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Development of ruthenium dioxide electrodes for pyroelectric devices based on lithium tantalate thin films
Thin Solid Films 515 (2007) 3971 – 3977 www.elsevier.com/locate/tsf
Development of ruthenium dioxide electrodes for pyroelectric devices based on lithium tantalate thin films Laurianne Nougaret, Philippe Combette ⁎, Richard Arinero, Jean Podlecki, Frédérique Pascal-Delannoy Centre d’Electronique et de Micro-optoélectronique de Montpellier, Unité mixte de Recherche du CNRS n° 5507, Université Montpellier II, Place E. Bataillon, 34095 Montpellier cedex 05, France Received 19 October 2005; received in revised form 29 August 2006; accepted 14 September 2006 Available online 20 October 2006
Abstract The aim of this paper is the study of ruthenium dioxide (RuO2) films, grown on low-stress silicon nitride on silicon (SiNx/Si), in order to develop thermal micro-sensors based on pyroelectric effect. The active part of these micro-sensors is constituted by a new arrangement : lithium tantalate (LiTaO3)/RuO2/SiNx/Si. Radio-frequency (RF) sputtering is employed to deposit RuO2 on SiNx/Si substrate. Morphology, crystallinity and resistivity of RuO2 are studied as function of growth parameters. Next, RF magnetron sputtering was used to deposit LiTaO3 on this electrode. Morphology studies, pyroelectric effect and dielectric parameters obtained, indicate that RuO2 material is a suitable candidate as back electrode for LiTaO3 thin films. © 2006 Elsevier B.V. All rights reserved. Keywords: Ruthenium oxide; Sputtering; Lithium tantalate; Pyroelectric properties
1. Introduction Earlier studies proved that for a given ferroelectric and pyroelectric material, electrical and reliability properties were strongly influenced by the bottom electrode material. Two principal groups could be identified for these electrodes. The first group consists of noble metals such Pt, Ir, Ru, the second group consist of rutile-type metal oxide such IrO2 and RuO2 [1]. The first group of noble metals usually forms a Schottky contact with many pyroelectrics and thus have a limiting effect on the leakage current. However, with these materials, interface charges play an increasingly dominant role as the thickness of the dielectric layer decreases below 100 nm [1], reducing the effective dielectric constant of the capacitor, and limiting the level of integration. For example, platinum has been one of the ⁎ Corresponding author. E-mail addresses: [email protected] (L. Nougaret), [email protected] (P. Combette), [email protected] (R. Arinero), [email protected] (J. Podlecki), [email protected] (F. Pascal-Delannoy). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.09.023
most widely used bottom electrode materials because it withstands high temperatures and oxidizing conditions. However, oxygen easily diffuses into the platinum films and oxidizes the adhesion layer (Ti, Cr, etc.) during deposition and thermal treatment, which is detrimental to device performances [1]. The second group of oxide electrodes does not form electrical blocking contact to many ferroelectrics and hence does not reduce the dielectric constant of the capacitor even when the capacitor thickness is about few tens of nanometers. RuO2 which belongs to this second group, is one of the most conductive oxides exhibiting a very low metallic resistivity in the range from 30–100 Fm cm [2–4]. RuO2 presents a high thermal stability, good corrosion resistance and diffusion barrier properties. This material is one of the most promising candidates for electrode in ferroelectric access memories and highdensity capacitors [5]. In particular, RuO2 has been considered as an alternative to noble metal due to the excellent fatigue properties of ferroelectric films on these electrodes [6,7]. Furthermore despite the fact, that the leakage current are higher with oxide electrodes than Pt electrodes, we chose to develop RuO2 electrodes for its relative good electrical properties and
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Table 1 Standard growth conditions for fabrication of RuO2 and LiTaO3 films by RF sputtering Target composition
Ru
Li2O2–Ta2O5
Target-substrate distance Total work pressure Gas content RF power Substrate temperature Annealing conditions Growth rate
10 cm 1.3 Pa Ar %–O2 %: 80–20 50 W 200–400 °C RTA 640 °C 1 min 20–40 nm/h
5 cm 0.7 Pa Ar %–O2 %: 50–50 50 W 100 °C RTA 640 °C 1 min 70 nm/h
because it prevents the inward diffusion of oxygen into underlying substrate. Most often, RuO2 thin films have been prepared by RadioFrequency (RF) reactive sputtering [8–12] on SiO2 or Si. Electrical, optical, microstructural and stress characteristics of RuO2 films have been also studied [13–19]. The aim of this paper is to study ruthenium dioxide (RuO2) films grown on low-stress silicon nitride on silicon (SiNx/Si) in order to develop thermal micro-sensors based on pyroelectric effect. This phenomenon allows the conversion of variable heat flux in an electrical current [20]. The active part of these microsensors will be constituted by a new arrangement: lithium
tantalate (LiTaO3)/RuO2/SiNx/Si. We focused our efforts on growth of RuO2 on SiNx/Si substrate by RF reactive sputtering. Morphology, crystallinity and resistivity of RuO2 are studied as function growth parameters. Next, first results on morphology, crystallinity and electrical characterizations of LiTaO3 grown on RuO2/SiNx/Si by RF magnetron sputtering are also presented. 2. Experimental details RF reactive sputtering equipment was used to elaborate RuO2 thin films. The target is 2-inch disk of Ru 99.99% pure. The sputtering conditions are shown in Table 1. The films were deposited on (100) Si recovered by 500 nm SiNx low stress grown by low pressure chemical vapour deposition. The substrates were first degreased in an organic solution and then in an acid solution of HNO3 followed by hydrofluoric acid chemical etching. Substrates were next annealed in oxygen at 400 °C directly in the RF reactive sputtering equipment before the deposition. This operation allows to remove the organic compounds and to improve the RuO2 growth [21]. After a pre-sputtering of the target, RuO2 is then grown during 2 h. After deposition, RuO2 thin films are annealed in nitrogen or oxygen ambient in Rapid Thermal Annealing (RTA) furnace Jipelec (Jetfirst 100) at 640 °C during 1 min. Thermal annealing
Fig. 1. SEM surface morphology of RuO2 films grown by RF sputtering at 1.3 Pa and 400 °C, with 50% O2 ratio annealed in oxygen a) or nitrogen b), with 20% O2 ratio annealed in oxygen c) or nitrogen d).
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Table 2 Different annealing conditions of RuO2 Annealing
Thermal ramp
Temperature annealing
Oxygen annealing 1 Oxygen annealing 2 Nitrogen annealing
4 °C/s 4 °C/s 4 °C/s
640 °C during 5 min 640 °C during 1 min 640 °C during 5 min
improves crystallinity and stabilizes material before the growth of LiTaO3. The film thickness was measured with a Dektac profilometer. The crystalline structure of the layers is characterized by a standard X-ray diffraction (X-Pert Philips equipment with CuKM radiation and 40 kV operating voltage). The surface morphology is observed by Scanning Electrons Microscopy (SEM) (Cambridge S200) and Atomic Force Microscopy (AFM) (Autoprobe cp, Park Scientific Instruments) in tapping mode with a silicon tip (59.18 kHz resonance frequency) and a 1 Hz scanning frequency. In order to obtain particles sizes, we perform images analysis with WsxM software a scanning probe microscopy freeware. Finally, we studied the resistivity of RuO2 layers with a four probes method provided by a HL 5500 system (Bio-Rad). 3. Results and discussion 3.1. Morphology
Fig. 2. SEM images of RuO2 after different annealing in oxygen: a) 5 min annealed in O2, b) 5 min annealed in O2 with a low ramp during the cooling, c) 1 min annealed in O2. Growth conditions: substrate temperature: 300 °C, pressure: 1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm.
The surface of as-deposited RuO2 films have dark black colour. After heat treatment, the colour of the films is unchanged and no peeling off is found by adhesive tape tests. According to P. Gopal Ganesan and al. [22] we observed, after RTA treatment, an increasing of film thickness (about 20%) and a change of grain micro-structure which depends on the nature of gas annealing, on the annealing time process and on oxygen percentage during the growth. The SEM images of the deposited films annealed at 640 °C in nitrogen and oxygen are shown in Fig. 1a, b, c, d. The films were deposited at 400 °C with respectively 20 and 50% of oxygen content. The layers are composed of granular-type grains with a reduced porosity for the films annealed under oxygen (1a and 1c). The presence of holes is enhanced in nitrogen gas during the annealing. In fact, the loss of O2 during the annealing explains this phenomenon. This result is confirmed by the observations of RuO2 films deposited with different oxygen content (Fig. 1b and d). The probability of defects incorporation increases with pressure ratio of oxygen during the growth as seen in Fig. 5, and as [23]. Before annealing, the resistivity value is higher for 50% O2 than for 20%. Since RuO2 is the only stable oxide of Ru phase, the excess oxygen should be distributed along the grain boundaries [24]. The excess of O2 evaporates during the annealing in the form of RuO3, RuO4 [25,26], which are unstable and volatile gases. This explains why RuO2 presents a high etching rate in an O2 plasma [27]. This evaporation of RuO3 and RuO4 induces the apparition of holes at the surface layer (see SEM and AFM photos), this phenomenon has been already observed by Oh et al. [3].
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This phenomenon is emphasized by an increase of annealing time as shown in Fig. 2. The corresponding annealing conditions are presented in Table 2. As the annealing time increased, morphology of films started to be disturbed with apparition of holes at grain boundaries. The best result was obtained with the shortest annealing time tested (1 min) as it is confirmed by the Fig. 2c. AFM images reveal spherical micro-structures of various sizes at the surface of the films. For a constant oxygen ratio during the growth, the grain size increases when nitrogen is replaced by oxygen gas as annealing gas as it is depicted in Fig. 3a and b. For these 120 nm thick films, the average particle size observed on 2 Fm-side AFM images is 120 nm and 140 nm, respectively. When the total gas pressure was hanged from 0.25 to 0.1 and 1.3 Pa while the other growth conditions were kept similar as the samples described above (T = 400 °C, O2/Ar: 20–80%, annealing at 640 °C in oxygen gas), we also observed an increase in particle size from 95 to 130 nm. For the growth conditions). At 1.3 Pa, the grain size decreased from 95 to 85 nm when the oxygen content in
Fig. 4. X-ray diffraction patterns of RuO2 films deposited at: a) 300 °C, 1.3 Pa, 20%O2/80% Ar ratio, 50 W and b) 300 °C, 1.3 Pa, 50%O2/50% Ar ratio, 50 W.
the fed gas was increased from 20 to 50%, similar to observations made by W. T. Lim et al. [28]. 3.2. Crystallinity As expected, the typical X-ray diffraction pattern of RuO2 presented in Fig. 4 clearly shows the presence of RuO2 peaks. The films deposited at all substrate temperatures and after annealing at 640 °C in oxygen exhibit polycrystalline RuO2 phases with preferential orientations of (110), (002) and (101). The films deposited in a 20% (O2/Ar) ratio pressure exhibit (110) RuO2 preferred orientation [28], while those elaborated in a 50% (O2/Ar) ratio pressure show a (101) preferential orientation. This results has been already mentioned [24,28–30] and according to Lee et al. [24], (110) orientation is preferential orientation of bulk RuO2 near stoichiometry oxygen composition, while preferential (101) grain growth occurs in the film with oxygen concentration near saturation. 3.3. Resistivity
Fig. 3. AFM micrograph of RuO2 films grown at 1.3 Pa, 400 °C, with 20% O2 ratio; a) annealed in O2; b) annealed in N2.
The electrical conductivity in RuO2 films is linked to its grain structure and grain boundaries composition. According to Mayadas [31], when the grain size is much larger that the electron mean free path, the contribution to the resistivity by the grain boundaries scattering is small. However, when the distance between grain boundaries become lesser, the grain boundary contribution can no longer be regarded as negligible. Furthermore, the effect of grain boundaries scattering would even dominate as the grain size became sufficiently small in
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Fig. 6. X-ray diffraction pattern of LiTaO3 film.
Fig. 5. Resistivity versus growth temperature a) before and b) after O2 annealing of RuO2 grown with 50% and 20 % O2. Growth conditions: substrate temperature: 200–400 °C, pressure: 1.3 Pa, RF power: 50 W, thickness: 100 nm.
comparison with the mean free path of the carriers [31]. As shown in Fig. 5, before annealing the resistivity of film grown in a 50%/50% (Ar/O2) ratio decreases as the growth temperature increases whereas in a 80%/20% (Ar/O2), resistivity is almost constant. Indeed, the increasing of growth temperature increases grains size of the films, thus decreasing resistivity. This is confirmed for 50%/50% (Ar/O2) ratio, but not for 80%/20% (Ar/O2). Taking into account this result, we suggest that with these growth conditions, the mean grain size is near the mean free path of the carriers and thus the contribution to the resistivity by the grain boundaries is not negligible [32]. Thus, the influence of grain boundaries is crucial: grain boundaries scattering due to excess oxygen may contribute to the higher resistivity. As the growth temperature increases, desorption of oxygen in the films becomes more important [30], and defects such as RuO3 and RuO4 in the grains boundaries tends to decrease. These results are in agreement with L. Krusin-Elbaum et al. [8] who explain this phenomenon by the grain boundary scattering due to excess oxygen in these regions. The resistivity decreases when the growth temperature varies from 200 °C from 400 °C as seen in Fig. 5, this is in agreement with Lee et al. [23]. The size of grains increasing with growth temperature, inducing a decrease of resistivity [28]. This increase in conductivity with temperature has been observed for films on SiO2 or Si [33] and also interpreted in terms of grain
boundaries effects [16,20] which in the end produce apparent semi-conductivity [34]. The resistivity varies also with time, thermal ramp of annealing and gas nature. Its value is always lower after annealing reaching the limit value of 110–120 μR cm for a layer thickness of 100 nm as shown in Fig. 5. This result can be explained by an increasing of the grain size and by the decreasing grain boundary defects [31]. Furthermore, we could observe after annealing in oxygen gas that the final resistivity is quasi constant and independent of oxygen percentage during the growth. However, when the annealing time becomes too long (over 15 mn) the thermal stress is very high and we observed a peeling of the metallic layer. We observed that the resistivity measured after nitrogen annealing is lower than the value measured after oxygen annealing. This resistivity variation is certainly due to a structure modification of the ruthenium oxide observed previously in the morphological observations and by the growth of its grain size. However, this layer is not stable in time and partially peel off making it unusable for the next deposition of LiTaO3. The best value of resistivity is equal to 110 μR cm after 1 min annealing at 640 °C in oxygen (Growth conditions: substrate temperature: 400 °C, pressure:
Fig. 7. SEM images of LiTaO3 showing the surface morphology.
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1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm). 3.4. First results obtained on LiTaO3 films First layers of LiTaO3 grown by RF magnetron sputtering were electrically characterized. The growth conditions are presented in Table 1. After magnetron deposition, the pyroelectric films are annealed at 640 °C in O2 for 1 min. Typical thickness of lithium tantalate is about 500 nm. Fig. 6 shows a typical X-ray diffraction pattern of LiTaO3 thin film grown on RuO2 electrode which confirms polycrystalline character of film. SEM characterization displayed on Fig. 7 shows the morphology of film. In order to test the quality of LiTaO3, we measured different electrical properties of the material. Firstly, Fig. 8 presents the frequency dependence of relative dielectric constant and dielectric losses at room temperature, measurement were achieved from a low frequency impedance analyzer 4192 A Hewlett Packard. Relative dielectric constant of the thin film is relatively closed to the value of bulk crystal (εr = 44 [35]). However, the relative high losses could be attributed to the electrode nature and need to be improved. Secondly, Fig. 9 shows the pyroelectric current of the structure Al/LiTaO3/RuO2 where Al and RuO2 are respectively the front and the back electrodes. Aluminium electrodes were thermally evaporated, with a diameter of pad size equal to 500 μm.
Fig. 9. a) Temperature derivative versus time, b) Pyroelectric current generated by LiTaO3 layer.
Fig. 9a shows the evolution of temperature derivative versus time. The temperature cycle is created by a Peltier module and varies from 16 °C to 40 °C during a half period of 100 s and induces a temperature time derivative about 1.5 K/s. An electrometer 6514A Keithley Instruments, is used to measure pyroelectric current. Fig. 9b depicts the pyroelectric current produced by the lithium tantalate following the temperature signal and varies from — 3 to 3 pA. This result is according to the results of Lang [20]. Indeed the pyroelectric current depends on the temperature derivative as following this relation: i = pA (dΔ/dt), where p is the pyroelectric coefficient and A is the surface electrode. From this last equation, we could obtain the value of pyroelectric coefficient, which is about 9 μC/m2 K. This value is less than bulk value (pbulk = 190 μC/m2 K [20]) and need to be improved. 4. Conclusions
Fig. 8. Dielectric parameters of LiTaO3 layer.
We studied the morphology and the crystalline properties of RuO2 grown on SiNx/Si by RF sputtering by varying the parameters of growth and post-annealing. We obtained results in agreement with studies already published on RuO2 grown on SiO2. The oxygen content improves preferential crystalline orientations. For a sputtering gas composition with a 50%/50% Ar/O2 ratio, the peak (101) is preponderant whereas the decreasing of oxygen content promotes the (110) direction which the principal direction for the bulk material. This phenomenon has been related to the growth of RuO2 in a near stoichiometry oxygen composition for the 80%/20% Ar/O2 ratio.
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Then, we specially studied the effect of the sputtering gas composition on the electrical properties of RuO2. We observed that the resistivity decreases when oxygen content decreases during the growth. This has been attributed to the structure of grains and grain boundaries. The resistivity decreases also when growth temperature increases. After the growth of RuO2, annealing was necessary in order to optimize the crystallization of the films. The best value of resistivity is equal to 110 μΩ cm after 1 min annealing at 640 °C in oxygen (Growth conditions: substrate temperature: 400 °C, pressure: 1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm). First structures of Al/LiTaO3/RuO2 grown on SiNx were tested. Electrical properties of stacked layers confirm the good qualities of lithium tantalate thin films. A pyroelectric current varying from − 3 to 3 pA was observed during a thermal cycle from 16 °C to 40 °C, revealing a pyroelectric coefficient about 9 μC/m2 K. These results are comparable to those obtained by Kohli et al. [36] who studied another arrangement: LiTaO3/ RuO2/SiO2. References [1] C.A. Paz de Araujo, J.F. Scott, G.W. Taylor, Ferroelectric Thin Films: Synthesis and Basic Properties, Gordon and Breach, Amsterdam, 1996, p. 193. [2] D.P. Vijay, S.B. Desu, J. Electrochem. Soc. 140 (1993) 2640. [3] Y.-J. Oh, S.H. Moon, C.-H. Chung, J. Electrochem. Soc. 148 (4) (2001) 56. [4] A. Gril, D. Beach, C. Smart, W. Kane, in: E.R. Meyers, B.A. Tuttle, S.B. Besu, P.K. Larsen (Eds.), Ferroelectric thin films III, San Fransisco, U.S.A., 13–16 April, 1993, Mat. Res. Soc. Symp. Proc., vol. 310, 1993, p. 189. [5] D.S. Yoon, J.S. Roh, Semicond. Sci. Technol. 17 (2002) 599. [6] H.N. Al-Shareef, A.I. Kingon, X. Chen, K.R. Bellur, J. Mater. Res. 9 (1994) 2968. [7] H.N. Al-Shareef, O. Auciello, A.I. Kingon, J. Appl. Phys. 77 (1995) 2146. [8] L. Krusin-Elbaum, Thin Solid Films 169 (1989) 17. [9] E. Kolawa, F.C.T. So, E.T.S. Pan, M.A. Nicolet, Appl. Phys. Lett. 50 (1987) 854. [10] E. Kolawa, F.C.T. So, W. Flick, X.A. Zhao, E.T.S. Pan, M.A. Nicolet, Thin Solid Films 173 (1989) 217.
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[11] K. Sakiyama, S. Onishi, K. Ishihara, K. Orita, T. Kajiyama, N. Hosoda, T. Hara, J. Electrochem. Soc. 140 (3) (1993) 834. [12] C.W. Law, K.Y. Tong, J.H. Li, K. Li, M.C. Poon, Thin Solid Films 354 (1999) 162. [13] M.L. Green, M.E. Gross, L.E. Papa, K.J. Schnoes, D. Barsen, J. Electrochem. Soc. 132 (11) (1985) 2677. [14] K.C. Smith, Y.M. Sun, N.R. Mettlach, R.L. Hance, J.M. White, Thin Solid Films 376 (2000) 73. [15] L. Krusin-Elbaum, M. Wittmer, D.S. Yee, Appl. Phys. Lett. 50 (1989) 1879. [16] A. Belkind, Z. Orban, J.L. Vossen, J.A. Woollam, Thin Solid Films 207 (1992) 242. [17] E. Kolawa, C.W. Nieh, F.C.T. So, M. Nicolet, J. Electron. Mater. 17 (1988) 425. [18] S.K. Hong, H.J. Kim, H.G. Yang, J. Appl. Phys. 80 (2) (1996) 822. [19] S.B. Desu, D.P. Vijay, S. Ramanathan, H.D. Bhatt, S. Tirumala, Thin Solid Films 350 (1999) 21. [20] S.B. Lang, Sourcebook of Pyroelectricity, Gordon and Breach, New York, 1974. [21] Y.S. Yoon, J.H. Kim, A.M. Schmidt, D.L. Polla, Q. Wang, W.L. Gladfelter, Y.H. Shin, J. Mater. Sci., Mater. Electron. 9 (1998) 465. [22] P. Gopal Ganesan, M. Eizenberg, Microelectron. Eng. 64 (2002) 449. [23] J.-G. Lee, S.-K. Min, S.H. Choh, Jpn. J. Appl. Phys. 33 (1994) 7080. [24] J.G. Lee, Y.T. Kim, S. Min, S.H. Choh, J. Appl. Phys. 77 (10) (1995) 5473. [25] J.A. Rard, Chem. Rev. 85 (1) (1985) 1. [26] D.B. Rogers, R.D. Shannon, A.W. Sleight, J.L. Gillson, Inorg. Chem. 8 (1969) 841. [27] J. Baborowski, P. Muralt, N. Lederman, S. Hiboux, Vacuum 56 (1) (2000) 51. [28] W.T. Lim, K.R. Cho, C.H. Lee, Thin Solid Films 348 (1999) 56. [29] Y. Matsui, M. Hiratani, S. Kimura, Jpn. J. Appl. Phys. 39 (2000) 256. [30] J.H. Huang, J.S. Chen, Thin Solid Films 382 (2001) 139. [31] A.F. Mayadas, M. Shartzkes, Phys. Rev., B 1 (1970) 1382. [32] Y. Mar, J.S. Liang, C.Y. Sun, Y.S. Huang, Thin Solid Films 238 (1994) 158. [33] S. Pizzini, G. Buzzanca, C. Mari, L. Rossi, S. Torchio, Mater. Res. Bull. 7 (5) (1972) 449. [34] G. Lodi, C. de Asmundis, S. Ardizzone, E. Sivieri, S. Trasatti, Surf. Technol. 14 (4) (1981) 335. [35] I. Tomeno, S. Matsumura, Phys. Rev., B 38 (1) (1988) 606. [36] C.H. Kohli, J. Moser, K. Prasad, P.E. Schmid, F. Levy, G. Burri, Microelectron. Eng. 29 (1995) 201.
Development of ruthenium dioxide electrodes for pyroelectric devices based on lithium tantalate thin films Laurianne Nougaret, Philippe Combette ⁎, Richard Arinero, Jean Podlecki, Frédérique Pascal-Delannoy Centre d’Electronique et de Micro-optoélectronique de Montpellier, Unité mixte de Recherche du CNRS n° 5507, Université Montpellier II, Place E. Bataillon, 34095 Montpellier cedex 05, France Received 19 October 2005; received in revised form 29 August 2006; accepted 14 September 2006 Available online 20 October 2006
Abstract The aim of this paper is the study of ruthenium dioxide (RuO2) films, grown on low-stress silicon nitride on silicon (SiNx/Si), in order to develop thermal micro-sensors based on pyroelectric effect. The active part of these micro-sensors is constituted by a new arrangement : lithium tantalate (LiTaO3)/RuO2/SiNx/Si. Radio-frequency (RF) sputtering is employed to deposit RuO2 on SiNx/Si substrate. Morphology, crystallinity and resistivity of RuO2 are studied as function of growth parameters. Next, RF magnetron sputtering was used to deposit LiTaO3 on this electrode. Morphology studies, pyroelectric effect and dielectric parameters obtained, indicate that RuO2 material is a suitable candidate as back electrode for LiTaO3 thin films. © 2006 Elsevier B.V. All rights reserved. Keywords: Ruthenium oxide; Sputtering; Lithium tantalate; Pyroelectric properties
1. Introduction Earlier studies proved that for a given ferroelectric and pyroelectric material, electrical and reliability properties were strongly influenced by the bottom electrode material. Two principal groups could be identified for these electrodes. The first group consists of noble metals such Pt, Ir, Ru, the second group consist of rutile-type metal oxide such IrO2 and RuO2 [1]. The first group of noble metals usually forms a Schottky contact with many pyroelectrics and thus have a limiting effect on the leakage current. However, with these materials, interface charges play an increasingly dominant role as the thickness of the dielectric layer decreases below 100 nm [1], reducing the effective dielectric constant of the capacitor, and limiting the level of integration. For example, platinum has been one of the ⁎ Corresponding author. E-mail addresses: [email protected] (L. Nougaret), [email protected] (P. Combette), [email protected] (R. Arinero), [email protected] (J. Podlecki), [email protected] (F. Pascal-Delannoy). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.09.023
most widely used bottom electrode materials because it withstands high temperatures and oxidizing conditions. However, oxygen easily diffuses into the platinum films and oxidizes the adhesion layer (Ti, Cr, etc.) during deposition and thermal treatment, which is detrimental to device performances [1]. The second group of oxide electrodes does not form electrical blocking contact to many ferroelectrics and hence does not reduce the dielectric constant of the capacitor even when the capacitor thickness is about few tens of nanometers. RuO2 which belongs to this second group, is one of the most conductive oxides exhibiting a very low metallic resistivity in the range from 30–100 Fm cm [2–4]. RuO2 presents a high thermal stability, good corrosion resistance and diffusion barrier properties. This material is one of the most promising candidates for electrode in ferroelectric access memories and highdensity capacitors [5]. In particular, RuO2 has been considered as an alternative to noble metal due to the excellent fatigue properties of ferroelectric films on these electrodes [6,7]. Furthermore despite the fact, that the leakage current are higher with oxide electrodes than Pt electrodes, we chose to develop RuO2 electrodes for its relative good electrical properties and
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Table 1 Standard growth conditions for fabrication of RuO2 and LiTaO3 films by RF sputtering Target composition
Ru
Li2O2–Ta2O5
Target-substrate distance Total work pressure Gas content RF power Substrate temperature Annealing conditions Growth rate
10 cm 1.3 Pa Ar %–O2 %: 80–20 50 W 200–400 °C RTA 640 °C 1 min 20–40 nm/h
5 cm 0.7 Pa Ar %–O2 %: 50–50 50 W 100 °C RTA 640 °C 1 min 70 nm/h
because it prevents the inward diffusion of oxygen into underlying substrate. Most often, RuO2 thin films have been prepared by RadioFrequency (RF) reactive sputtering [8–12] on SiO2 or Si. Electrical, optical, microstructural and stress characteristics of RuO2 films have been also studied [13–19]. The aim of this paper is to study ruthenium dioxide (RuO2) films grown on low-stress silicon nitride on silicon (SiNx/Si) in order to develop thermal micro-sensors based on pyroelectric effect. This phenomenon allows the conversion of variable heat flux in an electrical current [20]. The active part of these microsensors will be constituted by a new arrangement: lithium
tantalate (LiTaO3)/RuO2/SiNx/Si. We focused our efforts on growth of RuO2 on SiNx/Si substrate by RF reactive sputtering. Morphology, crystallinity and resistivity of RuO2 are studied as function growth parameters. Next, first results on morphology, crystallinity and electrical characterizations of LiTaO3 grown on RuO2/SiNx/Si by RF magnetron sputtering are also presented. 2. Experimental details RF reactive sputtering equipment was used to elaborate RuO2 thin films. The target is 2-inch disk of Ru 99.99% pure. The sputtering conditions are shown in Table 1. The films were deposited on (100) Si recovered by 500 nm SiNx low stress grown by low pressure chemical vapour deposition. The substrates were first degreased in an organic solution and then in an acid solution of HNO3 followed by hydrofluoric acid chemical etching. Substrates were next annealed in oxygen at 400 °C directly in the RF reactive sputtering equipment before the deposition. This operation allows to remove the organic compounds and to improve the RuO2 growth [21]. After a pre-sputtering of the target, RuO2 is then grown during 2 h. After deposition, RuO2 thin films are annealed in nitrogen or oxygen ambient in Rapid Thermal Annealing (RTA) furnace Jipelec (Jetfirst 100) at 640 °C during 1 min. Thermal annealing
Fig. 1. SEM surface morphology of RuO2 films grown by RF sputtering at 1.3 Pa and 400 °C, with 50% O2 ratio annealed in oxygen a) or nitrogen b), with 20% O2 ratio annealed in oxygen c) or nitrogen d).
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Table 2 Different annealing conditions of RuO2 Annealing
Thermal ramp
Temperature annealing
Oxygen annealing 1 Oxygen annealing 2 Nitrogen annealing
4 °C/s 4 °C/s 4 °C/s
640 °C during 5 min 640 °C during 1 min 640 °C during 5 min
improves crystallinity and stabilizes material before the growth of LiTaO3. The film thickness was measured with a Dektac profilometer. The crystalline structure of the layers is characterized by a standard X-ray diffraction (X-Pert Philips equipment with CuKM radiation and 40 kV operating voltage). The surface morphology is observed by Scanning Electrons Microscopy (SEM) (Cambridge S200) and Atomic Force Microscopy (AFM) (Autoprobe cp, Park Scientific Instruments) in tapping mode with a silicon tip (59.18 kHz resonance frequency) and a 1 Hz scanning frequency. In order to obtain particles sizes, we perform images analysis with WsxM software a scanning probe microscopy freeware. Finally, we studied the resistivity of RuO2 layers with a four probes method provided by a HL 5500 system (Bio-Rad). 3. Results and discussion 3.1. Morphology
Fig. 2. SEM images of RuO2 after different annealing in oxygen: a) 5 min annealed in O2, b) 5 min annealed in O2 with a low ramp during the cooling, c) 1 min annealed in O2. Growth conditions: substrate temperature: 300 °C, pressure: 1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm.
The surface of as-deposited RuO2 films have dark black colour. After heat treatment, the colour of the films is unchanged and no peeling off is found by adhesive tape tests. According to P. Gopal Ganesan and al. [22] we observed, after RTA treatment, an increasing of film thickness (about 20%) and a change of grain micro-structure which depends on the nature of gas annealing, on the annealing time process and on oxygen percentage during the growth. The SEM images of the deposited films annealed at 640 °C in nitrogen and oxygen are shown in Fig. 1a, b, c, d. The films were deposited at 400 °C with respectively 20 and 50% of oxygen content. The layers are composed of granular-type grains with a reduced porosity for the films annealed under oxygen (1a and 1c). The presence of holes is enhanced in nitrogen gas during the annealing. In fact, the loss of O2 during the annealing explains this phenomenon. This result is confirmed by the observations of RuO2 films deposited with different oxygen content (Fig. 1b and d). The probability of defects incorporation increases with pressure ratio of oxygen during the growth as seen in Fig. 5, and as [23]. Before annealing, the resistivity value is higher for 50% O2 than for 20%. Since RuO2 is the only stable oxide of Ru phase, the excess oxygen should be distributed along the grain boundaries [24]. The excess of O2 evaporates during the annealing in the form of RuO3, RuO4 [25,26], which are unstable and volatile gases. This explains why RuO2 presents a high etching rate in an O2 plasma [27]. This evaporation of RuO3 and RuO4 induces the apparition of holes at the surface layer (see SEM and AFM photos), this phenomenon has been already observed by Oh et al. [3].
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This phenomenon is emphasized by an increase of annealing time as shown in Fig. 2. The corresponding annealing conditions are presented in Table 2. As the annealing time increased, morphology of films started to be disturbed with apparition of holes at grain boundaries. The best result was obtained with the shortest annealing time tested (1 min) as it is confirmed by the Fig. 2c. AFM images reveal spherical micro-structures of various sizes at the surface of the films. For a constant oxygen ratio during the growth, the grain size increases when nitrogen is replaced by oxygen gas as annealing gas as it is depicted in Fig. 3a and b. For these 120 nm thick films, the average particle size observed on 2 Fm-side AFM images is 120 nm and 140 nm, respectively. When the total gas pressure was hanged from 0.25 to 0.1 and 1.3 Pa while the other growth conditions were kept similar as the samples described above (T = 400 °C, O2/Ar: 20–80%, annealing at 640 °C in oxygen gas), we also observed an increase in particle size from 95 to 130 nm. For the growth conditions). At 1.3 Pa, the grain size decreased from 95 to 85 nm when the oxygen content in
Fig. 4. X-ray diffraction patterns of RuO2 films deposited at: a) 300 °C, 1.3 Pa, 20%O2/80% Ar ratio, 50 W and b) 300 °C, 1.3 Pa, 50%O2/50% Ar ratio, 50 W.
the fed gas was increased from 20 to 50%, similar to observations made by W. T. Lim et al. [28]. 3.2. Crystallinity As expected, the typical X-ray diffraction pattern of RuO2 presented in Fig. 4 clearly shows the presence of RuO2 peaks. The films deposited at all substrate temperatures and after annealing at 640 °C in oxygen exhibit polycrystalline RuO2 phases with preferential orientations of (110), (002) and (101). The films deposited in a 20% (O2/Ar) ratio pressure exhibit (110) RuO2 preferred orientation [28], while those elaborated in a 50% (O2/Ar) ratio pressure show a (101) preferential orientation. This results has been already mentioned [24,28–30] and according to Lee et al. [24], (110) orientation is preferential orientation of bulk RuO2 near stoichiometry oxygen composition, while preferential (101) grain growth occurs in the film with oxygen concentration near saturation. 3.3. Resistivity
Fig. 3. AFM micrograph of RuO2 films grown at 1.3 Pa, 400 °C, with 20% O2 ratio; a) annealed in O2; b) annealed in N2.
The electrical conductivity in RuO2 films is linked to its grain structure and grain boundaries composition. According to Mayadas [31], when the grain size is much larger that the electron mean free path, the contribution to the resistivity by the grain boundaries scattering is small. However, when the distance between grain boundaries become lesser, the grain boundary contribution can no longer be regarded as negligible. Furthermore, the effect of grain boundaries scattering would even dominate as the grain size became sufficiently small in
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Fig. 6. X-ray diffraction pattern of LiTaO3 film.
Fig. 5. Resistivity versus growth temperature a) before and b) after O2 annealing of RuO2 grown with 50% and 20 % O2. Growth conditions: substrate temperature: 200–400 °C, pressure: 1.3 Pa, RF power: 50 W, thickness: 100 nm.
comparison with the mean free path of the carriers [31]. As shown in Fig. 5, before annealing the resistivity of film grown in a 50%/50% (Ar/O2) ratio decreases as the growth temperature increases whereas in a 80%/20% (Ar/O2), resistivity is almost constant. Indeed, the increasing of growth temperature increases grains size of the films, thus decreasing resistivity. This is confirmed for 50%/50% (Ar/O2) ratio, but not for 80%/20% (Ar/O2). Taking into account this result, we suggest that with these growth conditions, the mean grain size is near the mean free path of the carriers and thus the contribution to the resistivity by the grain boundaries is not negligible [32]. Thus, the influence of grain boundaries is crucial: grain boundaries scattering due to excess oxygen may contribute to the higher resistivity. As the growth temperature increases, desorption of oxygen in the films becomes more important [30], and defects such as RuO3 and RuO4 in the grains boundaries tends to decrease. These results are in agreement with L. Krusin-Elbaum et al. [8] who explain this phenomenon by the grain boundary scattering due to excess oxygen in these regions. The resistivity decreases when the growth temperature varies from 200 °C from 400 °C as seen in Fig. 5, this is in agreement with Lee et al. [23]. The size of grains increasing with growth temperature, inducing a decrease of resistivity [28]. This increase in conductivity with temperature has been observed for films on SiO2 or Si [33] and also interpreted in terms of grain
boundaries effects [16,20] which in the end produce apparent semi-conductivity [34]. The resistivity varies also with time, thermal ramp of annealing and gas nature. Its value is always lower after annealing reaching the limit value of 110–120 μR cm for a layer thickness of 100 nm as shown in Fig. 5. This result can be explained by an increasing of the grain size and by the decreasing grain boundary defects [31]. Furthermore, we could observe after annealing in oxygen gas that the final resistivity is quasi constant and independent of oxygen percentage during the growth. However, when the annealing time becomes too long (over 15 mn) the thermal stress is very high and we observed a peeling of the metallic layer. We observed that the resistivity measured after nitrogen annealing is lower than the value measured after oxygen annealing. This resistivity variation is certainly due to a structure modification of the ruthenium oxide observed previously in the morphological observations and by the growth of its grain size. However, this layer is not stable in time and partially peel off making it unusable for the next deposition of LiTaO3. The best value of resistivity is equal to 110 μR cm after 1 min annealing at 640 °C in oxygen (Growth conditions: substrate temperature: 400 °C, pressure:
Fig. 7. SEM images of LiTaO3 showing the surface morphology.
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1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm). 3.4. First results obtained on LiTaO3 films First layers of LiTaO3 grown by RF magnetron sputtering were electrically characterized. The growth conditions are presented in Table 1. After magnetron deposition, the pyroelectric films are annealed at 640 °C in O2 for 1 min. Typical thickness of lithium tantalate is about 500 nm. Fig. 6 shows a typical X-ray diffraction pattern of LiTaO3 thin film grown on RuO2 electrode which confirms polycrystalline character of film. SEM characterization displayed on Fig. 7 shows the morphology of film. In order to test the quality of LiTaO3, we measured different electrical properties of the material. Firstly, Fig. 8 presents the frequency dependence of relative dielectric constant and dielectric losses at room temperature, measurement were achieved from a low frequency impedance analyzer 4192 A Hewlett Packard. Relative dielectric constant of the thin film is relatively closed to the value of bulk crystal (εr = 44 [35]). However, the relative high losses could be attributed to the electrode nature and need to be improved. Secondly, Fig. 9 shows the pyroelectric current of the structure Al/LiTaO3/RuO2 where Al and RuO2 are respectively the front and the back electrodes. Aluminium electrodes were thermally evaporated, with a diameter of pad size equal to 500 μm.
Fig. 9. a) Temperature derivative versus time, b) Pyroelectric current generated by LiTaO3 layer.
Fig. 9a shows the evolution of temperature derivative versus time. The temperature cycle is created by a Peltier module and varies from 16 °C to 40 °C during a half period of 100 s and induces a temperature time derivative about 1.5 K/s. An electrometer 6514A Keithley Instruments, is used to measure pyroelectric current. Fig. 9b depicts the pyroelectric current produced by the lithium tantalate following the temperature signal and varies from — 3 to 3 pA. This result is according to the results of Lang [20]. Indeed the pyroelectric current depends on the temperature derivative as following this relation: i = pA (dΔ/dt), where p is the pyroelectric coefficient and A is the surface electrode. From this last equation, we could obtain the value of pyroelectric coefficient, which is about 9 μC/m2 K. This value is less than bulk value (pbulk = 190 μC/m2 K [20]) and need to be improved. 4. Conclusions
Fig. 8. Dielectric parameters of LiTaO3 layer.
We studied the morphology and the crystalline properties of RuO2 grown on SiNx/Si by RF sputtering by varying the parameters of growth and post-annealing. We obtained results in agreement with studies already published on RuO2 grown on SiO2. The oxygen content improves preferential crystalline orientations. For a sputtering gas composition with a 50%/50% Ar/O2 ratio, the peak (101) is preponderant whereas the decreasing of oxygen content promotes the (110) direction which the principal direction for the bulk material. This phenomenon has been related to the growth of RuO2 in a near stoichiometry oxygen composition for the 80%/20% Ar/O2 ratio.
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Then, we specially studied the effect of the sputtering gas composition on the electrical properties of RuO2. We observed that the resistivity decreases when oxygen content decreases during the growth. This has been attributed to the structure of grains and grain boundaries. The resistivity decreases also when growth temperature increases. After the growth of RuO2, annealing was necessary in order to optimize the crystallization of the films. The best value of resistivity is equal to 110 μΩ cm after 1 min annealing at 640 °C in oxygen (Growth conditions: substrate temperature: 400 °C, pressure: 1.3 Pa, gas ratio Ar/02: 80%/20%, RF power: 50 W, thickness: 100 nm). First structures of Al/LiTaO3/RuO2 grown on SiNx were tested. Electrical properties of stacked layers confirm the good qualities of lithium tantalate thin films. A pyroelectric current varying from − 3 to 3 pA was observed during a thermal cycle from 16 °C to 40 °C, revealing a pyroelectric coefficient about 9 μC/m2 K. These results are comparable to those obtained by Kohli et al. [36] who studied another arrangement: LiTaO3/ RuO2/SiO2. References [1] C.A. Paz de Araujo, J.F. Scott, G.W. Taylor, Ferroelectric Thin Films: Synthesis and Basic Properties, Gordon and Breach, Amsterdam, 1996, p. 193. [2] D.P. Vijay, S.B. Desu, J. Electrochem. Soc. 140 (1993) 2640. [3] Y.-J. Oh, S.H. Moon, C.-H. Chung, J. Electrochem. Soc. 148 (4) (2001) 56. [4] A. Gril, D. Beach, C. Smart, W. Kane, in: E.R. Meyers, B.A. Tuttle, S.B. Besu, P.K. Larsen (Eds.), Ferroelectric thin films III, San Fransisco, U.S.A., 13–16 April, 1993, Mat. Res. Soc. Symp. Proc., vol. 310, 1993, p. 189. [5] D.S. Yoon, J.S. Roh, Semicond. Sci. Technol. 17 (2002) 599. [6] H.N. Al-Shareef, A.I. Kingon, X. Chen, K.R. Bellur, J. Mater. Res. 9 (1994) 2968. [7] H.N. Al-Shareef, O. Auciello, A.I. Kingon, J. Appl. Phys. 77 (1995) 2146. [8] L. Krusin-Elbaum, Thin Solid Films 169 (1989) 17. [9] E. Kolawa, F.C.T. So, E.T.S. Pan, M.A. Nicolet, Appl. Phys. Lett. 50 (1987) 854. [10] E. Kolawa, F.C.T. So, W. Flick, X.A. Zhao, E.T.S. Pan, M.A. Nicolet, Thin Solid Films 173 (1989) 217.
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