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Cross-sectional structure of tetrahedral amorphous carbon thin films
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HCIINILOGY Surface and Coatings Technology 76-77 (1995) 316-321
Cross-sectional structure of tetrahedral amorphous carbon thin films c.A. Davis, K.M. Knowles, G.A.l. Amaratunga University of Cambridge, Department ofMaterials Science and Metallurgy, Pembroke Street, Cambridge CB2 3QZ, UK University of Cambridge, Department of Engineering, Trumpington Street, Cambridge CB2 1PZ, UK
Abstract A cross-sectional transmission electron microscope study of the low density layers at the surface and at the substrate-film interface of tetrahedral amorphous carbon (ta-C) films grown on (001) silicon substrates is presented. Spatially resolved electron energyloss spectroscopy is used to determine the bonding and compositionof a tetrahedral amorphous carbon film with nanometre spatial resolution. For a ta-C film grown with a substrate bias of - 300 V, an interfacial region approximately 5 nm wideis present in which the carbon is Sp2 bonded and is mixed with silicon and oxygen from the substrate. An Sp2 bonded layer observed at the surface of the film is 1.3 ± OJ nm thick and contains no detectable impurities. It is argued that the Sp2 bonded surface layer is intrinsic to the growth process, but that the Sp2 bonding in the interfacial layer at the substrate may be related to the presence of oxygen from the substrate. Keywords: Tetrahedral amorphous carbon (ta-C); Thin films; Electron energy loss spectroscopy (EELS); Filtered cathodic arc deposition process
1. Introduction Highly tetrahedral amorphous carbon (ta-C) is a predominantly tetrahedrally (Sp3) bonded carbon material which can be produced from a plasma of carbon ions [1]. The growth mechanism of ta-C films is a subject of continuing debate. On the one hand, it has been proposed that the high Sp3 content arises directly from subsurface incorporation (subplantation) of energetic carbon ions [2]. On the other hand, it is argued that the large compressive stresses in the films lead to pressures sufficiently high that Sp3 bonding is favoured energetically [l]. In the subplantation model a very thin layer of Sp2 bonded carbon is expected, with its thickness depending on the subplantation depth and, therefore, on the ion energy. The compressive stress model also predicts a thin surface layer of Sp2 carbon arising from stress relaxation at the surface of the film. In addition to consideration of the growth mechanism, the surface structure of ta-C films is likely to have a strong influence on their usefulness in electronic device applications, such as thin film transistors and field emission displays. Wear and adhesion properties of the films are also strongly dependent on the properties of the surface and interface. Previous structural studies of ta-C have concentrated on volume averaged properties, although there is some indirect evidence for the existence 0257-8972/95/$09.50 © 1995 Elsevier Science SA All rights reserved SSDI 0257-8972(95)02553-7
of a 1 nm thick layer of Sp2 bonded carbon at the surface of ta-C films [3]. In this paper the cross-sectional structure of ta-C films determined using transmission electron microscopy (TEM) and spatially resolved electron energy loss spectroscopy (EELS) is presented. The implications of these results for the film growth mechanism and for electronic applications are discussed.
2. Experimental details The ta-C films studied here were deposited in a filtered cathodic arc apparatus which has been described in detail elsewhere [4,5]. Briefly, a cathodic arc discharge on a 99.995% pure graphite cathode is used to produce a highly ionised carbon plasma. A curved solenoidal magnetic field guides the plasma around a 90° bend in order to eliminate macroscopic particles of graphite which are also emitted by the arc. The films were deposited on silicon (001) wafers which had been ultrasonically cleaned in acetone. The base pressure of the vacuum system was better than 10- 4 Pa, although the pressure rose to around 10- 3 Pa during deposition. The substrate was masked during a transient pressure increase, which occurred in the initial seconds of arcing, by turning off the magnetic field coils on the 90° bend. In order to produce cross-sectional specimens for
CA. Davis et al.jSurface and Coatings Technology
transnnssion electron microscopy TEM and scanning transmission electron microscopy (STEM), the uncoated side of the silicon substrate was ground mechanically to a thickness of 50~ 100 urn and polished to a 6 urn finish. The thinned silicon was then cleaved, with the coated side in tension, along two orthogonal (110) and (l TO) planes. The sharply pointed 90 wedge thus formed was then mounted sideways on a copper support using silverloaded epoxy resin, so that a cross-sectional view of the film would be visible in the thinnest part of the wedge [6]. This preparation procedure was used, rather than ion beam etching, to avoid structural changes which might be introduced by ion bombardment. The crosssections can be accurately oriented edge-on to the optic axis of the microscope using the known (001) orientation of the silicon substrate. Fresnel contrast and high resolution bright field images of the cross-sections were obtained using a Jeol 4000EXII TEM which has a point-to-point resolution of 0.17 nm when operated at 400 kV. Spatially resolved EELS was performed using a Gatan 678 DRV GIF in combination with a 100 kV VG501 STEM. A 7.5 mrad objective aperture was used, giving a spherical aberration-limited probe size of 0.4 nm. Beam spreading through a specimen around 30 nm thick and a delocalisation of 0.4 nm for the carbon K-edge [7] combine with the spherical aberration to give a theoretically achievable spatial resolution of around 0.7 nm. EELS data were recorded with a 0.5 s acquisition time and a collection angle of approximately 15 mrad. In order to obtain a profile of chemical composition and bonding configuration through the film a series of 80 EELS spectra were recorded at 0.46 nm intervals along a line parallel to the growth direction. The specimen drift was found to change depending on the illumination conditions, but was likely to be less than 1 nm min- 1 under the conditions used to acquire the EELS profile. A 1 nm min - 1 drift would give a negligible effect on the spatial resolution but would cause a 10% uncertainty in the position scale of the profile. The spectra were, therefore, recorded from approximately 70 eV to 580 eV so that all analysis could be performed from a single series of spectra to avoid problems with alignment of the various signals. Standard procedures [8] were used to determine the amount of each element present in the specimen from the areas of the absorption edges in the EELS spectra. The close proximity of the calcium L-edges and the carbon K-edge precluded an accurate determination of the calcium concentration, although detection of this distinctive edge was unambiguous. The relative size of the carbon ls-n* pre-edge peak compared with the area of the carbon K-edge was determined by least squares fitting using a Gaussian peak shape. The Sp3 fraction was then determined by comparison to a polycrystalline graphite standard, assuming that the transition matrix 0
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317
elements were invariant and that no Sp1 bonding is present [9]. Noticeable changes in the EELS spectra and bright field contrast occurred in the STEM after approximately 1 min illumination by a focused probe. Care was taken to minimise exposure of the area of interest to the electron beam prior to data collection. Therefore, EELS spectra acquired with 0.5 s illumination are expected to be free of the effects of beam damage.
3. Results
A bright field TEM image of a ta-C film which was grown with a-50 V substrate bias is shown in Fig. 1. The film consists of five distinct layers which are primarily shown by Fresnel contrast, although some absorption contrast is also present. In this underfocused image, lighter regions in the ta-C film correspond to regions of lower density. That is, low density regions are observed
Fig. 1. Cross-sectional bright field TEM image of a ta-C double layer acquired with a 6.5 nm -1 objective aperture and -120 nm defocus. The wedge thickness increases from bottom to top and the silicon substrate is at the left of the image. The deposition was carried out in two stages, between which the film was exposed to atmosphere. The substrate bias was - 50 V. The three low density layers are indicated by arrows.
318
CA. Daviset al.ISurface and Coatings Technology 76-77 (1995) 316-321
at the upper and lower surfaces of the film. These low density layers have been observed repeatedly in films deposited over a wide range of conditions. In addition to the surface and interface layers seen in all ta-C films, the film shown in Fig. 1 also shows a low density layer in the middle of the film. This particular film was deposited in two stages between which it was exposed to the atmosphere for 10 min. Films in which the deposition was stopped for a comparable time of around 2 h but left under vacuum also show a buried layer, but with considerably weaker contrast. We speculate that the buried layer is caused by exposure to atmospheric gases. The fact that a weak layer is also observed for specimens left in the vacuum system is likely to be caused by residual gases present at the base pressure of 10- 4 Pa. Analysis of Fig. 1 is made difficult by the fact that amorphous silicon dioxide, amorphous silicon carbide and spz amorphous carbon all have very similar forward scattering potentials, especially when compared with likely values for the very dense ta-C layers. They would, therefore, give similar contrast when compared to the ta-C layer. For example, the low density layer at the silicon surface could be an SiOz layer or a low density a-C layer, or a combination of the two. Therefore, spatially resolved EELS has been used to give direct information on the composition and bonding of the layers in a ta-C film. Fig. 2 shows the variation of the Sp3 fraction and the carbon, silicon, oxygen and calcium concentrations across a cross-section of a ta-C film which was deposited in a single stage with a substrate bias of - 300 V. For convenience of discussion, the specimen is divided into seven zones labelled A to G which are shown in Fig. 2. Zone A is the silicon substrate. The silicon concentration is at its maximum and the only other element present is a small amount of oxygen in the top 1 nm. A silicon L-edge from this region, curve (a) in Fig. 3, shows an onset around 99 eV, with a clear peak at 101 eV. This is typical of pure silicon [10]. Zone B mainly contains silicon and oxygen, with the oxygen concentration forming a peak approximately 2 nm wide. Through this zone, the carbon concentration rapidly rises to almost half its maximum value. A silicon L-edge obtained in this region is shown as curve (b) in Fig. 3. Peaks at 107 eV and 115 eV are typical of silicon dioxide [lOJ, although a broad onset below 105 eV indicates that some unoxidised silicon is also present. Although the carbon K-edge is too weak for accurate evaluation of the Sp3 content, a clear n* peak is observed, indicating that the carbon in this zone is primarily Sp2 bonded. Across zone C, the silicon concentration drops below the detection limit whilst the carbon concentration increases and the oxygen concentration decreases slowly. The onset of the silicon L-edge (curve (c) in Fig. 3) has
moved to around 103 eV and the 115 eV peak is not present, suggesting that the silicon is bonded to carbon rather than oxygen [10]. The carbon remains low in density, with an Sp3 fraction of only 10-20% Zone D is primarily carbon, although a small amount of oxygen is also present through at least part of this zone. The Sp3 content increases to more than 80% and there is an accompanying increase in the apparent carbon concentration as the carbon atom density increases. Zone E contains 100% carbon, with no detectable impurities and around 90% Sp3 bonding. In zone F, the Sp3 fraction decreases rapidly, with a corresponding decrease in the apparent carbon concentration. No impurities are detected. This zone corresponds to the low-density layer observed on the surface of ta-C films in TEM. Measurement of the distance from the halfway point of the decrease in the Sp3 fraction to the halfway point of the increase in the oxygen concentration yields a layer thickness of 1.3 ± OJ nm. The same layer thickness is obtained from the distance from the halfway point of the Sp3 fraction to the halfway point in the carbon concentration if a 40% decrease in density from Sp3 to spz carbon is taken into account. This is consistent with the 1.3 ± 0.4 nm width of the Fresnel contrast in TEM images. Zone G consists of oxygen, silicon, calcium and spz_ bonded carbon. The extent of this layer varies greatly from specimen to specimen and it is clearly contamination introduced after the film growth. The presence of calcium is consistent with contamination from a biological source such as fingerprints. The large amount of oxygen is consistent with the incorporation of oxygen from the atmosphere into the porous carbon contamination.
4. Discussion
In order to interpret the results in Fig. 2, it is necessary to have confidence that the existence of broad features in the data is not due to poor spatial resolution. A lower bound on the resolution can be obtained from the width of the sharpest feature in the data. The oxygen concentration in zone G increases with a full width at half maximum (FWHM) of 0.7 nm, which is close to the theoretically attainable resolution. The full width at 80% of maximum of the oxygen onset is 1.3 nm and the oxygen concentration does not show a broad tail into zone F (reaching 5% of its maximum value with a half width of 0.8 nm), indicating that the probe does not have extended wings containing a large fraction of the probe current. The magnitude of Fresnel contrast is dependent on the abruptness with which the forward scattering potential changes. The relatively gradual increase in density
319
CA. Davis et al./Surface and Coatings Technology 76-77 (1995) 316-321 4
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Fig. 2. The Sp3 fraction and the carbon, silicon, oxygen and calcium concentrations obtained from EELS spectra as a function of distance across a cross-section of a ta-C film. The film was deposited with a substrate bias of - 300 V. See text for details.
through zone D therefore leads this layer to contribute weakly to the Fresnel contrast, so that the layer thickness of around 3 nm determined from a simple measurement of the width of the Fresnel contrast is significantly less than the actual width of the Sp2 bonded layer. The existence of a 5 nm interface layer of Sp2 amorphous carbon is not predicted by the subplantation model, since it is considerably greater than the expected penetration depth of 300 eV carbon ions [11]. The compressive stress model, however, can accommodate an sp2-bonded layer if the film stress is low in the initial stages of growth. Alternatively, the correlation of the
increase in Sp3 bonding with the decrease in the oxygen concentration in zone D could indicate that the presence of oxygen may act to make Sp3 bonding less stable. This explanation is supported by recent work which has shown a decrease in the Sp3 content for ta-C films containing more than 2 at. % of nitrogen [12]. Although the mechanism which causes oxygen and nitrogen to prevent Sp3 bonding is not clear, the fact that only 2 at. % oxygen is present shows that the effect of the oxygen atoms must extend beyond their nearest neighbours. The existence of the low density buried layer shown
CA, Daviset al,/Surfaceand Coatings Technology 76-77 ( 1995) 316-321
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in Fig. 1 is consistent with the proposition that Sp3 bonding is prevented by the incorporation of oxygen. On exposure to atmosphere after the first deposition stage, the surface of the ta-C film is likely to be covered by adsorbed material including oxygen and water. Bombardment by carbon ions at the beginning of the second deposition stage will remove most of the adsorbed species by sputtering, but some oxygen may be mixed into the film, preventing the formation of Sp3 bonding in a layer a few nanometres thick. Spatially resolved EELS analysis of the two-stage film in Fig. 1 will be performed in the near future to test the prediction that the buried layer is spz carbon containing at least 1-2 at. % oxygen. In contrast to the interface layer formed in the early stages of growth, the spz layer observed at the top of the film (zone F) contains no detectable impurities. The fact that no oxygen is incorporated in this layer even after exposure to atmospheric gases over a period of weeks indicates that it is not porous and that oxygen is not the cause of the spz bonding. That a clear buried layer is not formed in two-stage depositions when the film is left under vacuum for up to 2 h suggests that the surface spz layer does not form by a surface relaxation after deposition. Thus, it appears that the spz-bonded layer is present during growth and so is intrinsic to the growth process. That is, throughout growth the surface layer is continuously converted into the underlying dense sp3-bonded material. The same process then removes the surface layer at the beginning of the second stage of a two-stage deposition, unless sufficientoxygen is present to prevent the formation of the dense Sp3 material. Both TEM images and the spatially resolved EELS data give a 1.3 ± 0.3 nm thickness for the spz surface layer. This thickness is consistent with the subplantation
model since a penetration depth of 1.3 nm is not unreasonable for 300 eV carbon ions, particularly if the surface is low density amorphous carbon. The compressive stress model has also predicted an spz layer approximately 1 nm thick, so that the two models are both consistent with the present data. An investigation of the dependence of the surface layer thickness on the energy of the carbon ions used to deposit the film will be carried out in the near future and may shed further light on the growth mechanism of ta-C films. Despite demonstrations of n-type doping [13] and photoconductivity [14], useful electronic devices based on ta-C films have yet to be achieved. For example, thin film transistor (TFT) devices based on ta-C have been fabricated but show negligible gain. This has been attributed to the presence of a highly defective layer at the ta-C-SiO z interface and is consistent with the 1.3 nm layer of spz amorphous carbon reported in this paper. Failure to observe transistor action in devices in which the ta-C film is grown 'on top of' the SiOz gate insulator is consistent with the wide spz layer observed at the SiOz-ta-C interface. If the spz surface layer is intrinsic to the growth process, as argued above, then subsequent etching of the ta-C film should remove this layer without it reforming. This may enable successful fabrication of ta-C TFT devices. The use of ta-C films as an electron emitting layer in flat screen displays is an exciting new area of research. The field emission properties of ta-C are likely to be strongly effected by the properties of the spz-bonded surface layer. Control of the surface layer may therefore be an important strategy for reducing the electron affinity of ta-C films.
5. Conclusions
The bonding and composition of a ta-C film has been determined with nanometre spatial resolution through the thickness of the film. It has been shown that for 300 eV carbon ions there is significant mixing of the carbon with silicon and oxygen from the native oxide on the surface of the silicon substrate. An interfacial layer approximately 5 nm thick is observed. The prevalence of Sp2 bonding in this layer is argued to be due to the presence of oxygen from the substrate. In contrast, the observed layer of spz carbon at the surface is 1.3 ± 0.3 nm thick and no contaminants were detected. Further work is required to determine the formation processes of the interfacial and surface layers in ta-C films unambiguously. The results presented here, however, suggest several strategies which may lead to ta-C films with improved properties for electronic applications.
CA. Daviset al.jSurfaceand Coatings Technology 76-77 (1995) 316-321
Acknowledgement The authors gratefully acknowledge useful discussions with Dr. CB, Boothroyd and would also like to thank the EPSRC for funding this work.
References [1] D.R. McKenzie, D.A. Muller and B.A. Pailthorpe, Phys. Rev. Lett., 67 (1991) 773. [2] r.Robertson, Phi/os. Trans. R. Soc.LondonSer.A, 342(1993) 277. [3] K.W.R. Gilkes, P.H. Gaskell and J. Yuan, J. Non-Cryst. Solids, 164-166 (1993) 1107. [4] V.S. Veerasamy, GAl Amaratunga, W.I. Milne, P. Hewitt, PJ. Fallon, D.R McKenzie and CA. Davis, Diamond Relat. Mater., 2 (1993) 782.
321
[5] CA. Davis, V.S. Veerasamy, GAl Amaratunga, W.I. Milne and D.R McKenzie, Phi/os. Mag. B, 69 (1994) 1121. [6] lM. Cowley, Acta Crystallogr. A, 25 (1969) 129. [7J H. Kohl and H. Rose, Adv. Electronics Electron Phys., 65 (1985) 173. [8J R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, Plenum, London, 1986, pp. 262-278. [9] S.D. Berger, D.R. McKenzie and PJ. Martin, Philos. Mag. Lett; 57 (1988) 285. [10] W.M. Skiff, RW. Carpenter and S.H. Lin, J. Appl. Phys., 62 (1987) 2439. [11] Y. Lifshitz, S.R Kasi, lW. Rabalais and W. Eckstein, Phys. Rev. B, 41 (1990) 10468. [12] C.A. Davis, D.R McKenzie, Y. Yin, E. Kravtchinskaia, GAJ. Amaratunga and V.S. Veerasamy, Phi/os. Mag. B, 69 (1994) 1133. [13] V.S. Veerasamy, GAl Amaratunga, C.A. Davis, A.E. Timbs, W.L Milne and D.R. McKenzie, J. Phys.: Condens. Matter, 5 (1993) Ll69. [14] GAJ. Amaratunga, V.S. Veerasamy, W.I. Milne, CA. Davis, S.RP. Silva and H.S. McKenzie, Appl. Phvs. Lett., 63 (1993) 370.
&COIRNIiS
ELSEVIER
HCIINILOGY Surface and Coatings Technology 76-77 (1995) 316-321
Cross-sectional structure of tetrahedral amorphous carbon thin films c.A. Davis, K.M. Knowles, G.A.l. Amaratunga University of Cambridge, Department ofMaterials Science and Metallurgy, Pembroke Street, Cambridge CB2 3QZ, UK University of Cambridge, Department of Engineering, Trumpington Street, Cambridge CB2 1PZ, UK
Abstract A cross-sectional transmission electron microscope study of the low density layers at the surface and at the substrate-film interface of tetrahedral amorphous carbon (ta-C) films grown on (001) silicon substrates is presented. Spatially resolved electron energyloss spectroscopy is used to determine the bonding and compositionof a tetrahedral amorphous carbon film with nanometre spatial resolution. For a ta-C film grown with a substrate bias of - 300 V, an interfacial region approximately 5 nm wideis present in which the carbon is Sp2 bonded and is mixed with silicon and oxygen from the substrate. An Sp2 bonded layer observed at the surface of the film is 1.3 ± OJ nm thick and contains no detectable impurities. It is argued that the Sp2 bonded surface layer is intrinsic to the growth process, but that the Sp2 bonding in the interfacial layer at the substrate may be related to the presence of oxygen from the substrate. Keywords: Tetrahedral amorphous carbon (ta-C); Thin films; Electron energy loss spectroscopy (EELS); Filtered cathodic arc deposition process
1. Introduction Highly tetrahedral amorphous carbon (ta-C) is a predominantly tetrahedrally (Sp3) bonded carbon material which can be produced from a plasma of carbon ions [1]. The growth mechanism of ta-C films is a subject of continuing debate. On the one hand, it has been proposed that the high Sp3 content arises directly from subsurface incorporation (subplantation) of energetic carbon ions [2]. On the other hand, it is argued that the large compressive stresses in the films lead to pressures sufficiently high that Sp3 bonding is favoured energetically [l]. In the subplantation model a very thin layer of Sp2 bonded carbon is expected, with its thickness depending on the subplantation depth and, therefore, on the ion energy. The compressive stress model also predicts a thin surface layer of Sp2 carbon arising from stress relaxation at the surface of the film. In addition to consideration of the growth mechanism, the surface structure of ta-C films is likely to have a strong influence on their usefulness in electronic device applications, such as thin film transistors and field emission displays. Wear and adhesion properties of the films are also strongly dependent on the properties of the surface and interface. Previous structural studies of ta-C have concentrated on volume averaged properties, although there is some indirect evidence for the existence 0257-8972/95/$09.50 © 1995 Elsevier Science SA All rights reserved SSDI 0257-8972(95)02553-7
of a 1 nm thick layer of Sp2 bonded carbon at the surface of ta-C films [3]. In this paper the cross-sectional structure of ta-C films determined using transmission electron microscopy (TEM) and spatially resolved electron energy loss spectroscopy (EELS) is presented. The implications of these results for the film growth mechanism and for electronic applications are discussed.
2. Experimental details The ta-C films studied here were deposited in a filtered cathodic arc apparatus which has been described in detail elsewhere [4,5]. Briefly, a cathodic arc discharge on a 99.995% pure graphite cathode is used to produce a highly ionised carbon plasma. A curved solenoidal magnetic field guides the plasma around a 90° bend in order to eliminate macroscopic particles of graphite which are also emitted by the arc. The films were deposited on silicon (001) wafers which had been ultrasonically cleaned in acetone. The base pressure of the vacuum system was better than 10- 4 Pa, although the pressure rose to around 10- 3 Pa during deposition. The substrate was masked during a transient pressure increase, which occurred in the initial seconds of arcing, by turning off the magnetic field coils on the 90° bend. In order to produce cross-sectional specimens for
CA. Davis et al.jSurface and Coatings Technology
transnnssion electron microscopy TEM and scanning transmission electron microscopy (STEM), the uncoated side of the silicon substrate was ground mechanically to a thickness of 50~ 100 urn and polished to a 6 urn finish. The thinned silicon was then cleaved, with the coated side in tension, along two orthogonal (110) and (l TO) planes. The sharply pointed 90 wedge thus formed was then mounted sideways on a copper support using silverloaded epoxy resin, so that a cross-sectional view of the film would be visible in the thinnest part of the wedge [6]. This preparation procedure was used, rather than ion beam etching, to avoid structural changes which might be introduced by ion bombardment. The crosssections can be accurately oriented edge-on to the optic axis of the microscope using the known (001) orientation of the silicon substrate. Fresnel contrast and high resolution bright field images of the cross-sections were obtained using a Jeol 4000EXII TEM which has a point-to-point resolution of 0.17 nm when operated at 400 kV. Spatially resolved EELS was performed using a Gatan 678 DRV GIF in combination with a 100 kV VG501 STEM. A 7.5 mrad objective aperture was used, giving a spherical aberration-limited probe size of 0.4 nm. Beam spreading through a specimen around 30 nm thick and a delocalisation of 0.4 nm for the carbon K-edge [7] combine with the spherical aberration to give a theoretically achievable spatial resolution of around 0.7 nm. EELS data were recorded with a 0.5 s acquisition time and a collection angle of approximately 15 mrad. In order to obtain a profile of chemical composition and bonding configuration through the film a series of 80 EELS spectra were recorded at 0.46 nm intervals along a line parallel to the growth direction. The specimen drift was found to change depending on the illumination conditions, but was likely to be less than 1 nm min- 1 under the conditions used to acquire the EELS profile. A 1 nm min - 1 drift would give a negligible effect on the spatial resolution but would cause a 10% uncertainty in the position scale of the profile. The spectra were, therefore, recorded from approximately 70 eV to 580 eV so that all analysis could be performed from a single series of spectra to avoid problems with alignment of the various signals. Standard procedures [8] were used to determine the amount of each element present in the specimen from the areas of the absorption edges in the EELS spectra. The close proximity of the calcium L-edges and the carbon K-edge precluded an accurate determination of the calcium concentration, although detection of this distinctive edge was unambiguous. The relative size of the carbon ls-n* pre-edge peak compared with the area of the carbon K-edge was determined by least squares fitting using a Gaussian peak shape. The Sp3 fraction was then determined by comparison to a polycrystalline graphite standard, assuming that the transition matrix 0
76~77
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317
elements were invariant and that no Sp1 bonding is present [9]. Noticeable changes in the EELS spectra and bright field contrast occurred in the STEM after approximately 1 min illumination by a focused probe. Care was taken to minimise exposure of the area of interest to the electron beam prior to data collection. Therefore, EELS spectra acquired with 0.5 s illumination are expected to be free of the effects of beam damage.
3. Results
A bright field TEM image of a ta-C film which was grown with a-50 V substrate bias is shown in Fig. 1. The film consists of five distinct layers which are primarily shown by Fresnel contrast, although some absorption contrast is also present. In this underfocused image, lighter regions in the ta-C film correspond to regions of lower density. That is, low density regions are observed
Fig. 1. Cross-sectional bright field TEM image of a ta-C double layer acquired with a 6.5 nm -1 objective aperture and -120 nm defocus. The wedge thickness increases from bottom to top and the silicon substrate is at the left of the image. The deposition was carried out in two stages, between which the film was exposed to atmosphere. The substrate bias was - 50 V. The three low density layers are indicated by arrows.
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CA. Daviset al.ISurface and Coatings Technology 76-77 (1995) 316-321
at the upper and lower surfaces of the film. These low density layers have been observed repeatedly in films deposited over a wide range of conditions. In addition to the surface and interface layers seen in all ta-C films, the film shown in Fig. 1 also shows a low density layer in the middle of the film. This particular film was deposited in two stages between which it was exposed to the atmosphere for 10 min. Films in which the deposition was stopped for a comparable time of around 2 h but left under vacuum also show a buried layer, but with considerably weaker contrast. We speculate that the buried layer is caused by exposure to atmospheric gases. The fact that a weak layer is also observed for specimens left in the vacuum system is likely to be caused by residual gases present at the base pressure of 10- 4 Pa. Analysis of Fig. 1 is made difficult by the fact that amorphous silicon dioxide, amorphous silicon carbide and spz amorphous carbon all have very similar forward scattering potentials, especially when compared with likely values for the very dense ta-C layers. They would, therefore, give similar contrast when compared to the ta-C layer. For example, the low density layer at the silicon surface could be an SiOz layer or a low density a-C layer, or a combination of the two. Therefore, spatially resolved EELS has been used to give direct information on the composition and bonding of the layers in a ta-C film. Fig. 2 shows the variation of the Sp3 fraction and the carbon, silicon, oxygen and calcium concentrations across a cross-section of a ta-C film which was deposited in a single stage with a substrate bias of - 300 V. For convenience of discussion, the specimen is divided into seven zones labelled A to G which are shown in Fig. 2. Zone A is the silicon substrate. The silicon concentration is at its maximum and the only other element present is a small amount of oxygen in the top 1 nm. A silicon L-edge from this region, curve (a) in Fig. 3, shows an onset around 99 eV, with a clear peak at 101 eV. This is typical of pure silicon [10]. Zone B mainly contains silicon and oxygen, with the oxygen concentration forming a peak approximately 2 nm wide. Through this zone, the carbon concentration rapidly rises to almost half its maximum value. A silicon L-edge obtained in this region is shown as curve (b) in Fig. 3. Peaks at 107 eV and 115 eV are typical of silicon dioxide [lOJ, although a broad onset below 105 eV indicates that some unoxidised silicon is also present. Although the carbon K-edge is too weak for accurate evaluation of the Sp3 content, a clear n* peak is observed, indicating that the carbon in this zone is primarily Sp2 bonded. Across zone C, the silicon concentration drops below the detection limit whilst the carbon concentration increases and the oxygen concentration decreases slowly. The onset of the silicon L-edge (curve (c) in Fig. 3) has
moved to around 103 eV and the 115 eV peak is not present, suggesting that the silicon is bonded to carbon rather than oxygen [10]. The carbon remains low in density, with an Sp3 fraction of only 10-20% Zone D is primarily carbon, although a small amount of oxygen is also present through at least part of this zone. The Sp3 content increases to more than 80% and there is an accompanying increase in the apparent carbon concentration as the carbon atom density increases. Zone E contains 100% carbon, with no detectable impurities and around 90% Sp3 bonding. In zone F, the Sp3 fraction decreases rapidly, with a corresponding decrease in the apparent carbon concentration. No impurities are detected. This zone corresponds to the low-density layer observed on the surface of ta-C films in TEM. Measurement of the distance from the halfway point of the decrease in the Sp3 fraction to the halfway point of the increase in the oxygen concentration yields a layer thickness of 1.3 ± OJ nm. The same layer thickness is obtained from the distance from the halfway point of the Sp3 fraction to the halfway point in the carbon concentration if a 40% decrease in density from Sp3 to spz carbon is taken into account. This is consistent with the 1.3 ± 0.4 nm width of the Fresnel contrast in TEM images. Zone G consists of oxygen, silicon, calcium and spz_ bonded carbon. The extent of this layer varies greatly from specimen to specimen and it is clearly contamination introduced after the film growth. The presence of calcium is consistent with contamination from a biological source such as fingerprints. The large amount of oxygen is consistent with the incorporation of oxygen from the atmosphere into the porous carbon contamination.
4. Discussion
In order to interpret the results in Fig. 2, it is necessary to have confidence that the existence of broad features in the data is not due to poor spatial resolution. A lower bound on the resolution can be obtained from the width of the sharpest feature in the data. The oxygen concentration in zone G increases with a full width at half maximum (FWHM) of 0.7 nm, which is close to the theoretically attainable resolution. The full width at 80% of maximum of the oxygen onset is 1.3 nm and the oxygen concentration does not show a broad tail into zone F (reaching 5% of its maximum value with a half width of 0.8 nm), indicating that the probe does not have extended wings containing a large fraction of the probe current. The magnitude of Fresnel contrast is dependent on the abruptness with which the forward scattering potential changes. The relatively gradual increase in density
319
CA. Davis et al./Surface and Coatings Technology 76-77 (1995) 316-321 4
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through zone D therefore leads this layer to contribute weakly to the Fresnel contrast, so that the layer thickness of around 3 nm determined from a simple measurement of the width of the Fresnel contrast is significantly less than the actual width of the Sp2 bonded layer. The existence of a 5 nm interface layer of Sp2 amorphous carbon is not predicted by the subplantation model, since it is considerably greater than the expected penetration depth of 300 eV carbon ions [11]. The compressive stress model, however, can accommodate an sp2-bonded layer if the film stress is low in the initial stages of growth. Alternatively, the correlation of the
increase in Sp3 bonding with the decrease in the oxygen concentration in zone D could indicate that the presence of oxygen may act to make Sp3 bonding less stable. This explanation is supported by recent work which has shown a decrease in the Sp3 content for ta-C films containing more than 2 at. % of nitrogen [12]. Although the mechanism which causes oxygen and nitrogen to prevent Sp3 bonding is not clear, the fact that only 2 at. % oxygen is present shows that the effect of the oxygen atoms must extend beyond their nearest neighbours. The existence of the low density buried layer shown
CA, Daviset al,/Surfaceand Coatings Technology 76-77 ( 1995) 316-321
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in Fig. 1 is consistent with the proposition that Sp3 bonding is prevented by the incorporation of oxygen. On exposure to atmosphere after the first deposition stage, the surface of the ta-C film is likely to be covered by adsorbed material including oxygen and water. Bombardment by carbon ions at the beginning of the second deposition stage will remove most of the adsorbed species by sputtering, but some oxygen may be mixed into the film, preventing the formation of Sp3 bonding in a layer a few nanometres thick. Spatially resolved EELS analysis of the two-stage film in Fig. 1 will be performed in the near future to test the prediction that the buried layer is spz carbon containing at least 1-2 at. % oxygen. In contrast to the interface layer formed in the early stages of growth, the spz layer observed at the top of the film (zone F) contains no detectable impurities. The fact that no oxygen is incorporated in this layer even after exposure to atmospheric gases over a period of weeks indicates that it is not porous and that oxygen is not the cause of the spz bonding. That a clear buried layer is not formed in two-stage depositions when the film is left under vacuum for up to 2 h suggests that the surface spz layer does not form by a surface relaxation after deposition. Thus, it appears that the spz-bonded layer is present during growth and so is intrinsic to the growth process. That is, throughout growth the surface layer is continuously converted into the underlying dense sp3-bonded material. The same process then removes the surface layer at the beginning of the second stage of a two-stage deposition, unless sufficientoxygen is present to prevent the formation of the dense Sp3 material. Both TEM images and the spatially resolved EELS data give a 1.3 ± 0.3 nm thickness for the spz surface layer. This thickness is consistent with the subplantation
model since a penetration depth of 1.3 nm is not unreasonable for 300 eV carbon ions, particularly if the surface is low density amorphous carbon. The compressive stress model has also predicted an spz layer approximately 1 nm thick, so that the two models are both consistent with the present data. An investigation of the dependence of the surface layer thickness on the energy of the carbon ions used to deposit the film will be carried out in the near future and may shed further light on the growth mechanism of ta-C films. Despite demonstrations of n-type doping [13] and photoconductivity [14], useful electronic devices based on ta-C films have yet to be achieved. For example, thin film transistor (TFT) devices based on ta-C have been fabricated but show negligible gain. This has been attributed to the presence of a highly defective layer at the ta-C-SiO z interface and is consistent with the 1.3 nm layer of spz amorphous carbon reported in this paper. Failure to observe transistor action in devices in which the ta-C film is grown 'on top of' the SiOz gate insulator is consistent with the wide spz layer observed at the SiOz-ta-C interface. If the spz surface layer is intrinsic to the growth process, as argued above, then subsequent etching of the ta-C film should remove this layer without it reforming. This may enable successful fabrication of ta-C TFT devices. The use of ta-C films as an electron emitting layer in flat screen displays is an exciting new area of research. The field emission properties of ta-C are likely to be strongly effected by the properties of the spz-bonded surface layer. Control of the surface layer may therefore be an important strategy for reducing the electron affinity of ta-C films.
5. Conclusions
The bonding and composition of a ta-C film has been determined with nanometre spatial resolution through the thickness of the film. It has been shown that for 300 eV carbon ions there is significant mixing of the carbon with silicon and oxygen from the native oxide on the surface of the silicon substrate. An interfacial layer approximately 5 nm thick is observed. The prevalence of Sp2 bonding in this layer is argued to be due to the presence of oxygen from the substrate. In contrast, the observed layer of spz carbon at the surface is 1.3 ± 0.3 nm thick and no contaminants were detected. Further work is required to determine the formation processes of the interfacial and surface layers in ta-C films unambiguously. The results presented here, however, suggest several strategies which may lead to ta-C films with improved properties for electronic applications.
CA. Daviset al.jSurfaceand Coatings Technology 76-77 (1995) 316-321
Acknowledgement The authors gratefully acknowledge useful discussions with Dr. CB, Boothroyd and would also like to thank the EPSRC for funding this work.
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[5] CA. Davis, V.S. Veerasamy, GAl Amaratunga, W.I. Milne and D.R McKenzie, Phi/os. Mag. B, 69 (1994) 1121. [6] lM. Cowley, Acta Crystallogr. A, 25 (1969) 129. [7J H. Kohl and H. Rose, Adv. Electronics Electron Phys., 65 (1985) 173. [8J R.F. Egerton, Electron Energy Loss Spectroscopy in the Electron Microscope, Plenum, London, 1986, pp. 262-278. [9] S.D. Berger, D.R. McKenzie and PJ. Martin, Philos. Mag. Lett; 57 (1988) 285. [10] W.M. Skiff, RW. Carpenter and S.H. Lin, J. Appl. Phys., 62 (1987) 2439. [11] Y. Lifshitz, S.R Kasi, lW. Rabalais and W. Eckstein, Phys. Rev. B, 41 (1990) 10468. [12] C.A. Davis, D.R McKenzie, Y. Yin, E. Kravtchinskaia, GAJ. Amaratunga and V.S. Veerasamy, Phi/os. Mag. B, 69 (1994) 1133. [13] V.S. Veerasamy, GAl Amaratunga, C.A. Davis, A.E. Timbs, W.L Milne and D.R. McKenzie, J. Phys.: Condens. Matter, 5 (1993) Ll69. [14] GAJ. Amaratunga, V.S. Veerasamy, W.I. Milne, CA. Davis, S.RP. Silva and H.S. McKenzie, Appl. Phvs. Lett., 63 (1993) 370.