Scanning probe measurements on luminescent Si nanoclusters in SiO2 films

Thin Solid Films 515 (2007) 6375 – 6380 www.elsevier.com/locate/tsf

Scanning probe measurements on luminescent Si nanoclusters in SiO2 films J. Mayandi a , T.G. Finstad a,b,⁎, A. Thøgersen b , S. Foss a , U. Serincan c , R. Turan c a

b

Department of Physics, University of Oslo, Pb 1048, Blindern, N-0316, Norway Centre for Materials Science and Nanotechnology, Pb 1126, Blindern, N-0318, Norway c Department of Physics, Middle East Technical University, 06531, Ankara, Turkey Available online 16 January 2007

Abstract Embedded Si nanocrystals in SiO2 have a large current interest due to the prospects for practical applications. For most of these it is essential to characterize and ultimately control the nanocrystal size, size distribution and spatial distribution. Here we present a study of Si nanocrystals and clusters in SiO2 studied by atomic force microscopy (AFM). Since it is an indirect method, it requires several other methods to establish a reliable description of the structure of the samples. We here compare the AFM results with photoluminescence (PL) and transmission electron microscopy (TEM). Si nanocrystals in thermal oxide films (∼ 250 nm) were fabricated by 100 keV Si ion implantation at a dose of 1 × 1017 cm− 2 followed by high temperature annealing. AFM micrographs were taken after different etching times of the oxide and compared to TEM measuerements of the nanocrystal size and distribution. The correlations observed strongly indicate AFM signatures connected to the nanocrystals. We have analyzed and modeled the etch sectioning technique. Comparisons with the experiments let us conclude that the sectioning technique has some memory effect, but yields a distribution of nanocrystals with depth. A dose of 5 × 1016 cm− 2 yields a PL blue shift of about 100 nm relative to the higher dose. No nanocrystals are observed with TEM in this latter case. However distinct signatures can be observed with AFM and is tentatively attributed to the presence of non-crystalline Si-rich nanoclusters. © 2006 Elsevier B.V. All rights reserved. Keywords: AFM; Luminescence; Si QD; Nanocrystal; Nanocluster

1. Introduction There has been much interest in Si nanocrystals (NCs) after Canham's report of luminescence yield in porous Si [1] which was attributed to quantum confinement and many studies followed with different methods for Si NC fabrication in particular for NCs embedded in oxides [2]. It is now established that the luminescence contains effects of quantum confinement and of interface states of the NCs to the oxide [2,3]. In the current study we have used SiO2 films implanted with Si followed by annealing to segregate Si NCs, which is a common way of producing Si NCs [4–6]. A number of experimental analytical techniques have been used to study NCs embedded in oxides. Among these Transmission Electron Microscopy (TEM) is perhaps the most intuitive while one should not ignore the highly refined skills and the effort that are needed to

⁎ Corresponding author. Department of Physics, University of Oslo, Pb 1048, Blindern, N-0316, Norway. Tel.: +47 2 285 2855; fax: +47 2 285 6422. E-mail address: [email protected] (T.G. Finstad). 0040-6090/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2006.11.174

do high quality TEM. We have employed many techniques in our studies and are currently exploring the feasibility of using Atomic Force Microscopy (AFM) as a simple technique to study NC formation, distributions and their evolution. We will present some of the measurements and their analysis in the present work. The aim of the study, in a broader sense, is to understand better the light emission from Si rich oxide and from Si NCs. There has been much experimental and theoretical work discussing the origin of luminescence in Si rich oxides and related systems [2,7], but still there is more room to explore especially the lower super saturation case regarding the origin of luminescence debate. [8,9] In the study of this low saturation case we have previously [9] introduced the use of AFM in combination with TEM, correlating luminescence with structure. We have also argued elsewhere that the local damage may influence the nucleation process of Si nanocrystals sufficiently that simple extrapolations and expectations based upon only the super saturation of excess Si in SiO2 may lead to incorrect conclusions [10], which becomes important for low supersaturation.

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Fig. 1. Photoluminescence spectra of SiO2 films on Si implanted with 100 keV Si for the doses indicated and annealed at 1050 °C for 2 h.

Here we will present a study of different techniques with emphasis on AFM for studying NCs. It should be noted that AFM has been used to study NCs before. Heng et al. [11] studied Ge NCs in SiO2 using that technique, while Sharp et al. [12] recently reported on the complete etching of a SiO2 layer containing NCs where a sedimentation layer of Si NCs remained. We will present measurements on two different implantation doses of Si in SiO2. The highest provides a case for studying the depth distribution of nanocrystals, and present AFM as an alternative method for such studies. The lower implantation dose represent lower Si supersaturation and offer a case where AFM could prove useful as few techniques give indications of the microstructure for this case. After presenting the experimental details, we will summarize the main PL and TEM data in Section 3. Then the AFM measurements are presented in Section 4.1, Section 4.2 presents an analysis and some modeling of etching for revealing NC distribution and Section 4.3 explicitly draws the conclusion about the AFM data by comparison to the model. The lower dose case will be discussed separately in Section 5.

excited by an Ar ion laser (488 nm, 2.54 eV) and detected by a Si detector. Specimens for cross-sectional TEM were prepared of the samples using standard techniques. The structures of the specimens were examined at 200 keV using an analytical JEOL 2000 FX TEM and a field emission analytical JEOL 2010 F TEM system. AFM measurements were performed on the samples measuring the topography with a Digital Instruments Dimension 3100 Scanning Probe Microscope operated in the tapping mode™. This reveals features we attribute to nanocrystals on the surface. The distribution of nanocrystals with depth was probed by thinning the oxide by etching the annealed samples for periods of time in a 1% HF aqueous solution. This etches the SiO2 while not Si. The remaining oxide was measured by ellipsometry in order to find the thickness. However we consider that Si NC formation in the present cases alters the optical properties of the SiO2 films, while we do not have a reliable model of this. Thus the exact etching depth is unknown. Note that we can clearly observe when we have etched through the oxide though, since SiO2 is hydrophilic while Si is hydrophobic. 3. Summary of PL and TEM observations It is necessary to review some of the PL and TEM measurements that have been made on the samples in order to state the foundation for the AFM measurements and interpretations. Fig. 1 shows the PL intensity spectra from two samples implanted with different dose of 1 × 10 17 cm − 2 and 5 × 1016 cm− 2 respectively and annealed at 1050 °C for 2 h. A broad PL peak with center around 880 nm was observed for the higher dose whereas the dose of 5 × 1016 cm− 2 yields a considerable blue shift of about 100 nm relative to the higher

2. Experimental details SiO2 films with a thickness of 250 nm were prepared in a standard process for Si integrated circuits by wet thermal oxidation of p-type single (100) Si wafers with a resistivity of 25–30 Ωcm. The films were implanted with 28Si ions to doses of 1 × 1017 and 5 × 1016 cm− 2 respectively at energy of 100 keV using a Varian DF4 implanter. These doses will be referred to as the high dose case and the lower dose case respectively. The Si depth distribution expected from the SRIM program [13] is near Gaussian with a mean projected range of 140 nm and standard deviation of 43 nm. The samples were annealed at 1050 °C in a nitrogen flow for 2 h. PL spectra of Si+-implanted samples before and after annealing were measured at room temperature;

Fig. 2. High resolution TEM micrograph of the SiO2 film implanted with 1 ×1017 Si cm− 2 and annealed 1050 °C 2 h. The inset is a diffraction pattern showing that the film has Si nanocrystals.

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dose. The higher dose corresponds to a peak concentration of 18 at.% while the lower one corresponds to 9 at.%. This low dose case corresponds to an interesting case in that no NCs are observed by TEM. NCs are clearly identified for the high dose case as can be seen from Fig. 2 giving an example of lattice fringes observed as well as diffraction rings from Si. The NC size is approximately 3.2 ± 0.2 nm with a standard deviation of 1.1 nm. A contrast corresponding to Si NCs are also observed by dark field TEM and by energy filtered imaging. Neither of these techniques yields any contrast in the images from the low dose case, which will be briefly discussed in Section 5. 4. AFM characterization 4.1. AFM results after different etch times Fig. 3 shows a series of AFM topography images of a sample implanted with 28 Si doses of a) 1 × 10 17 cm − 2 and b) 5 × 1016 cm− 2 ions and annealed at 1050 °C for 2 h. The samples were etched in a 1% aqueous HF solution for different times as indicated on the images. It is seen that the surface is smooth before any etching in both cases. Also for the longest etching time the surface appears to be smooth where the oxide has been completely away. It is seen from Fig. 3a that particlelike features appear in the AFM image for the different depths reached by the etching while the etch rate may not be linear with time as explained in Section 2. We argue that the particle-like features are signatures of regions with nanocrystals for Fig. 3a based upon the correlations with PL and TEM [9]. There are no similar features observed in unimplanted etched thermal oxide etched similarly. Fig. 4 shows the height distribution of Si particles for the higher dose case depending upon the etch time. The measurements in Fig. 3b will be treated in Section 5. 4.2. Analysis and modeling of AFM measurements of NCs We will analyze what is measured with the AFM technique by going back to what we could expect to observe in typical situations. We consider the case shown schematically shown in

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Fig. 5 which is Si NCs in SiO2, but the analysis could apply to any particles embedded in any matrix that can be selectively etched. The NCs have a certain distribution in size and they have a distribution in space. The spatial dependence of these distributions is assumed to be fully described by one spatial coordinate, y, taken as the distance (depth) from the surface. The distributions are assumed to be given by a function f(x,y) representing the density of NCs with a size x at a distance y from the surface. A simple example of this function is shown in Fig. 6a where the distribution has a narrow range of sizes, but where the NCs are uniformly distributed in depth. This situation should resemble the case of NCs formed by co-sputtering of Si and SiO2. Another case is shown in Fig. 6d. There the distribution is a function of depth with a maximum at a certain depth from the surface. At any depth the distribution of NC sizes is Gaussian. ! ðx0 ðyÞ−xÞ2 f ðx; yÞ~exp − : Dx0 ðyÞ

ð1Þ

The most probable size, x0, as well as the width of the distribution, Δx0, is a function of depth, y. This situation could apply to the case of ion-implanted Si in SiO2. When a film with NCs, as sketched in Fig. 5a is etched for a period of time we can in an idealized case imagine that we have a surface where the NCs are ‘protruding out of’ the surface as sketched in Fig. 5b. The AFM is now measuring the topography of this surface. Usually the ability of the AFM to accurately reproduce the topography is limited by tip convolution effects [14]. The lateral spread of the image of a NC will appear larger than the actual particle (at least for sizes in the range 1–10 nm) while the height of the NCs extension out of the surface plane can be accurately reproduced as long as the NCs are not too close together. The AFM image would then produce an image where the height of the extension of the NC out of an assumed flat surface can be measured. From the AFM images we have analyzed one can clearly identify the particle-like features and their height can be measured. This could also be done

Fig. 3. AFM tapping mode topography scans of the surface after etching for the times indicated on the graphs, probing different depths of the sample. The Si implantation dose was a) 1 × 1017 cm− 2 and b) 5 × 1016 cm− 2 respectively.

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note that the height distribution of these cases have a negative slope; dh / dz ≤ 0. In fact this is a characteristic of all cases where the AFM height distributions have no ‘memory effect’ from previously etched layers. On the other hand, when etching away a layer of SiO2, there is a possibility that some or all the NCs of that layer become attached to the surface. In such a case the height distribution of NCs on the surface will be Z hðz; yÞ ¼ z

l

Z f ðt; yÞdt þ

y

kðu; yÞf ðz; uÞdu;

ð3Þ

0

where k(u,y) is the fraction of NCs from the etched layer at depth u that becomes attached to the surface at y. This is the memory effect. Fig. 5c and d shows schematically the etching without and with memory respectively. Fig. 6c and f shows the case for having an incorporation of 10% NCs from each etched layers, i.e. k = 0.1 when etching films with size distributions as in Fig. 6a and d respectively. We see in those cases (Fig. 6c and f) that h shows a maximum and dh / dz is changing sign. The height distribution h will then give information about the changing size distribution with depth, but the depth resolution will be worse than that given by the etching rate due to this memory effect. 4.3. Interpretations of AFM measurements after different etch times Comparing the observations to the analysis, we see that the distributions in Fig. 4 have a clear peak, thus the etching has a

Fig. 4. Histograms of height ranges for the different particle-like features of Fig. 3a at the indicated etching times.

automatically on the image allowing data containing the statistics of large numbers to be processed rapidly. Returning to the idealized cases under consideration, we realize that the distribution of heights, h(z,y) of NCs measured with an AFM for the etching case of Fig. 5b should simply be the integral of the size distribution of NCs: Z

l

hðz; yÞ ¼

f ðt; yÞdt:

ð2Þ

z

So h(z,y) is a function of the depth, y, and yields the area density of NCs with a certain height, z, out from the surface. Fig. 6b and e shows the height distributions calculated from the size distributions for the cases in Fig. 6a and d respectively. We

Fig. 5. Schematic drawing of the etching of a SiO2 film containing Si nanocrystals for increasing times of etching; the results of memory effect upon topography. The etch attacks SiO2 but not Si. a) The SiO2 layer with NCs, the surface is smooth. b) After etching a short time. The etched away layer, (light gray), has a thickness y corresponding to the depth below the original surface. NCs ‘protrude’ the surface having a height z out from the surface. c) After further etching when all NCs from the etched layers are dispersed into the etch and transported away. d) The same situation as c) but the NCs of previous etched layers all stick to the surface.

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Fig. 6. Models of NC size distributions, f(x,y), in a matrix for two situations and the resulting AFM topography height distributions, h(z,y), after etching to various depths, y. The size of the NC is called x and z is its height out of the surface. a)–c) are for a narrow Gaussian size distribution which does not vary with depth. This could be expected for example when there is a constant supersaturation as in a co-sputtered film; d)–f) are for a Gaussian distribution of sizes where the most probable size and the standard deviation vary with depth; a) and d) show the size distributions, b) and e) show the corresponding height distributions when there is no memory effect while c) and f) show the corresponding height distributions when there is a memory effect.

memory effect. There is a statistical significant difference in this peak with depth, thus there is a change of average nanocrystal size with depth. It is smallest towards the surface goes through a maximum and then becomes smaller. Since the implanted Si has approximately Gaussian shape this is qualitatively what is expected if the nanocrystal size scales with the Si supersaturation, which is also expected. We should notice our TEM investigation was unable to reveal any statistically significant variation of particle size with depth, also there has been reports

about no variation of crystal size with depth, albeit without any evidence. This shows that AFM can be quite useful for studying nanocrystals. 5. The lower dose case The observations on the low dose case have several interesting aspects to it. No nanocrystals were observed in TEM but PL is observed and it blue shifted with respect to the

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higher dose case. It is not yet resolved whether the luminescence is fundamentally different in the lighter doped case. However, we make some reflections on the case. If the luminescence wavelength is determined by quantum confinement, then the nanocrystal size should be larger than 2 nm and this large size should be detectable under the current TEM conditions [10] as we observe NCs with a size of 1 nm in the 1 × 1017 cm− 2 implantation case. However the issue should be considered controversial. It is in this connection particularly interesting that the AFM method described yields features which resemble those observed for Si nanocrystals in the low dose case. It is possible to think that sufficiently small Si precipitates might be amorphous, especially in an amorphous matrix like SiO2 that could help to stabilize the amorphous phase. Such amorphous clusters might also absorb as well as emit electromagnetic radiation. Since we with certainty have observed Si NCs with sizes down to 1 nm, we deduct that if there is a critical size a cluster has to attain to be crystalline, it is smaller than 1 nm. The energy levels of an amorphous cluster and the energy gap would have to be quite different than those of a nanocrystal if the luminescent properties of the film should be assigned to amorphous clusters of sizes smaller than 1 nm. Of course, as is widely recognized [2,7,15], surface states or silanol associated states could be coupled to the states of the particles and reduce the optical band gap substantially. Other Si–O related bonds may have a similar effect. Another interesting possibility is that at the lower supersaturation, the precursors to precipitated nanocrystals are clusters or chains of Si atoms. They may not give sufficient contrast to be detected by the TEM techniques we have used, while it is possible that they could yield luminescence. It is also possible that they can be attacked differently by the etch and leave a topography signature seen by AFM. While the low supersaturation regime is very interesting by itself, it also indicates that one should approach studies of nanocrystals with AFM with care if it is suspected that one is approaching the parameter space we have discussed here. It is also likely to prevail at the ends of the concentration profiles for higher dose implantations. 6. Summary and conclusion We have applied AFM for the study of Si nanocrystals in SiO2. It can be applied beneficially for qualitative Si nanocrystal

presence and distribution within a parameter space for which the presence has been firmly established. It can be applied semi quantitatively for measurement of nanocrystal size by measuring their height. In diluted HF layer removal, there is a ‘memory’ effect. This should be better quantized and the effect controlled. HF etching and AFM is sensitive to the precursor stage of Si nanocrystals, where their atomic arrangement is unknown, and there is PL associated with the state. Acknowledgments This work has been partially supported by FUNMAT@UiO, the Norwegian Research Council under the NANOMAT program and the European Commission through the FP6 project called SEMINANO under Contract No. NMP4-CT-2004505285. References [1] L.T. Canham, Appl. Phys. Lett. 57 (1990) 1046. [2] L. Pavesi, D.J. Lockwood, Silicon Photonics, Springer–Verlag, Berlin, 2004. [3] J. Heitmann, F. Muller, M. Zacharias, U. Gosele, Adv. Mater. 17 (2005) 795. [4] D.I. Tetelbaum, A.N. Mikhaylov, O.N. Gorshkov, A.P. Kasatkin, A.I. Belov, D.M. Gaponova, S.V. Morozov, Vacuum 78 (2005) 519. [5] J. Grisolia, M. Shalchian, G.Ben Assayag, H. Coffin, C. Bonafos, C. Dumas, S.M. Atarodi, A. Claverie, Solid State Phenom. 108–109 (2005) 71. [6] U.S. Sias, E.C. Moreira, E. Ribeiro, H. Boudinov, L. Amaral, M. Behar, J. Appl. Phys. 95 (2004) 5053. [7] M. Luppi, S. Ossicini, Phys. Rev., B 71 (2005) 035340. [8] F. Iacona, G. Franzo, C. Spinella, J. Appl. Phys. 87 (2000) 1298. [9] J. Mayandi, T.G. Finstad, S. Foss, A. Thøgersen, U. Serincan, R. Turan, Phys. Scr. T126 (2006) 77. [10] J. Mayandi, T.G. Finstad, S. Foss, A. Thøgersen, U. Serincan, R. Turan, Surf. Coat. Technol. (In Press). [11] C.L. Heng, Y.J. Liu, A.T.S. Wee, T.G. Finstad, J. Cryst. Growth 262 (2004) 95. [12] I.D. Sharp, Q. Xu, C.Y. Liao, D.O. Yi, J.W. Beeman, Z. Liliental-Weber, K. M. Yu, et al., J. Appl. Phys. 97 (2005) 124316. [13] J.F. Ziegler, SRIM: The Stopping and Range of Ions in Matter (Program) http://www.srim.org. [14] P. Markiewicz, M.C. Goh, J. Vac. Sci. Technol., B 13 (1995) 1115. [15] T.S. Iwayama, K. Fujita, S. Nakao, K. Saitoh, T. Fujita, N. Itoh, J. Appl. Phys. 75 (1994) 7779.