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〈Si〉 multilayers
Surface and Coatings Technology 176 (2003) 115–123
Properties of interfaces in CuyTi1yxAlxNyNSiM multilayers Ding-Fwu Liia,*, Jow-Lay Huangb, Jenn-Fuh Linb a
Department of Electrical Engineering, Cheng Shiu Institute of Technology, Niau-sung shiang, Kaohsiung County, 833, Taiwan, ROC b Department of Material Science and Engineering, National Cheng-Kung University, Tainan 701, Taiwan, ROC Received 8 November 2002; accepted in revised form 11 February 2003
Abstract Ti1yxAlxN films have been grown onto Si(1 0 0) substrate by d.c. reactive magnetron sputtering as a diffusion barrier between Cu and Si. The residual stresses of Ti1yxAlxN films depended on the microstructure and composition. Higher residual stresses were obtained for Ti1yxAlxN films deposited with higher bias voltage. The minimum residual stress was 11.4 MPa, found in Ti1yxAlxN films deposited with a nitrogen flow rate of 8 mlymin and a bias voltage of y50 V. Residual stresses of Cu films decreased with the increase of surface roughness of the Ti1yx Alx N films. The good adhesion of Ti1yxAlxN films on Si substrate was due to the chemical reactions, however, the bad adhesion of Cu films on Ti1yx Alx N films was attributed to the large residual tensile stresses in Cu film. 䊚 2003 Elsevier Science B.V. All rights reserved. Keywords: Reactive magnetron sputtering; Adhesion; Surface roughness; Residual stress
1. Introduction Ultra large scale integrated circuits (ULSI) with corresponding small contact areas has gotten more stringent demands for the metallization in electronic packing technology. Due to its low resistivity (1.67 mV cm for bulk), high reliability against electromigration and the feasibility of physical vapor deposition, copper is an attractive material for interconnection metallization in such ULSI devices w1,2x. However, copper could be an impurity in semiconductor fabrication because of its fast diffusivity in silicon. It is known that copper atoms dissolve in the silicon crystal on interstitial sites and become deep level dopants. These deep levels can act as generation-recombination centers, hence decreasing the carrier lifetime and induce disorder of the device w3x. Therefore, there is a need to develop a thin diffusion barrier film to prevent copper diffusion into silicon in order to increase the device reliability. TiN is presently one of the most widely used barrier materials in copper and aluminum based metallization since it has one of the lowest electrical resistivity among the transition-metal mononitrides w3–7x. However, TiN *Corresponding author. Fax: q886-7-5556789. E-mail address: [email protected] (D.-F. Lii).
can be oxidized to form TiO2 and degrades above 500 8C w8–10x. In a previous study w11x, a potential alternative material with metastable NaCl structure w12x, Ti1yxAlxN, was deposited on boron-doped Si substrates by magnetron reactive sputtering using titanium–aluminum alloy (TiyAls90y10 at.%). The effects of heat treatment on sheet resistance and the results of the AES depth analysis in CuyTi1yxAlxNyNSiM specimens indicated the feasibility of using Ti1yxAlxN films as a diffusion barrier between CuySi up to 700 8C w11x. In this study, the properties of interfaces in Cuy Ti1yxAlxNyNSiM multilayers were further investigated. The effects of nitrogen flow rate and bias voltage on the internal stresses and adhesive strength of the films were evaluated. 2. Experimental procedure 2.1. Sample preparation Titanium aluminum nitride films were deposited onto boron-doped Si(1 0 0) substrates with a resistivity of 40 V cm, by d.c. magnetron reactive sputtering. The titanium–aluminum target alloy (⭋7.62=0.635 cm2, Plasmaterial Inc. USA) had a composition of TiyAl (at.%)s
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90y10. The substrates were degreased, ultrasonically cleaned in trichloroethylene for 10 min, and subsequently in acetone for 10 min. The substrates were then cleaned in isopropyl alcohol for 10 min, rinsed in deionized water for 3 min and cleaned in a diluted HF solution (HF:H2Os1:10 by volume) for 30 s. Finally, the substrates were rinsed in de-ionized water for 3 min and blown dry in flowing N2 gas before deposition. Sputtering was conducted in a mixture of Ar (99.999%, Lien Hwa gas Co. Hsin Chu, Taiwan)yN2 (99.9995%, Lien Hwa gas Co.) atmosphere at a targetsubstrate distance of 4 cm. A diffusion pump coupled with a rotary pump was used to achieve an ultimate pressure of 2.67=10y4 Pa before introducing gas mixtures of Ar and N2. The substrates were pre-sputtered by r.f. argon discharge at 0.8 Pa (6 mTorr) with an energy density of 0.4 Wycm for 10 min prior to film deposition for cleaning purpose. The TiAl target was cleaned by sputtering in Ar gas at 2.67 Pa for 5 min prior to the deposition of films. The total pressure and flow rate of Ar were maintained at 0.8 Pa and 20 mlymin, respectively, during the deposition process. The cathode current was maintained at 0.5 A and the substrate temperature was kept at 350 8C. The nitrogen flow rate varied between 0 and 10 mly min. An r.f. power supply was used for establishing negative bias voltages (0–80 V) on the substrate during deposition. The thickness of the TiAlN films was maintained at 100 nm. Deposited samples were cooled down in Ar atmosphere to less than 100 8C before venting the system. In addition, the copper layer (250 nm) over Ti1yxAlxNySi was deposited at a current of 0.5 A and a vacuum pressure of 0.8 Pa under an Ar flow rate of 20 mlymin, with no bias voltage or heating on the substrate. 2.2. Characterization of Ti1yxAlxN films The film thickness was measured by an a-step apparatus (Surface profiler, Tencor, USA). The deposition rate was calculated from the film thickness and deposition time. A surface analyzer, atomic force microscope (AFM, Nanoscope IIIa, Digital Instrument) was used to analyze the three dimensional space morphology of the film surface. Composition of deposited films was determined by electron spectroscopy for chemical analysis (VG Scientific 210, X-ray photoelectron spectroscopy, UK), using Mg Ka radiation (1253.6 eV) with an accelerating voltage of 12 keV and current of 17 mA under a vacuum of 2=10y5 Pa. Field emission scanning electron microscopy (SEM, XL-40FEG, Philips) was used for microstructural analysis and examination of cross-sections of the films. Residual stresses were determined by the bending technique (Tencor FLX-2320, Tencor Instruments,
Fig. 1. The NyTi ratio of Ti1yx Alx N films vs. nitrogen flow rates under various bias voltages.
Mountain View, CA), which compares the curvatures of the substrates before and after coated with a film. The residual stress (s) of the film was calculated according to Stoney’s equation w13x, ss
1 w Es z ts2 w 1 1z x | x y | 6 y 1yns ~ tf y R R0 ~
(1)
where Es, ts and ns are the Young’s modulus, thickness and Poisson ratio of the substrate, respectively. tf is the thickness of the film. R0 and R are the radius of curvature, measured on the substrate before and after deposition. The adhesive strength was determined using a scratch tester apparatus (Romulus II Universal Tester, Quad, USA), using a Rockwell diamond stylus with a tip radius of 300 mm. Testing was conducted at a loading rate of 1 Nys and a scratch length of 10 mm. Both acoustic emission and light microscopy were used to correlate the friction, appearance of failures and critical scratch loads. Interfacial adherence was determined by the pull-off test. An epoxy cement was applied to the films and then pulled off by a ball bond shear module (Sebastian FiveA, Quard Group, Inc.). The load at which the film was removed indicates the bonding strength between substrate and film. 3. Results and discussion 3.1. Compositions and surface morphology Fig. 1 shows the NyTi ratio of Ti1yxAlxN films vs. nitrogen flow rates under various bias voltages. NyTi ratios increased with the nitrogen flow rate as observed. It also can be seen that all the films have not completely
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approached stoichiometric composition, even the maximum NyTi ratio being only equal to 0.7 instead of being higher than 1.5 for stoichiometric Ti1yxAlxN films, w14x especially there deposited at a nitrogen flow rate equal to 2 mlymin, probably due to the insufficient nitrogen concentration for the formation of Ti1yxAlxN films. Typical SEM micrographs illustrating the fracture cross-sections of films deposited under a bias voltage of y50 V at different nitrogen flow rates are shown in Fig. 2. A columnar texture with porous grain boundary was observed in the Ti1yxAlxN films deposited at a nitrogen flow rate of 2 mlymin (Fig. 2a); as the nitrogen flow rate was increased to 6 mlymin, the structure became coarser and denser (Fig. 2b); while the nitrogen flow rate was increased to 10 mlymin, a dense Ti1yxAlxN film with no columnar texture was obtained (Fig. 2c). Similar variation in grain size and morphology was observed for reactive sputtered TiN films w15x. The morphology depends strongly on the mobility of the impinging atoms on the substrate’s surface, which is considerably influenced by the chemical potential (the free energy of formation) of the nitride w15x. According to Sundgren et al., as the number of non-metal atoms (e.g. nitrogen) or ions on the surfaces is smaller than the number of metal atoms (e.g. titanium), being a nonstoichiometry structure, composition gradients and thus chemical potential may develop w15x. A directed migration of the pre-condensed species may occur leading to the formation of relatively large grains (columnar texture) with voided boundaries (Fig. 2a). Subsequently, while the amount of nitrogen on the surfaces increases, the composition gradients and the directed migration will decrease. Therefore, smaller grains with denser grain boundaries forms (Fig. 2c). In order to analyze three dimensional space morphology and roughness of the film surface, AFM (3D) was used. Fig. 3 shows the micrographs of Ti1yxAlxN films deposited under bias voltage of y50 V with various nitrogen flow rates. All surface morphologies were essentially uniformly distributed. Some change of the surface roughness of the films deposited under different nitrogen flow rate could be observed. Fig. 4 shows both the root-mean-square ‘RMS’ roughness and the average surface roughness ‘Ave.’, which all obtained from the AFM measurements of Ti1yxAlxN films (Fig. 3), vs. nitrogen flow rates under bias voltage of y50 V. The roughness appears to decrease slightly with the increase of nitrogen flow rate. This could be explained by the increase of the reactive nitrogen atoms with uniform distribution, which subsequently reduce the under-stoichiometry region and the porous grain boundary within the film. Fig. 5 represents SEM micrographs of TiAlN films deposited at a nitrogen flow rate of 6 mlymin under various bias voltages. As the substrate bias was
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Fig. 2. SEM micrographs of Ti1yxAlxN films deposited under bias voltage of y50 V at various nitrogen flow rates.
increased, the columnar structure of the films became coarser and denser. Formation and disruption of the columnar morphology under bias has been proposed by Bland et al. w16x Supposedly, the surface morphology of the film is generated either by replication of the substrate’s surface or by the preferential growth of crystallographic planes
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Fig. 3. AFM micrographs (3D) of Ti1yxAlxN films deposited under bias voltage of y50 V at nitrogen flow rate of (a) 2; (b) 6 and (c) 10 mlymin.
during deposition. During deposition, the geometrical shadowing will prevent deposition in the valleys of the substrate’s surface while its peaks will grow under no bias. Thus, a columnar morphology (Fig. 5a) was developed. As the bias increases, the ion bombardment erodes the peaks and forward sputtering will fill the valleys. Simultaneously, the backsputtered material may be returned by scattering process in the gas, which further randomizing the deposition direction and making the structure denser (as shown in Fig. 5c). Fig. 6 is the AFM micrographs (3D) of Ti1yxAlxN films deposited at nitrogen flow rate of 6 mlymin under various negative bias voltages (0–80 V). All surface morphologies were essentially uniformly distributed. An obvious change of the surface roughness of the films deposited under different bias voltages has been observed. Fig. 7 shows the corresponding roughness (RMS and Ave.) of films vs. bias voltages. The roughness (RMS) rapidly decreased from 1.414 nm (unbi-
Fig. 4. The roughness of Ti1yxAlxN films vs. nitrogen flow rates under bias voltage of y50 V.
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3.2. Residual stress
Fig. 5. SEM micrographs of TiAlN films deposited at nitrogen flow rate of 6 mlymin under various bias voltages.
ased) to about a constant of 0.65 nm (y30 to y80 V). The decrease of roughness is probably due to the enhanced mobility of deposited atoms with higher ion bombardment conditions according to the bias voltage. In addition, the slight increase of roughness at a bias voltage of y80 V was possibly caused by the ion irradiation on the sputtered surface due to the excessive bias voltage.
3.2.1. Residual stress of Ti1yxAlxN film on the Si substrate Fig. 8 shows the positive residual tensile stress of Ti1yxAlxN films deposited at various nitrogen flow rates under y50 V bias voltage. The residual stress decreased with the increase of nitrogen flow rate from the maximum of 218.0 MPa at 2 mlymin to the minimum of 11.4 MPa at 8 mlymin. It was probably due to the porous grain boundary between columnar texture of the film deposited under lower nitrogen flow rate (Fig. 2a), which was easily to transmit the tensile stress. Therefore, it is supposed to characterize with a bad mechanical behavior. A similar result of PVD coatings with lower residual compressive stress (i.e. bad mechanical behavior), corresponding to the microstructure of pronounced columns and cavities has been reported before w17x. It was previously discussed that the residual stress could be generated in the PVD films w13,17x upon cooling (from 350 8C to room temperature) due to the mismatch in linear thermal expansion coefficient (TEC) between the layers of Ti0.9Al0.1N (as9.0=10y6 yK) and Si (as2.6=10y6 yK). A positive tensile stress was theoretically due to the higher TEC value of the coating layer (Ti0.9Al0.1N). Similar positive thermal stress (;200–550 MPa) has been reported for AlN coatings deposited on silica, silicon and tungsten carbide substrate w18x. Essentially, the positive tensile stresses are possible but comparable smaller (O1 GPa) than the negative compressive intrinsic stresses ()5 GPa) w17,19x. According to Cheng et al. w13x and Oettel and Wiedemann w17x, the total internal stresses in the PVD films are mainly equal to the sum of the thermal stresses and the intrinsic stresses (if neglecting the priori negligibly small stress comes from hetero-epitaxy). Conversely, all compressive stresses with a monotonical increase to a maximum (7 GPa) at nitrogen pressure of 1.98=10y1 Pa was reported for a filtered cathodic vacuum arc deposited (Ti, Al)N films w13x. It was presumably due to the thermal stress component is negligible since the film deposition and stress measurement were carried out at room temperature in the previous study w13x. The residual stress of Ti1yxAlxN films deposited under various bias voltages at a nitrogen flow rate of 6 mly min was depicted in Fig. 9. The residual stress was approximately equal to 50 MPa as the bias voltage was smaller than 50 V. Afterwards, the residual stress increased with the increase of bias voltage, and it attained a value of 80.3 MPa as the bias voltage was equal to 80 V, presumably due to the sufficient particle energy to generate different types of defects in the volume of the grown layer (the result of atomic or ion shoot-peening effect) w20x.
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Fig. 6. AFM micrographs (3D) of Ti1yxAlxN films deposited at nitrogen flow rate of 6 mlymin under various bias voltages of (a) 0; (b) y30; (c) y50 and (d) y80 V.
According to Oettel and Wiedemann, the residual stresses in PVD hard coatings could be possibly determined by the specific energy transfer from the incoming particles to the surface atoms, which mainly controlled by the bias voltage in the case of magnetron sputtering w17x. A similar trend of residual stresses (although being negative compressive stress) increasing with the bias voltages in TiN, (Ti, Al)N and ZrN coatings has been reported before w17x. 3.2.2. Residual stress of Cu layer on the Ti1yxAlxNySi film The residual stresses of Cu layer on the Ti1yxAlxN films, which was deposited at a nitrogen flow rate of 6 mlymin under various negative bias voltages are shown in Fig. 10. A minimum residual stress of 350 MPa was obtained for Cu layer on Ti1yxAlxN film deposited with no bias. The residual stress greatly increased to a constant value of 480 MPa, approximately under various biases (30–80 V). It was presumably due to the different microstructure (e.g. surface roughness, Fig. 7) of the coatings caused by the bias voltage w17x.
Fig. 7. The roughness of Ti1yxAlxN films vs. bias voltages at a nitrogen flow rate of 6 mlymin.
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Fig. 8. The residual stress of Ti1yxAlxN films deposited at various nitrogen flow rates under y50 V bias voltage.
Fig. 9. The residual stress of Ti1yxAlx N films deposited under various bias voltages at a nitrogen flow rate of 6 mlymin.
3.3. Adhesive strength
Obviously, the Cu layer has been completely detached from TiAlN. In addition, chipping failure w22x has also been observed on the SEM micrograph of the scratch channel of Cu layer on the TiAlNySi layers (Fig. 13). It indicates that a bad adhesion existed between the layer of Cu and TiAlN. Oettel and Wiedemann w17x investigated the residual stress in PVD hard coatings and reported that the compressive stress could prevent
3.3.1. TiAlN film on the Si substrate Fig. 11 shows the SEM micrographs of the Ti1yxAlxN films after pull-test. Samples were deposited at a nitrogen flow rate of 6 mlymin without bias voltage. Both the line scan (a) and the mapping (b) of carbon on the tested specimen indicate that the adhesive failure for the Ti1yxAlxNySi film was along the epoxy cement. It is suggested that a strong adhesive strength of TiAlN film on the Si substrate, was higher than the strength of the epoxy cement. Additionally, no detachment on the coating was observed for a normal force of 60 N (defined as a critical adhesive load) in a scratch test before the crack of Si substrate, which also conformed the pulltest result. Similar to the previous result of AES depth profiles of the CuyTi1yxAlxNyNSiM multilayer w11x, titanium and oxygen atoms were both detected among the interphase of Ti1yxAlxNyNSiM. With the essence of the non-stoichiometric composition of Ti1yxAlxN (Fig. 1), the excessive Ti atoms might react with SiO2 to form a stable TiO phase. Actually, the stable chemical TiO bonding was reported w21,22x to be a major reason for a coating with higher adhesive strength. 3.3.2. Cu on the TiAlNySi layers The SEM micrographs (Fig. 12) show the Cu line scan analysis of the Cu (250 nm)yTi1yxAlxN (100 nm)y Si multilayer structure. Samples were deposited at a nitrogen flow rate of 6 mlymin under y50 V. Asdeposited Cu layer has been shown in grey area, however Cu was absent from the dark area after a pull-test.
Fig. 10. The residual stresses of Cu layer on the Ti1yxAlxN films, which was deposited at nitrogen flow rate of 6 mlymin under various bias voltage.
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Fig. 13. The SEM micrographs showing the scratch channel of the Cu layer on the Ti1yxAlxNySi film.
Fig. 11. The SEM micrographs showing the Ti1yxAlx N film after pulltest. Samples were deposited at a nitrogen flow rate of 6 mlymin without bias voltage. (a) The line scan of carbon and (b) the mapping of carbon.
cracking, if the external load acting alone generates extended regions with tensile stresses. Higher compressive residual stresses explain the higher critical load, (the higher adhesive strength) w17,23x. Inversely, a tensile residual stress in the layers will theoretically reduce the adhesive strength of the coating. The low adhesive strength of Cu layer on the TiAlNySi film could be ascribed to its larger tensile stress, comparatively to the high adhesive strength of TiAlN film on the Si substrate with lower tensile stress (as depicted in Section 3.2). 4. Summary and conclusions (1) The residual tensile stress within Ti1yxAlxN films decreased with the increase of nitrogen flow rate. Higher residual stresses were obtained within Ti1yxAlxN films deposited under higher bias voltages. The minimum residual stress was 11.4 MPa within the Ti1yxAlxN films deposited at nitrogen flow rate of 8 mlymin under bias voltage of y50 V. (2) The residual stress of Cu films decreased with the increase of the surface roughness of the Ti1yxAlxN films. The adhesive strength was rather high between Ti1yxAlxN and Si substrate due to the chemical bonding. However, because of the high residual tensile stress among Cu film, the adhesive strength between Cu and Ti1yxAlxN was comparatively low. Acknowledgments
Fig. 12. The SEM micrographs showing the Cu line scan analysis of the Cu (250 nm)yTi1yxAlxN (100 nm)ySi multilayer structure. Samples were deposited at a nitrogen flow rate of 6 mlymin under y50 V.
The authors would like to thank the National Science Council of the ROC for its financial support under the contract no. NSC 89-2216-E-012-001.
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References w1x L. Jian, Seidel, E. Tom, Mayer, W. Jim, MRS Bull. August (1994) 15. w2x E. Kolawa, J.S. Chen, J.S. Reid, P.J. Pokela, M.A. Nicolet, J. Appl. Phys. 70 (1991) 1369. w3x S.K. Rha, W.J. Lee, S.Y. Lee, D.W. Kim, C.O. Park, S.S. Chun, J. Mater. Res. 12 (1997) 3367. w4x M. Wittmer, B. Studer, H. Melchior, J. Appl. Phys. 52 (1981) 5722. w5x G. Lemperiere, J.M. Poitevin, Thin Solid Films 111 (1984) 339. w6x M. Wittmer, H. Melchior, Thin Solid Films 93 (1982) 397. w7x A. Sherman, J. Electrochem. Soc. 137 (1892) 1990. w8x Y. Tanaka, T.M. Gur, ¨ M. Kelly, S.B. Hagstrom, T. Ikeda, Thin Solid Films 228 (1993) 238–241. w9x D. Mcintyre, J.E. Greene, G. Hakansson, J.E. Sundgren, W.D. Munz, J. Appl. Phys. 67 (1990) 1542. w10x D.F. Lii, J. Mater. Sci. 33 (1998) 2137. w11x D.F. Lii, J.L. Huang, W.K. Tsai, W.T. Lo, B.S. Yau, Surf. Eng. 17 (4) (2001) 295–299.
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w12x O. Knotek, T. Leyendecker, J. Solid State Chem. 70 (1987) 318. w13x Y.H. Cheng, B.K. Tay, S.P. Lau, X. Shi, H.C. Chua, Thin Solid Films 379 (2000) 76–82. w14x D.F. Lii, J.L. Huang, B.S. Yau, Thin Solid Films 335 (1998) 122–126. w15x J.-E. Sundgren, B.-O. Johansson, S.-E. Karlsson, Thin Solid Films 105 (1983) 367–384. w16x R.D. Bland, G.J. Kominiak, D.M. Mattox, J. Vac. Sci. Technol. 11 (4) (1974) 671–674. w17x H. Oettel, R. Wiedemann, Surf. Coat. Technol. 76–77 (1995) 265–273. w18x G. Kleer, R. Kassner, E.-M. Meyer, M.G. Schinker, W. Doell, Coat. Technol. 54–55 (1992) 167. w19x S.J. Bull, A.M. Jones, A.R. Mecabe, Surf. Coat. Technol. 54– 55 (1992) 173. w20x F.M. d’Heurle, Metall. Trans. 1 (1970) 725. w21x P. Anderson, M. Moske, K. Dyrbye, J. Bottiger, Thin Solid Films 340 (1999) 205–209. w22x P.J. Burnett, D.S. Rickerby, Thin Solid Films 154 (1987) 403–416. w23x D.F. Lii, B.S. Yau, J.L. Huang, C.Y. Chen, W.T. Lo, J. Ceramic Soc. Jpn. 109 (1) (2001) 9–11.
Properties of interfaces in CuyTi1yxAlxNyNSiM multilayers Ding-Fwu Liia,*, Jow-Lay Huangb, Jenn-Fuh Linb a
Department of Electrical Engineering, Cheng Shiu Institute of Technology, Niau-sung shiang, Kaohsiung County, 833, Taiwan, ROC b Department of Material Science and Engineering, National Cheng-Kung University, Tainan 701, Taiwan, ROC Received 8 November 2002; accepted in revised form 11 February 2003
Abstract Ti1yxAlxN films have been grown onto Si(1 0 0) substrate by d.c. reactive magnetron sputtering as a diffusion barrier between Cu and Si. The residual stresses of Ti1yxAlxN films depended on the microstructure and composition. Higher residual stresses were obtained for Ti1yxAlxN films deposited with higher bias voltage. The minimum residual stress was 11.4 MPa, found in Ti1yxAlxN films deposited with a nitrogen flow rate of 8 mlymin and a bias voltage of y50 V. Residual stresses of Cu films decreased with the increase of surface roughness of the Ti1yx Alx N films. The good adhesion of Ti1yxAlxN films on Si substrate was due to the chemical reactions, however, the bad adhesion of Cu films on Ti1yx Alx N films was attributed to the large residual tensile stresses in Cu film. 䊚 2003 Elsevier Science B.V. All rights reserved. Keywords: Reactive magnetron sputtering; Adhesion; Surface roughness; Residual stress
1. Introduction Ultra large scale integrated circuits (ULSI) with corresponding small contact areas has gotten more stringent demands for the metallization in electronic packing technology. Due to its low resistivity (1.67 mV cm for bulk), high reliability against electromigration and the feasibility of physical vapor deposition, copper is an attractive material for interconnection metallization in such ULSI devices w1,2x. However, copper could be an impurity in semiconductor fabrication because of its fast diffusivity in silicon. It is known that copper atoms dissolve in the silicon crystal on interstitial sites and become deep level dopants. These deep levels can act as generation-recombination centers, hence decreasing the carrier lifetime and induce disorder of the device w3x. Therefore, there is a need to develop a thin diffusion barrier film to prevent copper diffusion into silicon in order to increase the device reliability. TiN is presently one of the most widely used barrier materials in copper and aluminum based metallization since it has one of the lowest electrical resistivity among the transition-metal mononitrides w3–7x. However, TiN *Corresponding author. Fax: q886-7-5556789. E-mail address: [email protected] (D.-F. Lii).
can be oxidized to form TiO2 and degrades above 500 8C w8–10x. In a previous study w11x, a potential alternative material with metastable NaCl structure w12x, Ti1yxAlxN, was deposited on boron-doped Si substrates by magnetron reactive sputtering using titanium–aluminum alloy (TiyAls90y10 at.%). The effects of heat treatment on sheet resistance and the results of the AES depth analysis in CuyTi1yxAlxNyNSiM specimens indicated the feasibility of using Ti1yxAlxN films as a diffusion barrier between CuySi up to 700 8C w11x. In this study, the properties of interfaces in Cuy Ti1yxAlxNyNSiM multilayers were further investigated. The effects of nitrogen flow rate and bias voltage on the internal stresses and adhesive strength of the films were evaluated. 2. Experimental procedure 2.1. Sample preparation Titanium aluminum nitride films were deposited onto boron-doped Si(1 0 0) substrates with a resistivity of 40 V cm, by d.c. magnetron reactive sputtering. The titanium–aluminum target alloy (⭋7.62=0.635 cm2, Plasmaterial Inc. USA) had a composition of TiyAl (at.%)s
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90y10. The substrates were degreased, ultrasonically cleaned in trichloroethylene for 10 min, and subsequently in acetone for 10 min. The substrates were then cleaned in isopropyl alcohol for 10 min, rinsed in deionized water for 3 min and cleaned in a diluted HF solution (HF:H2Os1:10 by volume) for 30 s. Finally, the substrates were rinsed in de-ionized water for 3 min and blown dry in flowing N2 gas before deposition. Sputtering was conducted in a mixture of Ar (99.999%, Lien Hwa gas Co. Hsin Chu, Taiwan)yN2 (99.9995%, Lien Hwa gas Co.) atmosphere at a targetsubstrate distance of 4 cm. A diffusion pump coupled with a rotary pump was used to achieve an ultimate pressure of 2.67=10y4 Pa before introducing gas mixtures of Ar and N2. The substrates were pre-sputtered by r.f. argon discharge at 0.8 Pa (6 mTorr) with an energy density of 0.4 Wycm for 10 min prior to film deposition for cleaning purpose. The TiAl target was cleaned by sputtering in Ar gas at 2.67 Pa for 5 min prior to the deposition of films. The total pressure and flow rate of Ar were maintained at 0.8 Pa and 20 mlymin, respectively, during the deposition process. The cathode current was maintained at 0.5 A and the substrate temperature was kept at 350 8C. The nitrogen flow rate varied between 0 and 10 mly min. An r.f. power supply was used for establishing negative bias voltages (0–80 V) on the substrate during deposition. The thickness of the TiAlN films was maintained at 100 nm. Deposited samples were cooled down in Ar atmosphere to less than 100 8C before venting the system. In addition, the copper layer (250 nm) over Ti1yxAlxNySi was deposited at a current of 0.5 A and a vacuum pressure of 0.8 Pa under an Ar flow rate of 20 mlymin, with no bias voltage or heating on the substrate. 2.2. Characterization of Ti1yxAlxN films The film thickness was measured by an a-step apparatus (Surface profiler, Tencor, USA). The deposition rate was calculated from the film thickness and deposition time. A surface analyzer, atomic force microscope (AFM, Nanoscope IIIa, Digital Instrument) was used to analyze the three dimensional space morphology of the film surface. Composition of deposited films was determined by electron spectroscopy for chemical analysis (VG Scientific 210, X-ray photoelectron spectroscopy, UK), using Mg Ka radiation (1253.6 eV) with an accelerating voltage of 12 keV and current of 17 mA under a vacuum of 2=10y5 Pa. Field emission scanning electron microscopy (SEM, XL-40FEG, Philips) was used for microstructural analysis and examination of cross-sections of the films. Residual stresses were determined by the bending technique (Tencor FLX-2320, Tencor Instruments,
Fig. 1. The NyTi ratio of Ti1yx Alx N films vs. nitrogen flow rates under various bias voltages.
Mountain View, CA), which compares the curvatures of the substrates before and after coated with a film. The residual stress (s) of the film was calculated according to Stoney’s equation w13x, ss
1 w Es z ts2 w 1 1z x | x y | 6 y 1yns ~ tf y R R0 ~
(1)
where Es, ts and ns are the Young’s modulus, thickness and Poisson ratio of the substrate, respectively. tf is the thickness of the film. R0 and R are the radius of curvature, measured on the substrate before and after deposition. The adhesive strength was determined using a scratch tester apparatus (Romulus II Universal Tester, Quad, USA), using a Rockwell diamond stylus with a tip radius of 300 mm. Testing was conducted at a loading rate of 1 Nys and a scratch length of 10 mm. Both acoustic emission and light microscopy were used to correlate the friction, appearance of failures and critical scratch loads. Interfacial adherence was determined by the pull-off test. An epoxy cement was applied to the films and then pulled off by a ball bond shear module (Sebastian FiveA, Quard Group, Inc.). The load at which the film was removed indicates the bonding strength between substrate and film. 3. Results and discussion 3.1. Compositions and surface morphology Fig. 1 shows the NyTi ratio of Ti1yxAlxN films vs. nitrogen flow rates under various bias voltages. NyTi ratios increased with the nitrogen flow rate as observed. It also can be seen that all the films have not completely
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approached stoichiometric composition, even the maximum NyTi ratio being only equal to 0.7 instead of being higher than 1.5 for stoichiometric Ti1yxAlxN films, w14x especially there deposited at a nitrogen flow rate equal to 2 mlymin, probably due to the insufficient nitrogen concentration for the formation of Ti1yxAlxN films. Typical SEM micrographs illustrating the fracture cross-sections of films deposited under a bias voltage of y50 V at different nitrogen flow rates are shown in Fig. 2. A columnar texture with porous grain boundary was observed in the Ti1yxAlxN films deposited at a nitrogen flow rate of 2 mlymin (Fig. 2a); as the nitrogen flow rate was increased to 6 mlymin, the structure became coarser and denser (Fig. 2b); while the nitrogen flow rate was increased to 10 mlymin, a dense Ti1yxAlxN film with no columnar texture was obtained (Fig. 2c). Similar variation in grain size and morphology was observed for reactive sputtered TiN films w15x. The morphology depends strongly on the mobility of the impinging atoms on the substrate’s surface, which is considerably influenced by the chemical potential (the free energy of formation) of the nitride w15x. According to Sundgren et al., as the number of non-metal atoms (e.g. nitrogen) or ions on the surfaces is smaller than the number of metal atoms (e.g. titanium), being a nonstoichiometry structure, composition gradients and thus chemical potential may develop w15x. A directed migration of the pre-condensed species may occur leading to the formation of relatively large grains (columnar texture) with voided boundaries (Fig. 2a). Subsequently, while the amount of nitrogen on the surfaces increases, the composition gradients and the directed migration will decrease. Therefore, smaller grains with denser grain boundaries forms (Fig. 2c). In order to analyze three dimensional space morphology and roughness of the film surface, AFM (3D) was used. Fig. 3 shows the micrographs of Ti1yxAlxN films deposited under bias voltage of y50 V with various nitrogen flow rates. All surface morphologies were essentially uniformly distributed. Some change of the surface roughness of the films deposited under different nitrogen flow rate could be observed. Fig. 4 shows both the root-mean-square ‘RMS’ roughness and the average surface roughness ‘Ave.’, which all obtained from the AFM measurements of Ti1yxAlxN films (Fig. 3), vs. nitrogen flow rates under bias voltage of y50 V. The roughness appears to decrease slightly with the increase of nitrogen flow rate. This could be explained by the increase of the reactive nitrogen atoms with uniform distribution, which subsequently reduce the under-stoichiometry region and the porous grain boundary within the film. Fig. 5 represents SEM micrographs of TiAlN films deposited at a nitrogen flow rate of 6 mlymin under various bias voltages. As the substrate bias was
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Fig. 2. SEM micrographs of Ti1yxAlxN films deposited under bias voltage of y50 V at various nitrogen flow rates.
increased, the columnar structure of the films became coarser and denser. Formation and disruption of the columnar morphology under bias has been proposed by Bland et al. w16x Supposedly, the surface morphology of the film is generated either by replication of the substrate’s surface or by the preferential growth of crystallographic planes
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Fig. 3. AFM micrographs (3D) of Ti1yxAlxN films deposited under bias voltage of y50 V at nitrogen flow rate of (a) 2; (b) 6 and (c) 10 mlymin.
during deposition. During deposition, the geometrical shadowing will prevent deposition in the valleys of the substrate’s surface while its peaks will grow under no bias. Thus, a columnar morphology (Fig. 5a) was developed. As the bias increases, the ion bombardment erodes the peaks and forward sputtering will fill the valleys. Simultaneously, the backsputtered material may be returned by scattering process in the gas, which further randomizing the deposition direction and making the structure denser (as shown in Fig. 5c). Fig. 6 is the AFM micrographs (3D) of Ti1yxAlxN films deposited at nitrogen flow rate of 6 mlymin under various negative bias voltages (0–80 V). All surface morphologies were essentially uniformly distributed. An obvious change of the surface roughness of the films deposited under different bias voltages has been observed. Fig. 7 shows the corresponding roughness (RMS and Ave.) of films vs. bias voltages. The roughness (RMS) rapidly decreased from 1.414 nm (unbi-
Fig. 4. The roughness of Ti1yxAlxN films vs. nitrogen flow rates under bias voltage of y50 V.
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3.2. Residual stress
Fig. 5. SEM micrographs of TiAlN films deposited at nitrogen flow rate of 6 mlymin under various bias voltages.
ased) to about a constant of 0.65 nm (y30 to y80 V). The decrease of roughness is probably due to the enhanced mobility of deposited atoms with higher ion bombardment conditions according to the bias voltage. In addition, the slight increase of roughness at a bias voltage of y80 V was possibly caused by the ion irradiation on the sputtered surface due to the excessive bias voltage.
3.2.1. Residual stress of Ti1yxAlxN film on the Si substrate Fig. 8 shows the positive residual tensile stress of Ti1yxAlxN films deposited at various nitrogen flow rates under y50 V bias voltage. The residual stress decreased with the increase of nitrogen flow rate from the maximum of 218.0 MPa at 2 mlymin to the minimum of 11.4 MPa at 8 mlymin. It was probably due to the porous grain boundary between columnar texture of the film deposited under lower nitrogen flow rate (Fig. 2a), which was easily to transmit the tensile stress. Therefore, it is supposed to characterize with a bad mechanical behavior. A similar result of PVD coatings with lower residual compressive stress (i.e. bad mechanical behavior), corresponding to the microstructure of pronounced columns and cavities has been reported before w17x. It was previously discussed that the residual stress could be generated in the PVD films w13,17x upon cooling (from 350 8C to room temperature) due to the mismatch in linear thermal expansion coefficient (TEC) between the layers of Ti0.9Al0.1N (as9.0=10y6 yK) and Si (as2.6=10y6 yK). A positive tensile stress was theoretically due to the higher TEC value of the coating layer (Ti0.9Al0.1N). Similar positive thermal stress (;200–550 MPa) has been reported for AlN coatings deposited on silica, silicon and tungsten carbide substrate w18x. Essentially, the positive tensile stresses are possible but comparable smaller (O1 GPa) than the negative compressive intrinsic stresses ()5 GPa) w17,19x. According to Cheng et al. w13x and Oettel and Wiedemann w17x, the total internal stresses in the PVD films are mainly equal to the sum of the thermal stresses and the intrinsic stresses (if neglecting the priori negligibly small stress comes from hetero-epitaxy). Conversely, all compressive stresses with a monotonical increase to a maximum (7 GPa) at nitrogen pressure of 1.98=10y1 Pa was reported for a filtered cathodic vacuum arc deposited (Ti, Al)N films w13x. It was presumably due to the thermal stress component is negligible since the film deposition and stress measurement were carried out at room temperature in the previous study w13x. The residual stress of Ti1yxAlxN films deposited under various bias voltages at a nitrogen flow rate of 6 mly min was depicted in Fig. 9. The residual stress was approximately equal to 50 MPa as the bias voltage was smaller than 50 V. Afterwards, the residual stress increased with the increase of bias voltage, and it attained a value of 80.3 MPa as the bias voltage was equal to 80 V, presumably due to the sufficient particle energy to generate different types of defects in the volume of the grown layer (the result of atomic or ion shoot-peening effect) w20x.
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Fig. 6. AFM micrographs (3D) of Ti1yxAlxN films deposited at nitrogen flow rate of 6 mlymin under various bias voltages of (a) 0; (b) y30; (c) y50 and (d) y80 V.
According to Oettel and Wiedemann, the residual stresses in PVD hard coatings could be possibly determined by the specific energy transfer from the incoming particles to the surface atoms, which mainly controlled by the bias voltage in the case of magnetron sputtering w17x. A similar trend of residual stresses (although being negative compressive stress) increasing with the bias voltages in TiN, (Ti, Al)N and ZrN coatings has been reported before w17x. 3.2.2. Residual stress of Cu layer on the Ti1yxAlxNySi film The residual stresses of Cu layer on the Ti1yxAlxN films, which was deposited at a nitrogen flow rate of 6 mlymin under various negative bias voltages are shown in Fig. 10. A minimum residual stress of 350 MPa was obtained for Cu layer on Ti1yxAlxN film deposited with no bias. The residual stress greatly increased to a constant value of 480 MPa, approximately under various biases (30–80 V). It was presumably due to the different microstructure (e.g. surface roughness, Fig. 7) of the coatings caused by the bias voltage w17x.
Fig. 7. The roughness of Ti1yxAlxN films vs. bias voltages at a nitrogen flow rate of 6 mlymin.
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Fig. 8. The residual stress of Ti1yxAlxN films deposited at various nitrogen flow rates under y50 V bias voltage.
Fig. 9. The residual stress of Ti1yxAlx N films deposited under various bias voltages at a nitrogen flow rate of 6 mlymin.
3.3. Adhesive strength
Obviously, the Cu layer has been completely detached from TiAlN. In addition, chipping failure w22x has also been observed on the SEM micrograph of the scratch channel of Cu layer on the TiAlNySi layers (Fig. 13). It indicates that a bad adhesion existed between the layer of Cu and TiAlN. Oettel and Wiedemann w17x investigated the residual stress in PVD hard coatings and reported that the compressive stress could prevent
3.3.1. TiAlN film on the Si substrate Fig. 11 shows the SEM micrographs of the Ti1yxAlxN films after pull-test. Samples were deposited at a nitrogen flow rate of 6 mlymin without bias voltage. Both the line scan (a) and the mapping (b) of carbon on the tested specimen indicate that the adhesive failure for the Ti1yxAlxNySi film was along the epoxy cement. It is suggested that a strong adhesive strength of TiAlN film on the Si substrate, was higher than the strength of the epoxy cement. Additionally, no detachment on the coating was observed for a normal force of 60 N (defined as a critical adhesive load) in a scratch test before the crack of Si substrate, which also conformed the pulltest result. Similar to the previous result of AES depth profiles of the CuyTi1yxAlxNyNSiM multilayer w11x, titanium and oxygen atoms were both detected among the interphase of Ti1yxAlxNyNSiM. With the essence of the non-stoichiometric composition of Ti1yxAlxN (Fig. 1), the excessive Ti atoms might react with SiO2 to form a stable TiO phase. Actually, the stable chemical TiO bonding was reported w21,22x to be a major reason for a coating with higher adhesive strength. 3.3.2. Cu on the TiAlNySi layers The SEM micrographs (Fig. 12) show the Cu line scan analysis of the Cu (250 nm)yTi1yxAlxN (100 nm)y Si multilayer structure. Samples were deposited at a nitrogen flow rate of 6 mlymin under y50 V. Asdeposited Cu layer has been shown in grey area, however Cu was absent from the dark area after a pull-test.
Fig. 10. The residual stresses of Cu layer on the Ti1yxAlxN films, which was deposited at nitrogen flow rate of 6 mlymin under various bias voltage.
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Fig. 13. The SEM micrographs showing the scratch channel of the Cu layer on the Ti1yxAlxNySi film.
Fig. 11. The SEM micrographs showing the Ti1yxAlx N film after pulltest. Samples were deposited at a nitrogen flow rate of 6 mlymin without bias voltage. (a) The line scan of carbon and (b) the mapping of carbon.
cracking, if the external load acting alone generates extended regions with tensile stresses. Higher compressive residual stresses explain the higher critical load, (the higher adhesive strength) w17,23x. Inversely, a tensile residual stress in the layers will theoretically reduce the adhesive strength of the coating. The low adhesive strength of Cu layer on the TiAlNySi film could be ascribed to its larger tensile stress, comparatively to the high adhesive strength of TiAlN film on the Si substrate with lower tensile stress (as depicted in Section 3.2). 4. Summary and conclusions (1) The residual tensile stress within Ti1yxAlxN films decreased with the increase of nitrogen flow rate. Higher residual stresses were obtained within Ti1yxAlxN films deposited under higher bias voltages. The minimum residual stress was 11.4 MPa within the Ti1yxAlxN films deposited at nitrogen flow rate of 8 mlymin under bias voltage of y50 V. (2) The residual stress of Cu films decreased with the increase of the surface roughness of the Ti1yxAlxN films. The adhesive strength was rather high between Ti1yxAlxN and Si substrate due to the chemical bonding. However, because of the high residual tensile stress among Cu film, the adhesive strength between Cu and Ti1yxAlxN was comparatively low. Acknowledgments
Fig. 12. The SEM micrographs showing the Cu line scan analysis of the Cu (250 nm)yTi1yxAlxN (100 nm)ySi multilayer structure. Samples were deposited at a nitrogen flow rate of 6 mlymin under y50 V.
The authors would like to thank the National Science Council of the ROC for its financial support under the contract no. NSC 89-2216-E-012-001.
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