- Email: [email protected]
Effect of oxygen on residual stress and structural properties of tungsten nitride films grown by reactive magnetron sputtering
Materials Science and Engineering B76 (2000) 107 – 115 www.elsevier.com/locate/mseb
Effect of oxygen on residual stress and structural properties of tungsten nitride films grown by reactive magnetron sputtering Y.G. Shen *, Y.W. Mai Centre for Ad6anced Materials Technology (CAMT), Department of Mechanical and Mechatronic Engineering, Uni6ersity of Sydney, Sydney, NSW 2006, Australia Received 5 January 2000; accepted 28 January 2000
Abstract Thin films of W –O–N were produced by reactive d.c. magnetron sputtering of tungsten in an Ar – N2 –O2 gas mixture. The effects of oxygen incorporation on the residual stress and structural properties of these films as well as the influence of post-deposition annealing have been studied. The films were analyzed in situ by a cantilever beam technique, and ex situ by X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), transmission electron microscopy (TEM), and electron energy-loss spectroscopy (EELS). It was found that the stoichiometric W2N films deposited under oxygen-free conditions had a high compressive stress of 1.45 GPa. The compressive stress in W–O–N films decreased significantly with an increase in the oxygen concentration and became slightly tensile for films near 10 – 15 at.% oxygen. These results can be ascribed to the decrease in the lattice parameter caused by incorporating small oxygen atoms in the lattice sites and the development of an amorphous network in the W–O–N films as the incorporation of oxygen was increased. By high temperature annealing the structural conversion from W2N to W in oxygen-free films was observed using XRD and the microstructure evolution after conversion was demonstrated using cross-sectional TEM. The effect of oxygen in stabilizing the W2N structure was also elucidated and explained on the basis of structural and thermodynamic stability. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Oxygen; Residual stress; Structural properties; Tungsten nitride films; Reactive magnetron sputtering
1. Introduction Transition metal nitrides are well known for their extreme hardness, high melting point, chemical inertness and good thermodynamic stability [1]. For these reasons, they are used for diffusion barriers in microelectronics, electrodes in semiconductor devices, and protective coatings in industry. Tungsten nitride films are of particular interest and have emerged as a very promising candidate for use in all of the above mentioned applications. However, the preparation and investigation of tungsten nitrides [2 – 6] are not as intensive as other transition metal nitrides such as titanium nitrides. It has been observed that titanium nitride films form a better diffusion barrier between silicon and alu-
* Corresponding author. Tel.: + 61-2-93517142; fax: +61-293517060. E-mail address: [email protected] (Y.G. Shen)
minium when the film is exposed to air before the aluminium deposition [7–9], or when the film contains significant amounts of oxygen [10,11]. Mandl et al. [12] comments that these oxygen impurities stuffed grain boundaries and blocked the rapid diffusion paths between aluminium and silicon, thus improving the thermal stability of diffusion barriers. However, the effect of oxygen on the properties of tungsten nitride films remains almost totally unexplored. In this paper we investigate the effect of oxygen on the residual stress and structural properties of W–O–N thin films deposited by reactive magnetron sputtering of tungsten in an Ar–N2 –O2 gas mixture. This work was motivated by our concern that oxygen incorporation and residual stress in W–O–N films may have a great influence on their structural properties, and our hope that a detailed understanding of the relationship between residual stress and microstructure in the films would be useful in optimizing these materials for diffusion barrier applications. The evolution of the film
0921-5107/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 1 0 7 ( 0 0 ) 0 0 4 2 4 - 4
108
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
stress was characterized by in situ real time stress measurement. The atomic concentration (at.%) of elements was determined by X-ray photoelectron spectroscopy (XPS). The phase of the films was identified by X-ray diffraction (XRD). The microstructure of the films was demonstrated using cross-sectional transmission electron microscopy (TEM). The nitrogen-loss behavior of the films during heating was characterized in situ by electron energy-loss spectroscopy (EELS).
2. Experimental details
2.1. Film deposition Thin films of W2N and W – O – N were grown on unheated Si(100) substrates in a reactive d.c. magnetron sputtering system at a base pressure of 2×10 − 4 Pa, which has been described previously [13,14]. A currentregulated d.c. supply was used to provide a discharge current of 280 mA and a target voltage of 300 V with a target-to-substrate separation of 70 mm. Prior to deposition, the target was sputter cleaned for 10 min while the substrates were isolated from the plasma by a shutter. The tungsten disk target, which is 50 mm in diameter and 10 mm in thickness with 99.99% purity, was sputtered in either an Ar – N2 or an Ar – N2 –O2 gas mixture to deposit W2N or W – O – N films. The sputtering conditions were first adjusted at a constant total pressure of 0.8 Pa at N2 and Ar flow rates of 4.0 and 10.2 sccm, respectively, to obtain a single-phase W2N, and then O2 was added to the gas mixture for depositing W –O –N films. The ratio Of O2 to Ar+N2 could be changed, so as to vary the oxygen concentration between 0 and 18 at.%. In all cases, the films with oxygen-free conditions are stoichiometric with W/N ratios, determined by XPS, of 2.0 90.03. Consequently, the stoichiometry of all oxygen-free nitride films is hereafter referred to as W67N33. Steady-state operation and reproducible stoichiometries on the target were ensured using a sequence of low-to-high N2 concentrations. The thickness of the films was determined using a stylus-type surface profilometer (Tencor P-10) and confirmed by cross-sectional TEM measurements. Substrate temperature during film growth was estimated to be 5200°C (accounting for plasma heating). These samples will be denoted as as-deposited samples. Unless otherwise illustrated, post-deposition annealing of the samples was carried out in a vacuum of better than 6 ×10 − 5 Pa up to 900°C. Annealing duration time for each temperature cycle was 30 min.
2.2. Film characterization The films were analyzed by XPS using a Kratos Axis/800 hemispherical energy analyzer equipped with
an unmonochromatized Mg Ka X-ray source (hv = 1253.6 eV). The binding energy scale was calibrated against Ag 3d5/2 at 368.25 eV and the instrumental energy resolution, measured using the same peak, was 0.9 eV (pass energy set at 20 eV). Charging corrections were made using the as measured Fermi level of Ag and Au pieces on the metallic substrate holder using the Mg Ka radiation as a reference. The elemental concentration of the films was derived from the area under each elemental spectrum using published sensitivity factors [15]. The crystallographic structure of the films was determined by XRD using a Siemens D5000 diffractometer operated at 40 kV and 30 mA. The measurements were carried out using Cu Ka radiation with a Ni filter to remove Cu Kb reflections, a powder diffractometer scanned over the range up to 150° 2u at steps of 0.01°, a 0.5° divergence slit, and a 0.2° receiving slit. The films were also characterized by EELS in a Phillips EM430 TEM equipped with a GATAN 666 parallel electron energy-loss spectrometer. All electron energy-loss measurements were done in diffraction mode, 300 keV electrons, and a collection angle of approximately 8 mrad. A Phillips CM12 microscope with a LaB6 filament operated at 120 kV was used for plan-view TEM analyses. The cross-sectional TEM images were obtained with a Phillips EM430 microscope operated at 300 kV. Plan-view specimens were prepared by mechanical thinning from the backside to 30 mm followed by 3 keV Ar+-ion milling at 8° with respect to the specimen surface. Cross-sectional specimens were prepared by gluing two samples face-to-face with epoxy resin and then mechanically thinning, followed by ion-beam milling in a manner similar to that of the plan-view samples, but from both sides.
3. Results
3.1. In situ stress measurements The development of film stress during deposition was determined with a vacuum cantilever beam apparatus using an optical two-beam deflection method to detect changes in the substrate curvature. The film stress s is obtained by the well-known Stoney’s equation revised for biaxial stress: s=
n
1 Esh 2s 1 1 − 6 (1−ns)hf r r0
(1)
where Es(1 −ns) is the biaxial modulus of the substrate, hs and hf are the thicknesses of the substrate and film, and r0, and r denote the curvature radii before and after film deposition, respectively. The substrate curvature was measured on a beam-shaped sample by deflection
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
109
of a split laser beam from both ends. The sample stripes were 2 mm wide and 18 mm long. They were cut from the Si(100) wafers (hs =85 mm, Es/(1 −ns) =180.5 GPa). The in situ measurement reveals how the oxygen incorporation influences the formation of residual stresses in W –O – N films. Fig. 1 shows representative examples of the instantaneous film force as a function of the mean film thickness. The slope of the force curves corresponds to the incremental film stress, with the negative slope being compressive. The W67N33 film without oxygen dosing shows an increasing instantaneous compressive stress with film thickness. A remarkable evolution of the instantaneous stress was observed in the film containing 10 at.% oxygen, where a state of tensile instantaneous stress develops with increasing film thickness. At a high oxygen concentration, approximately 18 at.% oxygen, the film is almost free of residual stresses up to 200 nm. Fig. 2 shows the evolution of the stress s for as-deposited 150-nm-thick W – O – N films as a function of the oxygen concentration. It is noted that the film deposited without oxygen dosing also contains approximately 1 at.% oxygen. Close examination of the deposi-
Fig. 2. Residual stresses of as-deposited 150-nm-thick nitride as a function of oxygen concentration. The inset shows the TEM brightfield micrograph with selected-area electron diffraction pattern of the film containing 18 at.% oxygen.
Fig. 1. Representative examples of the film force versus mean film thickness in as-deposited W –O–N films with several different oxygen concentrations.
tion processing reveals that the most likely source of oxygen contamination was either from residual oxygen in the deposition chamber or from the tungsten sputter target under the present experimental conditions. In our study, the total Ar+ N2 pressure and the power density were fixed, and the only variable was the O2 partial pressure. No delamination of the films was observed in all cases. It is known that the residual stress in thin films is composed of two parts, i.e. the thermal stress and the intrinsic stress. It seems reasonable to assume that the thermal stress can be neglected under the present conditions because the contribution of the thermal (tensile) stress to the total stress for 150-nmthick films at growth temperature of approximately 150°C did not exceed 10%. It is seen that the W67N33 film (containing approximately 1 at.% oxygen) is under high compression (1.45 GPa). The stress decreases significantly with an increase in the oxygen concentration. The film stress changes from compression to tension at 10 at.% oxygen, where the lattice parameter of the film comes slightly below the bulk value of W2N, 0.4126 nm [16] (see Fig. 3). With further increase in oxygen concentration the film is almost free of residual stresses.
110
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
This full relaxation of stresses may, therefore, be understood in terms of the development of a disordered amorphous network as the oxygen incorporation is increased. The amorphous nature of the films at high oxygen incorporation was evidenced by TEM measurements. The inset of Fig. 2 shows the typical bright-field (first diffused ring paraxial) image of the film containing 18 at.% oxygen and the selected area electron diffraction pattern obtained from this area. The image reveals the largest discernible details being smaller than approximately 1– 2 nm, indicating an amorphous structure. The observation was further supported by the fact that the dark-field image in TEM exhibits no microcrystallites (not shown).
3.2. Film structure and XPS depth profiles Our XRD measurements indicate that the crystallographic structure of sputter-deposited W67N33 (nominally 150 nm in thickness) is W2N in structure. The XRD u –2u scan data in the Bragg – Brentano geometry show a linear decrease in the lattice parameter of as-deposited films from 0.4220 to 0.4116 nm (calculated from the (111) diffraction peak position) while the oxygen concentration increases from 1 to 10 at.% as
Fig. 4. Typical XPS (a) N 1s and (b) O 1s spectra measured from the W56O15N29 film at a depth of approximately of 8 nm before and after annealing to 800°C. The curves have been offset vertically for clarity.
Fig. 3. Variation of lattice parameters in W –O–N films before and after annealing to 700°C as a function of oxygen concentration.
shown in Fig. 3. The observation that the lattice parameter of oxygen-contained films shifts toward lower values as the oxygen concentration increases indicates the contraction of the lattice by the incorporated oxygen atoms. This can be explained as increasing the incorporation of oxygen into the lattice sites without the formation of detectable amounts of W–O phases. When the oxygen concentration reaches 15 at.% or above, amorphous films are formed. Fig. 3 also shows that, after annealing to 700°C, the lattice parameter approaches to the reported bulk value
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
Fig. 5. XPS depth profiles from an as-deposited 150-nm-thick W56O15N29 film.
of 0.4126 nm, indicating that the layers are nearly fully relaxed. It should be pointed out that we could not detect any loss of both nitrogen and oxygen from any of the films after vacuum annealing up to 800°C. XPS N 1s and O 1s spectra shown in Fig. 4 measured on as-deposited films and after annealing unambiguously reveal this fact. XPS depth profile measurements from oxygen-contained films provide evidence of oxygen enrichment in the near-surface region during deposition. Fig. 5 shows typical depth profiles from a 150-nm-thick W56O15N29 film. In order to avoid oxygen contamination from air exposure, the film was capped in situ with 10-nm-thick carbon overlayers prior to air exposure. It can be seen from Fig. 5 that the profile shows a high concentration of oxygen near the surface. Below this surface layer, however, the oxygen concentration is reduced to a steady-state value of about 13 – 15 at.%. This can be explained in terms of original oxygen enrichment in the near-surface region. The shape of the oxygen profile can be understood as follows. When a new nitride monolayer is deposited on a surface partially covered with oxygen, a significant fraction of the adsorbed oxygen can diffuse to the new surface. The diffusion of oxygen toward the near-surface region can be explained by a so called floating phenomenon [17,18]. The floating oxygen reduces the overall energy of the film/substrate composite by lowering the surface free energy. Thus with a continuous arrival of oxygen to the surface of the film, oxygen retained in the film will show a monotonic increase toward the surface as shown in Fig. 5.
3.3. XRD, TEM and EELS Example plots of XRD u – 2u scans exhibiting the structural changes in an oxygen-free W67N33 film after
111
high temperature annealing are shown in Fig. 6. The inset shows the XRD pattern obtained from the as-deposited film. At a low growth temperature atomic displacement, local atomic arrangements, and enhancement of adatom mobilities are limited. Thus, the (111) and (200) diffraction peaks observed in XRD u–2u measurements are broader and their intensities are relatively weak, indicating that the degree of ordering for the as-deposited W67N33 film is lower, probably due to small crystallite size and distortion of the fcc lattice structure. After annealing to 700°C, adatom mobilities in the film are sufficiently high to form a well ordered single-phase structure of W2N. The major peaks observed at angles 2u equal to 37.68, 43.70, 63.76, 76.12, and 80.16°. These peaks are assigned to (111), (200), (220), (311), and (222) reflections of fcc W2N This structure is thermally stable and all of W2N reflection peaks remain up to 800°C. After annealing to 825°C, the W(110) reflection peak appears at 40.3°, indicating that W2N has partially converted to W. With further increase in annealing temperature, the ratio of W to W2N increases. After annealing to 900°C, the full conversion from W2N to W in the film is observed.
Fig. 6. XRD u – 2u scans from the 150-nm-thick W67N films after high temperature annealing. The fresh film deposited simultaneously was used for each temperature annealing. The inset shows the XRD pattern from the as-deposited film.
112
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
some fine W2N grains. This is in good agreement with the XRD results shown in Fig. 6, where the film has been converted into a two-phase structure consisting of both W2N and bcc W after 850°C annealing. For the bcc W structure of the film formed after 900°C annealing (Fig. 7c), we see that this dense structure is still columnar in nature, but it does not exhibit the voids. This is indicative of the presence of large crystallites. It is also noted that alignment of the columns in W is less pronounced than that in W2N. Fig. 8 shows a detailed XRD pattern for 2u angles up to 150° obtained from a W56O15N29 film after annealing to 900°C. The as-deposited W56O15N29 film is amorphous in structure (see the inset). When heated at 650°C or above, a well ordered single-phase structure of W2N appears (not shown). This means that crystallization had progressed. The fact that no other phase except W2N was observed after annealing is due to no detectable phase of W–O formed for oxygen-contained films by annealing. The W2N structure of W56O15N29 is thermodynamically very stable and all of W2N diffraction peaks still remain up to 900°C. These observations are in good agreement with the results reported by Lin et al. [19], where W2N was stable up to 900°C if a thin WO3 film was formed above W2N. It is also noted that a very broad peak is observed at low 2u angles. This may be ascribed to the contributions from the surface W–O overlayers.
Fig. 7. TEM cross-sectional micrographs of theW67N33 films after annealing to three different temperatures. (a) 700°C, (b) 850°C, and (c) 900°C. All the images have the same scale. The films are those used for Fig. 6.
In order to investigate the microstructure evolution during the structural conversion, the cross-sectional TEM measurements were carried out from the same samples used in Fig. 6. Typical TEM micrographs are shown in Fig. 7. Columnar microstructure was observed in the crystailineW67N33film after annealing to 700°C (Fig. 7a). It contains a well-defined columnar microstructure with uniform columns from substrate to the top of the films. The column boundaries appear to be decorated by a relatively low density of voids. The average column diameter is about 15 – 20 nm. After annealing to 850°C, the microstructure is also a columnar with voids decorating the column boundaries (Fig. 7b). The average column diameter becomes larger, on the order of 30– 40 nm. It is clear that most of the microstructure for the film after annealing to 850°C are composed of large course bcc W grains together with
Fig. 8. XRD u– 2u scans from a 150-nm-thick W56O15N29 film after annealing to 900°C. The inset shows the XRD pattern from the as-deposited film.
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
Fig. 9. EELS N K-edge intensity as a function of sample temperature for W67N33 and W56O15N29. From each point, data were taken during in situ heating at a heating rate of 20°C min − 1, held for 5 min and then the N K-edge spectrum was taken. Error bars are based on 4–6 measurements.
Measurements using EELS and N K-edge loss spectra, which is more sensitive to the bulk [20], show that the signal intensity of N K-edge for W56O15N29 is much stronger than that for W67N33 as the sample temperature increases to 900°C (Fig. 9). This result directly indicates that nitrogen in the oxygen-free W67N33 film evaporates to vacuum after high temperature annealing, in agreement with the XRD and TEM measurements. A detailed description of EELS and K-edge spectra analyses is covered in our previous paper [13].
4. Discussion The stress-versus-oxygen concentration curve was obtained after repeated deposition over many days (Fig. 2). The results are reproducible. The compressive stress observed at a low working-gas pressure was attributed to the bombardment of the growing film surface by energetic particles such as sputtered atoms and working-gas atoms or ions through the atomic peening process [21]. Thus the crystal lattice is deformed and excess atoms are forced into interstitial sites to expand the crystal lattice, resulting in compressive stress in the films. The stress became smaller by introducing a small amount of O2 in the sputtering chamber. It can be suggested that the stress depends mainly on the oxygen content in the W – O – N films. It seems reasonable to assume that some of the oxygen atoms replace nitrogen to constitute the crystal lattice, resulting in the decrease in compressive stresses. At higher oxygen concentrations the structure collapses into a completely disordered amorphous network, leading to the full relaxation of stresses. This relaxation of compressive stresses may, therefore, be understood in terms of a loss
113
in intercolumnar coupling as a result of increased porosity in the grain boundaries. The kinetics of growth is controlled by the mobility of the impinging atoms on the surface before they condense and become entrapped in the film. This mobility can be enhanced by increasing temperature or by supplying impact energy through ion bombardment. For deposition at a low pressure of 0.8 Pa, pure reactively sputtered films of W67N33 grew in crystalline form. This is in agreement with previous reported results [4,5], where the crystallite structure favors sputterdeposited films whose composition is in the vicinity of W2N. The lattice parameter for as-deposited W67N33 films is higher than the corresponding bulk value of 0.4126 nm. A continuous shift of the lattice parameter toward lower values with increasing oxygen concentration is also evidenced (Fig. 3). The crystal structure of W2N is a B1–NaCl type where the W atoms occupy the positions of the fcc lattice points and N atoms occupy 50% of the total octahedral sites. In the case of oxygencontained films, some of the incorporated oxygen could possibly be occupying these nitrogen sites, resulting in a contraction of the lattice. This is clarified by the fact that the covalent radius of oxygen (0.066 nm) is smaller than that of nitrogen (0.070 nm). However, the possibility of the presence of oxygen impurities at the interstitial sites and grain boundaries also cannot be completely ruled out. As the oxygen is introduced in the deposition chamber, the mobility of both W and N particles decreases. At higher O2 partial pressures the incorporation of oxygen in the films is significant, resulting in the amorphous growth. We suggest that the oxygen impurities incorporated in the film can prevent crystallization growth in several ways: Oxygen can act as a roadblock to the diffusing W and N or can trap diffusing W and N, or can serve as a nucleation site for lattice defects. Therefore, the growing W and N particles may not be sufficiently mobile to migrate to the preferred sites for crystallization growth. When the concentration of permanent defects induced by the incorporated oxygen exceeds the equilibrium concentration of the mobile defects, failure of crystallization growth occurs and amorphization begins to take place. In addition, the decreasing grain size, observable with increasing O2 partial pressure during deposition can be explained in terms of oxygen enrichment at the grain surface, thus preventing further growth and establishing new grains. The oxygen is present as a W–O oxide, since the O1 s binding energy is approximately 530 eV (Fig. 4b). This leads to the picture that the incorporated oxygen mainly bonds to W, eventually leading to the partial formation of tungsten oxide. From thermodynamic considerations the formation of a WOx oxide is energetically strongly favored over the formation of a NO layer because the heat of formation for WO3 (−842.9
114
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
M mol − 1) and for WO2 ( −589.7 M mol − 1) is much larger than the corresponding value for NO (91.3 M mol − 1) [22]. Such an oxide formation is favored by the more stable WO bonding. However, the present XPS data provide no distinction between WO2 or WO3, or WO. The annealing used in the present investigation presumably serves only to activate atom diffusion, phase ordering, and structure conversion [23,24]. It is of interest to determine the thermal range of structure stability for both oxygen-free and oxygen-contained nitride films. It was found that the W2N structure of W67N33 was stable up to 800°C (Fig. 6), whereas the W2N structure of W56O15N29 was thermodynamically stable up to 900°C (Fig. 8). This can be explained in terms of original oxygen enrichment at the surface (Fig. 5) as well as oxygen species in the W2N matrix in the oxygencontained films. With increasing temperature the oxygen atoms move increasingly outward to the near-surface region to form thicker W – O (concentration dominant) overlayers (not shown here). This can be obviously ascribed to two causes: first, oxygen atoms are minority impurities; and second, the W – O bond is more stable than the WN bond. Both effects favor oxygen segregation. In addition, part of the bulk oxygen during annealing may also be possible to diffuse into the grain boundaries. We believe that thin W–O overlayers in the near-surface region and oxygen species in the W2N grain boundaries can be helpful for preventing the desorption of nitrogen from W2N matrix. Based on the results obtained and the above discussion, the thermal stability of the W2N phase is believed to be governed by two main trends: (1) From an energetic point of view, the bonding energy required to stabilize the W2N structure must be overcome in order for breaking WN bonds to occur. (2) From a structural point of view, the transformation to bcc W occurs by desorption of nitrogen and localized rearrangement of the lattice. No long range diffusion or atomic movement is required. Further study with ab initio total-energy calculations for establishing the origin of the stable W–O and W–N bonding structures is in progress in our Materials Research Center.
5. Conclusions The effect of oxygen on residual stress and structural properties of tungsten nitride thin films prepared by reactive d.c. magnetron sputtering in an Ar – N2 –O2 gas mixture has been investigated. Film stresses were analyzed in situ by a cantilever beam technique during the film formation. The films were also characterized ex situ by using XRD, XPS, TEM, and EELS. To the authors’ knowledge, this is the first report on sputter deposited W –O –N films.
The film stress was found to strongly depend on the oxygen incorporation in the W–O–N films under the present experimental conditions. The oxygen-free films had high compressive stresses. As the oxygen concentration was increased, the stress in the films decreased markedly and became slightly tensile for films near 10–15 at.% oxygen. It is suggested that the relaxation of the compressive stress is due to the decrease in the lattice parameter caused by incorporating oxygen atoms in the lattice sites (since the oxygen atom is smaller than the nitrogen atom) and the development of an amorphous network in the W–O–N films as the incorporation of oxygen was increased. The role of oxygen in stabilizing the W2N structure was also elucidated herein by XRD, XPS, TEM, and EELS. The oxygen-contained nitride films are thermodynamically more stable than the oxygen-free films. We believe that thin W–O overlayers in the near-surface region and oxygen species in the grain boundaries can be helpful for preventing the desorption of nitrogen from W2N matrix during high temperature annealing. We hope that our new results will stimulate further experimental as well as theoretical studies.
Acknowledgements The authors gratefully acknowledge the financial support of this work by the Australian Research Grants Scheme. The authors also appreciate the use of the facilities in the Australian Key Center for Microscopy and Microanalysis, University of Sydney, which is supported by the Australia Research Council.
References [1] L.E. Toth, Transition Metal Carbides and Nitrides, Academic Press, New York, 1971. [2] S.D. Marcus, R.F. Foster, Thin Solid Films 263 (1993) 330. [3] K.M. Chang, T.H. Yeh, I.C. Deng, J. Appl. Phys. 81 (1997) 3670. [4] F.C.T. So, E. Kolawa, X.A. Zhao, M.A. Nicolet, Thin Solid Films 153 (1988) 507. [5] L. Boukhris, J.M. Poitevin, Thin Solid Films 310 (1997) 222. [6] H.T. Chiu, S.H. Chuang, J. Mater. Res. 8 (1993) 1353. [7] W. Sinke, G.P.A. Friflink, F.W Saris, Appl. Phys. Lett. 47 (1985) 471. [8] N. Kumar, K. Pourrezaei, B. Lee, E.C. Douglas, Thin Solid Films 164 (1988) 417. [9] P. Jin, S. Maruno, Jpn. J. Appl. Phys. 30 (1991) 2059. [10] K.J. Huber, C.R. Aita, J. Vac. Sci. Technol. A6 (1988) 1717. [11] H.G. Tompkins, J. Appl. Phys. 70 (1991) 3876. [12] M. Mandi, H. Hoffrnann, P. Kucher, J. Appl. Phys. 68 (1990) 2127. [13] Y.G. Shen, Y.W. Mai, Q.C. Zhang, D.R. McKenzie, W.D. FcFall, W.E. McBride, J. Appl. Phys. 87 (2000) 177. [14] Y.G. Shen, Y.W. Mai, W.E. McBride, D.R. McKenzie, Q.C. Zhang, Appl. Phys. Lett. 75 (1999) 2211.
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115 [15] J.F. Moulder, W.F. Stichle, P.E. Sobol, K.D. Bomben Jr, in: J. Chastain, R.C. King Jr (Eds.), Handbook of X-Ray Photoelectron Spectroscopy, Physical Electronics, Eden Prairie, MN, 1995. [16] M. Bauccio Jr (Ed.), ASM Metals Reference Book, American Society for Metals, Metals Park, OH, 1994. [17] W.F. Egelhoff Jr, D.A. Steigerwald, J. Vac. Sci. Technol. A7 (1989) 2167. [18] Y.G. Shen, D.J. O’Connor, H. van Zee, K. Wandelt, R.J. MacDonald, Thin Solid Films 263 (1995) 72. [19] J. Lin, A. Tsukune, T. Suzuki, M. Yamada, J. Vac. Sci. Technol.
.
115
A16 (1998) 611. [20] R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, 2nd edn, Plenum, New York, 1996. [21] J.A. Thornton, D.W. Hoffman, Thin Solid Films 171 (1989) 5. [22] D.R. Lide (Ed.), CRC Handbook of Chemistry and Physics, 72nd edn, CRC Press, Boca Raton, FL, 1991/1992. [23] Y.G. Shen, J. Yao, D.J. O’Connor, B.V. King, R.J. MacDonald, Phys. Rev. B56 (1997) 9894. [24] Y.G. Shen, D.J. O’Connor, J. Yao, Appl. Surf. Sci. 125 (1998) 300.
.
Effect of oxygen on residual stress and structural properties of tungsten nitride films grown by reactive magnetron sputtering Y.G. Shen *, Y.W. Mai Centre for Ad6anced Materials Technology (CAMT), Department of Mechanical and Mechatronic Engineering, Uni6ersity of Sydney, Sydney, NSW 2006, Australia Received 5 January 2000; accepted 28 January 2000
Abstract Thin films of W –O–N were produced by reactive d.c. magnetron sputtering of tungsten in an Ar – N2 –O2 gas mixture. The effects of oxygen incorporation on the residual stress and structural properties of these films as well as the influence of post-deposition annealing have been studied. The films were analyzed in situ by a cantilever beam technique, and ex situ by X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), transmission electron microscopy (TEM), and electron energy-loss spectroscopy (EELS). It was found that the stoichiometric W2N films deposited under oxygen-free conditions had a high compressive stress of 1.45 GPa. The compressive stress in W–O–N films decreased significantly with an increase in the oxygen concentration and became slightly tensile for films near 10 – 15 at.% oxygen. These results can be ascribed to the decrease in the lattice parameter caused by incorporating small oxygen atoms in the lattice sites and the development of an amorphous network in the W–O–N films as the incorporation of oxygen was increased. By high temperature annealing the structural conversion from W2N to W in oxygen-free films was observed using XRD and the microstructure evolution after conversion was demonstrated using cross-sectional TEM. The effect of oxygen in stabilizing the W2N structure was also elucidated and explained on the basis of structural and thermodynamic stability. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Oxygen; Residual stress; Structural properties; Tungsten nitride films; Reactive magnetron sputtering
1. Introduction Transition metal nitrides are well known for their extreme hardness, high melting point, chemical inertness and good thermodynamic stability [1]. For these reasons, they are used for diffusion barriers in microelectronics, electrodes in semiconductor devices, and protective coatings in industry. Tungsten nitride films are of particular interest and have emerged as a very promising candidate for use in all of the above mentioned applications. However, the preparation and investigation of tungsten nitrides [2 – 6] are not as intensive as other transition metal nitrides such as titanium nitrides. It has been observed that titanium nitride films form a better diffusion barrier between silicon and alu-
* Corresponding author. Tel.: + 61-2-93517142; fax: +61-293517060. E-mail address: [email protected] (Y.G. Shen)
minium when the film is exposed to air before the aluminium deposition [7–9], or when the film contains significant amounts of oxygen [10,11]. Mandl et al. [12] comments that these oxygen impurities stuffed grain boundaries and blocked the rapid diffusion paths between aluminium and silicon, thus improving the thermal stability of diffusion barriers. However, the effect of oxygen on the properties of tungsten nitride films remains almost totally unexplored. In this paper we investigate the effect of oxygen on the residual stress and structural properties of W–O–N thin films deposited by reactive magnetron sputtering of tungsten in an Ar–N2 –O2 gas mixture. This work was motivated by our concern that oxygen incorporation and residual stress in W–O–N films may have a great influence on their structural properties, and our hope that a detailed understanding of the relationship between residual stress and microstructure in the films would be useful in optimizing these materials for diffusion barrier applications. The evolution of the film
0921-5107/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 1 0 7 ( 0 0 ) 0 0 4 2 4 - 4
108
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
stress was characterized by in situ real time stress measurement. The atomic concentration (at.%) of elements was determined by X-ray photoelectron spectroscopy (XPS). The phase of the films was identified by X-ray diffraction (XRD). The microstructure of the films was demonstrated using cross-sectional transmission electron microscopy (TEM). The nitrogen-loss behavior of the films during heating was characterized in situ by electron energy-loss spectroscopy (EELS).
2. Experimental details
2.1. Film deposition Thin films of W2N and W – O – N were grown on unheated Si(100) substrates in a reactive d.c. magnetron sputtering system at a base pressure of 2×10 − 4 Pa, which has been described previously [13,14]. A currentregulated d.c. supply was used to provide a discharge current of 280 mA and a target voltage of 300 V with a target-to-substrate separation of 70 mm. Prior to deposition, the target was sputter cleaned for 10 min while the substrates were isolated from the plasma by a shutter. The tungsten disk target, which is 50 mm in diameter and 10 mm in thickness with 99.99% purity, was sputtered in either an Ar – N2 or an Ar – N2 –O2 gas mixture to deposit W2N or W – O – N films. The sputtering conditions were first adjusted at a constant total pressure of 0.8 Pa at N2 and Ar flow rates of 4.0 and 10.2 sccm, respectively, to obtain a single-phase W2N, and then O2 was added to the gas mixture for depositing W –O –N films. The ratio Of O2 to Ar+N2 could be changed, so as to vary the oxygen concentration between 0 and 18 at.%. In all cases, the films with oxygen-free conditions are stoichiometric with W/N ratios, determined by XPS, of 2.0 90.03. Consequently, the stoichiometry of all oxygen-free nitride films is hereafter referred to as W67N33. Steady-state operation and reproducible stoichiometries on the target were ensured using a sequence of low-to-high N2 concentrations. The thickness of the films was determined using a stylus-type surface profilometer (Tencor P-10) and confirmed by cross-sectional TEM measurements. Substrate temperature during film growth was estimated to be 5200°C (accounting for plasma heating). These samples will be denoted as as-deposited samples. Unless otherwise illustrated, post-deposition annealing of the samples was carried out in a vacuum of better than 6 ×10 − 5 Pa up to 900°C. Annealing duration time for each temperature cycle was 30 min.
2.2. Film characterization The films were analyzed by XPS using a Kratos Axis/800 hemispherical energy analyzer equipped with
an unmonochromatized Mg Ka X-ray source (hv = 1253.6 eV). The binding energy scale was calibrated against Ag 3d5/2 at 368.25 eV and the instrumental energy resolution, measured using the same peak, was 0.9 eV (pass energy set at 20 eV). Charging corrections were made using the as measured Fermi level of Ag and Au pieces on the metallic substrate holder using the Mg Ka radiation as a reference. The elemental concentration of the films was derived from the area under each elemental spectrum using published sensitivity factors [15]. The crystallographic structure of the films was determined by XRD using a Siemens D5000 diffractometer operated at 40 kV and 30 mA. The measurements were carried out using Cu Ka radiation with a Ni filter to remove Cu Kb reflections, a powder diffractometer scanned over the range up to 150° 2u at steps of 0.01°, a 0.5° divergence slit, and a 0.2° receiving slit. The films were also characterized by EELS in a Phillips EM430 TEM equipped with a GATAN 666 parallel electron energy-loss spectrometer. All electron energy-loss measurements were done in diffraction mode, 300 keV electrons, and a collection angle of approximately 8 mrad. A Phillips CM12 microscope with a LaB6 filament operated at 120 kV was used for plan-view TEM analyses. The cross-sectional TEM images were obtained with a Phillips EM430 microscope operated at 300 kV. Plan-view specimens were prepared by mechanical thinning from the backside to 30 mm followed by 3 keV Ar+-ion milling at 8° with respect to the specimen surface. Cross-sectional specimens were prepared by gluing two samples face-to-face with epoxy resin and then mechanically thinning, followed by ion-beam milling in a manner similar to that of the plan-view samples, but from both sides.
3. Results
3.1. In situ stress measurements The development of film stress during deposition was determined with a vacuum cantilever beam apparatus using an optical two-beam deflection method to detect changes in the substrate curvature. The film stress s is obtained by the well-known Stoney’s equation revised for biaxial stress: s=
n
1 Esh 2s 1 1 − 6 (1−ns)hf r r0
(1)
where Es(1 −ns) is the biaxial modulus of the substrate, hs and hf are the thicknesses of the substrate and film, and r0, and r denote the curvature radii before and after film deposition, respectively. The substrate curvature was measured on a beam-shaped sample by deflection
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
109
of a split laser beam from both ends. The sample stripes were 2 mm wide and 18 mm long. They were cut from the Si(100) wafers (hs =85 mm, Es/(1 −ns) =180.5 GPa). The in situ measurement reveals how the oxygen incorporation influences the formation of residual stresses in W –O – N films. Fig. 1 shows representative examples of the instantaneous film force as a function of the mean film thickness. The slope of the force curves corresponds to the incremental film stress, with the negative slope being compressive. The W67N33 film without oxygen dosing shows an increasing instantaneous compressive stress with film thickness. A remarkable evolution of the instantaneous stress was observed in the film containing 10 at.% oxygen, where a state of tensile instantaneous stress develops with increasing film thickness. At a high oxygen concentration, approximately 18 at.% oxygen, the film is almost free of residual stresses up to 200 nm. Fig. 2 shows the evolution of the stress s for as-deposited 150-nm-thick W – O – N films as a function of the oxygen concentration. It is noted that the film deposited without oxygen dosing also contains approximately 1 at.% oxygen. Close examination of the deposi-
Fig. 2. Residual stresses of as-deposited 150-nm-thick nitride as a function of oxygen concentration. The inset shows the TEM brightfield micrograph with selected-area electron diffraction pattern of the film containing 18 at.% oxygen.
Fig. 1. Representative examples of the film force versus mean film thickness in as-deposited W –O–N films with several different oxygen concentrations.
tion processing reveals that the most likely source of oxygen contamination was either from residual oxygen in the deposition chamber or from the tungsten sputter target under the present experimental conditions. In our study, the total Ar+ N2 pressure and the power density were fixed, and the only variable was the O2 partial pressure. No delamination of the films was observed in all cases. It is known that the residual stress in thin films is composed of two parts, i.e. the thermal stress and the intrinsic stress. It seems reasonable to assume that the thermal stress can be neglected under the present conditions because the contribution of the thermal (tensile) stress to the total stress for 150-nmthick films at growth temperature of approximately 150°C did not exceed 10%. It is seen that the W67N33 film (containing approximately 1 at.% oxygen) is under high compression (1.45 GPa). The stress decreases significantly with an increase in the oxygen concentration. The film stress changes from compression to tension at 10 at.% oxygen, where the lattice parameter of the film comes slightly below the bulk value of W2N, 0.4126 nm [16] (see Fig. 3). With further increase in oxygen concentration the film is almost free of residual stresses.
110
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
This full relaxation of stresses may, therefore, be understood in terms of the development of a disordered amorphous network as the oxygen incorporation is increased. The amorphous nature of the films at high oxygen incorporation was evidenced by TEM measurements. The inset of Fig. 2 shows the typical bright-field (first diffused ring paraxial) image of the film containing 18 at.% oxygen and the selected area electron diffraction pattern obtained from this area. The image reveals the largest discernible details being smaller than approximately 1– 2 nm, indicating an amorphous structure. The observation was further supported by the fact that the dark-field image in TEM exhibits no microcrystallites (not shown).
3.2. Film structure and XPS depth profiles Our XRD measurements indicate that the crystallographic structure of sputter-deposited W67N33 (nominally 150 nm in thickness) is W2N in structure. The XRD u –2u scan data in the Bragg – Brentano geometry show a linear decrease in the lattice parameter of as-deposited films from 0.4220 to 0.4116 nm (calculated from the (111) diffraction peak position) while the oxygen concentration increases from 1 to 10 at.% as
Fig. 4. Typical XPS (a) N 1s and (b) O 1s spectra measured from the W56O15N29 film at a depth of approximately of 8 nm before and after annealing to 800°C. The curves have been offset vertically for clarity.
Fig. 3. Variation of lattice parameters in W –O–N films before and after annealing to 700°C as a function of oxygen concentration.
shown in Fig. 3. The observation that the lattice parameter of oxygen-contained films shifts toward lower values as the oxygen concentration increases indicates the contraction of the lattice by the incorporated oxygen atoms. This can be explained as increasing the incorporation of oxygen into the lattice sites without the formation of detectable amounts of W–O phases. When the oxygen concentration reaches 15 at.% or above, amorphous films are formed. Fig. 3 also shows that, after annealing to 700°C, the lattice parameter approaches to the reported bulk value
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
Fig. 5. XPS depth profiles from an as-deposited 150-nm-thick W56O15N29 film.
of 0.4126 nm, indicating that the layers are nearly fully relaxed. It should be pointed out that we could not detect any loss of both nitrogen and oxygen from any of the films after vacuum annealing up to 800°C. XPS N 1s and O 1s spectra shown in Fig. 4 measured on as-deposited films and after annealing unambiguously reveal this fact. XPS depth profile measurements from oxygen-contained films provide evidence of oxygen enrichment in the near-surface region during deposition. Fig. 5 shows typical depth profiles from a 150-nm-thick W56O15N29 film. In order to avoid oxygen contamination from air exposure, the film was capped in situ with 10-nm-thick carbon overlayers prior to air exposure. It can be seen from Fig. 5 that the profile shows a high concentration of oxygen near the surface. Below this surface layer, however, the oxygen concentration is reduced to a steady-state value of about 13 – 15 at.%. This can be explained in terms of original oxygen enrichment in the near-surface region. The shape of the oxygen profile can be understood as follows. When a new nitride monolayer is deposited on a surface partially covered with oxygen, a significant fraction of the adsorbed oxygen can diffuse to the new surface. The diffusion of oxygen toward the near-surface region can be explained by a so called floating phenomenon [17,18]. The floating oxygen reduces the overall energy of the film/substrate composite by lowering the surface free energy. Thus with a continuous arrival of oxygen to the surface of the film, oxygen retained in the film will show a monotonic increase toward the surface as shown in Fig. 5.
3.3. XRD, TEM and EELS Example plots of XRD u – 2u scans exhibiting the structural changes in an oxygen-free W67N33 film after
111
high temperature annealing are shown in Fig. 6. The inset shows the XRD pattern obtained from the as-deposited film. At a low growth temperature atomic displacement, local atomic arrangements, and enhancement of adatom mobilities are limited. Thus, the (111) and (200) diffraction peaks observed in XRD u–2u measurements are broader and their intensities are relatively weak, indicating that the degree of ordering for the as-deposited W67N33 film is lower, probably due to small crystallite size and distortion of the fcc lattice structure. After annealing to 700°C, adatom mobilities in the film are sufficiently high to form a well ordered single-phase structure of W2N. The major peaks observed at angles 2u equal to 37.68, 43.70, 63.76, 76.12, and 80.16°. These peaks are assigned to (111), (200), (220), (311), and (222) reflections of fcc W2N This structure is thermally stable and all of W2N reflection peaks remain up to 800°C. After annealing to 825°C, the W(110) reflection peak appears at 40.3°, indicating that W2N has partially converted to W. With further increase in annealing temperature, the ratio of W to W2N increases. After annealing to 900°C, the full conversion from W2N to W in the film is observed.
Fig. 6. XRD u – 2u scans from the 150-nm-thick W67N films after high temperature annealing. The fresh film deposited simultaneously was used for each temperature annealing. The inset shows the XRD pattern from the as-deposited film.
112
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
some fine W2N grains. This is in good agreement with the XRD results shown in Fig. 6, where the film has been converted into a two-phase structure consisting of both W2N and bcc W after 850°C annealing. For the bcc W structure of the film formed after 900°C annealing (Fig. 7c), we see that this dense structure is still columnar in nature, but it does not exhibit the voids. This is indicative of the presence of large crystallites. It is also noted that alignment of the columns in W is less pronounced than that in W2N. Fig. 8 shows a detailed XRD pattern for 2u angles up to 150° obtained from a W56O15N29 film after annealing to 900°C. The as-deposited W56O15N29 film is amorphous in structure (see the inset). When heated at 650°C or above, a well ordered single-phase structure of W2N appears (not shown). This means that crystallization had progressed. The fact that no other phase except W2N was observed after annealing is due to no detectable phase of W–O formed for oxygen-contained films by annealing. The W2N structure of W56O15N29 is thermodynamically very stable and all of W2N diffraction peaks still remain up to 900°C. These observations are in good agreement with the results reported by Lin et al. [19], where W2N was stable up to 900°C if a thin WO3 film was formed above W2N. It is also noted that a very broad peak is observed at low 2u angles. This may be ascribed to the contributions from the surface W–O overlayers.
Fig. 7. TEM cross-sectional micrographs of theW67N33 films after annealing to three different temperatures. (a) 700°C, (b) 850°C, and (c) 900°C. All the images have the same scale. The films are those used for Fig. 6.
In order to investigate the microstructure evolution during the structural conversion, the cross-sectional TEM measurements were carried out from the same samples used in Fig. 6. Typical TEM micrographs are shown in Fig. 7. Columnar microstructure was observed in the crystailineW67N33film after annealing to 700°C (Fig. 7a). It contains a well-defined columnar microstructure with uniform columns from substrate to the top of the films. The column boundaries appear to be decorated by a relatively low density of voids. The average column diameter is about 15 – 20 nm. After annealing to 850°C, the microstructure is also a columnar with voids decorating the column boundaries (Fig. 7b). The average column diameter becomes larger, on the order of 30– 40 nm. It is clear that most of the microstructure for the film after annealing to 850°C are composed of large course bcc W grains together with
Fig. 8. XRD u– 2u scans from a 150-nm-thick W56O15N29 film after annealing to 900°C. The inset shows the XRD pattern from the as-deposited film.
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
Fig. 9. EELS N K-edge intensity as a function of sample temperature for W67N33 and W56O15N29. From each point, data were taken during in situ heating at a heating rate of 20°C min − 1, held for 5 min and then the N K-edge spectrum was taken. Error bars are based on 4–6 measurements.
Measurements using EELS and N K-edge loss spectra, which is more sensitive to the bulk [20], show that the signal intensity of N K-edge for W56O15N29 is much stronger than that for W67N33 as the sample temperature increases to 900°C (Fig. 9). This result directly indicates that nitrogen in the oxygen-free W67N33 film evaporates to vacuum after high temperature annealing, in agreement with the XRD and TEM measurements. A detailed description of EELS and K-edge spectra analyses is covered in our previous paper [13].
4. Discussion The stress-versus-oxygen concentration curve was obtained after repeated deposition over many days (Fig. 2). The results are reproducible. The compressive stress observed at a low working-gas pressure was attributed to the bombardment of the growing film surface by energetic particles such as sputtered atoms and working-gas atoms or ions through the atomic peening process [21]. Thus the crystal lattice is deformed and excess atoms are forced into interstitial sites to expand the crystal lattice, resulting in compressive stress in the films. The stress became smaller by introducing a small amount of O2 in the sputtering chamber. It can be suggested that the stress depends mainly on the oxygen content in the W – O – N films. It seems reasonable to assume that some of the oxygen atoms replace nitrogen to constitute the crystal lattice, resulting in the decrease in compressive stresses. At higher oxygen concentrations the structure collapses into a completely disordered amorphous network, leading to the full relaxation of stresses. This relaxation of compressive stresses may, therefore, be understood in terms of a loss
113
in intercolumnar coupling as a result of increased porosity in the grain boundaries. The kinetics of growth is controlled by the mobility of the impinging atoms on the surface before they condense and become entrapped in the film. This mobility can be enhanced by increasing temperature or by supplying impact energy through ion bombardment. For deposition at a low pressure of 0.8 Pa, pure reactively sputtered films of W67N33 grew in crystalline form. This is in agreement with previous reported results [4,5], where the crystallite structure favors sputterdeposited films whose composition is in the vicinity of W2N. The lattice parameter for as-deposited W67N33 films is higher than the corresponding bulk value of 0.4126 nm. A continuous shift of the lattice parameter toward lower values with increasing oxygen concentration is also evidenced (Fig. 3). The crystal structure of W2N is a B1–NaCl type where the W atoms occupy the positions of the fcc lattice points and N atoms occupy 50% of the total octahedral sites. In the case of oxygencontained films, some of the incorporated oxygen could possibly be occupying these nitrogen sites, resulting in a contraction of the lattice. This is clarified by the fact that the covalent radius of oxygen (0.066 nm) is smaller than that of nitrogen (0.070 nm). However, the possibility of the presence of oxygen impurities at the interstitial sites and grain boundaries also cannot be completely ruled out. As the oxygen is introduced in the deposition chamber, the mobility of both W and N particles decreases. At higher O2 partial pressures the incorporation of oxygen in the films is significant, resulting in the amorphous growth. We suggest that the oxygen impurities incorporated in the film can prevent crystallization growth in several ways: Oxygen can act as a roadblock to the diffusing W and N or can trap diffusing W and N, or can serve as a nucleation site for lattice defects. Therefore, the growing W and N particles may not be sufficiently mobile to migrate to the preferred sites for crystallization growth. When the concentration of permanent defects induced by the incorporated oxygen exceeds the equilibrium concentration of the mobile defects, failure of crystallization growth occurs and amorphization begins to take place. In addition, the decreasing grain size, observable with increasing O2 partial pressure during deposition can be explained in terms of oxygen enrichment at the grain surface, thus preventing further growth and establishing new grains. The oxygen is present as a W–O oxide, since the O1 s binding energy is approximately 530 eV (Fig. 4b). This leads to the picture that the incorporated oxygen mainly bonds to W, eventually leading to the partial formation of tungsten oxide. From thermodynamic considerations the formation of a WOx oxide is energetically strongly favored over the formation of a NO layer because the heat of formation for WO3 (−842.9
114
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115
M mol − 1) and for WO2 ( −589.7 M mol − 1) is much larger than the corresponding value for NO (91.3 M mol − 1) [22]. Such an oxide formation is favored by the more stable WO bonding. However, the present XPS data provide no distinction between WO2 or WO3, or WO. The annealing used in the present investigation presumably serves only to activate atom diffusion, phase ordering, and structure conversion [23,24]. It is of interest to determine the thermal range of structure stability for both oxygen-free and oxygen-contained nitride films. It was found that the W2N structure of W67N33 was stable up to 800°C (Fig. 6), whereas the W2N structure of W56O15N29 was thermodynamically stable up to 900°C (Fig. 8). This can be explained in terms of original oxygen enrichment at the surface (Fig. 5) as well as oxygen species in the W2N matrix in the oxygencontained films. With increasing temperature the oxygen atoms move increasingly outward to the near-surface region to form thicker W – O (concentration dominant) overlayers (not shown here). This can be obviously ascribed to two causes: first, oxygen atoms are minority impurities; and second, the W – O bond is more stable than the WN bond. Both effects favor oxygen segregation. In addition, part of the bulk oxygen during annealing may also be possible to diffuse into the grain boundaries. We believe that thin W–O overlayers in the near-surface region and oxygen species in the W2N grain boundaries can be helpful for preventing the desorption of nitrogen from W2N matrix. Based on the results obtained and the above discussion, the thermal stability of the W2N phase is believed to be governed by two main trends: (1) From an energetic point of view, the bonding energy required to stabilize the W2N structure must be overcome in order for breaking WN bonds to occur. (2) From a structural point of view, the transformation to bcc W occurs by desorption of nitrogen and localized rearrangement of the lattice. No long range diffusion or atomic movement is required. Further study with ab initio total-energy calculations for establishing the origin of the stable W–O and W–N bonding structures is in progress in our Materials Research Center.
5. Conclusions The effect of oxygen on residual stress and structural properties of tungsten nitride thin films prepared by reactive d.c. magnetron sputtering in an Ar – N2 –O2 gas mixture has been investigated. Film stresses were analyzed in situ by a cantilever beam technique during the film formation. The films were also characterized ex situ by using XRD, XPS, TEM, and EELS. To the authors’ knowledge, this is the first report on sputter deposited W –O –N films.
The film stress was found to strongly depend on the oxygen incorporation in the W–O–N films under the present experimental conditions. The oxygen-free films had high compressive stresses. As the oxygen concentration was increased, the stress in the films decreased markedly and became slightly tensile for films near 10–15 at.% oxygen. It is suggested that the relaxation of the compressive stress is due to the decrease in the lattice parameter caused by incorporating oxygen atoms in the lattice sites (since the oxygen atom is smaller than the nitrogen atom) and the development of an amorphous network in the W–O–N films as the incorporation of oxygen was increased. The role of oxygen in stabilizing the W2N structure was also elucidated herein by XRD, XPS, TEM, and EELS. The oxygen-contained nitride films are thermodynamically more stable than the oxygen-free films. We believe that thin W–O overlayers in the near-surface region and oxygen species in the grain boundaries can be helpful for preventing the desorption of nitrogen from W2N matrix during high temperature annealing. We hope that our new results will stimulate further experimental as well as theoretical studies.
Acknowledgements The authors gratefully acknowledge the financial support of this work by the Australian Research Grants Scheme. The authors also appreciate the use of the facilities in the Australian Key Center for Microscopy and Microanalysis, University of Sydney, which is supported by the Australia Research Council.
References [1] L.E. Toth, Transition Metal Carbides and Nitrides, Academic Press, New York, 1971. [2] S.D. Marcus, R.F. Foster, Thin Solid Films 263 (1993) 330. [3] K.M. Chang, T.H. Yeh, I.C. Deng, J. Appl. Phys. 81 (1997) 3670. [4] F.C.T. So, E. Kolawa, X.A. Zhao, M.A. Nicolet, Thin Solid Films 153 (1988) 507. [5] L. Boukhris, J.M. Poitevin, Thin Solid Films 310 (1997) 222. [6] H.T. Chiu, S.H. Chuang, J. Mater. Res. 8 (1993) 1353. [7] W. Sinke, G.P.A. Friflink, F.W Saris, Appl. Phys. Lett. 47 (1985) 471. [8] N. Kumar, K. Pourrezaei, B. Lee, E.C. Douglas, Thin Solid Films 164 (1988) 417. [9] P. Jin, S. Maruno, Jpn. J. Appl. Phys. 30 (1991) 2059. [10] K.J. Huber, C.R. Aita, J. Vac. Sci. Technol. A6 (1988) 1717. [11] H.G. Tompkins, J. Appl. Phys. 70 (1991) 3876. [12] M. Mandi, H. Hoffrnann, P. Kucher, J. Appl. Phys. 68 (1990) 2127. [13] Y.G. Shen, Y.W. Mai, Q.C. Zhang, D.R. McKenzie, W.D. FcFall, W.E. McBride, J. Appl. Phys. 87 (2000) 177. [14] Y.G. Shen, Y.W. Mai, W.E. McBride, D.R. McKenzie, Q.C. Zhang, Appl. Phys. Lett. 75 (1999) 2211.
Y.G. Shen, Y.W. Mai / Materials Science and Engineering B76 (2000) 107–115 [15] J.F. Moulder, W.F. Stichle, P.E. Sobol, K.D. Bomben Jr, in: J. Chastain, R.C. King Jr (Eds.), Handbook of X-Ray Photoelectron Spectroscopy, Physical Electronics, Eden Prairie, MN, 1995. [16] M. Bauccio Jr (Ed.), ASM Metals Reference Book, American Society for Metals, Metals Park, OH, 1994. [17] W.F. Egelhoff Jr, D.A. Steigerwald, J. Vac. Sci. Technol. A7 (1989) 2167. [18] Y.G. Shen, D.J. O’Connor, H. van Zee, K. Wandelt, R.J. MacDonald, Thin Solid Films 263 (1995) 72. [19] J. Lin, A. Tsukune, T. Suzuki, M. Yamada, J. Vac. Sci. Technol.
.
115
A16 (1998) 611. [20] R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, 2nd edn, Plenum, New York, 1996. [21] J.A. Thornton, D.W. Hoffman, Thin Solid Films 171 (1989) 5. [22] D.R. Lide (Ed.), CRC Handbook of Chemistry and Physics, 72nd edn, CRC Press, Boca Raton, FL, 1991/1992. [23] Y.G. Shen, J. Yao, D.J. O’Connor, B.V. King, R.J. MacDonald, Phys. Rev. B56 (1997) 9894. [24] Y.G. Shen, D.J. O’Connor, J. Yao, Appl. Surf. Sci. 125 (1998) 300.
.