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Mechanisms of void formation in Ge implanted SiO2 films
Nuclear Instruments and Methods in Physics Research B 207 (2003) 424–433 www.elsevier.com/locate/nimb
Mechanisms of void formation in Ge implanted SiO2 films E.S. Marstein a, A.E. Gunnæs a, U. Serincan b, S. Jørgensen c, A. Olsen a, R. Turan b, T.G. Finstad a,* a
c
Department of Physics, University of Oslo, P.O. Box 1048 Blindern, N-0316 Oslo, Norway b Department of Physics, Middle East Technical University, Ankara 06531, Turkey Department of Chemistry, University of Oslo, P.O. Box 1033 Blindern, N-0315 Oslo, Norway Received 8 August 2002; received in revised form 27 December 2002
Abstract The present paper reports on annealing of Ge implanted SiO2 films and emphasize the observation of voids and the mechanism behind their formation which is considered new. Samples were prepared by ion implanting fluences of 3 1016 and 1 1017 cm2 respectively of 100 keV Ge into amorphous SiO2 films which were subsequently annealed up to 1000 °C in a N2 atmosphere. The structure of the films was studied by transmission electron microscopy (TEM), energy dispersive X-ray spectroscopy (EDS), secondary ion mass spectrometry (SIMS) and X-ray photoelectron spectroscopy (XPS). The most striking of the observations is that spherical voids with diameters up to tens of nanometers are observed in the films after annealing at 1000 °C for 1 h. The volume fraction of voids increases with the Ge fluence. The mechanism behind the void formation is indicated by the evolution of the sample structure after increasing annealing time or temperature; Ge first segregates into nanocrystals which then increase in size by diffusion and Oswald ripening. Ge is quite mobile in SiO2 , and as oxygen or moisture from the annealing atmosphere diffuse in from the surface, the Ge will be bonded in an oxide closer to the surface than the precipitate. There is thus a net flux of Ge out of the nanoprecipitate into an oxide closer to the surface. The volume occupied by the Ge precipitate becomes a void. This model is discussed and it is concluded that it fits the observations. We also report on the filling of the voids by beam induced migration under TEM electron beam exposure. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 81.07.Bc; 66.30.Pa; 68.37.Lp; 68.55.Ln Keywords: Voids; Ion-beam synthesis; Silicon dioxide; Germanium; Nanocrystals
1. Introduction The formation of cavities in crystalline materials after ion implantation is a well known phenomenon
*
Corresponding author. Tel.: +47-2285-6109; fax: +47-22856422. E-mail address: terje.fi[email protected] (T.G. Finstad).
[1,2]. Such cavities have been divided into two groups, namely voids and bubbles. Bubbles are gas filled cavities which are formed after implanting materials with inert gases. Voids are empty cavities and are formed by vacancy clustering. During ion implantation, damage is introduced into the material. Both vacancies and interstitials are generated. Voids can form when more vacancies than interstitials arrive at the void surface. In crystalline
0168-583X/03/$ - see front matter Ó 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0168-583X(03)00965-0
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materials dislocations are suggested to act as a stronger sink for interstitial atoms than for vacancies, resulting in a higher concentration of vacancies than that of interstitials, which again favor growth of voids when these point defects are mobile [1]. Bubbles can be considered to be more closely related to voids than they are related to other precipitates since the binding energy between the constituents of the cavity is small in the former case. Voids and bubbles have been observed in several crystalline materials. To our knowledge there are so far no reported observations of voids in amorphous SiO2 as a result of ion implantation and the mechanisms described above. One can thus say that voids are not expected in SiO2 . Bubbles on the other hand, may have been assumed in SiO2 by some researchers but a direct observation is unknown to us. The present paper reports on the observation of cavities in Ge implanted SiO2 films after annealing at 1000 °C for 1 h. Ge implanted SiO2 films have been the subject of many recent studies [3–12]. None of those have reported void formation to our knowledge. Many of the investigations of Ge and other semiconductor forming species that have been implanted into SiO2 films onto Si substrates have relevance to possible applications utilizing the luminescence [5,13–15] and charge storage properties [4] of nanocrystals in SiO2 . We have previously reported the formation of Ge nanocrystals in SiO2 films implanted with fluences of 3 1016 and 1 1017 cm2 Ge ions at 100 keV after annealing at 800 °C [16]. The present paper focuses on the void formation and the mechanisms leading to the voids. We will present the main experimental observations related to the evolution of the structure of the sample after first presenting the experimental methods. After presentation of the main evidences in Sections 3.1 and 3.2, we will present a qualitative model of the void formation in Section 3.3. In Sections 3.4 and 3.5 we include interesting observations that are off the main focus, which is void formation.
2. Experimental SiO2 films with a thickness of 200 nm were prepared by standard wet thermal oxidation
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(steam oxidation at 1000 °C) of p-type single crystal (1 0 0) Si wafers with resistivities of 25–30 X cm. Two sample sets of SiO2 films were implanted with fluences of 3 1016 and 1 1017 cm2 respectively of 74 Ge ions using an implantation energy of 100 keV using a 0–200 keV Varian DF4 implanter. TRIM calculations [17] give a projected range for the distribution of the implanted Ge at approximately 69 nm and estimated peak Ge concentrations of, respectively, 2.7 and 9.0 at.% for the two different fluences. Annealing of the samples was performed at temperatures between 800 and 1000 °C in a N2 gas flow in a quartz tube furnace using procedures for semiconductor device processing. The gas purity during annealing was not measured. Thus we consider it unknown. Samples for cross-sectional TEM were prepared using standard techniques. The structure of the different samples was examined at 200 keV using an analytical JEOL 2000 FX TEM and a field emission analytical JEOL 2010 F TEM equipped with a Noran Vantage DI+ energy dispersive X-ray spectroscopy (EDS) system. The composition of selected samples was measured using a Cameca 4F secondary ion mass spectrometer (SIMS) and the chemical state of elements was analyzed by X-ray photoelectron spectroscopy (XPS) combined with sputtering (4 keV Arþ ) in a VG Micro Lab III instrument for combining XPS with Auger Electron Spectroscopy at a residual pressure of 109 mbar with Al Ka radiation as excitation source. The binding energies are calibrated against the Si 2p peak from SiO2 at 103.3 eV [18].
3. Results and discussion 3.1. Cavity signature observations Fig. 1 shows a TEM micrograph of a SiO2 film implanted with a fluence of 1 1017 cm2 74 Ge ions that have been annealed at 1000 °C for 1 h. One of the striking features of the micrograph is the band of circular disks which are brighter than the surroundings and located at a depth of approximately 100 nm from the surface. One should also notice that the disks have the brightest color where several of them overlap. These features are
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Fig. 1. TEM micrograph of an SiO2 film on Si implanted with 1 1017 Ge ions per cm2 and annealed at 1000 °C for 1 h. Note the bright disks argued (see text) to be cavities.
interpreted as spherical voids whose projections along the electron beam direction overlap. This interpretation is supported by many other observations which together make a plausible and consistent description of the sample structure and itÕs evolution. It is the main purpose of this paper to present these. The direct evidence from Fig. 1 is not sufficient to identify the disks as voids. The picture of Fig. 1 is a bright field image, and thus the disks corresponds to areas of less electron scattering than the surroundings. This reduced scattering could be due to a material of lighter average atomic number than the surroundings or it could be caused by areas of thinner effective sample thickness than the surroundings. These features could be associated with segregation of light elements or void formation. Whether the disks are precipitates can be deduced from an analysis of the elements present in the samples. The EDS signals from inside the disks are not significantly different from those from the surroundings. We have also considered the possibility of foreign elements being present in the samples but EDS has very low sensitivity for the elements we primarily expect being present. By
SIMS analysis we have observed that the sample after implantation contains excess H and F atoms. The peak concentration of H and F are both qualitatively significant. The actual concentration is uncertain, but this uncertainty is not considered to be critical for the experiments at hand. For each of H and F, the peak concentration is estimated to be around 1020 cm3 by applying sensitivity factors for a Si matrix, which can yield quantitatively incorrect results. After annealing such a sample at 1000 °C, the SIMS signals of F and H are practically absent in the SiO2 film. This indicates that F and H has diffused out of the SiO2 film. The observations are consistent with the bright disks of Fig. 1 being voids, but so far the SIMS observation cannot rule out that they could be gas bubbles containing H or F as it is uncertain how efficient SIMS would detect elements in such a case. (We believe these elements would be detected in such a hypothetical structure, because in similar cases believed to have bubbles of Xe and Ar, those elements have been detected with SIMS.) The SIMS measurements referred to above can be used to further support the formation of voids, but only together with some other results. The origin of F in the samples is from contamination of the beam during ion implantation. GeF4 was used as the gas source and Fþ 4 molecules can evidently be implanted simultaneously with 74 Geþ . The contamination level in the beam was too low to be detected by the beam profile meter, though. (The implanter has later been improved by adding a customized pair of slits after the analyzer.) The excess H in the SiO2 could partly arise from the implantation or could be due to an increased uptake of H or water vapour by the heavily damaged oxide network which has been suggested by Schmidt et al. [6], who measured H concentrations in the range 5–10 atomic percent in their ion implanted SiO2 films after storage. Another important feature to note from Fig. 1 is the dark band located between the surface and the disks. This dark band is rich in Ge, as determined by EDS, and has an amorphous character as determined by selected area diffraction (SAD). The XPS spectrum of Fig. 2(b) corresponds to the Ge signal from the dark amorphous layer. The spectrum is obtained by combining XPS with
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Fig. 2. XPS spectra around the 2p3=2 peak from a sample implanted with a Ge fluence of 1017 cm2 . Both spectra are taken after sputtering in from the surface: (a) is from a sample annealed for 15 min at 1000 °C and is interpreted as arising from a region with Ge crystalline precipitates, (b) from a sample annealed for 1 h at 1000 °C and is interpreted as coming from a region with Ge in an oxidized state probably a SiGe glass.
sputtering. XPS spectra were collected at a series of depths and a depth profile could be established. The depth of the dark band of Fig. 1 corresponds well to a region enriched in Ge and the XPS spectrum indicates that the Ge is in an oxidized state. This would have been observed if the dark layer of Fig. 1 consisted of a SiGe-glass. Of course one should allow for the possibility that ion beam mixing during the sputtering can influence the binding properties of the elements and thus influence the XPS results. The interpretation of the dark layer of Fig. 1 consisting of a SiGe-glass is based upon all the evidences presented and the plausibility of the model. From TEM we also observe features interpreted as voids in the case of the lower fluence of 3 1016 cm3 , as can be seen from Fig. 3. Similarly to the case for higher fluence, we also observe by SAD and EDS an amorphous band that is enriched in Ge towards the surface. The area covered with disks is correlated with the Ge implant fluence. The interpretation is that the void volume is related to the Ge fluence. This would be expected for different interpretations of the disks and their origin. The amount of damage to the SiO2 network is expected to increase with the implantation fluence. An analogy with agglomeration of vacancies is conceivable for an SiO2 network and the present
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Fig. 3. TEM micrograph of a SiO2 film implanted with a fluence of 3 1016 Ge ions cm2 after annealing for 1 h at 1000 °C. Note the small white ÔdisksÕ interpreted as cavities (see text).
observation is qualitatively in agreement with such a model. Also if the disks had been H or F precipitations or bubbles we would expect a scaling of precipitate volume with the implanted fluence. 3.2. Evolution of the structure upon annealing After annealing the Ge implanted SiO2 films at 800 or 900 °C we observe Ge nanocrystals for both fluences. Figs. 4 and 5 show examples of crosssectional TEM micrographs observed after annealing for 1 h at 800 and 900 °C, respectively. We have previously [16] reported that the precipitates formed after an 800 °C anneal in this kind of samples are crystalline as determined from SAD and Raman spectroscopy. We now also observe nanocrystals after 900 °C annealing. From Figs. 4 and 5 it is evident that the precipitates are largest for the 900 °C anneal. These precipitates are also nearly spherical in shape, and we observe that they can contain several domains of different crystal orientations. It should also be noted that the size and positions of the precipitates in Fig. 4 corresponds well to the depth and size of the disks/voids seen in Fig. 1. The observation of nanocrystals is expected and has been reported by several research groups
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Fig. 4. TEM micrographs of a SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h at 800 °C.
Fig. 6 shows cross-sectional TEM micrographs of the high fluence Ge implanted SiO2 film after annealing for different times at 1000 °C. The figure indicates that the general trend is that precipitates form after annealing for short times. We have confirmed that the precipitates are Ge nanocrystals by SAD and EDS. The XPS spectrum of Fig. 2(a) is taken from the depth of the precipitates and is similar to that from crystalline Ge. Fig. 6 shows further that the precipitates grow in size upon extending the annealing time. Then bright disks appear where the Ge crystal precipitates were located and this links the disks/voids directly to the Ge nanocrystal volume. The last stage of the evolution can be referred to as converting the Ge crystal volume into a void, while we still need to argue further that the observed disks are voids. We have referred to our SIMS measurements from the high fluence sample annealed at 1000 °C. If we look at the Ge profile of Fig. 7 and compare it to the series of TEM micrographs of the sample shown in Fig. 6 or to the insert of Fig. 7, we see that the situation of the SIMS profile is that of Ge nanocrystals being located in the SiO2 film. Since this situation occurs prior to the disks/voids we can argue strongly that the disks are not F or H filled cavities since at this stage there is very little F and H in the film. We have now argued that the observation is that of voids. We will argue further that the evolution and formation of voids is a reasonable one. 3.3. Qualitative model for void formation
Fig. 5. TEM micro graphs of a SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h at 900 °C.
[3,4,9,19] using similar fabrication procedure to ours and can be readily understood; The solubility of Ge in SiO2 is lower than the implanted concentration. Therefore the Ge will segregate and form crystalline precipitates. The size distribution of precipitates will change over time with Oswald ripening [20]. The growth rate of the nanocrystals will be faster the higher the temperature.
We can now present a qualitative model for the void formation, and then justify it afterwards. We suggest the voids form in the same volume that was previously occupied by Ge precipitates. There is a net flux of Ge out of the precipitate. The Ge reacts to form strong oxygen bonds in an oxide layer created towards the surface of the sample. The oxygen containing molecules for this reaction are supplied from the surface by diffusion of residual oxygen or moisture in the annealing atmosphere. In order for our model to be correct it is required that the Ge atoms are quite mobile. That the Ge atoms are quite mobile is evidenced in our
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Fig. 6. TEM micrographs of SiO2 films implanted with a fluence of 1 1017 Ge ions cm2 after annealing for at 1000 °C for (a) 15 min, (b) 30 min, (c) 45 min and (d) 60 min.
Fig. 7. SIMS elemental signals versus sputtering time in the sample implanted with a fluence of 1 1017 Ge ions cm2 after annealing at 1000 °C. The signals from Si, Ge and F are indicated by open circles, closed circles and triangles, respectively. A small portion af a TEM cros-section micrograph is shown for evaluation of the depth scale. One broad Ge peak is seen where the Ge nanocrystals are located. There is also a Ge peak at the Si/SiO2 interface. There is F impurities at this interface, while most of the F has been annealed out of the remaining SiO2 film.
experiments by additional observations to those of voids. One of those is the fast Oswald ripening
evidenced for example by Fig. 6 at 1000 °C, but also from Fig. 5 at 900 °C if a smaller initial nucleus is implied. The other observation implying a high mobility is the observation of accumulation of Ge at the Si/SiO2 interface by Ge diffusing through the SiO2 layer [16]. These observations of Ge being quite mobile at elevated temperatures in SiO2 are qualitatively similar to those reported by others [19,21]. The size of nanocrystals of Figs. 5 and 6 are at the large end of the size range reported by others [19,21]. The average size is reported to vary with annealing conditions. A large size implies a large flux of Ge atoms during ripening. Large sizes have been observed under conditions believed to be influenced by H as the annealing atmosphere contained several percent of H. In conclusion, it seems reasonable that we can have a sufficiently high flux in SiO2 to move most of the Ge out of the precipitate to create a void during the time available. We cannot quantify what the required diffusivity would be unless we make a detailed model of the diffusion which would include the Ge concentration, solubility and reaction location. These are presently not known in detail. In order for our model to be reasonable there has to be a driving force for the diffusion. The driving force is suggested to be the change in free energy for Ge bonded to Ge in the Ge precipitate and the Ge bonded to oxygen in the SiGe glass
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network. The oxide bond is of course stronger than that of Ge–Ge. At 1000 °C the Ge precipitate is expected to be in the liquid state since the melting temperature of bulk Ge is 937 °C, so the driving force should really be the free energy change of Ge in the liquid and Ge in the SiGe-glass network. That Ge implanted in SiO2 can oxidize during annealing has been reported before [9,10,16]. In the model of Heinig et al. [9] the oxidation reaction rate of Ge atoms is largest at a location between the surface and the projected range of the implanted Ge. The void formation mechanism described by the presented model can be a general phenomenon. The parameter range over which voids can form by this mechanism is unknown, but we see no reason why the parameter range cannot be wider than what is reported here, nor that it is limited to the local experimental procedures. Of course, as unintentional oxidants such as moisture in the annealing atmosphere enter into the model, and this can be difficult to control accurately, there could be variations in the parameters such as dose, time and temperature leading to void formation. In many experiments the aim is to make nanocrystals. Then void formation is unwanted. The mechanism of the model leading to the voids is new in the sense that it is not a mechanism normally considered for the formation of voids [1,2]. The mechanism itself is similar to that of Kirkendall voids. In the common description thereof the diffusing species are assumed to diffuse by a vacancy mechanism and the net flux of vacancies is opposite the net flux of atoms. In the present case, Ge is probably not being incorporated in the SiO2 random network before it reacts with oxidizing species diffusing in from the surface, and then it appears to us that the diffusion in the present case is not similar to a vacancy diffusion. Another factor which could contribute to the void formation is evaporation of Ge from the surface during annealing. Evaporation from the surface would make another driving force for the net transport of Ge out of the precipitate. Evaporation of Ge from the surface of ion implanted SiO2 during annealing has been reported before [11,12] and should thus be considered as a possible contributing process in our case. Mark-
witz et al. [11] has attributed the Ge loss to the presence of H in their film. In our experiments, we do not know how much Ge that has evaporated. Our experimental observations indicate that Ge evaporation is not the main driving force for the Ge flux out of the Ge precipitate, as we observe a large amount of Ge which is redistributed to the dark belt between the void and the surface as seen for example in Fig. 1. It is likely that the amount of Ge evaporated will depend upon annealing temperature and time and also upon the partial pressure of oxidizing species during the oxidation, as well as the distance between the Ge precipitate and the surface. The Ge loss could be less if much of the Ge is strongly bonded in a SiGe glass than if the Ge was more free to diffuse. Yet another driving force for the Ge diffusion out of the nanoprecipitate is trapping at the Si/ SiO2 interface. We have observed Ge accumulation at this interface by EDS measurements and SIMS measurements [16] in agreement with what others have observed [9]. The model presented is reasonable from our observations and also those by others. Other models for the void formation that are reasonable can be constructed, but at least some of the obvious ones do not fit our experimental observations very well. Precipitation of H or F is already ruled out. Etching of SiO2 by F or HF as the void mechanism does not fit with the voids being most likely linked to the Ge precipitate. Direct reaction of H with Ge, as a possible void formation mechanism, does not fit with the low concentration of H found in the samples at a stage where we have large Ge precipitates but before void formation. H may very well still have an effect on the diffusivity and the diffusion mechanism. The details of chemistry and diffusion related to Ge in SiO2 are not very well known. The current experiments are not designed to shine new light on that topic, but rather focuses on the observation of voids in the system. It is worth mentioning that another kind of mechanism leading to void formation within CdS nanoprecipitates in SiO2 , formed by ion implantation and annealing, has been reported by Meldrum et al. [22–24]. In that case voids form within the CdS precipitate and was explained partly by
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the volume contraction of the implanted material and partly by the out diffusion of the implanted species. 3.4. Spherical shape of nanocrystals We draw the attention to the spherical shape of nanocrystals formed after annealing at 900 and 1000 °C. We think that this shape is peculiar and worth discussing. This shape minimizes the surface or interface energy of a body which has no preferred orientation and which does not have surface planes with different surface energies in contrast to what the surface of a crystal, and certainly that of Ge, tends to have. The examples of faceting of Ge crystal islands are many within the field of surface science and MBE growth [25,26]. At 1000 °C Ge is expected to be in a molten state as the melting temperature of bulk Ge is 927 °C. Then since both SiO2 and Ge is amorphous it is natural that their interface adopts a spherical shape if the mobility allows. For the 900 °C anneal it can also be rationalized by the Ge precipitate being molten when it is formed. The actual temperature could be different from the measured one in the furnace and the melting temperature of a small crystallite is expected to be slightly lower than the melting
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point of bulk crystals. The latter phenomena has been reported for several systems [27]. An alternative (to Ge precipitate being molten) is that it is the SiO2 which controls the shape and SiO2 has no preferred orientations. We think it is interesting to consider the deformation of the SiO2 which must occur and how rigid the SiO2 network appears to be; When Ge is precipitating and nanocrystals are growing, the SiO2 network needs to rearrange to yield space for the crystals. This rearrangement constitute a rigid deformation of the network. When Ge atoms escape the precipitate and are trapped by reaction with oxygen closer to the surface, the deformation made necessary by the nanocrystal essentially remains. The SiO2 network thus reacts quite differently during creation of the Ge nanocrystal and during creation of the void. Part of this difference might be due to implantation damage being present during creation of Ge crystals, but at the same time being annealed out. 3.5. Electron beam filling of voids When the implanted SiO2 films are exposed to intense electron beam irradiation in the TEM, structural changes occur. Fig. 8 shows TEM
Fig. 8. TEM micrograph of the same area of an SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h: (a) before intense electron beam irradiation and (b) after intense electron beam irradiation in the TEM.
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micrographs of the same region of a SiO2 film implanted with a fluence of 1 1017 cm2 74 Ge ions after annealing at 1000 °C for 1 h before and after prolonged exposure to an intense electron beam for Fig. 8(a) and (b), respectively. During exposure, we observe that the voids are gradually filled by the surrounding material. Further, darker micro-regions appear in the Ge-rich belt. This speckled pattern evident in Fig. 8(b) indicates that segregation occurs. The migration of O, Si and Ge in Ge implanted SiO2 films induced by intense electron beam irradiation in a TEM has previously been reported [7,8,12]. Regarding the segregation, we believe the amorphous dark belt in Fig. 8(a) is a SiGe glass. Given sufficient energy to mobilize the atoms in the network, this glass will segregate into a Si rich and Ge rich oxide respectively. It should be noted that this is a well known phenomenon. SiGe glass is commonly used in fiber optics, and this segregation is used to write gratings in the glass or the fiber. The segregation can be induced by energetic photons [28] or by electrons [12]. We observe filling of the voids. This is believed to occur by mobilizing the atoms allowing the system to lower its free energy. The voids should be unstable because of the associated high surface energy. By observation of voids that are interpreted to be located in the TEM specimen without intersecting the TEM specimen surface, one can conclude whether the electron beam induced migration is a surface phenomenon or also a bulk phenomenon. Our preliminary observations indicate it is a bulk phenomenon.
4. Conclusions and summary We have studied Ge implantation in SiO2 films. Our implantations were done at 100 keV and with fluences of 1 1017 and 3 1016 cm2 , respectively. We observe an unexpected microstructure after annealing at 1000 °C consisting of disks in the TEM micrographs. We have argued and concluded that these are voids, and we believe that this is the first time voids have been reported in this context and for SiO2 films. A plausible description for the void formation is that they are the result of
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Mechanisms of void formation in Ge implanted SiO2 films E.S. Marstein a, A.E. Gunnæs a, U. Serincan b, S. Jørgensen c, A. Olsen a, R. Turan b, T.G. Finstad a,* a
c
Department of Physics, University of Oslo, P.O. Box 1048 Blindern, N-0316 Oslo, Norway b Department of Physics, Middle East Technical University, Ankara 06531, Turkey Department of Chemistry, University of Oslo, P.O. Box 1033 Blindern, N-0315 Oslo, Norway Received 8 August 2002; received in revised form 27 December 2002
Abstract The present paper reports on annealing of Ge implanted SiO2 films and emphasize the observation of voids and the mechanism behind their formation which is considered new. Samples were prepared by ion implanting fluences of 3 1016 and 1 1017 cm2 respectively of 100 keV Ge into amorphous SiO2 films which were subsequently annealed up to 1000 °C in a N2 atmosphere. The structure of the films was studied by transmission electron microscopy (TEM), energy dispersive X-ray spectroscopy (EDS), secondary ion mass spectrometry (SIMS) and X-ray photoelectron spectroscopy (XPS). The most striking of the observations is that spherical voids with diameters up to tens of nanometers are observed in the films after annealing at 1000 °C for 1 h. The volume fraction of voids increases with the Ge fluence. The mechanism behind the void formation is indicated by the evolution of the sample structure after increasing annealing time or temperature; Ge first segregates into nanocrystals which then increase in size by diffusion and Oswald ripening. Ge is quite mobile in SiO2 , and as oxygen or moisture from the annealing atmosphere diffuse in from the surface, the Ge will be bonded in an oxide closer to the surface than the precipitate. There is thus a net flux of Ge out of the nanoprecipitate into an oxide closer to the surface. The volume occupied by the Ge precipitate becomes a void. This model is discussed and it is concluded that it fits the observations. We also report on the filling of the voids by beam induced migration under TEM electron beam exposure. Ó 2003 Elsevier Science B.V. All rights reserved. PACS: 81.07.Bc; 66.30.Pa; 68.37.Lp; 68.55.Ln Keywords: Voids; Ion-beam synthesis; Silicon dioxide; Germanium; Nanocrystals
1. Introduction The formation of cavities in crystalline materials after ion implantation is a well known phenomenon
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Corresponding author. Tel.: +47-2285-6109; fax: +47-22856422. E-mail address: terje.fi[email protected] (T.G. Finstad).
[1,2]. Such cavities have been divided into two groups, namely voids and bubbles. Bubbles are gas filled cavities which are formed after implanting materials with inert gases. Voids are empty cavities and are formed by vacancy clustering. During ion implantation, damage is introduced into the material. Both vacancies and interstitials are generated. Voids can form when more vacancies than interstitials arrive at the void surface. In crystalline
0168-583X/03/$ - see front matter Ó 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0168-583X(03)00965-0
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materials dislocations are suggested to act as a stronger sink for interstitial atoms than for vacancies, resulting in a higher concentration of vacancies than that of interstitials, which again favor growth of voids when these point defects are mobile [1]. Bubbles can be considered to be more closely related to voids than they are related to other precipitates since the binding energy between the constituents of the cavity is small in the former case. Voids and bubbles have been observed in several crystalline materials. To our knowledge there are so far no reported observations of voids in amorphous SiO2 as a result of ion implantation and the mechanisms described above. One can thus say that voids are not expected in SiO2 . Bubbles on the other hand, may have been assumed in SiO2 by some researchers but a direct observation is unknown to us. The present paper reports on the observation of cavities in Ge implanted SiO2 films after annealing at 1000 °C for 1 h. Ge implanted SiO2 films have been the subject of many recent studies [3–12]. None of those have reported void formation to our knowledge. Many of the investigations of Ge and other semiconductor forming species that have been implanted into SiO2 films onto Si substrates have relevance to possible applications utilizing the luminescence [5,13–15] and charge storage properties [4] of nanocrystals in SiO2 . We have previously reported the formation of Ge nanocrystals in SiO2 films implanted with fluences of 3 1016 and 1 1017 cm2 Ge ions at 100 keV after annealing at 800 °C [16]. The present paper focuses on the void formation and the mechanisms leading to the voids. We will present the main experimental observations related to the evolution of the structure of the sample after first presenting the experimental methods. After presentation of the main evidences in Sections 3.1 and 3.2, we will present a qualitative model of the void formation in Section 3.3. In Sections 3.4 and 3.5 we include interesting observations that are off the main focus, which is void formation.
2. Experimental SiO2 films with a thickness of 200 nm were prepared by standard wet thermal oxidation
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(steam oxidation at 1000 °C) of p-type single crystal (1 0 0) Si wafers with resistivities of 25–30 X cm. Two sample sets of SiO2 films were implanted with fluences of 3 1016 and 1 1017 cm2 respectively of 74 Ge ions using an implantation energy of 100 keV using a 0–200 keV Varian DF4 implanter. TRIM calculations [17] give a projected range for the distribution of the implanted Ge at approximately 69 nm and estimated peak Ge concentrations of, respectively, 2.7 and 9.0 at.% for the two different fluences. Annealing of the samples was performed at temperatures between 800 and 1000 °C in a N2 gas flow in a quartz tube furnace using procedures for semiconductor device processing. The gas purity during annealing was not measured. Thus we consider it unknown. Samples for cross-sectional TEM were prepared using standard techniques. The structure of the different samples was examined at 200 keV using an analytical JEOL 2000 FX TEM and a field emission analytical JEOL 2010 F TEM equipped with a Noran Vantage DI+ energy dispersive X-ray spectroscopy (EDS) system. The composition of selected samples was measured using a Cameca 4F secondary ion mass spectrometer (SIMS) and the chemical state of elements was analyzed by X-ray photoelectron spectroscopy (XPS) combined with sputtering (4 keV Arþ ) in a VG Micro Lab III instrument for combining XPS with Auger Electron Spectroscopy at a residual pressure of 109 mbar with Al Ka radiation as excitation source. The binding energies are calibrated against the Si 2p peak from SiO2 at 103.3 eV [18].
3. Results and discussion 3.1. Cavity signature observations Fig. 1 shows a TEM micrograph of a SiO2 film implanted with a fluence of 1 1017 cm2 74 Ge ions that have been annealed at 1000 °C for 1 h. One of the striking features of the micrograph is the band of circular disks which are brighter than the surroundings and located at a depth of approximately 100 nm from the surface. One should also notice that the disks have the brightest color where several of them overlap. These features are
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Fig. 1. TEM micrograph of an SiO2 film on Si implanted with 1 1017 Ge ions per cm2 and annealed at 1000 °C for 1 h. Note the bright disks argued (see text) to be cavities.
interpreted as spherical voids whose projections along the electron beam direction overlap. This interpretation is supported by many other observations which together make a plausible and consistent description of the sample structure and itÕs evolution. It is the main purpose of this paper to present these. The direct evidence from Fig. 1 is not sufficient to identify the disks as voids. The picture of Fig. 1 is a bright field image, and thus the disks corresponds to areas of less electron scattering than the surroundings. This reduced scattering could be due to a material of lighter average atomic number than the surroundings or it could be caused by areas of thinner effective sample thickness than the surroundings. These features could be associated with segregation of light elements or void formation. Whether the disks are precipitates can be deduced from an analysis of the elements present in the samples. The EDS signals from inside the disks are not significantly different from those from the surroundings. We have also considered the possibility of foreign elements being present in the samples but EDS has very low sensitivity for the elements we primarily expect being present. By
SIMS analysis we have observed that the sample after implantation contains excess H and F atoms. The peak concentration of H and F are both qualitatively significant. The actual concentration is uncertain, but this uncertainty is not considered to be critical for the experiments at hand. For each of H and F, the peak concentration is estimated to be around 1020 cm3 by applying sensitivity factors for a Si matrix, which can yield quantitatively incorrect results. After annealing such a sample at 1000 °C, the SIMS signals of F and H are practically absent in the SiO2 film. This indicates that F and H has diffused out of the SiO2 film. The observations are consistent with the bright disks of Fig. 1 being voids, but so far the SIMS observation cannot rule out that they could be gas bubbles containing H or F as it is uncertain how efficient SIMS would detect elements in such a case. (We believe these elements would be detected in such a hypothetical structure, because in similar cases believed to have bubbles of Xe and Ar, those elements have been detected with SIMS.) The SIMS measurements referred to above can be used to further support the formation of voids, but only together with some other results. The origin of F in the samples is from contamination of the beam during ion implantation. GeF4 was used as the gas source and Fþ 4 molecules can evidently be implanted simultaneously with 74 Geþ . The contamination level in the beam was too low to be detected by the beam profile meter, though. (The implanter has later been improved by adding a customized pair of slits after the analyzer.) The excess H in the SiO2 could partly arise from the implantation or could be due to an increased uptake of H or water vapour by the heavily damaged oxide network which has been suggested by Schmidt et al. [6], who measured H concentrations in the range 5–10 atomic percent in their ion implanted SiO2 films after storage. Another important feature to note from Fig. 1 is the dark band located between the surface and the disks. This dark band is rich in Ge, as determined by EDS, and has an amorphous character as determined by selected area diffraction (SAD). The XPS spectrum of Fig. 2(b) corresponds to the Ge signal from the dark amorphous layer. The spectrum is obtained by combining XPS with
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Fig. 2. XPS spectra around the 2p3=2 peak from a sample implanted with a Ge fluence of 1017 cm2 . Both spectra are taken after sputtering in from the surface: (a) is from a sample annealed for 15 min at 1000 °C and is interpreted as arising from a region with Ge crystalline precipitates, (b) from a sample annealed for 1 h at 1000 °C and is interpreted as coming from a region with Ge in an oxidized state probably a SiGe glass.
sputtering. XPS spectra were collected at a series of depths and a depth profile could be established. The depth of the dark band of Fig. 1 corresponds well to a region enriched in Ge and the XPS spectrum indicates that the Ge is in an oxidized state. This would have been observed if the dark layer of Fig. 1 consisted of a SiGe-glass. Of course one should allow for the possibility that ion beam mixing during the sputtering can influence the binding properties of the elements and thus influence the XPS results. The interpretation of the dark layer of Fig. 1 consisting of a SiGe-glass is based upon all the evidences presented and the plausibility of the model. From TEM we also observe features interpreted as voids in the case of the lower fluence of 3 1016 cm3 , as can be seen from Fig. 3. Similarly to the case for higher fluence, we also observe by SAD and EDS an amorphous band that is enriched in Ge towards the surface. The area covered with disks is correlated with the Ge implant fluence. The interpretation is that the void volume is related to the Ge fluence. This would be expected for different interpretations of the disks and their origin. The amount of damage to the SiO2 network is expected to increase with the implantation fluence. An analogy with agglomeration of vacancies is conceivable for an SiO2 network and the present
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Fig. 3. TEM micrograph of a SiO2 film implanted with a fluence of 3 1016 Ge ions cm2 after annealing for 1 h at 1000 °C. Note the small white ÔdisksÕ interpreted as cavities (see text).
observation is qualitatively in agreement with such a model. Also if the disks had been H or F precipitations or bubbles we would expect a scaling of precipitate volume with the implanted fluence. 3.2. Evolution of the structure upon annealing After annealing the Ge implanted SiO2 films at 800 or 900 °C we observe Ge nanocrystals for both fluences. Figs. 4 and 5 show examples of crosssectional TEM micrographs observed after annealing for 1 h at 800 and 900 °C, respectively. We have previously [16] reported that the precipitates formed after an 800 °C anneal in this kind of samples are crystalline as determined from SAD and Raman spectroscopy. We now also observe nanocrystals after 900 °C annealing. From Figs. 4 and 5 it is evident that the precipitates are largest for the 900 °C anneal. These precipitates are also nearly spherical in shape, and we observe that they can contain several domains of different crystal orientations. It should also be noted that the size and positions of the precipitates in Fig. 4 corresponds well to the depth and size of the disks/voids seen in Fig. 1. The observation of nanocrystals is expected and has been reported by several research groups
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Fig. 4. TEM micrographs of a SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h at 800 °C.
Fig. 6 shows cross-sectional TEM micrographs of the high fluence Ge implanted SiO2 film after annealing for different times at 1000 °C. The figure indicates that the general trend is that precipitates form after annealing for short times. We have confirmed that the precipitates are Ge nanocrystals by SAD and EDS. The XPS spectrum of Fig. 2(a) is taken from the depth of the precipitates and is similar to that from crystalline Ge. Fig. 6 shows further that the precipitates grow in size upon extending the annealing time. Then bright disks appear where the Ge crystal precipitates were located and this links the disks/voids directly to the Ge nanocrystal volume. The last stage of the evolution can be referred to as converting the Ge crystal volume into a void, while we still need to argue further that the observed disks are voids. We have referred to our SIMS measurements from the high fluence sample annealed at 1000 °C. If we look at the Ge profile of Fig. 7 and compare it to the series of TEM micrographs of the sample shown in Fig. 6 or to the insert of Fig. 7, we see that the situation of the SIMS profile is that of Ge nanocrystals being located in the SiO2 film. Since this situation occurs prior to the disks/voids we can argue strongly that the disks are not F or H filled cavities since at this stage there is very little F and H in the film. We have now argued that the observation is that of voids. We will argue further that the evolution and formation of voids is a reasonable one. 3.3. Qualitative model for void formation
Fig. 5. TEM micro graphs of a SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h at 900 °C.
[3,4,9,19] using similar fabrication procedure to ours and can be readily understood; The solubility of Ge in SiO2 is lower than the implanted concentration. Therefore the Ge will segregate and form crystalline precipitates. The size distribution of precipitates will change over time with Oswald ripening [20]. The growth rate of the nanocrystals will be faster the higher the temperature.
We can now present a qualitative model for the void formation, and then justify it afterwards. We suggest the voids form in the same volume that was previously occupied by Ge precipitates. There is a net flux of Ge out of the precipitate. The Ge reacts to form strong oxygen bonds in an oxide layer created towards the surface of the sample. The oxygen containing molecules for this reaction are supplied from the surface by diffusion of residual oxygen or moisture in the annealing atmosphere. In order for our model to be correct it is required that the Ge atoms are quite mobile. That the Ge atoms are quite mobile is evidenced in our
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Fig. 6. TEM micrographs of SiO2 films implanted with a fluence of 1 1017 Ge ions cm2 after annealing for at 1000 °C for (a) 15 min, (b) 30 min, (c) 45 min and (d) 60 min.
Fig. 7. SIMS elemental signals versus sputtering time in the sample implanted with a fluence of 1 1017 Ge ions cm2 after annealing at 1000 °C. The signals from Si, Ge and F are indicated by open circles, closed circles and triangles, respectively. A small portion af a TEM cros-section micrograph is shown for evaluation of the depth scale. One broad Ge peak is seen where the Ge nanocrystals are located. There is also a Ge peak at the Si/SiO2 interface. There is F impurities at this interface, while most of the F has been annealed out of the remaining SiO2 film.
experiments by additional observations to those of voids. One of those is the fast Oswald ripening
evidenced for example by Fig. 6 at 1000 °C, but also from Fig. 5 at 900 °C if a smaller initial nucleus is implied. The other observation implying a high mobility is the observation of accumulation of Ge at the Si/SiO2 interface by Ge diffusing through the SiO2 layer [16]. These observations of Ge being quite mobile at elevated temperatures in SiO2 are qualitatively similar to those reported by others [19,21]. The size of nanocrystals of Figs. 5 and 6 are at the large end of the size range reported by others [19,21]. The average size is reported to vary with annealing conditions. A large size implies a large flux of Ge atoms during ripening. Large sizes have been observed under conditions believed to be influenced by H as the annealing atmosphere contained several percent of H. In conclusion, it seems reasonable that we can have a sufficiently high flux in SiO2 to move most of the Ge out of the precipitate to create a void during the time available. We cannot quantify what the required diffusivity would be unless we make a detailed model of the diffusion which would include the Ge concentration, solubility and reaction location. These are presently not known in detail. In order for our model to be reasonable there has to be a driving force for the diffusion. The driving force is suggested to be the change in free energy for Ge bonded to Ge in the Ge precipitate and the Ge bonded to oxygen in the SiGe glass
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network. The oxide bond is of course stronger than that of Ge–Ge. At 1000 °C the Ge precipitate is expected to be in the liquid state since the melting temperature of bulk Ge is 937 °C, so the driving force should really be the free energy change of Ge in the liquid and Ge in the SiGe-glass network. That Ge implanted in SiO2 can oxidize during annealing has been reported before [9,10,16]. In the model of Heinig et al. [9] the oxidation reaction rate of Ge atoms is largest at a location between the surface and the projected range of the implanted Ge. The void formation mechanism described by the presented model can be a general phenomenon. The parameter range over which voids can form by this mechanism is unknown, but we see no reason why the parameter range cannot be wider than what is reported here, nor that it is limited to the local experimental procedures. Of course, as unintentional oxidants such as moisture in the annealing atmosphere enter into the model, and this can be difficult to control accurately, there could be variations in the parameters such as dose, time and temperature leading to void formation. In many experiments the aim is to make nanocrystals. Then void formation is unwanted. The mechanism of the model leading to the voids is new in the sense that it is not a mechanism normally considered for the formation of voids [1,2]. The mechanism itself is similar to that of Kirkendall voids. In the common description thereof the diffusing species are assumed to diffuse by a vacancy mechanism and the net flux of vacancies is opposite the net flux of atoms. In the present case, Ge is probably not being incorporated in the SiO2 random network before it reacts with oxidizing species diffusing in from the surface, and then it appears to us that the diffusion in the present case is not similar to a vacancy diffusion. Another factor which could contribute to the void formation is evaporation of Ge from the surface during annealing. Evaporation from the surface would make another driving force for the net transport of Ge out of the precipitate. Evaporation of Ge from the surface of ion implanted SiO2 during annealing has been reported before [11,12] and should thus be considered as a possible contributing process in our case. Mark-
witz et al. [11] has attributed the Ge loss to the presence of H in their film. In our experiments, we do not know how much Ge that has evaporated. Our experimental observations indicate that Ge evaporation is not the main driving force for the Ge flux out of the Ge precipitate, as we observe a large amount of Ge which is redistributed to the dark belt between the void and the surface as seen for example in Fig. 1. It is likely that the amount of Ge evaporated will depend upon annealing temperature and time and also upon the partial pressure of oxidizing species during the oxidation, as well as the distance between the Ge precipitate and the surface. The Ge loss could be less if much of the Ge is strongly bonded in a SiGe glass than if the Ge was more free to diffuse. Yet another driving force for the Ge diffusion out of the nanoprecipitate is trapping at the Si/ SiO2 interface. We have observed Ge accumulation at this interface by EDS measurements and SIMS measurements [16] in agreement with what others have observed [9]. The model presented is reasonable from our observations and also those by others. Other models for the void formation that are reasonable can be constructed, but at least some of the obvious ones do not fit our experimental observations very well. Precipitation of H or F is already ruled out. Etching of SiO2 by F or HF as the void mechanism does not fit with the voids being most likely linked to the Ge precipitate. Direct reaction of H with Ge, as a possible void formation mechanism, does not fit with the low concentration of H found in the samples at a stage where we have large Ge precipitates but before void formation. H may very well still have an effect on the diffusivity and the diffusion mechanism. The details of chemistry and diffusion related to Ge in SiO2 are not very well known. The current experiments are not designed to shine new light on that topic, but rather focuses on the observation of voids in the system. It is worth mentioning that another kind of mechanism leading to void formation within CdS nanoprecipitates in SiO2 , formed by ion implantation and annealing, has been reported by Meldrum et al. [22–24]. In that case voids form within the CdS precipitate and was explained partly by
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the volume contraction of the implanted material and partly by the out diffusion of the implanted species. 3.4. Spherical shape of nanocrystals We draw the attention to the spherical shape of nanocrystals formed after annealing at 900 and 1000 °C. We think that this shape is peculiar and worth discussing. This shape minimizes the surface or interface energy of a body which has no preferred orientation and which does not have surface planes with different surface energies in contrast to what the surface of a crystal, and certainly that of Ge, tends to have. The examples of faceting of Ge crystal islands are many within the field of surface science and MBE growth [25,26]. At 1000 °C Ge is expected to be in a molten state as the melting temperature of bulk Ge is 927 °C. Then since both SiO2 and Ge is amorphous it is natural that their interface adopts a spherical shape if the mobility allows. For the 900 °C anneal it can also be rationalized by the Ge precipitate being molten when it is formed. The actual temperature could be different from the measured one in the furnace and the melting temperature of a small crystallite is expected to be slightly lower than the melting
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point of bulk crystals. The latter phenomena has been reported for several systems [27]. An alternative (to Ge precipitate being molten) is that it is the SiO2 which controls the shape and SiO2 has no preferred orientations. We think it is interesting to consider the deformation of the SiO2 which must occur and how rigid the SiO2 network appears to be; When Ge is precipitating and nanocrystals are growing, the SiO2 network needs to rearrange to yield space for the crystals. This rearrangement constitute a rigid deformation of the network. When Ge atoms escape the precipitate and are trapped by reaction with oxygen closer to the surface, the deformation made necessary by the nanocrystal essentially remains. The SiO2 network thus reacts quite differently during creation of the Ge nanocrystal and during creation of the void. Part of this difference might be due to implantation damage being present during creation of Ge crystals, but at the same time being annealed out. 3.5. Electron beam filling of voids When the implanted SiO2 films are exposed to intense electron beam irradiation in the TEM, structural changes occur. Fig. 8 shows TEM
Fig. 8. TEM micrograph of the same area of an SiO2 film implanted with a fluence of 1 1017 Ge ions cm2 after annealing for 1 h: (a) before intense electron beam irradiation and (b) after intense electron beam irradiation in the TEM.
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micrographs of the same region of a SiO2 film implanted with a fluence of 1 1017 cm2 74 Ge ions after annealing at 1000 °C for 1 h before and after prolonged exposure to an intense electron beam for Fig. 8(a) and (b), respectively. During exposure, we observe that the voids are gradually filled by the surrounding material. Further, darker micro-regions appear in the Ge-rich belt. This speckled pattern evident in Fig. 8(b) indicates that segregation occurs. The migration of O, Si and Ge in Ge implanted SiO2 films induced by intense electron beam irradiation in a TEM has previously been reported [7,8,12]. Regarding the segregation, we believe the amorphous dark belt in Fig. 8(a) is a SiGe glass. Given sufficient energy to mobilize the atoms in the network, this glass will segregate into a Si rich and Ge rich oxide respectively. It should be noted that this is a well known phenomenon. SiGe glass is commonly used in fiber optics, and this segregation is used to write gratings in the glass or the fiber. The segregation can be induced by energetic photons [28] or by electrons [12]. We observe filling of the voids. This is believed to occur by mobilizing the atoms allowing the system to lower its free energy. The voids should be unstable because of the associated high surface energy. By observation of voids that are interpreted to be located in the TEM specimen without intersecting the TEM specimen surface, one can conclude whether the electron beam induced migration is a surface phenomenon or also a bulk phenomenon. Our preliminary observations indicate it is a bulk phenomenon.
4. Conclusions and summary We have studied Ge implantation in SiO2 films. Our implantations were done at 100 keV and with fluences of 1 1017 and 3 1016 cm2 , respectively. We observe an unexpected microstructure after annealing at 1000 °C consisting of disks in the TEM micrographs. We have argued and concluded that these are voids, and we believe that this is the first time voids have been reported in this context and for SiO2 films. A plausible description for the void formation is that they are the result of
a process where first Ge is precipitating and then Ge effectively diffuses out of the precipitates. The main driving force for the Ge flux out of the precipitate is the reaction with oxidant molecules (oxygen or water molecules) diffusing in from the surface. The described void formation process could in principle be active over a wider range of processing parameters than we report here and possibly also for other systems. Acknowledgements This work has been supported by the Norwegian Research Council (NFR) and the Turkish Scientific and Technical Research Council (TUBITAK). We thank Drs. B.G. Svensson and J. Tibbals for stimulating discussions and Dr. M. Linnarsson for the SIMS measurements. References [1] J.H. Evans, Scanning Microscopy 7 (1993) 837. [2] H.S. Rosenbaum, Treatise on Materials Science and Technology, Microstructure of Irradiated Materials, 7, Academic press, 1975. [3] J. von Borany, R. Gr€ otzschel, K.H. Heinig, A. Markwitz, B. Schmidt, W. Skorupa, H.-J. Thees, Solid-State Electron 43 (1999) 1159. [4] Y.C. King, T.J. King, C. Hu, IEEE. Trans. Electron. Devices 48 (2001) 696. [5] K. Masuda, M. Yamamoto, M. Kanaya, Y. Kanemitsu, J. Non-Cryst. Solids 299–302 (2002) 1079. [6] B. Schmidt, D. Grambole, F. Herrmann, Nucl. Instr. and Meth. B 191 (2002) 482. [7] M. Klimenkov, W. Matz, J. von Borany, Nucl. Instr. and Meth. B 168 (2000) 367. [8] M. Klimenkov, W. Matz, S.A. Nepijko, M. Lehmann, Nucl. Instr. and Meth. B 179 (2001) 209. [9] K.H. Heinig, B. Schmidt, A. Markwitz, R. Gr€ otzschel, M. Strobel, S. Oswald, Nucl. Instr. and Meth. B 148 (1999) 969. [10] J. von Borany, R. Gr€ otzschel, K.H. Heinig, A. Markwitz, W. Matz, B. Schmidt, W. Skorupa, Appl. Phys. Lett. 71 (1997) 3215. [11] A. Markwitz, B. Schmidt, W. Matz, R. Grøtzchel, A. M€ ucklich, Nucl. Instr. and Meth. B 142 (1998) 338. [12] K. Beltsios, P. Normand, E. Kapetanakis, D. Tsoukalas, A. Travlos, Microelectron. Eng. 61 (2002) 631. [13] Y. Kanemitsu, H. Uto, Y. Masumoto, Y. Maeda, Appl. Phys. Lett. 61 (1992) 2187. [14] Y. Maeda, N. Tsukamoto, Y. Yazawa, K. Kanemitsu, Y. Masumoto, Appl. Phys. Lett. 59 (1992) 3168.
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