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Characterization by ion beams of surfaces and interfaces of alternative materials for future microelectronic devices
Applied Surface Science 212–213 (2003) 556–562
Characterization by ion beams of surfaces and interfaces of alternative materials for future microelectronic devices C. Kruga, F.C. Stedileb,*, C. Radtkea, E.B.O. da Rosaa, J. Moraisa, F.L. Freire Jr.c, I.J.R. Baumvola a Instituto de Fı´sica, Universidade Federal do Rio Grande do Sul, Av. Bento Gonc¸alves, 9500, Porto Alegre, RS 91509-900, Brazil Instituto de Quı´mica, Universidade Federal do Rio Grande do Sul, Av. Bento Gonc¸alves, 9500, Porto Alegre, RS 91509-900, Brazil c Departamento de Fı´sica, Pontifı´cia Universidade Cato´lica do Rio de Janeiro Rua Marqueˆs de Sa˜o Vicente, 225, Rio de Janeiro, RJ 22452-970, Brazil
b
Abstract We present the potential use of ion beam techniques such as nuclear reactions, channelling Rutherford backscattering spectrometry, and low energy ion scattering in the characterization of the surface and interface of materials thought to be possible substitutes to Si (like SiC, for example) and to SiO2 films (like Al2 O3 films, for example) in microelectronic devices. With narrow nuclear reaction resonance profiling the depth distribution of light elements such as Al and O in the films can be obtained non-destructively and with subnanometric depth resolution, allowing one to follow the mobility of each species under thermal treatments, for instance. Thinning of an amorphous layer at the surface of single-crystalline samples can be determined using channelling of Heþ ions and detection of the scattered light particles. Finally, the use of Heþ ions in the 1 keV range allows elemental analysis of the first monolayer at the sample surface. # 2003 Elsevier Science B.V. All rights reserved. PACS: 77.55.þf; 81.65.Mq; 68.49.h; 24.30.v Keywords: Narrow nuclear reaction resonance profiling; Channelling Rutherford backscattering spectrometry; Low energy ion scattering; Silicon carbide; Aluminum oxide
1. Introduction It is well recognized that the duo Si–SiO2 (meaning the physical properties of the materials and of the interface between them) enabled the microelectronics revolution. Moreover, Si and SiO2 allowed four decades of evolution that followed their original application. It is now realized that the evolutionary approach *
Corresponding author. Tel.: þ55-51-3316-7220; fax: þ55-51-3316-7304. E-mail address: [email protected] (F.C. Stedile).
that has been so successful up to now cannot be used for much longer. Materials properties are limiting further developments, and a new revolution is set to take place. On one side, for high performance digital integrated circuits, there is an ongoing search [1,2] for a substitute for silicon oxide, which becomes too leaky at the extremely reduced thicknesses (less than 2 nm) necessary to achieve the high gate dielectric capacitance demanded in submicron metal-oxide-semiconductor devices [3]. Having in sight direct tunnelling leakage current and gate dielectric capacitance, the search was started among materials with increased
0169-4332/03/$ – see front matter # 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0169-4332(03)00405-7
C. Krug et al. / Applied Surface Science 212–213 (2003) 556–562
dielectric constant as compared to SiO2 – so-called ‘‘high-k materials’’ – but other stringent requirements must be fulfilled [4]. One such high-k candidate is aluminum oxide (Al2 O3 ), whose dielectric constant is twice that of SiO2 . On the other side, it is silicon as semiconducting material that is to be substituted aiming at enhanced device performance and reliability for operation at high power and high temperature. The semiconductor of choice is silicon carbide (SiC), for a number of reasons [5] including wide bandgap and having SiO2 as the product of thermal oxidation. Much work has been done to reduce the density of electronic states at SiC–SiO2 interfaces, as that is one of the main issues preventing widespread application of SiC. It is thus clear that the tools of surface science – as ion beam analysis – have much to contribute to both highk and SiC issues. In essence, ion beam analysis (IBA) [6] yields nearsurface, depth-resolved elemental characterization. Common to the different techniques is the incidence of a nearly monoenergetic ion beam (between 1 keV and several megaelectronvolts) from a particle accelerator onto the sample to be studied. In narrow nuclear reaction resonance profiling (NRP), the energy of ion incidence is scanned step by step around the energy corresponding to a narrow, strong, and isolated resonance in the cross-section curve of a nuclear reaction induced by the particles in the beam on a nuclide of interest in the sample. For each beam energy, the number of nuclear reaction products after a fixed ion fluence is recorded with a suitable detector. A plot of the nuclear reaction product yield versus ion beam energy (called excitation curve) is the primary experimental result. We have conveniently used the SPACES [7] code to simulate excitation curves and so obtain nuclide profiles from the experimental data. In addition to NRP, a single ion energy corresponding to a plateau region in a nuclear reaction cross-section curve can be used to measure the total amount of a nuclide in a thin film; this corresponds to nuclear reaction analysis (NRA). In Rutherford backscattering spectrometry (RBS), a particle detector is used to count and register the energy of ions from the original beam scattered by atoms in the sample. Sensitivity to elements in an amorphous film on a crystalline substrate can be enhanced by aligning a major crystallographic direction to the ion beam, i.e. by using the channelling technique. In low energy ion
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scattering (LEIS), lower ion energies are used, and the particle detector is replaced by a hemispherical energy analyzer. Under experimental conditions given below, only the topmost atomic layer of the sample is probed. This paper presents the use of NRP, channelling RBS, and LEIS in the study of Al2 O3 films on Si and of the interface between thermally grown SiO2 and SiC, with attention given to experimental details. The application of NRP to reveal the depth distribution of Al and O atoms in thin Al2 O3 films deposited on Si is reported in Section 2. The technique was specially useful to study the diffusion of oxygen in Al2 O3 during thermal annealing in O2 . Section 3 shows the use of RBS to identify thinning of an amorphous layer at the surface of irradiated SiC, plus LEIS data evidencing C and N migration upon thermal oxidation of SiC. Section 4 summarizes and concludes the paper.
2. Nuclear reaction analysis of Al2 O3 films on Si Among the many possible high-k substitutes for SiO2 [2], we focus on the example of Al2 O3 , which has been considered a suitable candidate [8,9] but has also demonstrated weaknesses [10,11]. Here, the depth distribution of Al in 4 nm-thick Al2 O3 films deposited by remote plasma enhanced chemical vapor deposition on a 0.5 nm-thick SiO2 buffer layer grown in an oxygen-containing plasma on Si(0 0 1) substrates is shown [12]. In order to evaluate the stability of the thin films and of the interface with Si, the samples were characterized as-prepared and after thermal annealing mimicking device processing conditions. Annealing was performed in Ar using a custom-built rapid thermal processing module, for different times and at different temperatures. Additionally, 35 nm-thick Al2 O3 films deposited on HFlast Si(0 0 1) by atomic layer deposition were submitted to thermal annealing in O2 enriched to 98.5% in the 18O isotope (whose natural abundance is only 0.2%). As nuclear reactions are nuclide-specific, that enabled a distinction between oxygen incorporated upon annealing and oxygen originally in the samples. Such relatively thick films (as compared to those of interest for the microelectronics industry) were used to
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study the diffusion of oxygen in the high-k material apart from the Al2 O3 –Si interface. It was our aim to investigate how oxygen diffuses in Al2 O3 and determine if it is incorporated or not. Our ultimate goal is to find annealing conditions that result in optimum dielectric properties. The strong, narrow, and isolated resonances at the proton beam energies 404.9 and 151 keV for the nuclear reactions 27Al(p,g)28Si and 18O(p,a)15N, respectively, were used. Gamma rays of 7.3 MeV produced in the first are detected with a 3 in: 3 in: BGO crystal collinear with the proton beam, and alpha particles produced in the latter are recorded with a standard particle detector. A 6 mm-thick Mylar foil covers the particle detector to prevent scattered protons from reaching it without stopping the 3 MeV alpha particles produced in the nuclear reaction. Significant improvement in near surface depth resolution is gained through depth magnification associated to grazing ion incidence. The samples are typically tilted by 65 for the NRP measurements. This strategy has ultimate limits dictated by angular multiple scattering processes [13]. Fig. 1 shows 27Al profiling data for an ultrathin Al2 O3 film before and after rapid thermal annealing in Ar at 800 C for 30 s. Important features for technological applications of these films are constant
aluminum concentrations, corresponding to the stoichiometric oxide, and sharp interfaces with the SiO2 buffer layer (within depth resolution). The clear difference in the width of the excitation curves for films with thicknesses differing by 1 nm evidences the very high depth resolution of this technique. Fig. 1 also indicates a significant thinning of the dielectric film upon thermal annealing in an inert gas atmosphere. Instabilities such as the creation of voids and thickness inhomogeneities in thin Al2 O3 films have been related to the presence of hydrogen (in the form of hydroxyl groups) residual from film deposition [14,15]. This points to the necessity of extremely low system base pressures during Al2 O3 deposition if good thermal stability is to be achieved. Fig. 2 shows the evolution in time of 18O profiles in an 35 nm-thick Al2 O3 film upon thermal annealing in 18O2 . As stated above, these profiles correspond to oxygen incorporated during the annealing procedure. It can be seen that oxygen is progressively incorporated into the film as the annealing time evolves. In the sample submitted to the longest annealing, oxygen from the gas phase probably reaches the Al2 O3 –Si interface. As a rule, oxygen from the gas phase diffuses in the film strongly interacting with the Al2 O3 bulk’’ regions of the film and absence from the Al2 O3 – Si interface – at least before it is reached by the
Fig. 1. Experimental (symbols) and simulated (lines) excitation curves for the 27Al(p,g)28Si nuclear reaction around the resonance at Ep ¼ 404:9 keV for an Al2 O3 film on SiO2 on Si as-deposited (squares, dashed line) and after annealing (circles, solid line). The 27Al profiles used in the simulations are given in the inset, with Al concentration normalized to stoichiometric Al2 O3 .
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Fig. 2. Experimental (symbols) and simulated (lines) excitation curves for the 18O(p,a)15N nuclear reaction around the resonance at Ep ¼ 151 keV for an 35 nm-thick Al2 O3 film deposited on Si(0 0 1) after annealing at 750 C under 6 mbar of 18O2 for 60, 180, 300, and 720 s, respectively, from left to right. The 18O profiles used in the simulations are given in the inset, with O concentration normalized to stoichiometric Al2 O3 .
propagating front of 18O. Such behavior contrasts with that observed upon annealing of SiO2 films on Si, in which incoming oxygen is mainly incorporated at the SiO2 –Si interface, with minor amounts at the film surface and no incorporation (above natural isotopic abundance) to the bulk oxide [16]. We have also determined the total amounts of 18O in the annealed Al2 O3 films by NRA, using the 18O(p,a)15N nuclear reaction at the plateau region of the cross-section curve at the proton energy around 730 keV [16]. Besides a monotonic increase in the amount of incorporated oxygen with annealing time, the data indicates a non-linear relationship with annealing temperature between 650 and 800 C and a practically linear variation with 18O2 pressure (pO2 ) in the range 0.06–54 mbar [17]. Once again, the result differs from the observations regarding SiO2 on Si, for which the amount of oxygen incorporated to the near-surface 1=4 region follows a pO2 law [18]. Using NRP we have thus shown instabilities in thin Al2 O3 films on Si upon thermal annealing in inert gas and investigated the diffusion of oxygen. Evidences are that the instabilities occur due to the presence of hydrogen residual from the deposition process. Diffusing oxygen may also play a role [10]; this point continues under investigation.
3. Ion scattering characterization of the SiO2 –SiC interface While the physical and chemical characteristics of SiC make it a perfectly suitable semiconductor for use in microelectronic devices, many issues remain to be solved if such potential is to be effectively explored. Among the current difficulties are a poor understanding of the SiC–SiO2 interface and the slow oxidation kinetics of SiC. Both issues can be addressed by IBA, as shown below. The very first stages of the thermal oxidation of SiC in O2 were studied in situ using LEIS [19,20]. We use a Heþ ion beam at the primary energy of 1 keV in an ultra-high vacuum (UHV) system equipped with an Omicron EA 125 hemispherical energy analyzer. A 10 mm 5 mm piece of SiC was cut from a nitrogendoped 6H-SiC(0 0 0 1) wafer supplied by Cree, Inc. The sample was etched by dipping for 30 s in a 5% solution of HF in ethanol, then rinsed for additional 30 s in pure ethanol, and dried in Ar flow. It was then introduced in the UHV chamber and its Si-terminated surface was characterized by LEIS. The sample was then heated to 850 C in UHV for 10 min, and a new LEIS spectrum was acquired. Finally, the sample was submitted to in situ oxidation for 5.0 plus 3.5 plus
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Fig. 3. LEIS spectra obtained with Heþ ions at 1 keV of the SiC sample (a) as-introduced in the UHV chamber, (b) after heating in UHV, and (c) after in situ oxidation for a total time of 8.5 min in 105 mbar of O2 at 850 C. The dotted lines show the ion energy corresponding to each of the elements indicated at the top.
15 min in a dry O2 pressure of 105 mbar at 850 C. The LEIS characterization was repeated after each oxidation step. Fig. 3 shows LEIS spectra of the sample in different processing stages. Besides de Si signal expected for the Si-terminated SiC substrate, spectrum (a) indicates the presence of F and O at the sample surface (we recall that the LEIS analysis as performed here, i.e. with low energy noble gas ions, is practically restricted to the topmost atomic layer of the sample [21]). Spectrum (b) shows that after heating in UHV only Si and N remain, the strong N signal being probably a result of dopant segregation during the annealing step. Spectrum (c), taken after oxidation, indicates the presence of Si and O and, in minor amounts, N and C. While Si and O were expected at this point and N can be beneficial to the dielectric [16], contamination of the SiC–SiO2 interface with C could be one of the reasons for its relatively poor electrical properties.
The second issue related to SiC that was investigated using IBA (in this case, RBS) is acceleration of the oxidation kinetics through amorphization of the surface of the substrate [22–24]. Indeed, high temperatures and/or long oxidation times have to be used with standard single crystalline substrates to achieve the currently desired oxide thicknesses [25] (30 nm). In this experiment, a SiC sample was irradiated with an energetic ion beam in order to produce the amorphized region. The same SiC substrate described above was used. Immediately after the wet etching step, three samples were introduced in an ion implantation chamber, which was then pumped down to 5 107 mbar. Irradiation was performed with Arþ ions at the energy of 170 keV, to the fluence of 1:1 1015 cm2 . The samples were then separately submitted to oxidation at 1100 C in atmospheric pressure of N2 bubbled through deionized water kept at 95 C, for 1.0, 2.5, and 12.0 h. The samples were then analyzed by RBS with Heþ ions of 2 MeV. Detection of backscattered ions was performed at 170 relative to the direction of ion incidence. The samples were mounted on a goniometer so that the h0 0 0 1i crystallographic axis of the substrate could be aligned to the incident ion beam (channelling RBS technique). The amount of oxide grown on each sample was calculated from the signals corresponding to O in the samples and in a standard SiO2 film. The oxide layers were then etched away in 20% HF aqueous solution and the exposed SiC substrates were again analyzed by RBS in the conditions above. Fig. 4 shows the Si signal in RBS data for virgin, asirradiated, and irradiated, oxidized, and etched SiC samples. For spectra (b)–(d), the high yield plateau next to the high energy signal edge corresponds to a surface layer of amorphized SiC. As expected for complete amorphization, the plateau height is the same as that observed in spectrum (a). Also for spectra (b)–(d), the reduction of backscattered ion yield at lower energies is due to ion channelling in the deeper, crystalline portion of the substrate. This second plateau is still higher than that observed in spectrum (e) due to so-called ‘‘dechannelling’’ effects of the top amorphous layer. The reduction on the width of the plateau at the highest ion yield in spectra (b)–(d) is due to consumption of the top amorphized SiC layer in the oxidation process. Comparing the estimated amount of Si present in the oxide films (calculated
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Fig. 4. Si signal in (a) random RBS spectrum for virgin SiC substrate and channelling RBS spectra for (b) as-irradiated SiC; exposed irradiated substrate originally oxidized for (c) 1.0 h; (d) 2.5 h; and (e) 12 h. The latter reproduces the channelling RBS spectrum obtained for virgin SiC (not shown).
from the O signal in the channelling RBS spectra of the oxidized samples) with the amount of Si in the amorphous layer of the as-irradiated sample (obtained by simulation with the RUMP code [26]), one observes that the oxidation for 1.0, 2.5, and 12.0 h consumed, respectively, 15, 45, and 100% (within an error of 10%) of the amorphized layer; the estimated oxide thicknesses are 48, 160, and 307 nm. For comparison, a non-irradiated SiC substrate was submitted to oxidation for 6.0 h under analogous conditions, yielding a 25 nm-thick oxide overlayer. This result evidences a strong enhancement of the oxidation kinetics of SiC upon ion irradiation (i.e. amorphization). Evidently, for electronic device manufacturing the interface of SiO2 must be with single-crystalline SiC. If ion irradiation of SiC is to be used to enhance its oxidation kinetics, oxidation must proceed at least so as to consume the whole amorphized layer.
4. Summary and conclusion The use of the IBA techniques narrow nuclear reaction resonance profiling (NRP), channelling Rutherford backscattering spectrometry (RBS), and low energy ion scattering (LEIS) were shown in the
study of alternative materials for future microelectronic devices. These are well established techniques, which can be considered complementary as judged by the nature of the information that they provide. Of technological relevance, the paper shows (a) an extremely powerful method to calibrate deposition systems for Al2 O3 films on Si; (b) evidence of a possible segregation of dopant N upon thermal annealing of SiC; and (c) kinetic data on the oxidation of irradiated SiC. We understand that IBA – in close connection to other surface analysis techniques such as XPS, SIMS, FT-IR, STM, AFM, and others – will continue to play a key role in the evaluation of new materials for the electronics industry.
Acknowledgements Supported by the Brazilian agencies CNPq, FAPERGS, PADCT, and CAPES. References [1] A.I. Kingon, J.-P. Maria, S.K. Streiffer, Nature (London) 406 (2000) 1032. [2] G.D. Wilk, R.M. Wallace, J.M. Anthony, J. Appl. Phys. 87 (2000) 484.
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[3] M.L. Green, E.P. Gusev, R. Degraeve, E.L. Garfunkel, J. Appl. Phys. 90 (2001) 2057. [4] L. Manchanda, M.D. Morris, M.L. Green, R.B. van Dover, F. Klemens, T.W. Sorsch, P.J. Silverman, G. Wilk, B. Busch, S. Aravamudhan, Microelectron. Eng. 59 (2001) 351. [5] H. Morkoc¸, S. Strite, G.B. Gao, M.E. Lin, B. Sverdlov, M. Burns, J. Appl. Phys. 76 (1994) 1363. [6] J.R. Bird, J.S. Williams, Ion Beams for Materials Analysis, Academic Press, Sydney, 1989. [7] I. Vickridge, G. Amsel, Nucl. Instr. Meth. B 45 (1990) 6. [8] M. Copel, E. Cartier, E.P. Gusev, S. Guha, N. Bojarczuk, M. Poppeller, Appl. Phys. Lett. 78 (2001) 2670. [9] E.P. Gusev, M. Copel, E. Cartier, I.J.R. Baumvol, C. Krug, M.A. Gribelyuk, Appl. Phys. Lett. 76 (2000) 176. [10] C. Krug, E.B.O. da Rosa, R.M.C. de Almeida, J. Morais, I.J.R. Baumvol, T.D.M. Salgado, F.C. Stedile, Phys. Rev. Lett. 85 (2000) 4120. [11] M. Copel, Phys. Rev. Lett. 86 (2001) 4713. [12] R.S. Johnson, G. Lucovsky, I. Baumvol, J. Vac. Sci. Technol. A 19 (2001) 1353. [13] G. Battistig, G. Amsel, E. d’Artemare, A. L’Hoir, Nucl. Instr. Meth. B 85 (1994) 572. [14] L.G. Gosset, J.-J. Ganem, O. Renault, P. Holliger, J.-F. Damlencourt, G. Rolland, H.J. von Bardeleben, F. Pierre, D. Jalabert, I. Trimaille, et al., in: J. Morais, I.J.R. Baumvol
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(Eds.), Alternatives to SiO2 as Gate Dielectrics for Future SiBased Microelectronics, Materials Research Society, Warrendale, 2001, p. 6. J.M. Schneider, A. Anders, B. Hjo¨ rvarsson, I. Petrov, K. Maca´ k, U. Helmersson, J.-E. Sundgren, Appl. Phys. Lett. 74 (1999) 200. I.J.R. Baumvol, Surf. Sci. Rep. 36 (1999) 1. E.B.O. da Rosa, I.J.R. Baumvol, J. Morais, R.M.C. d Almeida, R.M. Papale´ o, F.C. Stedile, Phys. Rev. B 65 (2002) 121303. I. Trimaille, S. Rigo, Appl. Surf. Sci. 39 (1989) 65. C. Radtke, R.V. Branda˜ o, R.P. Pezzi, J. Morais, I.J.R. Baumvol, F.C. Stedile, Nucl. Instrum. Meth. B 190 (2002) 579. C. Radtke, I.J.R. Baumvol, J. Morais, F.C. Stedile, Appl. Phys. Lett. 78 (2001) 3601. D.G. Armour, Methods of Surface Analysis, Cambridge University Press, Cambridge, 1989, Chapter 8, p. 263. R. Nipoti, M. Madrigali, A. Sambo, Mater. Sci. Eng. B 61–62 (1999) 475. T. Yoneda, T. Nakata, M. Watanabe, M. Kitabatake, Mater. Sci. Eng. B 61–62 (1999) 502. D. Alok, B.J. Baliga, J. Electrochem. Soc. 144 (1997) 1135. C. Raynaud, J. Non-Crystal. Solids 280 (2001) 1. L.R. Doolittle, Nucl. Instrum. Meth. B 9 (1985) 334.
Characterization by ion beams of surfaces and interfaces of alternative materials for future microelectronic devices C. Kruga, F.C. Stedileb,*, C. Radtkea, E.B.O. da Rosaa, J. Moraisa, F.L. Freire Jr.c, I.J.R. Baumvola a Instituto de Fı´sica, Universidade Federal do Rio Grande do Sul, Av. Bento Gonc¸alves, 9500, Porto Alegre, RS 91509-900, Brazil Instituto de Quı´mica, Universidade Federal do Rio Grande do Sul, Av. Bento Gonc¸alves, 9500, Porto Alegre, RS 91509-900, Brazil c Departamento de Fı´sica, Pontifı´cia Universidade Cato´lica do Rio de Janeiro Rua Marqueˆs de Sa˜o Vicente, 225, Rio de Janeiro, RJ 22452-970, Brazil
b
Abstract We present the potential use of ion beam techniques such as nuclear reactions, channelling Rutherford backscattering spectrometry, and low energy ion scattering in the characterization of the surface and interface of materials thought to be possible substitutes to Si (like SiC, for example) and to SiO2 films (like Al2 O3 films, for example) in microelectronic devices. With narrow nuclear reaction resonance profiling the depth distribution of light elements such as Al and O in the films can be obtained non-destructively and with subnanometric depth resolution, allowing one to follow the mobility of each species under thermal treatments, for instance. Thinning of an amorphous layer at the surface of single-crystalline samples can be determined using channelling of Heþ ions and detection of the scattered light particles. Finally, the use of Heþ ions in the 1 keV range allows elemental analysis of the first monolayer at the sample surface. # 2003 Elsevier Science B.V. All rights reserved. PACS: 77.55.þf; 81.65.Mq; 68.49.h; 24.30.v Keywords: Narrow nuclear reaction resonance profiling; Channelling Rutherford backscattering spectrometry; Low energy ion scattering; Silicon carbide; Aluminum oxide
1. Introduction It is well recognized that the duo Si–SiO2 (meaning the physical properties of the materials and of the interface between them) enabled the microelectronics revolution. Moreover, Si and SiO2 allowed four decades of evolution that followed their original application. It is now realized that the evolutionary approach *
Corresponding author. Tel.: þ55-51-3316-7220; fax: þ55-51-3316-7304. E-mail address: [email protected] (F.C. Stedile).
that has been so successful up to now cannot be used for much longer. Materials properties are limiting further developments, and a new revolution is set to take place. On one side, for high performance digital integrated circuits, there is an ongoing search [1,2] for a substitute for silicon oxide, which becomes too leaky at the extremely reduced thicknesses (less than 2 nm) necessary to achieve the high gate dielectric capacitance demanded in submicron metal-oxide-semiconductor devices [3]. Having in sight direct tunnelling leakage current and gate dielectric capacitance, the search was started among materials with increased
0169-4332/03/$ – see front matter # 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0169-4332(03)00405-7
C. Krug et al. / Applied Surface Science 212–213 (2003) 556–562
dielectric constant as compared to SiO2 – so-called ‘‘high-k materials’’ – but other stringent requirements must be fulfilled [4]. One such high-k candidate is aluminum oxide (Al2 O3 ), whose dielectric constant is twice that of SiO2 . On the other side, it is silicon as semiconducting material that is to be substituted aiming at enhanced device performance and reliability for operation at high power and high temperature. The semiconductor of choice is silicon carbide (SiC), for a number of reasons [5] including wide bandgap and having SiO2 as the product of thermal oxidation. Much work has been done to reduce the density of electronic states at SiC–SiO2 interfaces, as that is one of the main issues preventing widespread application of SiC. It is thus clear that the tools of surface science – as ion beam analysis – have much to contribute to both highk and SiC issues. In essence, ion beam analysis (IBA) [6] yields nearsurface, depth-resolved elemental characterization. Common to the different techniques is the incidence of a nearly monoenergetic ion beam (between 1 keV and several megaelectronvolts) from a particle accelerator onto the sample to be studied. In narrow nuclear reaction resonance profiling (NRP), the energy of ion incidence is scanned step by step around the energy corresponding to a narrow, strong, and isolated resonance in the cross-section curve of a nuclear reaction induced by the particles in the beam on a nuclide of interest in the sample. For each beam energy, the number of nuclear reaction products after a fixed ion fluence is recorded with a suitable detector. A plot of the nuclear reaction product yield versus ion beam energy (called excitation curve) is the primary experimental result. We have conveniently used the SPACES [7] code to simulate excitation curves and so obtain nuclide profiles from the experimental data. In addition to NRP, a single ion energy corresponding to a plateau region in a nuclear reaction cross-section curve can be used to measure the total amount of a nuclide in a thin film; this corresponds to nuclear reaction analysis (NRA). In Rutherford backscattering spectrometry (RBS), a particle detector is used to count and register the energy of ions from the original beam scattered by atoms in the sample. Sensitivity to elements in an amorphous film on a crystalline substrate can be enhanced by aligning a major crystallographic direction to the ion beam, i.e. by using the channelling technique. In low energy ion
557
scattering (LEIS), lower ion energies are used, and the particle detector is replaced by a hemispherical energy analyzer. Under experimental conditions given below, only the topmost atomic layer of the sample is probed. This paper presents the use of NRP, channelling RBS, and LEIS in the study of Al2 O3 films on Si and of the interface between thermally grown SiO2 and SiC, with attention given to experimental details. The application of NRP to reveal the depth distribution of Al and O atoms in thin Al2 O3 films deposited on Si is reported in Section 2. The technique was specially useful to study the diffusion of oxygen in Al2 O3 during thermal annealing in O2 . Section 3 shows the use of RBS to identify thinning of an amorphous layer at the surface of irradiated SiC, plus LEIS data evidencing C and N migration upon thermal oxidation of SiC. Section 4 summarizes and concludes the paper.
2. Nuclear reaction analysis of Al2 O3 films on Si Among the many possible high-k substitutes for SiO2 [2], we focus on the example of Al2 O3 , which has been considered a suitable candidate [8,9] but has also demonstrated weaknesses [10,11]. Here, the depth distribution of Al in 4 nm-thick Al2 O3 films deposited by remote plasma enhanced chemical vapor deposition on a 0.5 nm-thick SiO2 buffer layer grown in an oxygen-containing plasma on Si(0 0 1) substrates is shown [12]. In order to evaluate the stability of the thin films and of the interface with Si, the samples were characterized as-prepared and after thermal annealing mimicking device processing conditions. Annealing was performed in Ar using a custom-built rapid thermal processing module, for different times and at different temperatures. Additionally, 35 nm-thick Al2 O3 films deposited on HFlast Si(0 0 1) by atomic layer deposition were submitted to thermal annealing in O2 enriched to 98.5% in the 18O isotope (whose natural abundance is only 0.2%). As nuclear reactions are nuclide-specific, that enabled a distinction between oxygen incorporated upon annealing and oxygen originally in the samples. Such relatively thick films (as compared to those of interest for the microelectronics industry) were used to
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study the diffusion of oxygen in the high-k material apart from the Al2 O3 –Si interface. It was our aim to investigate how oxygen diffuses in Al2 O3 and determine if it is incorporated or not. Our ultimate goal is to find annealing conditions that result in optimum dielectric properties. The strong, narrow, and isolated resonances at the proton beam energies 404.9 and 151 keV for the nuclear reactions 27Al(p,g)28Si and 18O(p,a)15N, respectively, were used. Gamma rays of 7.3 MeV produced in the first are detected with a 3 in: 3 in: BGO crystal collinear with the proton beam, and alpha particles produced in the latter are recorded with a standard particle detector. A 6 mm-thick Mylar foil covers the particle detector to prevent scattered protons from reaching it without stopping the 3 MeV alpha particles produced in the nuclear reaction. Significant improvement in near surface depth resolution is gained through depth magnification associated to grazing ion incidence. The samples are typically tilted by 65 for the NRP measurements. This strategy has ultimate limits dictated by angular multiple scattering processes [13]. Fig. 1 shows 27Al profiling data for an ultrathin Al2 O3 film before and after rapid thermal annealing in Ar at 800 C for 30 s. Important features for technological applications of these films are constant
aluminum concentrations, corresponding to the stoichiometric oxide, and sharp interfaces with the SiO2 buffer layer (within depth resolution). The clear difference in the width of the excitation curves for films with thicknesses differing by 1 nm evidences the very high depth resolution of this technique. Fig. 1 also indicates a significant thinning of the dielectric film upon thermal annealing in an inert gas atmosphere. Instabilities such as the creation of voids and thickness inhomogeneities in thin Al2 O3 films have been related to the presence of hydrogen (in the form of hydroxyl groups) residual from film deposition [14,15]. This points to the necessity of extremely low system base pressures during Al2 O3 deposition if good thermal stability is to be achieved. Fig. 2 shows the evolution in time of 18O profiles in an 35 nm-thick Al2 O3 film upon thermal annealing in 18O2 . As stated above, these profiles correspond to oxygen incorporated during the annealing procedure. It can be seen that oxygen is progressively incorporated into the film as the annealing time evolves. In the sample submitted to the longest annealing, oxygen from the gas phase probably reaches the Al2 O3 –Si interface. As a rule, oxygen from the gas phase diffuses in the film strongly interacting with the Al2 O3 bulk’’ regions of the film and absence from the Al2 O3 – Si interface – at least before it is reached by the
Fig. 1. Experimental (symbols) and simulated (lines) excitation curves for the 27Al(p,g)28Si nuclear reaction around the resonance at Ep ¼ 404:9 keV for an Al2 O3 film on SiO2 on Si as-deposited (squares, dashed line) and after annealing (circles, solid line). The 27Al profiles used in the simulations are given in the inset, with Al concentration normalized to stoichiometric Al2 O3 .
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Fig. 2. Experimental (symbols) and simulated (lines) excitation curves for the 18O(p,a)15N nuclear reaction around the resonance at Ep ¼ 151 keV for an 35 nm-thick Al2 O3 film deposited on Si(0 0 1) after annealing at 750 C under 6 mbar of 18O2 for 60, 180, 300, and 720 s, respectively, from left to right. The 18O profiles used in the simulations are given in the inset, with O concentration normalized to stoichiometric Al2 O3 .
propagating front of 18O. Such behavior contrasts with that observed upon annealing of SiO2 films on Si, in which incoming oxygen is mainly incorporated at the SiO2 –Si interface, with minor amounts at the film surface and no incorporation (above natural isotopic abundance) to the bulk oxide [16]. We have also determined the total amounts of 18O in the annealed Al2 O3 films by NRA, using the 18O(p,a)15N nuclear reaction at the plateau region of the cross-section curve at the proton energy around 730 keV [16]. Besides a monotonic increase in the amount of incorporated oxygen with annealing time, the data indicates a non-linear relationship with annealing temperature between 650 and 800 C and a practically linear variation with 18O2 pressure (pO2 ) in the range 0.06–54 mbar [17]. Once again, the result differs from the observations regarding SiO2 on Si, for which the amount of oxygen incorporated to the near-surface 1=4 region follows a pO2 law [18]. Using NRP we have thus shown instabilities in thin Al2 O3 films on Si upon thermal annealing in inert gas and investigated the diffusion of oxygen. Evidences are that the instabilities occur due to the presence of hydrogen residual from the deposition process. Diffusing oxygen may also play a role [10]; this point continues under investigation.
3. Ion scattering characterization of the SiO2 –SiC interface While the physical and chemical characteristics of SiC make it a perfectly suitable semiconductor for use in microelectronic devices, many issues remain to be solved if such potential is to be effectively explored. Among the current difficulties are a poor understanding of the SiC–SiO2 interface and the slow oxidation kinetics of SiC. Both issues can be addressed by IBA, as shown below. The very first stages of the thermal oxidation of SiC in O2 were studied in situ using LEIS [19,20]. We use a Heþ ion beam at the primary energy of 1 keV in an ultra-high vacuum (UHV) system equipped with an Omicron EA 125 hemispherical energy analyzer. A 10 mm 5 mm piece of SiC was cut from a nitrogendoped 6H-SiC(0 0 0 1) wafer supplied by Cree, Inc. The sample was etched by dipping for 30 s in a 5% solution of HF in ethanol, then rinsed for additional 30 s in pure ethanol, and dried in Ar flow. It was then introduced in the UHV chamber and its Si-terminated surface was characterized by LEIS. The sample was then heated to 850 C in UHV for 10 min, and a new LEIS spectrum was acquired. Finally, the sample was submitted to in situ oxidation for 5.0 plus 3.5 plus
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Fig. 3. LEIS spectra obtained with Heþ ions at 1 keV of the SiC sample (a) as-introduced in the UHV chamber, (b) after heating in UHV, and (c) after in situ oxidation for a total time of 8.5 min in 105 mbar of O2 at 850 C. The dotted lines show the ion energy corresponding to each of the elements indicated at the top.
15 min in a dry O2 pressure of 105 mbar at 850 C. The LEIS characterization was repeated after each oxidation step. Fig. 3 shows LEIS spectra of the sample in different processing stages. Besides de Si signal expected for the Si-terminated SiC substrate, spectrum (a) indicates the presence of F and O at the sample surface (we recall that the LEIS analysis as performed here, i.e. with low energy noble gas ions, is practically restricted to the topmost atomic layer of the sample [21]). Spectrum (b) shows that after heating in UHV only Si and N remain, the strong N signal being probably a result of dopant segregation during the annealing step. Spectrum (c), taken after oxidation, indicates the presence of Si and O and, in minor amounts, N and C. While Si and O were expected at this point and N can be beneficial to the dielectric [16], contamination of the SiC–SiO2 interface with C could be one of the reasons for its relatively poor electrical properties.
The second issue related to SiC that was investigated using IBA (in this case, RBS) is acceleration of the oxidation kinetics through amorphization of the surface of the substrate [22–24]. Indeed, high temperatures and/or long oxidation times have to be used with standard single crystalline substrates to achieve the currently desired oxide thicknesses [25] (30 nm). In this experiment, a SiC sample was irradiated with an energetic ion beam in order to produce the amorphized region. The same SiC substrate described above was used. Immediately after the wet etching step, three samples were introduced in an ion implantation chamber, which was then pumped down to 5 107 mbar. Irradiation was performed with Arþ ions at the energy of 170 keV, to the fluence of 1:1 1015 cm2 . The samples were then separately submitted to oxidation at 1100 C in atmospheric pressure of N2 bubbled through deionized water kept at 95 C, for 1.0, 2.5, and 12.0 h. The samples were then analyzed by RBS with Heþ ions of 2 MeV. Detection of backscattered ions was performed at 170 relative to the direction of ion incidence. The samples were mounted on a goniometer so that the h0 0 0 1i crystallographic axis of the substrate could be aligned to the incident ion beam (channelling RBS technique). The amount of oxide grown on each sample was calculated from the signals corresponding to O in the samples and in a standard SiO2 film. The oxide layers were then etched away in 20% HF aqueous solution and the exposed SiC substrates were again analyzed by RBS in the conditions above. Fig. 4 shows the Si signal in RBS data for virgin, asirradiated, and irradiated, oxidized, and etched SiC samples. For spectra (b)–(d), the high yield plateau next to the high energy signal edge corresponds to a surface layer of amorphized SiC. As expected for complete amorphization, the plateau height is the same as that observed in spectrum (a). Also for spectra (b)–(d), the reduction of backscattered ion yield at lower energies is due to ion channelling in the deeper, crystalline portion of the substrate. This second plateau is still higher than that observed in spectrum (e) due to so-called ‘‘dechannelling’’ effects of the top amorphous layer. The reduction on the width of the plateau at the highest ion yield in spectra (b)–(d) is due to consumption of the top amorphized SiC layer in the oxidation process. Comparing the estimated amount of Si present in the oxide films (calculated
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Fig. 4. Si signal in (a) random RBS spectrum for virgin SiC substrate and channelling RBS spectra for (b) as-irradiated SiC; exposed irradiated substrate originally oxidized for (c) 1.0 h; (d) 2.5 h; and (e) 12 h. The latter reproduces the channelling RBS spectrum obtained for virgin SiC (not shown).
from the O signal in the channelling RBS spectra of the oxidized samples) with the amount of Si in the amorphous layer of the as-irradiated sample (obtained by simulation with the RUMP code [26]), one observes that the oxidation for 1.0, 2.5, and 12.0 h consumed, respectively, 15, 45, and 100% (within an error of 10%) of the amorphized layer; the estimated oxide thicknesses are 48, 160, and 307 nm. For comparison, a non-irradiated SiC substrate was submitted to oxidation for 6.0 h under analogous conditions, yielding a 25 nm-thick oxide overlayer. This result evidences a strong enhancement of the oxidation kinetics of SiC upon ion irradiation (i.e. amorphization). Evidently, for electronic device manufacturing the interface of SiO2 must be with single-crystalline SiC. If ion irradiation of SiC is to be used to enhance its oxidation kinetics, oxidation must proceed at least so as to consume the whole amorphized layer.
4. Summary and conclusion The use of the IBA techniques narrow nuclear reaction resonance profiling (NRP), channelling Rutherford backscattering spectrometry (RBS), and low energy ion scattering (LEIS) were shown in the
study of alternative materials for future microelectronic devices. These are well established techniques, which can be considered complementary as judged by the nature of the information that they provide. Of technological relevance, the paper shows (a) an extremely powerful method to calibrate deposition systems for Al2 O3 films on Si; (b) evidence of a possible segregation of dopant N upon thermal annealing of SiC; and (c) kinetic data on the oxidation of irradiated SiC. We understand that IBA – in close connection to other surface analysis techniques such as XPS, SIMS, FT-IR, STM, AFM, and others – will continue to play a key role in the evaluation of new materials for the electronics industry.
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