Metal and composite nanocluster precipitate formation in silicon dioxide implanted with Sb ions

Transcription

1 JOURNAL OF APPLIED PHYSICS VOLUME 92, NUMBER 8 15 OCTOBER 2002 Metal and composite nanocluster precipitate formation in silicon dioxide implanted with Sb ions V. A. Ignatova a) MiTAC, Department of Chemistry, University of Antwerp, Universiteitsplein 1, B-2610 Wilrijk, Belgium O. I. Lebedev EMAT, Department of Physics, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerpen, Belgium U. Wätjen EC-JRC, Institute for Reference Materials and Measurements, Retieseweg, B-2440 Geel, Belgium L. Van Vaeck MiTAC, Department of Chemistry, University of Antwerp, Universiteitsplein 1, B-2610 Wilrijk, Belgium J. Van Landuyt EMAT, Department of Physics, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerpen, Belgium R. Gijbels and F. Adams MiTAC, Department of Chemistry, University of Antwerp, Universiteitsplein 1, B-2610 Wilrijk, Belgium Received 14 March 2002; accepted for publication 23 July 2002 Amorphous thin SiO 2 layers of different thickness on a Si 111 substrate were implanted with 12 kev 121 Sb ions at fluences between cm 2 and cm 2, followed by thermal annealing. Formation of nanocrystal precipitates was established by high-resolution electron microscopy. The size and the distribution of the nanocrystals formed around the initial depth of implantation were studied in the as-implanted samples. The crystalline structure of these nanocrystals was also explored and the presence of antimony oxide Sb 2 O 3 in the form of valentinite was proven. After the annealing step, the implanted material was observed to spread into a wider band and even to split into two discrete bands. The presence of antimony oxide in the intermediate phase as-implanted layer of Sb was confirmed using Fourier transform laser microprobe mass spectrometry. No definite signals of Sb 2 O 3 could be detected in the annealed samples American Institute of Physics. DOI: / I. INTRODUCTION The increasing interest in the formation of metal nanocrystals in barrier layers with low dielectric constants has moved the attention from fundamental studies to establishing the principles of nanotechnology and fabrication of different nanodevices over the past few years. One of the driving forces is their potential use as the necessary nanoscale structure in single-electron devices as elements of future microelectronics circuits, whereas their formation technique can be made compatible with conventional fabrication processes of semiconductor devices. Antimony is one of the typical doping species in conventional large-scale integrated devices. Low dose ion implantation followed by thermal annealing is one of the most promising methods for preparation, since, in principle, it allows control of the size and the depth of the Sb nanocrystals by changing the implantation dose and energy. However, the final position and size distribution are strongly affected by the annealing conditions and the thickness of the dioxide layer and are not fully understood. Although SiO 2 is commonly used in silicon integrated circuits, there is surprisingly little knowledge about the diffusion of dopant atoms in SiO 2. The previous studies 1 5 of Sb implantation in SiO 2 or in a Si substrate show that no appreciable diffusion of ion-implanted Sb can be observed by means of Rutherford backscattering RBS, the implanted Sb is immobile in SiO 2 upon annealing at temperatures at least as high as 1200 C in N 2,O 2,orO 2 /H 2 O ambient. To mobilize the implant, extra oxygen has to be supplied from the annealing ambient. A systematic study of the incorporation, diffusion, and formation of oxides of ion-implanted Sb in SiO 2 at a dose of atoms cm 2 and an energy of 190 kev and of other elements from groups V and III demonstrated the formation of Sb-O bonds due to the ion implantation in the presence of oxygen in the ambient or at high annealing temperatures 1000 C. 1 Without an external supply of oxygen, two thirds of the implanted Sb will become incorporated on O sites, and one third on the Si sites of the SiO 2 network. In other studies, 2 it has been shown by energy-dispersive x-ray spectrometry EDX and wavelength dispersive x-ray analysis that at Sb concentrations at. %, Sb tends to precipitate in nm size spherical inclusions after annealing with crystalline structure. The study of Sb precipitation during ion implantation has been continued by a series of papers of Nakajima et al. 3 5 concerning the formation of Sb nanocrystals in SiO 2 films. The formation of nanocrystals has been confirmed by transmission electron microscopy TEM and EDX analysis, but no antia Author to whom all correspondence should be addressed; electronic mail:velislav@uia.ua.ac.be /2002/92(8)/4336/6/$ American Institute of Physics

2 J. Appl. Phys., Vol. 92, No. 8, 15 October 2002 Ignatova et al mony oxides were reported. Small clusters, however, are expected to be very reactive and susceptible to oxidation due to their large ratio of the number of surface to volume atoms. In recent studies, 6 the concept of a chemical reaction between implant atoms and moisture atoms, diffusing in from the ambient during annealing, was introduced and the formation of GeO x clusters was observed by x-ray photoelectron spectroscopy. 7 But again, no proof of the presence of antimony oxides could be given, 6 and the problem of which physical and chemical interactions of implant atoms, host atoms, and possibly moisture atoms take place remains to be solved. The purpose of the present work is to investigate the distribution of the Sb ions implanted in SiO 2 layers and the possible formation of metal and composite oxide nanocrystals during/after the implantation at different thicknesses of the SiO 2 layer and at different fluences. We studied the size and the redistribution of the nanoclusters formed, after 12 kev Sb ions were implanted at fluences of cm 2 and cm 2 into 20 or 30 nm thick SiO 2. We combine high-resolution electron microscopy HREM for studying the microscopic structure of the formed clusters with Fourier transform laser microprobe mass spectrometry FTLMMS for molecular speciation, which can give direct proof of the presence of Sb oxides in the implanted samples. Since the size of nanocrystals is in the order of nanometers, HREM is the only way to visualize the crystal lattice and to achieve information on the structure and the distribution of the individual clusters. II. MATERIAL Wafers of n-type silicon in orientation 111 and specific resistance of the order of 1 to 10 cm were thermally oxidized to a nominal thickness of the SiO 2 layer of 20 and 30 nm. Subsequently, they were implanted with 12 kev 121 Sb ions mass separated in the implanter, rendering a mean projected range of 11 nm calculated by a dynamic TRIM code. The nominal retained doses were and Sb atoms cm 2. RBS analysis revealed that the effectively retained doses of the as-implanted samples were about 10% to 15% lower than the nominal values. Four chips of each implanted wafer were furnace-annealed at 900 C under dry nitrogen atmosphere for 10 min. III. HIGH-RESOLUTION ELECTRON MICROSCOPY Standard cross-sectional TEM specimens were cut parallel to a 110 plane of the Si substrate and mechanically ground to a thickness of about 20 m, followed by ion-beam milling under grazing incidence with respect to the surface until a small hole is formed. The area around this hole is sufficiently thin to transmit electrons in order to obtain TEM - HREM images. TEM investigations were carried out with a JEOL 4000EX microscope operated at 400 kv. The point resolution of the microscope was of the order of 0.17 nm. Since the size of the nanocrystals is very small in the order of a few nanometers, electron diffraction will not give a relevant structure information of nanocrystals. HREM images on the other hand reveal the local lattice spacing within a much smaller area of one separate nanocrystal. However, a reliable interpretation requires the measurement of a number of lattice spacings. This is facilitated by using the Fourier transformations FT of small local areas of the HREM image of representative nanocrystals, provided that the areas are large enough to produce sharp spots in the FT pattern. Instead of trying to measure distances in blurred lattice plane fringes, interdot spacings in the FT patterns are measured cf. Fig. 6. In the present study, we have used the combination of the HREM and FT techniques in order to determine the structure of nanocrystals. The image processing was performed with NIH Image 1.60 software. FT of lattice plane fringes in the digitized micrographs were calculated resulting in diffractograms. In this respect, the lattice image of the Si substrate and its corresponding FT pattern could be used for internal calibration of FT patterns and HREM images. This allowed us to measure the lattice spacing of the nanocrystals with high accuracy. Based on the geometry and d spacing of FT patterns, the structure and phase of the nanocrystals could be determined. IV. FOURIER-TRANSFORM LASER MICROPROBE MASS SPECTROMETRY Laser microprobe mass spectrometry LMMS employs focused laser-beam irradiation of solids and subsequent mass analysis by Fourier transform mass spectrometry FTMS for high mass resolution and mass accuracy. The technique is suitable for local analysis with spatial resolution in the m range by means of molecular information on both organic and inorganic constituents speciation. A significant advantage is the absence of surface charging of insulating samples. The FTLMMS instrument 8 has been developed from a Spectrospin CMS 47X FTMS Bruker Spectrospin, Billerica, MA. The system uses an Infinity Cell, 9 a 4.7 T magnet and an external ion source. Samples are irradiated under 45 in the reflection mode by a frequency quadrupled Nd:YAG laser Quanta-Ray DCR 2 10, Spectra Physics, Mountain View, CA delivering 20 mj per pulse at 266 nm with a pulse length of 4 5 ns. The beam is focused to a 5 m spot. The sample viewing system comprises a microscope Nikon M Plan, Japan, coupled to a charge-coupled device CCD camera CCD 500, Bischke, Switzerland. V. RESULTS AND DISCUSSION A. Distribution of the implanted material The HREM micrograph in Fig. 1 of the as-implanted sample with a dose of atoms cm 2 of 12 kev Sb shows the implant distribution at about the initial mean implantation depth, which is estimated as 13 nm better seen in Fig. 3. This is somewhat larger than the value of 11 nm for the projected range, predicted by calculations with the dynamic TRIM code. 10,11 Also, the observed thickness of the oxide layer is larger 23 nm than the nominal value of 20 nm. Concerning the annealed samples, the micrograph in Fig. 2 clearly shows a spatial spreading of the band of nanoparticles. From many images, the width of this spreading is estimated as 14 nm. The center of the band remains almost

3 4338 J. Appl. Phys., Vol. 92, No. 8, 15 October 2002 Ignatova et al. FIG. 1. Cross-sectional HREM image of an as-implanted SiO 2 film of nominal 20 nm thickness grown on a Si substrate. The layer marked by white arrows contains the implant ( atoms cm 2 of 12 kev Sb. unchanged at a depth of about 12 nm. The widening of the implantation band by more than a factor of 2 during the annealing implies a rather significant thermal diffusion, which differs from the observations in Ref. 1. Such widening of the implanted precipitates oxides or pure metal within a few nm would not be seen by using techniques like RBS. It is necessary to note that the RBS technique used in Ref. 1 has much less depth resolution, although at an impact energy of 190 kev of Sb used during implantation in Ref. 1, one should expect an effect of even increased mobility of the implant due to radiation-enhanced diffusion. In Figs. 3 and 4, the electron micrographs of samples with different SiO 2 thicknesses and implantation fluences are compared. The images in Fig. 3 represent as-implanted samples with atoms cm 2 in 30 nm SiO 2 Fig. 3 a, atoms cm 2 in 20 nm SiO 2 Fig. 3 b, and atoms cm 2 in 30 nm SiO 2 c. The images in Figs. 4 a - c show the corresponding annealed samples. The contrast differences at the SiO 2 /Si interface are due to the TEM specimen preparation and can be explained by the effects of different thicknesses. In the case of the thinner SiO 2 layer Fig. 3 b, the implant is positioned closer to the interface, since the primary energy is the same 12 kev. In all images FIG. 2. Cross-sectional HREM image of an annealed sample with 20 nm thick SiO 2 film. The implant layer ( atoms cm 2 of 12 kev Sb is wider than the one in Fig. 1. The arrows show the individual nanoclusters. FIG. 3. Comparison of HREM images for several as-implanted samples with 12 kev Sb ions: a atoms cm 2 in 30 nm SiO 2, b atoms cm 2 in 20 nm SiO 2, and c atoms cm 2 in 30 nm SiO 2. of the as-implanted sample with a 20 nm thick SiO 2 layer, the formation of a very narrow band close to the interface 2 nm distant is observed as in Fig. 3 b. The formation of such a narrow interface band before annealing is also documented by Nakajima et al. 12 after implantation of 10 kev Sn ions at a fluence of ions cm 2 into 15 nm thick SiO 2 on Si. After annealing, several studies with Sb, Sn, and Ge report such a layer, 4,6,7 but never as well defined as in the as-implanted case of the present work or Ref. 12. Different explanations have been proposed in literature. Radiation damage due to ion implantation causing ion-beaminduced mixing at the interface may result in an excess of Si in the dioxide layer close to the SiO 2 /Si interface. This excess Si may then act as a nucleation center for gettering implanted impurity atoms. 6 Another possible reason is the compressive strain which exists in the SiO 2 near the interface due to the density mismatch between Si and SiO 2. Introducing implanted ions into the transition region would reduce the strain. 12 The third suggested idea is the segregation of the implant about the interface plane, which may even cause an increase of the pure metal content at the side of the Si substrate as shown in 7,13. One should, however, bear in mind that the segregation cases have been observed under experimental conditions different from our study: For Ge in a thick

4 J. Appl. Phys., Vol. 92, No. 8, 15 October 2002 Ignatova et al FIG. 5. Size distribution of the nanoclusters in an annealed sample ( atoms cm 2 in 30 nm SiO 2 ). FIG. 4. Comparison of HREM images for the corresponding annealed samples 900 C, 10 min, the doses are as in Fig. 3. SiO 2 layer 480 nm with high energy 350 kev 7 and for Ga with implantation and self-sputtering at the same time at 6 and 10 kev. 13 All the images of our results after annealing show a large spreading of nanoclusters around the initial implantation depth, better expressed at higher fluences Figs. 4 a and 4 b. No interface bands are found. Again only with the thinner SiO 2 layer of 20 nm, splitting of the implantation band into two with maxima at depths of about 6 and 15 nm is observed Fig. 4 b. Note that this interesting phenomenon together with the narrow interface band in the asimplanted case is observed so far only in the thinnest silica layer and high fluence, but never for the 30 nm thick SiO 2 layer. An explanation may be the fact that the tail of the implantation profile reaches for the thinnest silica into the substrate, which is probably acting as a barrier as discussed herein. Decreasing the fluence clearly results in less and smaller Sb clusters, both for the as-implanted Fig. 3 c as compared to Fig. 3 a and for the annealed samples Fig. 4 c versus Fig. 4 a. This study also reveals that the thermal annealing is acting as a driving force for rearranging of the nanoclusters according to their size. Figure 5 shows the image of the same sample as in Fig. 4 a at higher magnification, and one can easily see that the particles with size around 5 6 nm are more often localized closer to the interface, while the smaller particles of about 2 nm size redistribute closer to the surface. This is in accordance with similar results found recently for implanted Sn ions after rapid thermal processing RTP and for Sb ions only after long 5 min RTP annealing or furnace annealing during 10 min as in our work, and which have been explained by Ostwald ripening. 6 In the case of the thin SiO 2 layer Fig. 4 b, the proximity to the interface may have inversed the process to result in inverse Ostwald ripening, making the small nanoclusters from the interface band grow at the expense of the large ones in the main implantation band, resulting in a monodisperse size distribution in two bands both redistributed about the middle of the SiO 2 layer. B. Nanocrystal formation It is important to note that even without thermal treatment the implant ( atoms cm 2 of 12 kev Sb into 20 nm thick SiO 2 ) forms discrete nanocrystals in a size range of 2.5 to 6.5 nm which are uniformly distributed in a narrow band of 6 nm at about the implantation depth. Most high-resolution images not all displayed here clearly show lattice plane fringes in some of the nanoclusters, indicating their crystalline structure. In order to determine the lattice parameters, FT of selected areas in the micrographs were performed. Figure 6 shows such a nanocrystal 5.3 nm in size, the masks used to limit the area chosen for the FT of parts of the nanoparticle and of the substrate, the latter serving as internal calibration, and the corresponding computed diffractograms. This particular nanocrystal exhibits two sets of lattice planes, which allow their lattice spacing d 0.62 and 0.27 nm, respectively, with an uncertainty of 0.02 nm and the orientation of the crystal axes to be determined. The nanoparticle is identified as antimony trioxide Sb 2 O 3 in its orthorhombic form valentinite with d nm and d nm. 14 Our work observed the crystalline structure of Sb nanoclusters in as-implanted samples and, beyond that, identified some of the nanocrystals

5 4340 J. Appl. Phys., Vol. 92, No. 8, 15 October 2002 Ignatova et al. FIG. 6. a Cross-sectional HREM image of an as-implanted sample ( Sb atoms cm 2 and 20 nm SiO 2 ) containing an Sb 2 O 3 nanocrystal close to the SiO 2 /Si interface. The two insets, b and c, are the FT diffractograms of the indicated circular areas of the micrograph in the substrate and the nanocrystal. FIG. 7. Cross-sectional HREM image of the same as-implanted sample as in Fig. 6 containing a nanocrystal with d 0.30 nm of an unidentified nature. FIG. 8. Positive ion mass spectrum from FTLMMS of an as-implanted sample 20 nm SiO 2 ) at a dose of Sb atoms cm 2, recorded in the high m/z range. to be of an oxidic form, namely the orthorhombic trioxide Sb 2 O 3. However, this oxide form valentinite has not been corroborated for most of the nanocrystals due to missing lattice distances 0.4 nm, which would distinguish them from metallic crystals. Many of such nanocrystals observed, e.g., the one shown in Fig. 7 from the same as-implanted sample with a lattice spacing of 0.30 nm, can be either metallic or oxidic antimony. In the annealed samples, the size of the nanoparticles is reduced to a range of 2 5 nm with the majority of them around 2 3 nm, some are crystalline. From the TEM images Fig. 2 of such small clusters, it is difficult to deduce their nature, some of them may also have been sputtered back from the crystalline Si substrate. The decreased size of most of the particles prevents recording the HREM micrographs, because the particles are so small that the electron beam used for the analysis could destroy the crystallinity by irradiation effects in this way, causing amorphization. The concentration and the small size of the particles embedded in the samples with lower implantation fluence ( atoms cm 2 ) and/or thinner silica layer 20 nm do not allow obtaining conclusions on the crystallinity of the nanoclusters. Nor can any distinction be made between oxidic or metallic form of such small nanoclusters. As outlined herein, it must be noted that the identification of metallic Sb nanocrystals such as claimed in Ref. 5 can not be based unambiguously on the observation of lattice spacings 0.4 nm because, in this range, all major lattice spacings of Sb metal are within the resolution of HREM images and their FT too close to those of Sb oxides. C. Molecular composition Figure 8 shows the mass spectra of the characteristic positive ions used to perform molecular speciation analysis of Sb implanted to a dose of atoms cm 2 in the 20 nm thick SiO 2 layers as-implanted sample. The spectra taken at low m/z contain only the ions of Sb at m/z 121 and SbO at m/z 137, while the spectrum in Fig. 8 shows the polymeric elemental ions Sb 2,Sb 3, and Sb 4 at m/z 242, 363, and 484, respectively, structural fragments such as SbO,Sb 2 O,and Sb 2 O 2 at m/z 137, 258, and 274, respectively, and the so-called adduct ions Sb 2 O 3.H m/z 291 and Sb 2 O 3.SbO m/z 427. The high mass resolution M/M 10 6 allows the elemental composition of the ions to be determined unambiguously. According to the desorptionionization model in LMMS, 15,16 the laser microbeam irradiation of the sample brings the originally present molecules as neutrals into gas phase, where they form adduct ions by recombination with co-desorbed ions. Hence, adduct ions are of main diagnostic interest for the molecular identification. Structural fragments are believed to originate from intact analyte ions but as they do not contain the entire molecule as a whole, they are only relevant to identify specific functionalities. Ions such as Sb 2 O 2,Sb 2 O, and SbO indicate that

6 J. Appl. Phys., Vol. 92, No. 8, 15 October 2002 Ignatova et al the analyte is an oxide, but that oxide can be Sb 2 O 3 or Sb 2 O 5. Therefore, the positive identification of the analyte as Sb 2 O 3 is based on the detected adducts Sb 2 O 3.H and Sb 2 O 3.SbO. The mass spectrum for the annealed sample after implantation at the same conditions exhibits the dimeric elemental ion Sb 2 at 242 as the main signal at high m/z. The fragment ion SbO indicates the presence of oxide molecules, but does not allow the molecular form to be derived. The FTLMMS is the only technique that can give us unambiguous information on the molecular speciation, but it does not allow depth profiling to be achieved. The diffusion, observed by HREM, can also explain the low signal and the difficulties in detection by FTLMMS for the annealed sample of the same fluence. Indeed, for the rest of the samples of lower fluences or larger thickness of the silica layer, the concentration of the accessible analyte is becoming very small because of the limited information depth of this method 10 nm. VI. CONCLUSIONS Amorphous SiO 2 layers of thickness between 20 and 30 nm on a Si 111 substrate were implanted with 12 kev 121 Sb ions at low fluences ( ions cm 2 ). Some of the chips were subsequently annealed at 900 C for 10 min and HREM and FTLMMS were used to study the implant distribution before and after annealing. All of the samples show considerable widening of the implantation band after the annealing step from 6 nm before to 14 nm afterward, most probably due to thermal diffusion. In the annealed thicker SiO 2 layers 30 nm, nanoclusters are redistributed according to their size with the larger clusters closer to the SiO 2 /Si interface. For the annealed SiO 2 layer of 20 nm thickness, splitting into two bands with two local maxima is observed. The thin SiO 2 layer 20 nm shows for the as-implanted case another peculiarity, a narrow interface band 2 nm distant from the SiO 2 /Si interface. The high-resolution images of HREM clearly show the formation of metal and/or composite oxide nanosized precipitates of Sb with crystalline structure. Nanocrystals were formed even before heat treatment of the samples. Orthorhombic antimony trioxide Sb 2 O 3 was detected in a few nanocrystals of the thin as-implanted SiO 2 layer. Annealing causes a decrease of the average nanocluster size, for most of the clusters from 5 6 nm before annealing to 2 3 nm afterward. A crystalline structure could be corroborated only for some of the annealed nanoclusters. The combined use of HREM and FTLMMS has demonstrated unambiguously, albeit in a few cases, the presence of nanocrystalline Sb 2 O 3 after implantation of 12 kev Sb at ions cm 2 into the 20 nm thick silicon dioxide layer on silicon. FTLMMS was used in nanocrystal research to give unambiguous information on the molecular form, whereas HREM allowed us to visualize individual nanoparticles and to determine their crystalline structure and location in the layer. ACKNOWLEDGMENTS The authors are grateful to R. GrWtzschel of Forschungszentrum Rossendorf for supplying us with the implanted and partly annealed samples. Part of the work has been conducted within the framework of the IUAP-5 program of the Belgian government. 1 A. H. van Ommen, J. Appl. Phys. 61, X. Huang, K. Terashima, Y. Anzai, E. Tokizaki, H. Sasaki, and S. Kimura, Appl. Phys. Lett. 64, A. Nakajima, T. Futatsugi, N. Horiguchi, and N. Yokoyama, Jpn. J. Appl. Phys., Part 2 36, L A. Nakajima, H. Nakao, H. Ueno, T. Futatsugi, and N. Yokoyama, Appl. Phys. Lett. 73, A. Nakajima, H. Nakao, H. Ueno, T. Futatsugi, and N. Yokoyama, J. Vac. Sci. Technol. B 17, S. Spiga, S. Ferrari, M. Fanciulli, B. Schmidt, K.H. Heinig, R. GrWtzschel, A. MXcklich, and G. Pavia, Mater. Res. Soc. Symp. Proc. 647, O S. Oswald, B. Schmidt, and K.H. Heinig, Surf. Interface Anal. 29, L. Van Vaeck, W. Van Roy, H. Struyf, F. Adams, and P. Caravatti, Rapid Commun. Mass Spectrom. 7, P. Caravatti and M. Alleman, Org. Mass Spectrom. 26, V.A. Ignatova, I. R. Chakarov, I. Katardjiev, and U. Wätjen unpublished. 11 S.S. Todorov and I.R. Chakarov, Vacuum 39, A. Nakajima, T. Futatsugi, N. Horiguchi, H. Nakao, and N. Yokoyama, IEEE Intern. Electron Devices Meet V. Ignatova, I. Chakarov, A. Torrisi, and A. Licciardello, Appl. Surf. Sci. 187, ASTM series , JSPDS - International Centre for Diffraction Data H. Struyf, L. Van Vaeck, and R. Van Grieken, Rapid Commun. Mass Spectrom. 10, H. Struyf, L. Van Vaeck, K. Poels, and R. Van Grieken, J. Am. Soc. Mass Spectrom. 9,