- Email: [email protected]
Physical characterization by valence electron energy loss spectroscopy
Microelectronic Engineering 83 (2006) 2364–2367 www.elsevier.com/locate/mee
Physical characterization by valence electron energy loss spectroscopy S. Pokrant
a,*
, R. Pantel b, M. Cheynet
c
a
Philips Semiconductors, 860 rue Jean Monnet, F-38920 Crolles, France ST Microelectronics, 850 rue Jean Monnet, F-38926 Crolles, France Laboratoire de Thermodynamique et Physico-Chimie Me´tallurgiques, UMR-CNRS 5614, ass.INPG-UJF – BP75, 38402 Saint-Martin d’He`res, France b
c
Available online 20 October 2006
Abstract The complexity and the size reduction of state of the art CMOS technologies generate a need for advanced physical characterization. The main difficulty is often to achieve the combination of both high spatial resolution and information on the chemical composition or on the physical properties like the k value or the electrical conductivity. One of the aptest methods to give the desired information is electron energy loss spectroscopy (EELS) in scanning transmission electron microscopy (STEM) mode, especially in the low loss region. This is demonstrated by performing the nickel silicide phase determination on an encroachment formed by Ni diffusion in active Si. 2006 Elsevier B.V. All rights reserved. Keywords: EELS; TEM; NiSi
1. Introduction With the downscaling of semiconductor devices, physical characterization faces two challenges to provide meaningful support for process development or defect analysis. First of all, the dimensions in advanced technology nodes, even at the backened level, are such that for many characterization needs the resolution of scanning electron microscopy (SEM) is not sufficient and is replaced by transmission electron microscopy (TEM). Well known examples are barrier layer conformity or seed layer deposition control. The second problem relates to the increasing demands on the physical properties of the involved materials like lower k values for the insulators and higher electrical conductivity values for the barrier layers. The consequence is very often an increase of the complexity of the employed materials. One of the best examples is the transition from the binary phase system Co–Si to the ternary phase system Ni–Si to obtain better conductivity between contact and active area. This example illustrates *
Corresponding author. Tel.: +33 0438 922620. E-mail address: [email protected] (S. Pokrant).
0167-9317/$ - see front matter 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.mee.2006.10.037
that a purely morphological characterization is not enough in most cases, but must be completed by phase determination or the measurement of physical properties like for example the conductivity. The difficulty is to find a technique which satisfies both needs, the need of local information and the need to determine the exact chemical composition and/or to measure the physical properties. EELS in combination with STEM is a very local technique. At each point of a predefined line or a two dimensional pattern an EEL spectrum is collected. The lateral resolution of the experiment is defined by both the diameter of the electron beam (P0.2 nm) and the sample thickness. Using very small low intensity STEM spot sizes, the spatial resolution of 0.5 nm in a 60 nm thick Si sample can be achieved [1]. As a function of the energy range either core electrons or valence electrons can be probed. For the purpose of chemical analysis core electron energy loss spectroscopy (CEELS) is usually preferred, since the presence of the core edge is linked directly to the presence of the chemical elements. Moreover, the stoichiometry of a compound can be deduced from the intensity ratio between two core edges (quantitative CEELS). The disadvantages of this technique have their origin in the weak scattering cross
S. Pokrant et al. / Microelectronic Engineering 83 (2006) 2364–2367
section of core edges. To obtain a reasonable signal to noise ratio large STEM spot sizes (= high intensity) have to be employed. This leads to reduction of the spatial resolution and in some cases to sample pollution and/or destruction due to long exposure times. Another issue which is very specific to semiconductor applications is the so called 3D effect. Some defects or structures are thinner than the sample thickness. A good example is the NixSiy encroachment in the active zone of a CMOS065 S-RAM cell. The TEM sample contains the entire active area (60 nm in width), but Ni diffusion occurs preferentially on the edges of the active zone and less in the centre, giving rise to a locally inhomogeneous sample. On the defect site the electrons traverse the NixSiy encroachment and the Si of the active area before and/or behind it. Hence the Si edge contains contributions from the defect and from the active zone. Since the Si core edge intensity of the defect cannot be measured separately from the active area contributions, quantitative CEELS cannot provide a satisfying solution to determine the nickel silicide phase. Moreover, the Si edge is not very intense. To obtain a reasonable signal to noise ratio high beam currents and long exposure times are required. This induces sample destruction, pollution and drift and hence affects the precision of the experimental data. An alternative approach is provided by valence electron energy loss spectroscopy (VEELS). In contrast to core electrons, valence electrons have a very large scattering cross section allowing for small spot sizes (=high spatial resolution) and small exposure times (=little sample damage). The signatures of the valence electrons in the low loss region (0–100 eV) can be considered as a finger print of the compound, not of the elements. At one hand this represents a disadvantage, because it is still very difficult to predict theoretically the position and the shape of plasmons for a given compound. Hence it is only possible to use VEELS for the purpose of chemical analysis after having acquired reference spectra and comparing to them. On the other hand the advantage is that in many cases 3D effects have less impact on the analysis, since we can select or separate the compound we are interested in and not a combination of elements as in CEELS. The encroachment is a good example. Crystalline Si as present in the active area gives rise to a peak at 16 eV, while the NixSiy plasmons are situated between 20 and 22 eV. This means that the Si contribution of the active can be separated thanks to the compound selectivity of the low loss signature. Moreover VEELS can provide insight into the physical properties of the material under study like the band gap, the k value or the optical conductivity [2]. In this contribution we show on the example of an encroachment (Fig. 1) formed by uncontrolled diffusion of Ni under a spacer in a failing S-RAM cell, how VEELS can provide local phase determination, which is necessary to understand and control the diffusion process. We will also demonstrate how the optical conductivity of NixSiy grains can be deduced.
2365
Fig. 1. TEM bright field image of a Ni encroachment (as indicated by the arrow) in an S-RAM cell.
2. Experimental The failing S-RAM cell is localised on a 300 mm wafer by the voltage-contrast method using scanning electron microscopy (SEM). A chunk containing the failing SRAM cell is cut out of the wafer by focused ion beam (FIB) sputtering, lifted with an in situ probe and welded to a TEM grid [3]. The chunk is thinned by FIB at 30 kV acceleration voltage and cleaned at 5 kV to get rid of surface amorphization. The final lamella has a thickness of about 50–60 nm. Additionally, two reference samples of the three process relevant Ni–Si phases (NiSi, NiSi2 and Ni2Si) are prepared by sawing and FIB thinning. One sample contains NiSi and NiSi2, the other sample NiSi and Ni2Si. EELS experiments are performed using a FEI TECNAI F20. This is a Schottky field emission gun transmission electron microscope (SFEG-TEM) equipped with a high resolution gatan imaging filter (HR-GIF 2000), a high angle annular dark field (HAADF) detector and a scanning module with spot sizes down to 0.2 nm. The instrumental energy resolution in spectroscopy mode defined as the full width at half maximum of the zero-loss peak in vacuum is 0.8 eV. The reference samples are studied by both CEELS in TEM mode with 4 mrad convergence angle and 20 mrad collection angle and by VEELS in STEM mode with convergence and collection angles both equal to 10 mrad. Core-loss spectra are recorded with an energy dispersion of 0.5 eV/channel in energy windows between 600 and 1100 eV for the Ni–L2,3 edge and between 1500 and 2000 eV for Si–K edge. Low loss spectra are recorded typically between 10 and 90 eV in a 100 eV energy window. On the encroachment sample STEM-EELS analysis in linescan mode is performed with high resolution in space and in energy (0.8 eV for spectra recorded with an energy dispersion of 0.1 eV/channel). The experimental conditions
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are given by the spot size (0.3 nm), the camera length (40 mm), the illumination and collection angles (16 and 4.76 mrad respectively) and the energy window (100 eV between 10 and 90 eV). All spectra are corrected for the dark current and gain variations of the detector. They are processed using the electronic structure tools software [4] developed within Gram 32 [5] to calculate the single scattering distribution using the Fourier-Log deconvolution method proposed by Egerton [2]. The goal of this procedure is to get rid of sample thickness effects before comparing the encroachment sample to the references. Up to this point, all intensities are on an arbitrary scale. To calculate the energy loss function (ELF) the absolute scaling of the data is achieved by applying the sum rule using the refractive index [2]. The refractive index for each nickel silicide phase is obtained from data bases (Mytrix: www.ioffe.ru/SVA/NSM/nk/) or from literature [6]. Kramers–Kronig analysis allows calculating the real (e1) and the imaginary part (e2) of the dielectric function. From there the optical conductivity r can be extrapolated [7]. 3. Results and discussion 3.1. Reference samples Fig. 2 shows a conventional TEM bright field image of the NiSi/Ni2Si reference sample. The diffraction contrast allows estimating the grain sizes to be in the order of 40– 140 nm. Since the reference samples are free of 3D effects and there are no limitations due to the need for high spatial resolution, their chemical composition can be determined by quantitative CEELS in TEM mode (see Ref. [7]). Then low loss spectra are recorded in STEM mode to provide reference spectra for the phase determination of the encroachment. Fig. 3 shows the reference single scattering distributions of the three nickel silicide phases. Since they are recorded between 10 eV and 90 eV with an energy dispersion of 0.1 eV/channel, both plasmons and Ni–M2,3 edges are observed in the same spectrum. The phases can be distinguished by their volume plasmon energy position (Ep). It is shifted from an energy position of 19.4 ± 0.1 eV for NiSi2, to 20 ± 0.1 eV for NiSi and 22.2 ± 0.1 eV for Ni2Si. Although the signal to noise ratio for the Ni–M2,3
Fig. 3. Single scattering distributions of the reference samples of NiSi2, NiSi and Ni2Si. The solid thin line represents Ni2Si, the dotted line NiSi and the thick solid line NiSi2.
core edge is low, since the experimental conditions are optimised to capture the plasmon response, Ni–M2,3 core edge shifts can be observed as a function of the phase: 65.8 eV for Ni2Si, 66.8 eV for NiSi and to 67.5 eV for NiSi2. In summary, taking into account the experimental energy resolution, Ni2Si can be distinguished easily by the plasmon energy and the Ni–M2,3 core edge position, while the difference between NiSi and NiSi2 is subtle, but sizeable in spectra with good signal to noise ratios. Starting from the single scattering distributions of the nickel silicides in the low loss region, the dielectric functions can be calculated. The optical conductivity (r) is related to the dielectric function by r(E) = 2pEee2/h. The optical resistivity q is deduced using the classical relation q = 1/r in analogy to the electrical resistivity. Since the intensive zero-loss peak hides fine structures below 1 eV, e(E), r(E) and q(E) are significant only above 1 eV. Thus, to determine the optical resistivity at E = 0 eV for each compound, the optical resistivity q(E) is extrapolated linearly in the range between 1 and 0 eV starting from 1 eV. Assuming that this extrapolation is correct, the optical resistivities obtained for NiSi, Ni2Si and NiSi2 are equal to 18, 27 and 58 lX cm, respectively. In good agreement with global dc electrical resistivity measurements [8] the optical resistivities determined by VEELS increase from NiSi, to Ni2Si and to NiSi2. 3.2. Encroachment
Fig. 2. TEM bright field image of the NiSi reference sample.
In Fig. 1 we show a bright field image of the encroachment. In agreement with the crystal grain sizes (40 nm– 140 nm) in the reference sample, no grain boundaries can be distinguished in the encroachment, since it is too small (30 nm · 80 nm). High resolution imaging (Fig. 4) shows that the silicide in the encroachment is not very well crystallized, since lattice planes in the reciprocal space can be
S. Pokrant et al. / Microelectronic Engineering 83 (2006) 2364–2367
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Since no shifts neither in the plasmon position nor in the Ni–M2,3 edge position are found throughout the scan, we deduce that the encroachment is built up of one single nickel silicide phase. The only variation we find concerns the width of the plasmon peak. It is a bit wider in the centre of the encroachment than in the middle. Since the width of the plasmon is related to the morphology of the sample, this can be interpreted as a consequence of slight changes in crystallinity, i.e. less crystalline in the middle of the encroachment. To determine the phase we compare the encroachment spectra with Ep = 19.8 ± 0.1 eV and Ni– M2,3 edge = 66.7 eV to the references. The position of the plasmon indicates that Ni2Si (Ep = 22 ± 0.1 eV) can be excluded as possible phase. Taking into account the plasmon peak position and the edge position, the phase of the encroachment is NiSi. As shown in Fig. 5 there is a difference between the NiSi reference and the encroachment single scattering distribution concerning the width of the plasmon peak, but not in the position. Again this can be attributed to the poor crystallinity of the encroachment in comparison to the reference NiSi. Fig. 4. High resolution TEM image of the encroachment. In the inset the Fourier transformation of the encroachment part is shown. Two weak lattice planes can be distinguished in the amorphous back ground.
4. Conclusion We have shown that VEELS is a valuable method for solving analytical problems in advanced semiconductor devices. The advantages over CEELS are better spatial resolution, less sample damage and, very importantly for small patterned structures, VEELS is in some cases less affected by the 3D effect. We have shown that physical properties like the optical resistivity can be successfully determined for full sheet (reference) samples. In future we would like being able to perform the same kind of analysis on patterned structures or defects like the encroachment. The analysis of the optical resistivity on the NiSi encroachment is hindered in this work by the small signal to noise ratio in the VEELS spectra. Working with new generation microscopes should lead to spectra which can be exploited. References
Fig. 5. Singe scattering distribution (solid line) of the encroachment linescan in comparison to the NiSi reference spectrum (dotted line). In the inset a HAADF STEM image of the encroachment is shown. The black line indicated the scan line profile.
hardly distinguished in the amorphous back ground. This represents a huge structural difference in comparison to the reference samples, where the silicide is very well crystallized. To determine the phase we collect 40 low loss spectra along the line shown in the inset in Fig. 5 with a step size of 2.9 nm. The single scattering distributions obtained for each point on the line are compared carefully one by one.
[1] S. Pokrant, M.C. Cheynet, S. Jullian, R. Pantel, Ultramicroscopy 104 (2005) 233. [2] R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, second ed., Plenum Press, New York, 1996. [3] N. Bicais-Lepinay, F. Andre´, R. Pantel, S. Julian, A. Margain, L.F.Tz. Kwakman, Microelectron. Reliab. 42 (2002) 1747. [4] L. Denoyer, G. Duscher, R. French, Electronic Structure Tool (EST) software, Tompkins Count Trust Company Inc., Ithaca, NY, 1996. [5] Gram 32, Galactic Industries, (1996), 325 Main Street, Salem NH 03079 USA. [6] H.W. Chen, J.T. Lue, J. Appl. Phys. 59 (1986) 2165. [7] M.C. Cheynet, R. Pantel, Micron, (2006) accepted. [8] M. Amiotti, A. Borghesi, G. Guizzetti, F. Nava, Phys. Rev. B 42 (1990) 8939.
Physical characterization by valence electron energy loss spectroscopy S. Pokrant
a,*
, R. Pantel b, M. Cheynet
c
a
Philips Semiconductors, 860 rue Jean Monnet, F-38920 Crolles, France ST Microelectronics, 850 rue Jean Monnet, F-38926 Crolles, France Laboratoire de Thermodynamique et Physico-Chimie Me´tallurgiques, UMR-CNRS 5614, ass.INPG-UJF – BP75, 38402 Saint-Martin d’He`res, France b
c
Available online 20 October 2006
Abstract The complexity and the size reduction of state of the art CMOS technologies generate a need for advanced physical characterization. The main difficulty is often to achieve the combination of both high spatial resolution and information on the chemical composition or on the physical properties like the k value or the electrical conductivity. One of the aptest methods to give the desired information is electron energy loss spectroscopy (EELS) in scanning transmission electron microscopy (STEM) mode, especially in the low loss region. This is demonstrated by performing the nickel silicide phase determination on an encroachment formed by Ni diffusion in active Si. 2006 Elsevier B.V. All rights reserved. Keywords: EELS; TEM; NiSi
1. Introduction With the downscaling of semiconductor devices, physical characterization faces two challenges to provide meaningful support for process development or defect analysis. First of all, the dimensions in advanced technology nodes, even at the backened level, are such that for many characterization needs the resolution of scanning electron microscopy (SEM) is not sufficient and is replaced by transmission electron microscopy (TEM). Well known examples are barrier layer conformity or seed layer deposition control. The second problem relates to the increasing demands on the physical properties of the involved materials like lower k values for the insulators and higher electrical conductivity values for the barrier layers. The consequence is very often an increase of the complexity of the employed materials. One of the best examples is the transition from the binary phase system Co–Si to the ternary phase system Ni–Si to obtain better conductivity between contact and active area. This example illustrates *
Corresponding author. Tel.: +33 0438 922620. E-mail address: [email protected] (S. Pokrant).
0167-9317/$ - see front matter 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.mee.2006.10.037
that a purely morphological characterization is not enough in most cases, but must be completed by phase determination or the measurement of physical properties like for example the conductivity. The difficulty is to find a technique which satisfies both needs, the need of local information and the need to determine the exact chemical composition and/or to measure the physical properties. EELS in combination with STEM is a very local technique. At each point of a predefined line or a two dimensional pattern an EEL spectrum is collected. The lateral resolution of the experiment is defined by both the diameter of the electron beam (P0.2 nm) and the sample thickness. Using very small low intensity STEM spot sizes, the spatial resolution of 0.5 nm in a 60 nm thick Si sample can be achieved [1]. As a function of the energy range either core electrons or valence electrons can be probed. For the purpose of chemical analysis core electron energy loss spectroscopy (CEELS) is usually preferred, since the presence of the core edge is linked directly to the presence of the chemical elements. Moreover, the stoichiometry of a compound can be deduced from the intensity ratio between two core edges (quantitative CEELS). The disadvantages of this technique have their origin in the weak scattering cross
S. Pokrant et al. / Microelectronic Engineering 83 (2006) 2364–2367
section of core edges. To obtain a reasonable signal to noise ratio large STEM spot sizes (= high intensity) have to be employed. This leads to reduction of the spatial resolution and in some cases to sample pollution and/or destruction due to long exposure times. Another issue which is very specific to semiconductor applications is the so called 3D effect. Some defects or structures are thinner than the sample thickness. A good example is the NixSiy encroachment in the active zone of a CMOS065 S-RAM cell. The TEM sample contains the entire active area (60 nm in width), but Ni diffusion occurs preferentially on the edges of the active zone and less in the centre, giving rise to a locally inhomogeneous sample. On the defect site the electrons traverse the NixSiy encroachment and the Si of the active area before and/or behind it. Hence the Si edge contains contributions from the defect and from the active zone. Since the Si core edge intensity of the defect cannot be measured separately from the active area contributions, quantitative CEELS cannot provide a satisfying solution to determine the nickel silicide phase. Moreover, the Si edge is not very intense. To obtain a reasonable signal to noise ratio high beam currents and long exposure times are required. This induces sample destruction, pollution and drift and hence affects the precision of the experimental data. An alternative approach is provided by valence electron energy loss spectroscopy (VEELS). In contrast to core electrons, valence electrons have a very large scattering cross section allowing for small spot sizes (=high spatial resolution) and small exposure times (=little sample damage). The signatures of the valence electrons in the low loss region (0–100 eV) can be considered as a finger print of the compound, not of the elements. At one hand this represents a disadvantage, because it is still very difficult to predict theoretically the position and the shape of plasmons for a given compound. Hence it is only possible to use VEELS for the purpose of chemical analysis after having acquired reference spectra and comparing to them. On the other hand the advantage is that in many cases 3D effects have less impact on the analysis, since we can select or separate the compound we are interested in and not a combination of elements as in CEELS. The encroachment is a good example. Crystalline Si as present in the active area gives rise to a peak at 16 eV, while the NixSiy plasmons are situated between 20 and 22 eV. This means that the Si contribution of the active can be separated thanks to the compound selectivity of the low loss signature. Moreover VEELS can provide insight into the physical properties of the material under study like the band gap, the k value or the optical conductivity [2]. In this contribution we show on the example of an encroachment (Fig. 1) formed by uncontrolled diffusion of Ni under a spacer in a failing S-RAM cell, how VEELS can provide local phase determination, which is necessary to understand and control the diffusion process. We will also demonstrate how the optical conductivity of NixSiy grains can be deduced.
2365
Fig. 1. TEM bright field image of a Ni encroachment (as indicated by the arrow) in an S-RAM cell.
2. Experimental The failing S-RAM cell is localised on a 300 mm wafer by the voltage-contrast method using scanning electron microscopy (SEM). A chunk containing the failing SRAM cell is cut out of the wafer by focused ion beam (FIB) sputtering, lifted with an in situ probe and welded to a TEM grid [3]. The chunk is thinned by FIB at 30 kV acceleration voltage and cleaned at 5 kV to get rid of surface amorphization. The final lamella has a thickness of about 50–60 nm. Additionally, two reference samples of the three process relevant Ni–Si phases (NiSi, NiSi2 and Ni2Si) are prepared by sawing and FIB thinning. One sample contains NiSi and NiSi2, the other sample NiSi and Ni2Si. EELS experiments are performed using a FEI TECNAI F20. This is a Schottky field emission gun transmission electron microscope (SFEG-TEM) equipped with a high resolution gatan imaging filter (HR-GIF 2000), a high angle annular dark field (HAADF) detector and a scanning module with spot sizes down to 0.2 nm. The instrumental energy resolution in spectroscopy mode defined as the full width at half maximum of the zero-loss peak in vacuum is 0.8 eV. The reference samples are studied by both CEELS in TEM mode with 4 mrad convergence angle and 20 mrad collection angle and by VEELS in STEM mode with convergence and collection angles both equal to 10 mrad. Core-loss spectra are recorded with an energy dispersion of 0.5 eV/channel in energy windows between 600 and 1100 eV for the Ni–L2,3 edge and between 1500 and 2000 eV for Si–K edge. Low loss spectra are recorded typically between 10 and 90 eV in a 100 eV energy window. On the encroachment sample STEM-EELS analysis in linescan mode is performed with high resolution in space and in energy (0.8 eV for spectra recorded with an energy dispersion of 0.1 eV/channel). The experimental conditions
2366
S. Pokrant et al. / Microelectronic Engineering 83 (2006) 2364–2367
are given by the spot size (0.3 nm), the camera length (40 mm), the illumination and collection angles (16 and 4.76 mrad respectively) and the energy window (100 eV between 10 and 90 eV). All spectra are corrected for the dark current and gain variations of the detector. They are processed using the electronic structure tools software [4] developed within Gram 32 [5] to calculate the single scattering distribution using the Fourier-Log deconvolution method proposed by Egerton [2]. The goal of this procedure is to get rid of sample thickness effects before comparing the encroachment sample to the references. Up to this point, all intensities are on an arbitrary scale. To calculate the energy loss function (ELF) the absolute scaling of the data is achieved by applying the sum rule using the refractive index [2]. The refractive index for each nickel silicide phase is obtained from data bases (Mytrix: www.ioffe.ru/SVA/NSM/nk/) or from literature [6]. Kramers–Kronig analysis allows calculating the real (e1) and the imaginary part (e2) of the dielectric function. From there the optical conductivity r can be extrapolated [7]. 3. Results and discussion 3.1. Reference samples Fig. 2 shows a conventional TEM bright field image of the NiSi/Ni2Si reference sample. The diffraction contrast allows estimating the grain sizes to be in the order of 40– 140 nm. Since the reference samples are free of 3D effects and there are no limitations due to the need for high spatial resolution, their chemical composition can be determined by quantitative CEELS in TEM mode (see Ref. [7]). Then low loss spectra are recorded in STEM mode to provide reference spectra for the phase determination of the encroachment. Fig. 3 shows the reference single scattering distributions of the three nickel silicide phases. Since they are recorded between 10 eV and 90 eV with an energy dispersion of 0.1 eV/channel, both plasmons and Ni–M2,3 edges are observed in the same spectrum. The phases can be distinguished by their volume plasmon energy position (Ep). It is shifted from an energy position of 19.4 ± 0.1 eV for NiSi2, to 20 ± 0.1 eV for NiSi and 22.2 ± 0.1 eV for Ni2Si. Although the signal to noise ratio for the Ni–M2,3
Fig. 3. Single scattering distributions of the reference samples of NiSi2, NiSi and Ni2Si. The solid thin line represents Ni2Si, the dotted line NiSi and the thick solid line NiSi2.
core edge is low, since the experimental conditions are optimised to capture the plasmon response, Ni–M2,3 core edge shifts can be observed as a function of the phase: 65.8 eV for Ni2Si, 66.8 eV for NiSi and to 67.5 eV for NiSi2. In summary, taking into account the experimental energy resolution, Ni2Si can be distinguished easily by the plasmon energy and the Ni–M2,3 core edge position, while the difference between NiSi and NiSi2 is subtle, but sizeable in spectra with good signal to noise ratios. Starting from the single scattering distributions of the nickel silicides in the low loss region, the dielectric functions can be calculated. The optical conductivity (r) is related to the dielectric function by r(E) = 2pEee2/h. The optical resistivity q is deduced using the classical relation q = 1/r in analogy to the electrical resistivity. Since the intensive zero-loss peak hides fine structures below 1 eV, e(E), r(E) and q(E) are significant only above 1 eV. Thus, to determine the optical resistivity at E = 0 eV for each compound, the optical resistivity q(E) is extrapolated linearly in the range between 1 and 0 eV starting from 1 eV. Assuming that this extrapolation is correct, the optical resistivities obtained for NiSi, Ni2Si and NiSi2 are equal to 18, 27 and 58 lX cm, respectively. In good agreement with global dc electrical resistivity measurements [8] the optical resistivities determined by VEELS increase from NiSi, to Ni2Si and to NiSi2. 3.2. Encroachment
Fig. 2. TEM bright field image of the NiSi reference sample.
In Fig. 1 we show a bright field image of the encroachment. In agreement with the crystal grain sizes (40 nm– 140 nm) in the reference sample, no grain boundaries can be distinguished in the encroachment, since it is too small (30 nm · 80 nm). High resolution imaging (Fig. 4) shows that the silicide in the encroachment is not very well crystallized, since lattice planes in the reciprocal space can be
S. Pokrant et al. / Microelectronic Engineering 83 (2006) 2364–2367
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Since no shifts neither in the plasmon position nor in the Ni–M2,3 edge position are found throughout the scan, we deduce that the encroachment is built up of one single nickel silicide phase. The only variation we find concerns the width of the plasmon peak. It is a bit wider in the centre of the encroachment than in the middle. Since the width of the plasmon is related to the morphology of the sample, this can be interpreted as a consequence of slight changes in crystallinity, i.e. less crystalline in the middle of the encroachment. To determine the phase we compare the encroachment spectra with Ep = 19.8 ± 0.1 eV and Ni– M2,3 edge = 66.7 eV to the references. The position of the plasmon indicates that Ni2Si (Ep = 22 ± 0.1 eV) can be excluded as possible phase. Taking into account the plasmon peak position and the edge position, the phase of the encroachment is NiSi. As shown in Fig. 5 there is a difference between the NiSi reference and the encroachment single scattering distribution concerning the width of the plasmon peak, but not in the position. Again this can be attributed to the poor crystallinity of the encroachment in comparison to the reference NiSi. Fig. 4. High resolution TEM image of the encroachment. In the inset the Fourier transformation of the encroachment part is shown. Two weak lattice planes can be distinguished in the amorphous back ground.
4. Conclusion We have shown that VEELS is a valuable method for solving analytical problems in advanced semiconductor devices. The advantages over CEELS are better spatial resolution, less sample damage and, very importantly for small patterned structures, VEELS is in some cases less affected by the 3D effect. We have shown that physical properties like the optical resistivity can be successfully determined for full sheet (reference) samples. In future we would like being able to perform the same kind of analysis on patterned structures or defects like the encroachment. The analysis of the optical resistivity on the NiSi encroachment is hindered in this work by the small signal to noise ratio in the VEELS spectra. Working with new generation microscopes should lead to spectra which can be exploited. References
Fig. 5. Singe scattering distribution (solid line) of the encroachment linescan in comparison to the NiSi reference spectrum (dotted line). In the inset a HAADF STEM image of the encroachment is shown. The black line indicated the scan line profile.
hardly distinguished in the amorphous back ground. This represents a huge structural difference in comparison to the reference samples, where the silicide is very well crystallized. To determine the phase we collect 40 low loss spectra along the line shown in the inset in Fig. 5 with a step size of 2.9 nm. The single scattering distributions obtained for each point on the line are compared carefully one by one.
[1] S. Pokrant, M.C. Cheynet, S. Jullian, R. Pantel, Ultramicroscopy 104 (2005) 233. [2] R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, second ed., Plenum Press, New York, 1996. [3] N. Bicais-Lepinay, F. Andre´, R. Pantel, S. Julian, A. Margain, L.F.Tz. Kwakman, Microelectron. Reliab. 42 (2002) 1747. [4] L. Denoyer, G. Duscher, R. French, Electronic Structure Tool (EST) software, Tompkins Count Trust Company Inc., Ithaca, NY, 1996. [5] Gram 32, Galactic Industries, (1996), 325 Main Street, Salem NH 03079 USA. [6] H.W. Chen, J.T. Lue, J. Appl. Phys. 59 (1986) 2165. [7] M.C. Cheynet, R. Pantel, Micron, (2006) accepted. [8] M. Amiotti, A. Borghesi, G. Guizzetti, F. Nava, Phys. Rev. B 42 (1990) 8939.