HfO2 nanolaminates

Thin Solid Films 518 (2010) 5057–5060

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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f

Combined spectroscopic ellipsometry and attenuated total reflection analyses of Al2O3/HfO2 nanolaminates M. Bonvalot a,⁎, M. Kahn a, C. Vallée a, E. Gourvest a, H. Abed a, C. Jorel a, C. Dubourdieu b a b

Laboratoire des Technologies de la Microélectronique LTM-CNRS, CEA-LETI, 17 avenue des martyrs, 38054 Grenoble Cedex 9, France Laboratoire des Matériaux et du Génie Physique, CNRS, Grenoble INP, 3 parvis L. Néel, BP 257, 38016 Grenoble Cedex 1, France

a r t i c l e

i n f o

Article history: Received 4 May 2009 Received in revised form 1 February 2010 Accepted 12 February 2010 Available online 19 February 2010 Keywords: HfO2 Al2O3 High κ oxide Nanolaminate

a b s t r a c t The microstructure of thin HfO2–Al2O3 nanolaminate high κ dielectric stacks grown by atomic vapor deposition has been studied by attenuated total reflection spectroscopy (ATR) and 8 eV spectroscopic ellipsometry (SE). The presence of Al2O3 below HfO2 prevents the crystallisation of HfO2 if an appropriate thickness is used, which depends on the HfO2 thickness. A thicker Al2O3 is required for thicker HfO2 layers. If crystallisation does occur, we show that the HfO2 signature in both ATR and 8 eV SE spectra allows the detection of monoclinic crystallites embedded in an amorphous phase. © 2010 Elsevier B.V. All rights reserved.

1. Introduction The ongoing aggressive downscaling of microelectronic devices has boosted researches on high κ insulating materials as possible candidates to replace the traditionally used SiO2 or SiOxNy insulators. Among the many binary oxides which have been put on trial for this purpose (Ta2O5, Al2O3, La2O3, HfO2, ZrO2, Y2O3), HfO2 is particularly suitable, due to its high chemical stability with respect to Si, large band gap and high κ value [1]. However, HfO2 shows a strong ability to favour charge defects, which in turn affect the intrinsic properties of metal oxide semiconductor and metal insulator metal (MIM) devices, such as threshold voltage and leakage currents [2]. From the structural point of view, it is generally preferred that the HfO2 material remains amorphous throughout a complete technological process to minimise leakage transport along grain boundaries. The HfO2 structure can be monoclinic, tetragonal, orthorhombic or cubic, each one of them having a specific κ value. The crystallisation temperature of HfO2 thin films has been reported to depend on the growth conditions and layer thickness [3,4]. For instance, it has been reported to be 360 °C in the case of films grown by metal organic chemical vapor deposition [5]. Thus, an important issue is how to control the formation of an amorphous HfO2 phase during deposition and subsequent annealing in order to guarantee reproducible electrical properties of the device formed thereafter. One way consists in alloying HfO2 with a compound that stays amorphous up to much higher temperatures than HfO2, such as SiO2

⁎ Corresponding author. Tel.: + 33 438 78 34 27; fax: +33 438 78 58 92. E-mail address: [email protected] (M. Bonvalot). 0040-6090/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2010.02.034

or Al2O3, as already reported in [6,7]. Another route consists in stacking these compounds in the form of nanolaminates. In this study, Al2O3 has been combined with HfO2. Indeed, Al2O3 has a high-energy band gap, which should limit leakage currents. Moreover, it remains amorphous in thin layers within the low thermal budget required for MIM device fabrication (T ≤ 450 °C) [8]. Thus, HfO2–Al2O3 nanolaminates are expected to remain amorphous up to higher annealing temperatures (as compared to HfO2), which has been shown to be of great interest for MIM capacitors with improved electrical properties [8,9]. We present an investigation on the amorphous phase stabilization in HfO2–Al2O3 nanolaminates intended for MIM applications. Since conventional X-ray diffraction is not appropriate for the determination of the crystalline structure of films with thickness less than 10 nm, we have used spectroscopic techniques such as attenuated total reflection infrared spectroscopy (ATR) and spectroscopic ellipsometry (SE). We discuss the stabilization of the HfO2 amorphous phase with respect to the relative thickness of the films in the HfO2–Al2O3 stacks. The impact of the amorphous phase stabilization on the electrical properties of MIM capacitors will be published in a subsequent paper. 2. Experimental details Several HfO2–Al2O3 nanolaminate structures (Fig. 1) have been prepared by atomic layer deposition (ALD) in a Pulsar 2000 ASM equipment, at 350 °C and with HfCl4, Al(CH3)3 (TMA) and H2O precursors. ALD offers the advantage of very precise thickness control during deposition. The overall nanolaminate thickness has been set to a fixed value of 13 nm and the intermediate thickness range has been evaluated with a 10% precision, the minimum Al2O3 layer thickness

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Fig. 1. Schematic representation of HfO2- and Al2O3-based thin films under study.

being 1 nm. WSi2.3 on Si(100) substrate has been chosen as bottom electrode, because we have previously shown that it improves the C (V) linearity behaviour of MIM capacitors, thanks to its high work function (ϕ = 4.5 eV) [10]. All these samples have been characterised by X-ray photoelectron spectroscopy (XPS) to check the Hf/Al relative stoichiometry in the stacks. Their microstructure has been investigated by spectroscopic ellipsometry (SE) up to 8 eV in a vacuum UV phase modulated spectroscopic ellipsometer from Horiba Jobin Yvon. Attenuated total reflection (ATR) spectroscopic measurements have been carried out in a homemade set-up consisting of a 65° germanium prism, a ppolarized infrared beam (from a Bruker IFS 55 Fourier transform IR spectrometer) operating at 71° incidence to insure total reflection, and a nitrogen-cooled HgCdTe detector [11]. Special attention has been paid on the crystallinity, microstructure and morphology of the stacking as a function of thickness and stacking sequence. 3. Results and discussion Fig. 2 shows the 8 eV SE spectra of pure aluminium and hafnium oxide layers, which will serve as references for discussion of HfO2–Al2O3 nanolaminate behaviour. A modified Tauc–Lorentz dispersion formalism with two oscillators has been applied to fit the optical dielectric function εR and εI (real and imaginary part, respectively), as developed in [12,13]. This leads to an optical index value of 1.9 for HfO2 and 1.7 for Al2O3, with an optical energy band gap of 5.5 eV for HfO2 and 6.5 eV for Al2O3. All these values are consistent with literature data [8,14–16]. Besides providing direct estimates of optical index and real energy band gap, the imaginary part of the dielectric function obtained from this measurement also allows structural change observations. In the case of monoclinic HfO2 films, we observe a first maximum in the imaginary part εI(ω) located at 7.2 eV. SE measurements up to 12 eV for such films indicate the existence of a second maximum slightly above 8 eV [14]. The ATR spectra of these two samples are shown in Fig. 3. HfO2 exhibits a large absorption band between 600 and 800 cm−1, which is

Fig. 2. 8 eV SE measurements of 13 nm-thick HfO2 and Al2O3 layers on WSi2.3 substrate.

attributed to Hf–O stretching vibrations. The presence of two maxima, centred around 690 cm−1 and 770 cm−1, corresponds to the monoclinic phase of HfO2 as proposed by Mc Devitt, back in 1964 [17] and reported for HfO2 thin films [5,11]. ATR spectra analyses of HfO2 are consistent with 8 eV SE results. Concerning Al2O3, a broad absorption peak centred at 900 cm−1 can be observed with a shoulder around 700 cm−1. This is attributed to stretching longitudinal and transverse modes in amorphous Al2O3, the longitudinal vibration mode being strongly enhanced by the P polarization of our ATR configuration. The portion located between 1000 and 1300 cm−1 on the ATR spectra originates from Si–O vibrations. The WSi2.3 bare substrate shows a large absorption with a relatively pronounced peak between 1200 and 1300 cm−1, corresponding to Si–O–Si longitudinal stretching vibrations and a bump at 1100 cm−1 for the transverse mode. The observation of these peaks is due to the presence of a native SiOx

Fig. 3. ATR spectra of 13 nm-thick HfO2 and Al2O3 and of a bare WSi2.3 substrate.

M. Bonvalot et al. / Thin Solid Films 518 (2010) 5057–5060

Fig. 4. ATR spectra of Al2O3 (3 nm)/HfO2 (10 nm) and HfO2 (3 nm)/Al2O3 (10 nm) bilayers deposited on WSi2.3 substrate.

(x ≤ 2) layer on top of the substrate originating from the oxidation of Si atoms at the surface, which are not fully coordinated in WSi2.3. Si-2p XPS spectra of bare WSi2.3 substrate corroborate the existence of this SiOx surface layer with an estimated thickness of 0.8 nm. Similar peaks in the 1100–1300 region can also be observed after HfO2 and Al2O3 deposition, with a slight shift of the longitudinal contribution attributed to the thickness increase of the SiOx interface during film deposition [18]. It should be noted here that the presence of a SiO2 interface has been proven to be beneficial to the linearization of the C(V) electrical characteristic of MIM capacitors, thanks to the reversed C(V) behaviour of SiO2 compared to most high κ materials [19,20]. ATR spectra of HfO2 (10 nm)/Al2O3 (3 nm) and Al2O3 (3 nm)/HfO2 (10 nm) bilayers are shown in Fig. 4. The absorption maximum in the 900 cm−1 region is due to the fully amorphous 3 nm-thick Al2O3 layer. The strong broad peak around 700 cm−1 indicates that HfO2 in both bilayers is amorphous. However, the peak located at 780 cm−1 for HfO2 (10 nm)/Al2O3 (3 nm) also indicates that the monoclinic HfO2 regions are present, whereas this peak does not appear for Al2O3 (3 nm)/HfO2 (10 nm). This clearly shows that the presence of a 3 nm-thick Al2O3 film in contact with the substrate prevents the crystallisation of HfO2 during subsequent deposition. Fig. 5 shows that the Al2O3 thickness required to prevent HfO2 crystallisation depends on HfO2 thickness: In the Al2O3 (1 nm)/HfO2 (11 nm)/Al2O3 (1 nm) trilayer indeed, the monoclinic structure of HfO2 is formed, but for a reduced thickness of 5 nm, HfO2 is amorphous in the Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm) nanolaminate. Thus, the crystallisation of HfO2 depends both on its thickness and Al2O3 thickness. In Fig. 6, we compare the 8 eV SE spectra of a 13 nm-thick monoclinic HfO2 film and of the Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm)/HfO2

Fig. 5. ATR spectra of Al2O3 (1 nm)/HfO2 (11 nm)/Al2O3 (1 nm) trilayer and of Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm) nanolaminate on WSi2.3 substrate.

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Fig. 6. 8 eV SE spectra (imaginary part) of a 13 nm-thick HfO2 film and of Al2O3 (1 nm)/ HfO2 (5 nm)/Al2O3 (1 nm)/HfO2 (5 nm)/Al2O3 (1 nm) nanolaminate deposited on WSi2.3 substrate.

(5 nm)/Al2O3 (1 nm) nanolaminate. One can note the disappearance of the maximum at 7.2 eV. In agreement with previous studies on amorphous HfO2 films [14,15], this confirms the overall amorphous structure of the nanolaminate. One point to be mentioned is that the real optical band gap of HfO2 is not altered by this change in the microstructure of the compound, as deduced from the imaginary part of the dielectric function. This indicates that 1 nm Al2O3 films do not alter the band structure of HfO2–Al2O3 nanolaminates. 4. Conclusion Deposition of a thin layer of Al2O3 prior to HfO2 inhibits the crystallisation of thin HfO2 films. The effect of the relative thicknesses of both oxides on HfO2 crystallisation has been investigated. When crystallisation is not fully inhibited, we have shown that a careful analysis of the HfO2 signature in ATR and 8 eV SE spectra allows the detection of monoclinic crystallites embedded in an amorphous phase. HfO2/Al2O3 nanolaminates are of particular interest for MIM applications. Al2O3 can help minimise leakage currents flowing through grain boundaries, and thus significantly improve electrical properties of devices. Acknowledgments This work has been partially supported by a BDI fellowship between STMicroelectronics and CNRS during the PhD work of M. Kahn. References [1] E.P. Gusev, C. Cabral Jr., M. Copel, C. D'Emic, M. Gribelyuk, Microelectron. Eng. 69 (2003) 145. [2] X. Yu, C. Zhu, H. Hu, A. Chin, M.F. Li, B.J. Cho, IEEE Electron. Dev. Lett. 24 (2003) 64. [3] B.H. Lee, L. Kang, W.J. Qi, R. Nieh, Y. Jeon, K. Onishi, J.C. Lee, Tech. Dig.-Int. Electron. Devices Meet. (1999) 133. [4] E.E. Hoppe, R.C. Aita, M. Gajdardziska-Josifovska, Appl. Phys. Lett. 91 (2007) 203105. [5] C. Dubourdieu, E. Rauwel, C. Millon, P. Chaudouët, F. Ducroquet, N. Rochat, S. Rushworth, V. Cosnier, Chem. Vap. Deposition 12 (2006) 187. [6] G. Molas, M. Bocquet, J. Buckley, H. Grampeix, M. Gély, J.P. Colonna, C. Licitra, N. Rochat, T. Veyront, X. Garros, F. Martin, P. Brianceau, V. Vidal, C. Bongiorno, S. Lombardo, B. De Salvo, S. Deleonibus, Solid State Electron. 51 (2007) 1540. [7] B. Govoreanu, R. Degraeve, M.B. Zahid, L. Nyns, M. Cho, B. Kaczer, M. Jurczak, J.A. Kittl, J. Van Houdt, Microelectron. Eng. 86 (2009) 1807. [8] V.V. Afanas'ev, A. Stesmans, J. Appl. Phys. 102 (2007) 081301. [9] S.J. Ding, D.W. Zhang, L.K. Wang, J. Phys. D: Appl. Phys. 40 (2007) 1072. [10] M. Kahn, C. Vallée, E. Defay, C. Dubourdieu, M. Bonvalot, S. Blonkowski, J.R. Plaussu, P. Garrec, T. Baron, Microelectron. Reliab. 47 (2007) 773. [11] N. Rochat, K. Dabertrand, V. Cosnier, S. Zoll, P. Besson, U. Weber, Phys. Status Solidi 8 (2003) 2965. [12] G.E. Jellison Jr, F.A. Modine, Appl. Phys. Lett. 69 (1996) 371. [13] G.E. Jellison Jr, F.A. Modine, Appl. Phys. Lett. 69 (1996) 2137.

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