APPLIED PHYSICS LETTERS 96, 162111 共2010兲
Depth-dependence of electrical conductivity of diamondlike carbon films A. Sikora,1,2 P. Paolino,3 H. Ftouni,1 C. Guerret-Piécourt,3 J.-L. Garden,1 A.-S. Loir,2 F. Garrelie,2 C. Donnet,2 and O. Bourgeois1,a兲 1
Institut NÉEL, CNRS-UJF, 25 avenue des Martyrs, 38042 Grenoble, France Laboratoire Hubert Curien, Université Jean Monnet, 42000 Saint Etienne, France 3 Ecole Centrale de Lyon, 36 avenue Guy de Collongue, 69134 Ecully, France 2
共Received 18 March 2010; accepted 30 March 2010; published online 23 April 2010兲 The electrical behavior of diamondlike carbon 共DLC兲 has been measured as a function of depth. The amorphous carbon 共a-C兲 films are deposited by pulsed laser deposition using two complementary setups: a femtosecond 共fs兲 and a nanosecond 共ns兲 pulse lasers. It is demonstrated through four probe resistance measurements and contact resistance mapping that the fs DLC are electrically heterogeneous in thickness. The presence of a thick sp2 rich layer on top is evidenced for fs a-C and is apparently away in the sp3 rich ns a-C. It is attributed to different subplantation processes between ns and fs a-C films. © 2010 American Institute of Physics. 关doi:10.1063/1.3407671兴 Electrical properties of the surface of diamondlike carbon 共DLC兲 films are of great importance for a large number of practical applications due to their unique properties:1 tribological, thermal, electrical and for their biocompatibility.2–6 The quality of the contact, for instance in the case of electrical switch, or for tribological purposes, depends only on the electrical conductivity of the top layers of the DLC films.2 Hence, the clear understanding of the electrical transport properties of the surface of DLC films is of great value for practical applications, especially in tomorrow’s microelectromechanical and nanoelectromechanical system technology.6 These resistive properties are strongly influenced by the different possible hybridizations of the carbon atoms. It has been reported before that, for DLC films deposited either by cathodic arc, mass selected ion beam deposition or by pulsed laser deposition 共PLD兲, a sp2 rich layer can form on the surface due to subplantation process.1 In this process, energetic ions coming from the target penetrate the surface. A small part of their kinetic energy is used to cross the surface, as the rest is dissipated into atom displacement and heat release. A fraction of the carbon atoms are relaxing toward the surface creating a sp2 layer and another fraction penetrates deeply in the films making denser the bottom layer which is generally rich in sp3 bonded atoms. In this letter we use the measurement of the resistive properties of hydrogen free DLC films as a unique tool to investigate the electrical properties of amorphous carbon films as function of the thickness. We show that the electrical properties of the femtosecond 共fs兲 laser-produced films are strongly depth-dependent. On the other hand, the nanosecond 共ns兲 laser-produced films do not exhibit such heterogeneity of resistance in thickness. This specific feature of fs laser produced film is attributed to the heterogeneous repartition in depth of sp2 / sp3 carbon bound of fs DLC films. On the other hand, the dominant sp3 bonded carbon content is found to be very homogeneous in thickness for the ns laser produced DLC film. A series of a-C films have been deposited onto sapphire substrates by ablating graphite 共purity 99.997%兲 targets in a a兲
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deposition chamber at a base pressure of 10−4 Pa.7 Two different lasers were used in this experiment. First, the fs laser 共 = 800 nm, pulse duration 150 fs, repetition rate 1 kHz, energy per pulse 1.2 mJ兲 has been used with an incident angle of 45° and an energy density set to 5 J / cm2 共⫾10%兲. Second, a ns KrF excimer laser 共 = 248 nm, pulse duration 17 ns, repetition rate 10 Hz, energy per pulse around 200 mJ兲 has been used with an incident angle of 45° and an energy density set to 60 J / cm2. The films will be referred to as fs a-C for the femtosecond amorphous carbon film and as ns a-C for the nanosecond one. A standard four probe technique has been used to measure the electrical resistance of the DLC films using four platinum electrical leads from 150 to 300 K. This geometry, as presented in the inset of the Fig. 1, has the major advantage of freeing the measure from contact resistances. The a-C films are deposited and structured using photolithographic process onto the platinum leads. The voltage lead distance is about 0.5 mm long and 2 mm wide. The typical thickness of the thin films is ranging between 200 and 300 nm. The surfaces of the DLC thin films have been also studied by conductive atomic-force-microscope 共C-AFM兲. This C-AFM was built by adding an extension to a commercial AFM
FIG. 1. Resistance vs temperature of a fs a-C before and after RIE etching of 50 nm on top of the film. The inset is giving the geometry of the sample, the contact to the electrical leads are underneath the DLC films.
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FIG. 2. Electrical resistivity plotted vs temperature for the fs and ns a-C films before and after RIE.
VEECO CPII able to measure contact resistance over nine decades 共from 102 to 1011 ⍀兲.8 The resistances are measured after a deep etching with an oxygen plasma of the top layers of the DLC film to identify the contribution of the surface to the electron transport. The etching is performed by reactive ion etching 共RIE兲 with an etching rate of 40 nm per minute. In the Fig. 1, the resistance of the fs a-C films before and after the etching of the top 50 nm is presented. The resistance is changing by a factor of 5.7, exceeding largely the expected geometrical factor of 1.3 due to the reduction in the thickness from 270 to 220 nm. The resistances changed from 1.47⫻ 106 to 8.42 ⫻ 106 ⍀. It can be concluded from this measurement that the top layers are much more conductive than the rest of the films, a signature of the electrical heterogeneity in depth of the fs a-C. The same procedure has been applied to DLC films deposited with the ns laser. A thickness of 65 nm has been etched by RIE out of 200 nm. In this case, the resistance change corresponds exactly to the change in the thickness of the film, an increase in resistance by a factor of 1.5. Indeed, the resistance change from 78.6⫻ 106 to 125.7⫻ 106 ⍀. First, it can be noted that the ns a-C film is much more resistive than the fs a-C film. Second, the change in the resistance is only due to a change in the geometry of the films, illustrating the homogeneity of such kind of a-C. No particular behavior of the surface can be evidenced. It is noteworthy that the change in resistance is identical whatever the temperature which gives good credit to the quality of the measurement. The temperature dependence of the resistance of the DLC, as already demonstrated, follows a regular variable range hopping behavior.7 To illustrate these changes and get rid of the geometrical parameter which are of no interest, the resistivity of the fs and ns a-C films are plotted in the Fig. 2. The resistivities, before and after RIE, are illustrated for the two different materials; a clear difference between fs and ns a-C is evidenced. As the fs DLC resistivity is changing a lot when removing the top layers, in the case of the ns DLC no change can be observed. This indicates that the subplantation growth mechanism1 if applied, will lead to a thinner sp2-rich layer in the ns a-C film. Indeed, the sp2-rich layer has been estimated by ellipsometric measurements around 1 nm for ns a-C films, and between 5 and 10 nm for fs a-C films.9 The induced variation in electrical transport properties can thus be disre-
Appl. Phys. Lett. 96, 162111 共2010兲
garded in the case of the ns a-C film whereas we show its great influence in the case of fs a-C film. The presence of such a sp2 rich layer reported by Robertson1 has been observed by energy filtering transmission electron microscopy of arc ion plating or chemical vapor deposition produced films or pointed out by optical measurement of PLD films,9 but here this property is demonstrated through a direct electrical measurement. Because this sp2 rich layer behaves like a short circuit for the circulation of the electrons, its effect on the resistance measurement is huge and can be easily evidenced. To confirm that the electrical properties of ns and fs a-C films originated from the different contents of sp2 / sp3 carbon bounds, Raman spectroscopy measurements at various wavelengths have been done on the same DLC films. The Raman spectra of the ns a-C and fs a-C have been acquired from an excitation wavelength of 325 nm, allowing to identify the T band characteristics of the carbon sp3 hybridization. Contrary to the ns a-C film, the fs a-C film exhibits no T band but it exhibits a D band characteristic of the stretching mode of aromatic clusters. From the fit of the spectra obtained at various wavelengths 共see Ref. 10兲, the ns a-C film contains a higher carbon sp3 content 共near 70%兲, compared to the fs a-C film 共near 40%兲. One of the origin of the different content of sp2 bound in fs and ns a-C films comes from the different kinetic energy of plasma species9 and the presence of stronger relaxation processes in fs laser ablation due to high kinetic energy ions.11 These results are entirely consistent with the conduction properties measured on both films. The conductive layers of the fs a-C can thus be ascribed to the presence of a sp2 rich layer on top. These layers can be removed by plasma etching if one wants to get rid of that electrical short circuit. One question emerges however from the use of RIE about the impact of the oxygen plasma. To show that the electrical behavior demonstrated for the fs a-C cannot be due to a reconstruction in depth of the DLC films induced by the collision of oxygen atoms on the carbon surface, we have mapped the contact resistance on the surface of the fs a-C films using the C-AFM technique before and after the RIE. The ns a-C films are too resistive 共R ⬎ 1012 ⍀兲 to be measured by C-AFM. In the Fig. 3, we show in parallel the topology and the resistance map of the fs a-C films before the RIE 共a兲 and after the RIE 共b兲. The repartition of resistance versus the surface is exploited in the graph 共c兲 where the percentage of pixel having a defined resistance 共extracted from the value of each pixel of the image兲 is given versus the logarithm of the resistance. The reading of this graph indicates clearly that before RIE the maximum resistance is located near 109 ⍀, while after the RIE the maximum is located near 1011 ⍀. These AFM measurements, sensitive to what happens on the surface, go in the same direction than the measurement of the resistance more sensitive to the volume properties. It confirms that the deep layers of a fs a-C films are much more resistive, a property certainly linked to the presence of more sp3 bounding at the bottom of the films. The possible presence on the surface of C‑O bonds after RIE, which may perturb the measurement done in this study, has been ruled out by x-ray photoelectron spectroscopy where no difference in the amount of oxygen has been detected between DLC films before and after RIE.
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bonding of the carbon atoms. One can notice here that these differences are not only due to the differences in wavelength of the laser used. Previous works related to graphite ablation using fs to ns lasers working both at 248 nm show that ns laser is more favorable to a rich sp3 coating.12 In summary, it was demonstrated that resistivity measurement as a function of thickness is an original tool for exploring electronic transport properties of the surface of DLC films. The presence of a sp2-rich layer on top of fs a-C films and its influence on electrical properties are clearly evidenced. On the contrary, ns a-C films do not show such thick sp2 rich top layer in good accordance with Raman spectroscopy data and C-AFM measurements. The influence of this sp2-rich layer has to be taken into account when planning to use DLC for applications involving the electrical properties of the surface of a-C films. FIG. 3. 共Color online兲 Topological and electrical characterizations of the surface of the fs a-C films before RIE 共a兲 and after RIE 共b兲 by scans of 2 m, the color scale bar indicates the value of the resistance measured by the C-AFM. In 共c兲, we give the repartition of resistance in percentage as a function of log共r兲, the blue curve corresponds to 共a兲 and the red one to 共b兲.
The resistance behavior of the fs a-C films can be described by a simple two channel model:
sp2rich共T兲 =
e1bulktot bulk共e1 + e2兲 − e2tot
where sp2rich is the resistivity of the top layer, e1 and e2 are, respectively, the thicknesses of the sp2-rich layer and the one of the bulk, bulk is the measured resistivity after RIE and tot the resistivity before the RIE. From this equation we can extract the resistivity of the etched layer: sp2rich = 32.8 ⍀ cm, which is 24 times smaller than the resistivity of the bottom layer at 300 K 共bulk = 800 ⍀ cm by itself ns =2 smaller than the resistivity of the ns a-C: a−C 3 ⫻ 10 ⍀ cm兲. This is in perfect agreement with what has been measured with the C-AFM. We ascribe these significant contrasted results to the different proportion of sp2 or sp3
This work was financially supported by Région RhôneAlpes and Agence Nationale de la Recherche through the Sensocarb project. J. Robertson, Mater. Sci. Eng. 37, 129 共2002兲. A. Erdemir and C. Donnet, J. Phys. D 39, R311 共2006兲. 3 A. A. Balandin, M. Shamsa, W. L. Liu, C. Casiraghi, W. I. Milne, and A. C. Ferrari, Appl. Phys. Lett. 93, 043115 共2008兲. 4 S. Egret, J. Robertson, W. I. Milne, and F. J. Clough, Diamond Relat. Mater. 6, 879 共1997兲. 5 J. C. Orlianges, A. Pothier, D. Mercier, P. Blondy, C. Chapeaux, A. Catherinot, M. I. de Barros, and S. Pavant, Thin Solid Films 482, 237 共2005兲. 6 J. K. Luo, Y. Q. Fu, H. R. Le, J. A. Williams, S. M. Spearing, and W. I. Milne, J. Micromech. Microeng. 17, S147 共2007兲. 7 A. Sikora, A. Berkesse, O. Bourgeois, J.-L. Garden, C. Guerret-Piécourt, A.-S. Loir, F. Garrelie, and C. Donnet, Appl. Phys. A: Mater. Sci. Process. 94, 105 共2009兲. 8 M. Gadenne, O. Schneegans, F. Houze, P. Chretien, C. Desmarest, J. Sztern, and P. Gadenne, Physica B 279, 94 共2000兲. 9 T. Katsuno, C. Godet, J. C. Orlianges, A. S. Loir, F. Garrelie, and A. Catherinot, Appl. Phys. A: Mater. Sci. Process. 81, 471 共2005兲. 10 A. C. Ferrari and J. Robertson, Phys. Rev. B 64, 075414 共2001兲. 11 F. Qian, V. Craciun, R. K. Singh, S. D. Dutta, and P. P. Pronko, J. Appl. Phys. 86, 2281 共1999兲. 12 F. Claeyssens, M. R. Ashfold, E. Sofoulakis, C. G. Ritoscu, D. Anglos, and C. Fotakis, J. Appl. Phys. 91, 6162 共2002兲. 1 2
