enhanced chemical vapor deposition [31–33]. Among these ⁎ Corresponding author. Tel./fax: +86 431 85168246. E-mail address: [email protected] (W.T. Zheng). 0257-8972/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2010.08.082

techniques, reactive sputtering is of special interest because it is an industrial process applicable to large-area deposition, and high quality films can be obtained even at a low substrate temperature. The microstructure and properties of WNx films deposited by reactive sputtering are strongly dependent on the nitrogen flow ratio (FN2) [15,26] and substrate bias voltage (Vb) [18]. Considering the difference in heats of formation, formation of WNx should be more difficult than that of TiN, NbN and ZrN (ΔH298 =−72 kJ/mol for W2N f [19], ΔH298 =−337 kJ/mol for TiN [34], ΔH298 =−220 kJ/mol for NbN f f [35], and ΔH298 =−365 kJ/mol for ZrN [34]). It is known that in the f TiN, NbN and ZrN systems, the stoichiometric TiN, NbN and ZrN phases can be easily obtained at reactive magnetron sputtering. However, many researchers [14,16,17,20,22,26] have reported that a phase transition from W to face centered-cubic (fcc)-W2N takes place in WNx films with increasing N2 partial pressure and no stoichiometric WN phase is formed, while other investigators [15,19] have observed a change in the structure for WNx films, from W to hexagonal WN through fcc-W2N with increasing N2 partial pressure. It is not clear when stoichiometric hexagonal WN appears. In addition, as the absolute value of Vb is increased, a phase transition from fccW2N structure to W or W + W2N has been reported [18], which has not been observed for TiN, NbN and ZrN films. The mechanisms about the effect of Vb on the phase evolution in WNx thin films have not yet been well understood. Furthermore, very little effect in the literature [19,26] has been focused on the relationship between stress, hardness, and microstructural properties of the WN film.

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In this work, the WNx films have been prepared on Si(100) substrates using direct current (DC) reactive magnetron sputtering in discharging a mixture of N2 and Ar gas, and the deposition rate, chemical bonding, phase configuration, intrinsic stress, and hardness as a function of FN2 and Vb for the obtained films have been investigated. The mechanisms about the effects of FN2 and Vb on the phase transition have been discussed, and the relationships between stress, hardness, and microstructural properties of the WNx film have also been built. 2. Experimental The WNx films were deposited on Si(100) wafers by DC reactive magnetron sputtering a 6-cm-diam W target (99.95%) in discharge of a mixing Ar (99.999%) and N2 (99.999%). The schematic drawing of magnetron sputter system was given in Rf. [36]. The distance between the target and substrate holder was fixed at 8 cm, and the chamber was evacuated by a turbomolecular pump to 4 × 10−4 Pa prior to film deposition, in which the pumping speed was fixed at a high value of 1000 l/s for all experiments. During the deposition, the applied current on the W target and substrate temperature were kept at 0.2 A and 200 °C, respectively. The flow rate of N2 and Ar was accurately controlled by mass flow controller, wherein FN2 was varied from 5 to 100 sccm (sccm denotes cubic centimeter per minute at STP), while Ar flow rate was always kept at 60 sccm. The total pressure was set at 0.8 Pa during all depositions by regulation of the throttling valve. Various DC negative substrate bias voltage ranging from −40 V to −200 V were applied to the substrate during deposition. The thickness of all WNx films was in the order of 1.0–1.2 μm by adjusting the deposition time. The structure, preferred orientation, and grain size were evaluated by X-ray diffraction (XRD) using a Bragg–Brentano diffractometer (D8_tools) in θ–2θ configuration with Cu Kα line at 0.15418 nm as a source. The chemical binding state and composition were analyzed by X-ray photoelectron spectroscopy (XPS) (Perkin-Elmer PHI-5702). The surface of all WNx films is cleaned by Ar ion beam for 2 min before measurement and all the surface spectra have been calibrated to the C 1s binding energy for adventitious carbon, known to be present at 284.8 eV. Since XPS was used as a semi-quantitative analysis technique, we mainly focused on the trend of composition variation with FN2 and Vb. The nanostructure was characterized by highresolution transmission electron microscopy (HRTEM) (field emission JEOL 2010F) operated at 200 kV. For the HRTEM specimen, a piece of film was peeled off from the substrate and then ultra-sonicated in acetone for 10 min. Droplets of the suspension were deposited on Cu grids fitted with a carbon membrane. The mechanical properties were evaluated by MTS Nanoindenter XP with continuous stiffness measurements (CSM) mode, in which a Berkovitch-type pyramidal diamond tip indented the films to a maximum depth of 500 nm. At least six indentations at different places on the film surface were made. Both the film thickness and curvature of the Si wafer were measured by an AMBIOS XP-2 surface profiler, and the residual stress was calculated using the Stoney equation [37]. Assuming that the film was at a biaxial stress state, the thermal contribution to the total measured stress was estimated as Ef  σ th = ΔαΔT  1−νf

substrate [39]). In this work, the intrinsic stress was obtained by subtracting the thermal stress from the total measured stress. 3. Results 3.1. Deposition rate and composition Fig. 1(a) shows the deposition rate and target voltage for WNx films as a function of FN2, in which the deposition rate decreases gradually from 17.4 to 8.8 nm/min as FN2 increases from 0 to 100 sccm, whereas at a constant current (0.2 A), target voltage increases with increasing FN2. Similar trend in deposition rate has been observed in reactive sputtering W [20], Ta [40], or Nb [41] target in discharging a mixture of N2 and Ar. Assuming that the ratio of the argon and nitrogen ion fluxes directed towards the target is kept constant, the sputtering yield of tungsten should increase with increasing target voltage. However, as FN2 increases, WNx formation takes place at the surface of the W target, and nitrogen ion flux directed towards the target increases while argon ion flux towards the target decreases. Therefore, the observed reduction in deposition rate for WNx films with an increase in FN2 is mainly due to the nitridation of the W target and the lower sputtering yield of WNx (W2N or WN) compared to W, as well as the reduced sputtering efficiency of nitrogen ions with regard to that of argon ions [20]. Occurrence of nitridation at the surface of the W target has been confirmed by the increment of the target voltage with increasing FN2 due to a lower emission coefficient of secondary electrons for WNx. As the target surface is covered with WNx, the target voltage must increase to maintain the supply of electrons and the constant discharging current. Fig. 1(b) shows atomic ratio (N/W) and oxygen content

ð1Þ

where Δα was the difference in thermal expansion coefficients between the film and substrate, ΔT the temperature change between before and after deposition, Ef the Young's modulus and νf Poisson's ratio of the film. Specially, the thermal expansion coefficients were reported as (αW = 4.5 × 10−6 K− 1 for the W phase [38], α W(N) = 5.0 × 10−6 K− 1 for the W(N) solid solution phase [38], α WN = 5.7 × 10−6 K− 1 for W2N and hexagonal WN phases [38], and α Si = 3.15 × 10−6 K− 1 for the Si

Fig. 1. The deposition rate and the voltage on W target (a), and atomic ratio (N/W) and oxygen composition (b) for WNx films as a function of FN2, in which the solid line is just guided to eyes.

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determined by XPS for WNx films as a function of FN2. The N/W atomic ratio increases gradually with increasing FN2 and reaches to 1.17 at FN2 = 100 sccm. The result indicates that the incorporation of nitrogen is approximately proportional to the FN2, which is consistent with the other reports [19,26]. Our XPS results also show the presence of 4–10 % oxygen in the WN films and the oxygen content decreases as FN2 increases. The most likely source of oxygen contamination is either from residual oxygen in the deposition chamber or from the tungsten sputter target under present experiment conditions. The deposition rate and N/W atomic ratio determined by XPS for WNx films as a function of Vb are shown in Fig.2. The variations of the deposition rate and N/W atomic ratio with Vb show a similar behavior. The deposition rate and N/W atomic ratio monotonously decrease when the absolute value of Vb increases. In addition, the corresponding oxygen content determined by XPS for WNx films deposited at Vb = −40, −80, −120, −160, and −200 V is about 5.4, 6.2, 6.9, 8.7, and 10.5%, respectively. It is noted that oxygen content increases as the absolute value of Vb increases and as FN2 decreases, wherein oxygen content is inversely proportional to N composition in the WNx film. The target voltage keeps constant for all Vb. The variations of deposition rate with Vb can be attributed to the back-scattering effect of incident ions and/or film re-sputtering [42]. Since the sputtering yield increases with increasing energy of bombarding ions, more film material is sputtered when Vb increases.

3.2. Chemical bonding The typical XPS survey spectra for WNx films deposited at FN2 = 60 sccm is shown in Fig. 3(a), in which spectral signatures of tungsten, nitrogen, oxygen and carbon are observed. Fig. 3(b) and (c) exhibit the high-resolution XPS core-level spectra in the energy region of W 4f and N 1s for WNx films as a function of FN2. For the W 4f spectra, a four-peak deconvolution using Gaussian functions is shown in Fig. 3(b). The WNx film deposited at 15 sccm exhibits two distinct doublet peaks for W4f7/2 and W4f5/2 spectra: one at 32.1 and 34.0 eV, and another at 35.3 and 37.5 eV. The low-energy side of W4f7/2– W4f5/2 doublet peaks is considered to be associated with the W–N bonding [26], since the W4f7/2–W4f5/2 doublet peaks relative to that (centered at 31.5 and 33.4 eV) of the W film deposited in pure Ar discharging were shifted to higher binding energies, reflecting a charge transfer from tungsten to nitrogen during the nitridation process. The high binding energy peaks fall within the reported binding range for WO3 (cubic structure) and correspond to a valence of VI for W [43]. It is noted that with increasing FN2 from 15 to 100 sccm, the W4f7/2–W4f5/2 doublet peaks at bonding energies of 32.1 and 34.0 eV were slightly shifted to high bonding energies of 32.4

Fig. 3. Typical XPS survey spectra of WNx films grown at FN2 = 60 sccm (a), and corelevel XPS spectra for (b) W 4f and (c) N 1s for WNx films as a function of FN2.

Fig. 2. The deposition rate and N/W atomic ratio for WNx films as a function of Vb, in which the solid line is just guided to eyes.

and 34.4 eV, which may be caused by an increase in the N/W atomic ratio. N 1s XPS spectra exhibit a number of N 1s peaks, as shown in Fig. 3(c). The N XPS peaks for deposition at FN2 = 15 sccm can be deconvoluted into four components with the corresponding binding energies at about 397.6 eV (N1), 398.9 eV (N2), 400.0 eV (N3), and 401.2 eV

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Fig. 5 shows the typical N 1s XPS spectra for WNx films deposited at Vb = −40, −80 and −120 V, respectively. It is noted that, compared to WNx films deposited at Vb = −40 and −120 V, a new peak at 395.3 eV appears for the film grown at Vb = −80 V. However, the reasons for the presence of this peak are still unclear.

(N4). The N1 peak component is contributed to W–N species, which is the typical of nitride ion with Nδ−, as reported in the literature [26,44]. Chang et al. [7] have reported that for the WNx film prepared by nitridation of fine grain chemical vapor deposited tungsten film, only one N XPS peak at 399.7 eV is observed, which is due to N atoms or molecules present in grain boundaries of WNx. Hence it is believed that the N3 peak at 400.0 eV is ascribed to the N atoms or molecules present in grain boundaries of WNx. The N4 peak at approximately 401.3 eV, analyzed by previous researchers, [45] is assigned to N atoms at tetrahedral interstitial sites. However, the origin of the observed N2 peak at a binding energy of about 398.9 eV is not clear. Given the presence of the interstitial species and presence of tungsten vacancy (formation of Frenkel defects), it seems that a strong possibility that the nitrogen coming from the growth atmosphere may be available to fill a tungsten vacancy. Intrinsic Frenkel defects have been observed widely in the sputtered InN [46,47] and ZnO [48] films, which is a consequence of energetical particle bombardment. Compared to W2N, InN prefers to form Frenkel defects with larger In3+ cations [46]. The consideration of larger W δ+ cations in WNx to form Frenkel defects is reasonable. The in situ formed tungsten vacancies are occupied by nitrogen coming from the growth atmosphere, and hence nitrogen-on-tungsten antisite defects are generated. As suggested in InN [46], the N1s bonding energy, higher than that of In–N species and lower than that of interstitial nitrogen, is contributed to nitrogen-on-indium antisites. Therefore, the observed peak at ~ 398.9 eV can be assigned as nitrogen-on-tungsten antisites. The N 1s peaks for WNx films deposited at FN2 = 30, 60, and 100 sccm can be deconvoluted into four peaks similarly. The relative integrated intensities of the four components in the N 1s spectra are plotted as a function of FN2, which is shown in Fig. 4, wherein the fraction of N2 is close to 50% for the WNx film deposition at low FN2 (15 sccm). This means that at low FN2, N atoms are more favorable to occupy the positions of tungsten vacancies created by ions bombardment because of low formation energy. As FN2 increases, once most tungsten vacancies are occupied, more nitrogen can enter octahedral interstitial sites to form W–N species since it is more difficulty for nitrogen to enter tetrahedral positions due to smaller interstitial radius. According to the cannonball rule, octahedral interstitials can have radii up to 0.414R and tetrahedral interstitials up to 0.225R, where R is the radius of the cannon balls at lattice sites [49]. Therefore with increasing FN2 from 15 to 100 sccm, the fraction of N1 increases rapidly and that of N2, N3 and N4 decreases (Fig. 4), accompanied by an increase in the N/W atomic ratio (Fig. 1b).

Fig. 6 shows the XRD patterns in symmetric θ–2θ configuration for sputtered WNx films deposited at (a) different FN2 with a constant Vb of −40 V and at (b) different Vb with a constant FN2 of 15 sccm. In this symmetric geometry, the X-ray scattering vector is confined along the film growth direction and such XRD measurement can provide information on the crystallographic texture and out-of-plane lattice parameters. The evaluated values of 2θ positions and out-of-plane interplanar spacing (d) for the WNx films are listed in Table 1. For the film grown at FN2 = 0 sccm, Fig. 6(a), sputtering tungsten occurs and the peaks at 2θ= 35.61° and 40.10° can be ascribed to the (200) and (210) planes of the cubic β-W phase (JCPDF 47-1319). In contrast, for the film deposited at FN2 = 5 sccm, the peak from β-W(200) disappears, while a small peak from W2N(111) appears due to some nitrogen atoms occupying the interstitial positions. It is believed that when inserting interstitial atoms, such as nitrogen, into the cubic β-W structure, the βW structure can be distorted and be transferred to fcc-like structure. When FN2 increases to 15 sccm, only the fcc-W2N phase is observed, wherein the diffraction peaks at 2θ = 37.03°, 42.79° and 62.41° are associated with (111), (200), and (220) reflections, respectively. The crystal structure of W2N is a B1-NaCl type where the W atoms occupy the positions of the fcc lattice sites and N atoms occupy 50% of the available octahedral interstitial positions. The peaks from fcc-W2N (111) and (200) shift to low diffraction angle with increasing FN2 up to 30 sccm, meaning that the lattice constant of the W2N phase dilates significantly. As FN2 further increases to 60 sccm, the XRD pattern dramatically changes and shows a hexagonal WN (100) peak. At FN2 of 100 sccm, the hexagonal WN (100) peak becomes broader and weaker, indicating the presence of an amorphous phase. [Fig. 6(b)] As Vb increases from 40 V to 80 V, the intensity of fcc-W2N (111) and (200) peaks drop dramatically and a broad W peak at 2θ ≈39.00o appears, indicating that film exhibits a mixed W + W2N phases and poor crystallinity. However, the only peak at 2θ ≈ 39.00o cannot be indexed unanimously since both bcc α-W(110) and cubic β-W(210) peaks appear at about 39.00o. In order to further clarify the phase structure, selected area electron diffraction (SAED) patterns for WNx film deposition at Vb = −40, −80, and −160 V are obtained in Fig. 7, from which the observed diffraction patterns from the

Fig. 4. The relative integrated intensities of the four components in the N1s spectra for WNx films as a function of FN2, in which the solid line is just guided to eyes.

Fig. 5. Typical XPS N 1s core-level spectra for WNx films disposition at Vb = −40, −80, and −120 V, respectively.

3.3. Crystalline structure (XRD and HRTEM)

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Table 1 Evaluated d, (hkl), and grain size (D) values, and corresponding standard values for the WNx thin films prepared at different nitrogen flow rates with a constant substrate bias of −40 V, and at different substrate biases with a constant nitrogen flow rate of 15 sccm. Deposition conditions

FN2 (sccm) 0 5 15 30 60 100 Vb (−V) 40 80

120

160 200

Fig. 6. The XRD patterns in the Bragg–Brentano (θ–2θ) configuration for WNx thin films grown at (a) different FN2 with a constant Vb of −40 V and at (b) different Vb with a constant FN2 of 15 sccm.

film grown at Vb = −80 V can be indexed to a mixture of β-W(N) + fccW2N. Hence the peak at 2θ ≈ 39.00o can be assigned as β-W(N) (210). At a Vb of −120 V, the intensity of β-W(N) (210) and fcc-W2N (111) peaks increases significantly and the film exhibits a better crystallinity compared to that deposited at Vb = −40 V. With further increasing the absolute value of Vb to 160 V, only the β-W(N) (210) peak remains. However, the W0.75 N phase is also observed in the SAED pattern for the film deposited at Vb = −160 V (Fig. 7c). The peak of the W0.75 N phase is not present in XRD for the film deposited at Vb = −160 V, which could be due to the detection limit of XRD. The evolution from a W2N structure to a mixed W(N) + W2N or W(N) phase with increasing the absolute value of Vb has been reported in Ref. [12], and the result is also consistent with the variation of the N/W atomic ratio in Fig. 2. The HRTEM analyses are further carried out to investigate the microstructure evolution as Vb varies. Fig. 7 shows the typical HRTEM images for the films formed at Vb = −40 V, −80 V and −160 V, respectively. It is clear that the film disposition at Vb = −40 V exhibits a good crystallinity and large grain size. Moreover, the locations of some dislocations can be identified from the inverse fast-Fouriertransformed (IFFT) image for the film deposited at Vb = −40 V, indicating that there exist a high density of defects in the film, which is consistent with the XPS results. As the absolute value of Vb increases from 40 V to 80 V, the grain size in the film decreases dramatically and an amorphous structure surrounding the nano-crystalline grain is

Evaluated values

Standard values (PCPDF)

2θ

d(Å)

hkl

D(nm)

2θ

d(Å)

hkl

35.61 40.10 39.95 37.03 42.79 36.56 42.56 35.47 35.47

2.519 2.247 2.255 2.426 2.112 2.456 2.122 2.529 2.529

200 210 110 111 200 111 200 100 100

57.5 18.3 9.7 15.2 13.8 11.1 13.4 5.1 4.0

35.52 39.88 39.88 37.73 43.85 37.73 43.85 35.82 35.82

2.525 2.258 2.258 2.382 2.063 2.382 2.063 2.505 2.505

200 210 210 111 200 111 200 100 100

37.03 42.79 37.13 39.01 42.87 37.54 39.85 42.82 39.45 39.63

2.426 2.112 2.419 2.308 2.108 2.394 2.261 2.110 2.282 2.272

111 200 111 110 200 111 110 200 110 110

15.2 13.8 12.1 2.2 10.7 8.7 4.9 6.8 5.0 4.7

37.73 43.85 37.73 39.88 43.85 37.73 39.88 43.85 39.88 39.88

2.382 2.063 2.382 2.258 2.063 2.382 2.258 2.063 2.258 2.258

111 200 111 210 200 111 210 200 210 210

observed. Further increasing the absolute value of Vb to 160 V, the content of the amorphous phase decreases and crystallinity becomes significantly enhanced. The HRTEM images confirm the XRD results. Table 1 lists the crystallite size (D) of the WNx films, which is roughly estimated from the full width at half maximum (FWHM) according to Scherrer's formula [50]. The crystallite size of the W(N) film reduces significantly with increasing FN2 from 0 to 5 sccm, which may be a result of that the growth of tungsten grain is limited by adding nitrogen as well as a transition from β-W(N) to W2 N. As FN2 increases to 15 sccm, the crystallite size of the W2 N phase increases. With further increasing FN2, crystallite sizes reduce again due to the presence of a hexagonal WN and addition of excess nitrogen. In contrast, as the absolute value of Vb increases, the crystallite size of the fcc-W2N phase decreases, while that of the W phase increases, indicating that the high absolute value of Vb promotes the growth of the W phase. 3.4. Residual stress and hardness Fig. 8 illustrates the intrinsic stress (IS) evolution as a function of FN2 and Vb. As Vb is fixed at −40 V, the compressive stress of WNx films is strongly dependent on FN2, as shown in Fig. 8(a). For the W film grown in pure Ar discharging gas, IS is a tensile stress, having a value of 0.8 GPa, while it becomes a compressive, having a value of 1.2 GPa, as FN2 = 5 sccm. The compressive stress increases from 1.2 to 5.6 GPa with further increasing FN2 from 5 to 60 sccm, and finally decreases as FN2 increases from 60 to 100 sccm. On the other hand, as FN2 is fixed at 15 sccm, the effect of Vb on IS is plotted in Fig. 8(b), wherein all films are under the compressive stress. The compressive stress deceases significantly as the absolute value of Vb increases from 40 to 80 V, and then increases gradually when that of Vb increases up to 160 V, and finally decreases again as that of Vb increases from 160 to 200 V. Fig. 9(a) displays the dependence of the hardness on FN2. The W film deposited at FN2 = 0 sccm exhibits the lowest hardness of about 14.4 GPa. With increasing FN2 to 15 sccm, the hardness sharply increases from 14.4 to 34.5 GPa. As FN2 increases up to 100 sccm, the film hardness drops gradually from 34.5 to 24.9 GPa. Fig. 9(b)

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Fig. 7. Typical HRTEM images for WNx films disposited at Vb of (a) −40 V, (b) −80 V, and (c) −160 V, in which the inset in (a) shows the corresponding inverse fast-Fouriertransformed image of area enclosed by the white square, while the position of the dislocation cores have been indicated by T's.

shows the variation of hardness for WNx films with Vb, in which the hardness is strongly dependent on Vb. As the absolute value of Vb increases from 40 to 80 V, hardness decreases dramatically from 34.5 to 19.2 GPa, and then increases to 38.9 GPa with a further increase in the absolute value of Vb to 120 V, and finally slightly decreases as the absolute value of Vb increases from 120 to 200 V.

4. Discussion 4.1. Composition and crystal structure The experimental results showed that the evolution of the phase structure is strongly correlated with N concentration in the film. As FN2 increases, the N/W atomic ratio gradually increases, accompanied by a phase transition from cubic β-W to hexagonal WN through fccW2N. On the other hand, the N/W atomic ratio decreases with increasing the absolute value of Vb, resulting in a change from fccW2N to a mixture of fcc-W2N + β-W(N) or β-W(N). FN2 and Vb play an opposite role in determining the variation of N concentration and the evolution of structure. It is known that the fraction of gas flow rate (or FN2) can affect the bombardment ion species, while the negative Vb can control the energy of the ionized bombarding species. How do FN2

and Vb affect the variation of N concentration and the evolution of structure in WNx films? During reactive bias sputter deposition of WNx, the relative N incorporation in the as-deposited film mainly results from the chemisorption of nitrogen ions and atomic N generated in the plasma, dissociative chemisorption of N2, direct implantation, and recoil implantation. Petrov et al. [51] have investigated the ion species in the ion flux directed towards the substrate, using sputtering Ti target in discharging a mixed Ar+ N2 gas. At a low nitrogen partial pressure Ar+ is the main ion species in the plasma. On the other hand, N+ 2 increases significantly with increasing the nitrogen partial pressure. The similar phenomenon might take place during sputtering W in discharging a mixed Ar+ N2 gas. The increase of nitrogen ions in the plasma with FN2 will enhance the chemisorption probability of nitrogen ions, resulting in an increase of the N/W atomic ratio. The ion bombarding on the growing WNx film on the surface of Si substrate should also be considered. Wang et al. [52] have reported that N+ 2 ions undergo dissociative neutralization, as they impinge on the target, giving rise to two N atoms which are then backscattered. Moreover, some researchers [53] have showed that the N species which are backscattered from the target have a significantly higher energy than the backscattered Ar species. Therefore an increase of FN2 leads to an increase in the flux of both N+ 2 species and backscattered N

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Fig. 8. Intrinsic stress for WNx films as a function of (a) FN2 and (b) Vb, in which the solid line is just guided to eyes.

Fig. 9. Hardness for WNx films as a function of (a) FN2 and (b) Vb, in which the solid line is just guided to eyes.

species incident at the growing film. Furthermore, as revealed from XPS analysis, the nitrogen-related defects are observed for the film grown at Vb = −40 V, implying that ion implantation has taken place since the defects generation is closely related to the subplantation process [54]. Trapping of N+ 2 species and backscattered N species in subsurface sites should also be an important factor for increasing N/W atomic ratio. Baker et al. [20] have reported that a maximum nitrogen concentration of only about 35 at.% is obtained for sputtering WNx films without bias, and it is not possible to further increase the nitrogen concentration of the film via increasing FN2. At FN2 = 0 sccm, a single β-W phase is obtained. At FN2 = 5 sccm, a mixed β-W(N) and fcc-W2N phases are acquired. Note that the N concentration for the film deposited at FN2 = 5 sccm is lower than that of fcc-W2N. With increasing FN2 to 15 sccm, as the N concentration almost reaches that of fcc-W2N, a single phase fcc-W2N with good crystallinity appears. As FN2 increases to 30 sccm, the N/W atomic ratio in the film increases to ~0.8 and the fcc-W2N phase still remains. The XPS results show that the fraction of N1 is large at FN2 = 30 sccm, implying that the excess nitrogen could accommodate the remaining available octahedral interstitial positions, which also results in the out-of-plane lattice dilation. It is expected that if all octahedral interstitial positions are occupied by nitrogen atoms, the fcc WN phase will form. However, when FN2 increases to 60 sccm, the N/W atomic ratio is close to 1 and a phase transition from fcc-W2N to hexagonal WN takes place. Suetin et al. [55]

have investigated the cohesive energies and the elastic properties for hexagonal WN and fcc WN with stoichiometric composition using the first-principle calculations, wherein the cohesive energy of hexagonal WN is larger than that of fcc WN. Meanwhile for hexagonal WN, their five independent elastic constants Cij satisfy the well-known Born stability criteria: C11 N 0, (C11 − C12) N 0, C44 N 0, and (C11 + C12) C33 − 2C212 N 0, confirming that the phase is mechanically stable. On the contrary, for fcc WN, C44 b 0 does not satisfy the generalized criteria for mechanically stable crystals: (C11 − C12) N 0; (C11 + 2C12) N 0; C44 N 0. Therefore, as the N/W atomic ratio increase to ~1, fcc WN becomes unstable and transforms into hexagonal WN. Bubenzer et al. [56] have proposed that ion energy can be determined by Vb and total pressure during deposition in the following expression: E=

KVb Pm

; 0≤m≤1;

ð2Þ

where K is constant, P the total pressure, and m the exponent. As other parameters remain constant, the ion energy will be proportional to Vb. As ion energy increases, the enhanced collisionally induced dissociative chemisorptions of N2 and N implantation-related processes should favor an increase in N concentration in the film. However, the observed N/W atomic ratio monotonously decreases when the

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M. Wen et al. / Surface & Coatings Technology 205 (2010) 1953–1961

absolute value of Vb increases. Hence the concurrent N loss processes should dominate as the absolute value of Vb increases. Note that with increasing the absolute value of Vb, deposition rate obviously decreases due to film re-sputtering, preferential re-sputtering of N may occur. Palmer et al. [57] have previously reported preferential loss of N from stoichiometric TiN targets due to sputtering with 1– 5 keV Ar+. In this work, ion implantation has taken place even at Vb = −40 V. Meanwhile, during the subplantation process the energetic ions participating in cascade collision can dissipate their excess energy to the lattice in a thermalization stage, which results in an increase in the thermal energy of the local lattice atoms. As Vb increases, the number of displaced atoms and the thermal energy of the lattice atoms should increase. A threshold of Vb could exist, above which lattice atoms can obtain enough thermal energy to break W–N atomic bonds, and hence some nitrogen atoms can escape from WNx film. As shown in Fig. 2, as the absolute value of Vb is large, an escape of nitrogen can be observed. However, compared to previous work [58], for the NbN films deposited under the same condition, such elimination of nitrogen was not observed, which may be due to that NbN is more stable than W2N, since the heat of formation for NbN is −220 kJ/mol, while that of W2N is −72 kJ/mol. For the film grown at Vb =−40 V, both XRD and HRTEM results show that the film has fcc structure and good crystallinity. However, as the absolute value of Vb increases to 80 V, a mixture of W + W2N occurs and poor crystallinity is observed. The N concentration begins to decrease for the film deposited at Vb = −80 V, indicating that N atoms escape from the film. The threshold thermal energy of breaking W–N atomic bonds may be surpassed for the film grown at Vb = −80 V since the bonding energy between W and N is relatively low. As a consequence, nitrogendeficiency would result in forming the W phase. However, the limited loss of the nitrogen atoms in the film deposited at Vb = −80 V makes the grain growth of the W phase difficult. HRTEM (Fig. 7b) result shows that an amorphous structure appears for the film grown at Vb = −80 V, which may be explained that a local structure collapses into a disordered amorphous network when the concentration of defects induced by bombarding ions exceeds the allowable equilibrium concentration of the mobile defects. As the absolute value of Vb increases to 120 V, N concentration reduces significantly, which promotes the grain growth of the W phase. As the absolute value of Vb increases to 160 V, the N concentration becomes low, and only the W(N) solid solution phase is observed. For a mixed W + W2N phase, the relative concentration of either W or W2N may be controlled by accurately adjusting Vb. 4.2. Residual stress From Fig. 8, the W film, deposited in discharging a pure Ar gas, exhibits a tensile stress, while the other WNx films are under compression. According to the models proposed by Windischmann [59] and Davis [60], the intrinsic compressive stress in the film deposited under ionic or atomic bombardment originates from the “atomic peening” mechanism. The bombarding ions or atoms with energy larger than the threshold energy will implant into the subsurface of the growing film and result in the increased defects. It is well known that the compressive stresses are caused by both the direct subplantation of bombarding species in the films and the displacement of atoms at lattice in the film to interstitial positions [59,60]. Therefore, the entrapped atoms and high density of defects will dilate the lattice and thus induce compressive stress. In contrast, the tensile stresses are generated due to the shrinkage of the grain boundaries [61]. For the W film deposited at FN2 = 0 sccm, the tensile stress term caused by the shrinkage of the grain boundaries cannot be compensated by the compressive stress term caused by the atomic peening. The reason may be that at FN2 = 0 sccm, under the relatively low energetic bombardment condition, the Ar entrapment in the W film is difficult, and only limited Ar atoms are trapped into the

subsurface since Ar is inert. However, at FN2 = 5 sccm, some N atoms are implanted into the subsurface, to occupy the interstitial positions of β-W, resulting in out-of-plane lattice dilation, and leading to a compressive stress. With increasing FN2 to 15 sccm, although the W2N phase with a nearly stoichiometric composition is obtained, the high density of defects in WNx film are observed, which may be responsible for the observed high compressive stress value of 3.3 GPa. As FN2 increases to 30 sccm, an overstoichiometric W2N phase is formed, resulting in a further out-of-plane lattice dilation and increase in compressive stress. The compressive stress reaches the maximum value of 5.6 GPa as hexagonal WN is formed. With a further increase in FN2, the stress decreases. This partial relaxation of the stress can be attributed to the formation of the amorphous phase. Therefore, the stress in the film is strongly dependent on the N concentration of the film. An analytical stress model [60] based on shallow implantation shows that the compressive stress is related to the ion impinging energyEi. At lowEi, with increasingEi, the number of displaced atoms (defect) increases, and thereby giving rise to an increase of compressive stress. As the threshold energy is reached, the relaxation of point defects, e.g., annihilation of Frankel pairs or annihilation of interstitials at the surface, will play a dominate role, and compressive stress begins to decrease. In the reactive sputtering process, the energy of the ions, as mentioned in Eq. (2), is proportional to Vb. Therefore, it is expected that the compressive stress should firstly increase with increasing Vb, as the absolute value of Vb is low. However, since the N concentration changes and a phase transition also takes place, as Vb varies, this stress model can only discuss the dependence of stress on Vb qualitatively. When the absolute value of Vb increases to 80 V, the compressive stress decrease dramatically to 0.6 GPa. It is noted that the amorphous structure is present in the film grown at Vb =−80 V, which could cause the stress relaxation in the WNx film deposited at Vb =−80 V. As the film is well crystallized, the compressive stress increases again. Although N concentration decreases with increasing the absolute value of Vb from 80 to 160 V, the mechanism of defect creation still plays a dominant role in the evolution of the compressive stress. The compressive stress begins to decrease again in the film grown at Vb =−200 V, which can be attributed to the relaxation of defects since the films deposited at Vb =−160 and −200 V have the same N concentration in the well-crystallized films. 4.3. Hardness There are many factors that affect the hardness of coatings [62] such as phase configuration, microstructure, residual stress, and grain size. In the range of 0 ≤ FN2 ≤ 30 sccm, the trend of the variation of residual stress and grain size with FN2 is not always in accord with that of the variation of hardness with FN2. This means that the stress hardening and grain size strengthening do not dominate the evolution of hardness at 0 ≤ FN2 ≤ 30 sccm. In Fig. 8, the pure W film at FN2 = 0 sccm exhibits a relatively low hardness value of 14.4 GPa. However, as FN2 = 5 sccm, an increase in hardness is observed, which can be ascribed to the solid solution strengthening since some N atoms occupy the β-W interstitial positions and the W(N) structure is formed. With increasing FN2 to 15 sccm, only fcc-W2N is observed and the hardness arrives at 34.5 GPa. When an overstoichiometric fccW2N phase is obtained at FN2 = 30 sccm, the hardness decreases. Hones et al. [19] have also reported that in fcc W–N, the hardness decreases when the film is deposited under overstoichiometric conditions, and they ascribe the decrease in hardness to the change in the electronic structure. With increasing N concentration, the valence electron concentration increases [63], and this leads to nonbonding and antibonding electronic states being filled, which are responsible for a decrease of the cohesive energy. It has been reported [15,19] that in the W–N system, the hexagonal phase exhibits a higher hardness than the fcc phase. However, in this work, as the hexagonal phase is obtained at FN2 ≥ 60 sccm, its hardness is not high. Noted that

M. Wen et al. / Surface & Coatings Technology 205 (2010) 1953–1961

as a phase transition from fcc to hexagonal WN takes place, the grain size decreases significantly, and reaches even to ~5 nm. Hence the reverse Hall–Petch relation [42,64] may go into effect, leading to softening. As the absolute value of Vb increases to 80 V, the hardness decreases dramatically. This can be ascribed to the existence of the amorphous phase. With increasing the absolute value of Vb to 120 V, a mixed wellcrystallized W+ fcc-W2N structure emerges, and hardness increases to 38.9 GPa. In a mixed W+ fcc-W2N structure, the interface between W and fcc-W2N phase may block the dislocation movement and give rise to the enhanced hardness. Shih and Dove [65] have investigated the mechanical properties of W/WN multilayer films and found that the coatings with an alternating layer of pure metal and metal nitride exhibit higher hardness than the single-layer film. Upon further increasing the absolute value of Vb, only the W(N) phase exists, and its hardness begins to decrease again. The absence of a mixed structure and the decrease in grain size (~5 nm) should be responsible for the drop of the hardness. In addition, comparing with the hardness of the W(N) film deposited at FN2 =15 sccm, Vb =−200 and that deposited at FN2= 5 sccn, Vb =−40, it is observed that both of films have about the same N/W atomic ratio, but the hardness is different. The higher hardness value for the W(N) film deposited at FN2 =15 sccm, Vb = −200, compared to that deposited at FN2 =5 sccm, Vb =−40, may be ascribed to the densification of the microstructure caused by ion bombardment and higher stress. 5. Conclusion The WNx thin films can be deposited on Si(100) using DC reactive magnetron sputtering in discharging a mixture of N2 and Ar gas, and both FN2 and Vb have a significant influence on the composition, phase structure, intrinsic stress and hardness for the obtained films. Keeping Vb at −40 V, as FN2 increases, the N/W atomic ratio gradually increases, accompanied by a phase transition from β-W to hexagonal WN through fcc-W2N. On the other hand, the N/W atomic ratio gradually decreases with increasing the absolute value of Vb, resulting in an evolution from the fcc-W2N structure to β-W(N) through a mixture of fcc-W2N + β-W (N). This means that a high Vb can induce an increase in the thermal energy of the local lattice atoms, thereby leading to breaking W–N bonds and the elimination of nitrogen due to the relatively low heat of formation for WNx. As FN2 varies, an increase of the nitrogen concentration and a high density of defects, created by ion bombarding, are responsible for the stress evolution with FN2. In contrast, for the film grown at Vb = −80 V, HRTEM image reveals that an amorphous structure is formed, resulting in a drop to 0.6 GPa in stress. Further increasing the absolute value of Vb from 80 to 200 V, the competition between the defect creation and defect relaxation determines the evolution of the compressive stress. Furthermore, the maximum hardness of 38.9 GPa is obtained for the film deposited at Vb = −120 V as a mixed fcc-W2N + β-W(N) structure is formed, which can be ascribed to that the interface between W and fcc-W2N phase blocks the dislocation movement. Acknowledgments The support from National Natural Science Foundation of China (Grant Nos. 50525204 and 50832001), the special Ph.D. program (Grant No. 200801830025) from MOE, and the “211” and “985” project of Jilin University, China, is highly appreciated. M.W. would like to thank the Chinese Scholarship Council for financial support. References [1] L.E. Toth, Transition Metal Carbides and Nitrides, Academic, New York, 1971. [2] B.L. Park, J. Electron. Mater. 26 (1997) L1. [3] M. Takeyama, A. Noya, Jpn. J. Appl. Phys. 36 (1997) 2261.

1961

[4] M. Uekubo, T. Oku, K. Nii, M. Murakami, K. Takahiro, S. Yamaguchi, T. Nakano, T. Ohta, Thin Solid Films 286 (1996) 170. [5] F.C.T. So, E. Kolawa, X.A. Zhao, M.A. Nicolet, J. Appl. Phys. 64 (1988) 2787. [6] M.H. Tsai, H.T. Chiu, S.H. Chuang, Appl. Phys. Lett. 68 (1996) 1412. [7] K.M. Chang, T.H. Yeh, I.C. Deng, J. Appl. Phys. 81 (1997) 3670. [8] S. Bystrova, A.A.I. Aarnink, J. Holleman, R.A.M. Wolters, J. Electrochem. Soc. 152 (2005) G522. [9] C.W. Lee, Y.T. Kim, Solid State Electron. 38 (1995) 679. [10] A.E. Gisserger, R.A. Sadler, F.A. Leyenaar, M.L. Balzan, J. Vac. Sci. Technol., A 4 (1986) 3091. [11] M. Moriwaki, T. Yamada, Y. Harada, S. Fujii, Jpn. J. Appl. Phys. 39 (2000) 2177. [12] B. Claflin, M. Binger, G. Lucovsky, J. Vac. Sci. Technol., A 16 (1998) 1757. [13] J.W. Lee, C.H. Han, J.S. Park, J.W. Park, J. Electrochem. Soc. 148 (2001) G95. [14] P.C. Jiang, Y.S. Lai, J.S. Chen, Appl. Phys. Lett. 89 (2006) 122107. [15] T. Yamamoto, M. Kawate, H. Hasegawa, T. Suzuki, Surf. Coat. Technol. 193 (2005) 372. [16] T. Polcar, N.M.G. Parreira, A. Cavaleiro, Wear 262 (2007) 655. [17] T. Polcar, N.M.G. Parreira, A. Cavaleiro, Wear 265 (2008) 319. [18] L. Boukhris, J.M. Poitevin, Thin Solid Films 310 (1997) 222. [19] P. Hones, N. Martin, M. Regula, F. Levy, J. Phys. D Appl. Phys. 36 (2003) 1023. [20] C.C. Baker, S.I. Shaha, J. Vac. Sci. Technol., A 20 (2002) 1699. [21] P.C. Jiang, J.S. Chen, Y.K. Lin, J. Vac. Sci. Technol., A 21 (2003) 616. [22] Y.G. Shen, Y.W. Mai, Surf. Coat. Technol. 127 (2000) 239. [23] Z.L. Wang, Z.H. Liu, Z.P. Yang, S. Shingubara, Microelectron. J. 85 (2008) 395. [24] S.H. Mohamed, Surf. Coat. Technol. 202 (2008) 2169. [25] S. Guruvenket, G.M. Rao, Mater. Sci. Eng., B 106 (2004) 172. [26] Y.G. Shen, Y.W. Mai, D.R. McKenzie, Q.C. Zhang, W.D. McFall, W.E. McBride, J. Appl. Phys. 88 (2000) 1380. [27] B.S. Suh, H.K. Cho, Y.J. Lee, W.J. Lee, C.O. Park, J. Appl. Phys. 89 (2001) 4128. [28] T. Migita, R. Kamei, T. Tanaka, K. Kawabata, Appl. Surf. Sci. 169 (2001) 362. [29] M. Bereznai, Z. Toth, A.P. Caricato, M. Fernandez, A. Luches, G. Majni, P. Mengucci, P.M. Nagy, A. Juhasz, L. Nanai, Thin Solid Films 473 (2005) 16. [30] M. Eizenberg, F. Meyer, A. Benhocine, D. Bouchier, J. Appl. Phys. 75 (1994) 3900. [31] C.W. Lee, Y.T. Kim, S.K. Min, Appl. Phys. Lett. 62 (1993) 3312. [32] C. Meunier, C. Monteil, C. Savall, F. Palmino, J. Weber, R. Berjoan, J. Durand, Appl. Surf. Sci. 125 (1998) 313. [33] K.K. Lai, A.W. Mak, T.P.H.F. Wendling, P. Jian, B. Hathcock, Thin Solid Films 332 (1998) 329. [34] J.S. Jeng, S.H. Wang, J.S. Chen, J. Vac. Sci. Technol., A 25 (2007) 651. [35] S. Motojima, N. Ueshima, J. Alloys Compd. 393 (2005) 307. [36] D.J. Li, M.X. Wang, J.J. Zhang, J. Yang, J. Vac. Sci. Technol., A 24 (2006) 966. [37] B.K. Gan, M.M.M. Bilek, D.R. McKenzie, M.B. Taylor, D.G. McCulloch, J. Appl. Phys. 95 (2004) 2130. [38] A. Lahav, K.A. Grim, I.A. Blech, J. Appl. Phys. 67 (1990) 734. [39] T.P. Drüsedau, K. Koppenhagen, J. Bläsing, T.M. John, Appl. Phys. A: Mater. Sci. Process. 72 (2001) 541. [40] T.C. Li, B.J. Lwo, N.W. Pu, S.P. Yu, C.H. Kao, Surf. Coat. Technol. 201 (2006) 1031. [41] M. Wen, C.Q. Hu, Q.N. Meng, Z.D. Zhao, T. An, Y.D. Su, W.X. Yu, W.T. Zheng, J. Phys. D Appl. Phys. 42 (2009) 035304. [42] P.K. Huang, J.W. Yeh, J. Phys. D Appl. Phys. 42 (2009) 115401. [43] C.D. Wagner, W.M. Riggs, L.E. Davis, J.F. Moulder, G.E. Mulinberg, Handbook of X-ray Photoelectron Spectroscopy, Perkin-Elmer Corporation, USA, 1979. [44] T. Nakajima, K. Watanabe, N. Watanabe, J. Electrochem. Soc. 134 (1987) 3175. [45] O.H. Gokcea, S. Amina, N.M. Ravindra, D.J. Szostakb, R.J. Paffb, J.G. Flemingc, C.J. Galewskid, J. Shallenbergere, R. Eby, Thin Solid Films 353 (1999) 149. [46] D.H. Kuo, C.H. Shih, Appl. Phys. Lett. 93 (2008) 164105. [47] K.S.A. Butcher, A.J. Fernandes, P.P.T. Chen, M. Wintrebert-Fouquet, J. Appl. Phys. 101 (2007) 123702. [48] Y.M. Chiang, D. Birnie, W.D. Kingery, Physical Ceramics: Principles for Ceramic Science and Engineering, Wiley, New York, 1997 Chap. 2. [49] H. Wulff, C. Eggs, J. Vac. Sci. Technol., A 15 (1997) 2938. [50] G. Abadias, Y.Y. Tse, P. Guérin, J. Appl. Phys. 99 (2006) 113519. [51] I. Petrov, A. Myers, J.E. Greene, J.R. Abelson, J. Vac. Sci. Technol., A 12 (1994) 2846. [52] Z. Wang, S.A. Cohen, D.N. Ruzic, M.J. Goeckner, Phys. Rev. E 61 (2000) 1904. [53] K. Sarakinos, J. Alami, P.M. Karimi, D. Severin, M. Wuttig, J. Phys. D Appl. Phys. 40 (2007) 778. [54] H. Ljungcrantz, L. Hultman, J.E. Sundgren, L. Karlsson, J. Appl. Phys. 78 (1995) 832. [55] D.V. Suetin, I.R. Shein, A.L. Ivanovskii, Phys. Status Solidi B 245 (2008) 1590. [56] A. Bubenzer, B. Dischler, G. Brandt, P.J. Koidl, Appl. Phys. 54 (1983) 4590. [57] W. Palmer, M. Huttinger, W. Bench, Thin Solid Films 174 (1989) 143. [58] M. Wen, C.Q. Hu, C. Wang, T. An, Y.D. Su, Q.N. Meng, W.T. Zheng, J. Appl. Phys. 104 (2008) 023527. [59] H. Windischmann, J. Appl. Phys. 62 (1987) 1800. [60] C.A. Davis, Thin Solid Films 226 (1993) 30. [61] R.W. Hoffman, Thin Solid Films 34 (1976) 185. [62] D.V. Shtansky, E.A. Levashov, A.N. Sheveiko, J.J. Moore, Metall. Mater. Trans. A 30 (1999) 2439. [63] H. Holleck, J. Vac. Sci. Technol., A 4 (1986) 2661. [64] L.U. Chunsheng, Y.W. Mai, Y.G. Shen, J. Mater. Sci. 41 (2006) 937. [65] K.K. Shih, D.B. Dove, Appl. Phys. Lett. 61 (1992) 654.