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Hardness, intrinsic stress, and structure of the a-C and a-C:H films prepared by magnetron sputtering
Diamond and Related Materials 10 Ž2001. 1076᎐1081
Hardness, intrinsic stress, and structure of the a-C and a-C:H films prepared by magnetron sputtering V. Kulikovsky a , P. Bohac b,U , F. Franc b, A. Deinekab, V. Vorlicek b, L. Jastrabik b a
b
Institute for Problems of Materials Science, Academy of Sciences of Ukraine, 3 Krzhyzhano¨ sky St., 252142 Kie¨ , Ukraine Institute of Physics, Academy of Sciences of the Czech Republic, Na Slo¨ ance 2, P.O. Box 61, CZ-18221 Prague 8, Czech Republic
Abstract The 1.5-m thick carbon films prepared by magnetron sputtering of a carbon target in Ar, Ar qCH 4 and Ar q O 2 gas mixtures show that the higher their intrinsic stress, the higher their microhardness and the lower their resistivity. The a-C films obtained without ion bombardment Žunbiased; the mean free path of sputtered C atoms larger or equal to the cathode᎐anode distance. show microhardness up to 25 GPa. The biased a-C films Ži.e. with ion bombardment. achieve the highest microhardness Žup to 50 GPa. and the lowest resistivity Ž0.01 ⍀ cm.. An increase in Ar pressure, or the optional addition of O 2 , results in a decrease in the microhardness and intrinsic stress and an increase in the film resistivity. In comparison to a-C films, by adding CH 4 to Ar up to certain limit, the microhardness and intrinsic stress of these a-C:H films increase and subsequently decrease steeply. It was specified by analysis of the electron diffraction patterns of thin films Ž30᎐60 nm. deposited under the same conditions that the radius of the first co-ordination sphere of C atoms for all the films is in a good agreement with the value for graphite. The prime interplanar distance for biased a-C films is considerably lower than that for unbiased ones and for graphite. Our data indicate the sp 2-bonded carbon structure of the deposited hard carbon films, in which the prime interplanar distance is reduced due to intrinsic stress. Thus, it is more suitable to explain the hardness origin as a consequence of the film nanostructure rather than the presence of sp 3 bonds . 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Diamond-like carbon; Stress; Hardness; Structure
1. Introduction The high hardness of amorphous carbon films Ža-C. is usually linked to the presence of a high percentage of sp 3 bonds. Such films are called diamond-like carbon ŽDLC.. Recently, it was shown that some hard films contain a high number of sp 2 bonds w1᎐5x. The structural origin of the very high hardness of a-C films deposited at ion bombardment by argon or any other noble gas is still a subject of discussion and investigation.
U
Corresponding author. Tel.: q420-2-6605-2959; fax: q420-28581448. E-mail address: [email protected] ŽP. Bohac..
In this paper, we attempt to find the connection between the hardness, internal stress, resistivity and structure of a-C films obtained by magnetron sputtering of carbon by varying the Ar pressure, the substrate bias and gas mixtures of Arq CH 4 or Ar q O 2 . To explain the origin of the hardness, the film nanostructure rather than existence of sp 3 bonds is taken into account. Direct experimental observation of the reduction of the prime interplanar distance for hard, compressed, sp 2-bonded films is also presented for the first time.
2. Experimental details All the films were prepared on a commercial Leybold
0925-9635r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 5 - 9 6 3 5 Ž 0 0 . 0 0 5 2 5 - 2
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Z 550M vacuum plant equipped with three planar magnetrons Žtype PK 150., which are situated 50 mm above the substrate carrier in the vacuum chamber w6x. DC sputtering of a carbon target Žof 99.5% purity. was performed in Ar Žpurity 99.999%., or in a mixture of argon and methane Žpurity 99.995%. to obtain a-C or a-C:H films, respectively. The argon flow rate was held at 13 sccm, while the CH 4 flow rate was varied from 0 to 24 sccm. The total working gas pressure was approximately 0.17᎐2.8 Pa for a-C and 0.17᎐0.5 Pa for a-C:H deposition. The discharge power was varied from 300 to 960 W, but was mainly at 960 W. Bias up to y150 V was induced due to a 13.56 MHz power source applied to the substrate carrier. Substrates were not specially heated during deposition process. Thick films on SiŽ111. substrates Žsize 20 = 8 = 0.5 mm. were used for measurement of the microhardness and determination of the intrinsic stress. Thin films Ž30᎐60 nm. were prepared on KCl substrates for investigation by transmission electron diffraction ŽTED.. The film thickness measured by Tencor ALPHA STEP 500 was approximately 1.5 m for the majority of films. The film structure was investigated by Raman spectroscopy, and also by TED. Electrical resistivity of the films on glass substrates was determined using standard four-point probe measurements. The microhardness measurements were carried out on a FISCHERSCOPE H100 apparatus at a growing load of up to 10 mN. Radii of curvature of the substrate before and after coating deposition were measured by observation of Newton’s rings on an optical interferometry system ŽZYGO Mark IV. and the intrinsic stresses were calcu-
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lated using Stoney’s classical equation w7x. All radius measurements and presented tendencies were reproducible.
3. Results 3.1. Microhardness, stress and resisti¨ ity Fig. 1 shows a strong dependency of the microhardness on intrinsic stress for all the thick Žapprox. 1.5 m. a-C and a-C:H films obtained at different values of discharge power Ž300᎐900 W., Ar or Arq CH 4 pressure Ž0.17᎐3.3 Pa. and negative substrate bias Žfrom 0 up to y100 V.. Similar dependencies were reported in w8,9x. Both hardness and intrinsic stress are closely related to the film microstructure. The increase in Ar pressure from 0.17 to 2.8 Pa leads to a decrease in the microhardness and intrinsic stress and an increase in resistivity ŽFig. 2.. Such behavior of the film properties could be explained by an increasing amount of absorbed gas on the boundaries of clusters, which leads to their imperfect interconnection. Note that at low Ar pressure, when the mean free path of sputtered C atoms is equal to or larger than the cathode᎐substrate distance L, the energy, not only of the Ar ions, but of every carbon atom arriving at the substrate, is of several eV. This is already sufficient to remove Žre-sputter. the gas impurities weakly bound to the growing film. In such a case, the value of the microhardness can be approximately 25 GPa without any ion bombardment ŽFig. 2.. A decrease in rL
Fig. 1. The microhardness of ; 1.5-m thick a-C and a-C:H films, sputtered under different conditions, depending on the intrinsic stress.
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V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Fig. 2. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films on the Ar pressure at discharge power of 960 W; L, cathode᎐substrate distance; and , mean free path of carbon atoms.
Fig. 3. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films on the O 2 partial pressure at discharge power of 960 W; pAr s 0.17 Pa.
leads to a reduction in the average energy of condensing C atoms and as a result, to a decrease in bonds between adjacent clusters. Similar, but more pronounced changes in microhardness, intrinsic stress and resistivity are observed when even a small amount of oxygen is added to the argon ŽFig. 3.. Raman spectra of the films deposited at pAr s 0.17 and 2.8 Pa and at different oxygen additions to Ar Ž pAr s 0.17 Pa, pO s 0.02 and 0.11 Pa. are shown in 2
Fig. 4. The increase in Ar pressure, and especially in the partial O 2 pressure, leads to an increase in the intensity of both the D- Ž; 1350 cmy1 . and G-band Ž; 1580 cmy1 .. It is particularly visible for the D-band. An increase in the ratio of intensity I D rIG is usually observed at a rise in substrate temperature w10x, and is due to an increase in the size andror number of the sp 2 clusters, according to w11x. Together with the increase in resistivity and decrease in intrinsic stress, the Raman spectra support the aforementioned idea of the
Fig. 4. Raman spectra of a-C films sputtered at discharge power of 960 W.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Fig. 5. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films Ž; 1.5 m. on the negative substrate bias at discharge power of 960 W; pAr s 0.17 Pa.
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against previous cases can mainly be explained by the re-sputtering of weakly bound gas impurities on the boundaries of clusters. This leads to enhancement of the contacts between adjacent clusters and a decrease in the electron scattering on numerous cluster boundaries. In addition, it leads to an increase in the intrinsic stress due to formation of the dense, incoherent boundaries of clusters. An increase in the film hardness can be caused by this growth process and also by a displacement of C atoms by Ar ions into subsurface positions. Fig. 6 shows the dependence of the intrinsic stress and microhardness of a-C:H films prepared without Žand partially with. ion bombardment on CH 4rAr flow ratio. A small amount of CH 4 added to Ar leads to an increase in the intrinsic stress and microhardness of a-C:H films in comparison to a-C ones. This can be caused by a decrease in the short-range order of the film, due to the formation of additional sp 3 bonds and a decrease in cluster size in the amorphous matrix. Further addition of CH 4 in Ar leads to the growth of numerous clusters with hydrogen terminating the carbon bonds at their boundaries, thereby decreasing the number of eventual cross-links in the carbon network. This results in the steep decrease in the film hardness and intrinsic stress. 3.2. Structure
clusters disconnecting due to absorbed gas impurities, which can lead to an increase in the number of sp 2 clusters. A microhardness of approximately 50 GPa was obtained for the 1.5-m thick film deposited at a discharge power of 960 W and substrate bias of y100 V ŽFig. 5.. The same result was achieved by decreasing these values to 300 W and y60 V, respectively. This increase in microhardness and intrinsic stress and simultaneous decrease in resistivity to a very low value
The electron diffraction patterns of all the investigated films are very similar to each other and consist of three visible haloes. The first small-angle halo appears only if the film contains fragments of a layer structure. It corresponds to the diffraction from prime interplanes wanalogous to the diffraction from Ž002. graphite planesx. The prime interplanar distance d and radius r of the first co-ordination sphere for C᎐C bonds de-
Fig. 6. Dependence of the microhardness and intrinsic stress of a-C:H films on the CH 4 rAr flow ratio at discharge power of 960 W; pAr s 0.17 Pa.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
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Table 1 Dependence of the prime interplanar distance Ž d . and the radius Ž r1 . of the first co-ordination sphere for a-C and a-C:H films on the depositions parameters Ar flow rate wsccmx
CH4 flow rate wsccmx
Pressure wPax
Substrate bias wVx
d wnmx
r1 " 0.001 wnmx
13 13 13 13 13 13
᎐ ᎐ ᎐ 1 12 24
0.17 0.17 0.17 0.19 0.32 0.44
y150 y100 Floating Ž; 5. Floating Ž; 5. Floating Ž; 5. Floating Ž; 5.
0.298᎐0.305 0.298᎐0.301 0.326᎐0.337 Not determined 0.389᎐0.405 0.405᎐0.410
0.142 0.142 0.142 0.143 0.143 0.143
termined are given in Table 1. The radius r was calculated by Ehrenfest’s equation w12x: 2 r sin s 1.23 ew in which 2 is scattering angle of the maximum of corresponding Žthird. visible halo and ew is the electron wavelengths. This procedure for all investigated a-C and a-C:H films gives a radius value r s 0.142᎐0.143 nm, corresponding closely to that of graphite. This means that carbon atoms in all our films, including the hardest a-C ones, are bonded to each other mainly through sp 2 bonding. The interplanar distance d can be determined from the Bragg’s formula: 2 dsin s n ew For a-C films obtained without ion bombardment, the interplanar distance d was found to be 0.326᎐0.337 nm Ž0.335 nm is the interlayer spacing in graphite.. Bias Žy100, y150 V. during deposition shifted these values to 0.298᎐0.305 nm ᎏ see Table 1. This con-
siderable decrease in d can be explained by the highly stressed sp 2 structure of the films. Such a structure is known for bulk graphite under high pressure, e.g. w13x. The main source of pressure in our films was intrinsic stress induced by ion bombardment. An intrinsic stress of approximately 10 GPa was obtained for a-C film with a thickness of approximately 50 nm deposited onto SiŽ111. at a substrate bias of y100 V ŽFig. 7.. The CH 4 addition to Ar led to a considerable increase in distance d ŽTable 1.. This agrees with the data for nanostructured graphite prepared by mechanical milling under a hydrogen atmosphere w14x. The main peak corresponding to this interplanar diffraction strongly decreases and shifts to lower angles, indicating an expansion of the graphite interlayer with increasing hydrogen in the samples. Part of the hydrogen probably occupied sites between the planes, as supported by neutron diffraction measurements w14x.
4. Discussion In the case of amorphous carbon film deposited by
Fig. 7. Dependence of the intrinsic stress of a-C films on the film thickness at discharge power of 960 W; pAr s 0.17 Pa; substrate bias y100 V.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
magnetron sputtering, it would be possible to reduce or almost avoid the influence of gas impurities on the boundaries between the growing grains or clusters. There are two ways to achieve this: Ž1. by the preferential re-sputtering of gas adsorbed, mainly at the joints of growing film islands, through enhanced ion bombardment; Ž2. through a decrease in Ar pressure during sputtering to the value at which the mean free path of carbon atoms is larger than the cathode᎐substrate distance. In the latter case, every depositing carbon atom has an energy of approximately several eV. This is sufficient for both sputtering the weakly bound gas impurity, and occupying a stable position without significant heating of the surrounding atoms. The structure formed in this way can be very dense with strong interactions between atoms at grain boundaries, even if they occupy incoherent positions. This interaction can be strongly enhanced due to the biaxial compressive stress which accompanies such a process of film growth. The possible sliding along prime planes within clusters is difficult, due to the high disordering within clusters and high compressive stress, which, in turn, can bind these planes, see for example w15x. In any case, if such sliding is possible, it is realized partially during film growth under compressive stress. That is why these films demonstrate high elasticity Ž89᎐94%.. If the substrate temperature is low during film deposition and the recrystallization process is not developed Žfor carbon film, the recrystallization temperature is over 900 K., a nanostructure having a high hardness can be obtained. The experimental results presented support the aforementioned concept, according to the references cited.
5. Conclusions The superhardness of amorphous carbon films is usually explained by the presence of sp 3 bonds in the film. The higher the sp 3rsp 2 ratio is, the higher the film hardness. However, our results Ždata on electron diffraction and the low level of film resistivity. and those reported in w1᎐4x indicate that very hard carbon films with sp 2 C᎐C bonds can be obtained. It has been shown on our 1.5-m thick carbon films Ža-C and a-C:H. that with increasing negative substrate bias andror decreasing Ar, O 2 or CH 4 pressure, the resistivity decreases with a simultaneous increase in microhardness and compressive stress. The a-C films condensed under ion bombardment showed a microhardness of 50 GPa and a very low level of resistivity. Analysis of the electron diffraction patterns of thin films Ž30᎐60 nm. prepared under the same conditions
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showed that the radius of the first co-ordination sphere of C atoms for all the films Ža-C and a-C:H. was 0.142᎐0.143 nm, which is in good agreement with the value for sp 2 bonded carbon. The interplanar distance d shifted from approximately 0.326᎐0.337 nm Žfor unbiased a-C films. to 0.298᎐0.305 Žfor those deposited at a substrate bias of y150 V.. This can be caused by strong compression due to the intrinsic stress. All the data indicate an sp 2 bonding configuration of the hard carbon films deposited. The sp 2 hard carbon films presumably consist of cross-linked randomly oriented dense clusters, bonded to each other through incoherent boundaries on which high stresses are accumulated. The fewer the gas impurities on these boundaries are, the stronger is the bonding between the clusters. In this way, a continuous rigid network is formed. The high compressive stress can result in strong bonding between the cluster boundaries, on which a small concentration of sp 3 sites may be occurring.
Acknowledgements This work has been supported by the Concerned European Action on Tribology COST 516 under contract no. 516.50 and by grant GA CR 202r00r1592. References w1x G.A.J. Amaratunga, M. Chhowalla, C.J. Kiely, I. Alexandrou, R. Aharonov, R.M. Devenish, Nature 383 Ž1996. 321᎐323. w2x D. Camino, A.H.S. Jones, D. Mercs, D.G. Teer, Vacuum V52 Ž1999. 125. w3x R.G. Lacerda, P. Hammer, F. Alvarez, F.C. Marques, F.L. Freire, Jr, Diamond Relat. Mater. 9 Ž2000. 796. w4x R.G. Lacerda, F.C. Marques, Appl. Phys. Lett. 73 Ž5. Ž1998. 617. w5x I. Alexandrou, H.-J. Scheibe, C.J. Kiely, A.J. Papworth, G.A.J. Amaratunga, B. Schultrich, Phys. Rev. B 60 Ž1999. 10903. w6x V.Y. Kulikovsky, F. Fendrych, L. Jastrabik, D. Chvostova, Surf. Coat. Technol. 91 Ž1997. 122. w7x G.G. Stoney, Proc. R. Soc. Lond. Ser. A 82 Ž1909. 172. w8x J.W. Zou, K. Reichelt, K. Schmidt, B. Dischler, J. Appl. Phys. 65 Ž1989. 3914. w9x D.R. McKenzie, Y. Yin, N.A. Marks et al., Diamond Relat. Mater. 3 Ž1994. 353. w10x R.S. Rubino, E.S. Takcuchi, J. Power Sources 81r82 Ž1999. 373. w11x R.O. Dillon, J.A. Woollam, V. Katkanant, Phys.Rev. B 29 Ž1984. 3482. w12x P. Ehrenfest, Prog. Amst. Acad. 17 Ž1915. 1132. w13x R.W. Lynch, H.G. Drickamer, J. Chem. Phys. 44 Ž1966. 181. w14x S. Orimo, G. Majer, T. Fukunaga, A. Zuttel, L. Schlapbach, H. Fujii, Appl. Phys. Lett. 75 Ž1999. 3093. w15x J. Schwan, S. Ulrich, T. Theel et al., J. Appl. Phys. 82 Ž1997. 6025.
Hardness, intrinsic stress, and structure of the a-C and a-C:H films prepared by magnetron sputtering V. Kulikovsky a , P. Bohac b,U , F. Franc b, A. Deinekab, V. Vorlicek b, L. Jastrabik b a
b
Institute for Problems of Materials Science, Academy of Sciences of Ukraine, 3 Krzhyzhano¨ sky St., 252142 Kie¨ , Ukraine Institute of Physics, Academy of Sciences of the Czech Republic, Na Slo¨ ance 2, P.O. Box 61, CZ-18221 Prague 8, Czech Republic
Abstract The 1.5-m thick carbon films prepared by magnetron sputtering of a carbon target in Ar, Ar qCH 4 and Ar q O 2 gas mixtures show that the higher their intrinsic stress, the higher their microhardness and the lower their resistivity. The a-C films obtained without ion bombardment Žunbiased; the mean free path of sputtered C atoms larger or equal to the cathode᎐anode distance. show microhardness up to 25 GPa. The biased a-C films Ži.e. with ion bombardment. achieve the highest microhardness Žup to 50 GPa. and the lowest resistivity Ž0.01 ⍀ cm.. An increase in Ar pressure, or the optional addition of O 2 , results in a decrease in the microhardness and intrinsic stress and an increase in the film resistivity. In comparison to a-C films, by adding CH 4 to Ar up to certain limit, the microhardness and intrinsic stress of these a-C:H films increase and subsequently decrease steeply. It was specified by analysis of the electron diffraction patterns of thin films Ž30᎐60 nm. deposited under the same conditions that the radius of the first co-ordination sphere of C atoms for all the films is in a good agreement with the value for graphite. The prime interplanar distance for biased a-C films is considerably lower than that for unbiased ones and for graphite. Our data indicate the sp 2-bonded carbon structure of the deposited hard carbon films, in which the prime interplanar distance is reduced due to intrinsic stress. Thus, it is more suitable to explain the hardness origin as a consequence of the film nanostructure rather than the presence of sp 3 bonds . 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Diamond-like carbon; Stress; Hardness; Structure
1. Introduction The high hardness of amorphous carbon films Ža-C. is usually linked to the presence of a high percentage of sp 3 bonds. Such films are called diamond-like carbon ŽDLC.. Recently, it was shown that some hard films contain a high number of sp 2 bonds w1᎐5x. The structural origin of the very high hardness of a-C films deposited at ion bombardment by argon or any other noble gas is still a subject of discussion and investigation.
U
Corresponding author. Tel.: q420-2-6605-2959; fax: q420-28581448. E-mail address: [email protected] ŽP. Bohac..
In this paper, we attempt to find the connection between the hardness, internal stress, resistivity and structure of a-C films obtained by magnetron sputtering of carbon by varying the Ar pressure, the substrate bias and gas mixtures of Arq CH 4 or Ar q O 2 . To explain the origin of the hardness, the film nanostructure rather than existence of sp 3 bonds is taken into account. Direct experimental observation of the reduction of the prime interplanar distance for hard, compressed, sp 2-bonded films is also presented for the first time.
2. Experimental details All the films were prepared on a commercial Leybold
0925-9635r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 5 - 9 6 3 5 Ž 0 0 . 0 0 5 2 5 - 2
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Z 550M vacuum plant equipped with three planar magnetrons Žtype PK 150., which are situated 50 mm above the substrate carrier in the vacuum chamber w6x. DC sputtering of a carbon target Žof 99.5% purity. was performed in Ar Žpurity 99.999%., or in a mixture of argon and methane Žpurity 99.995%. to obtain a-C or a-C:H films, respectively. The argon flow rate was held at 13 sccm, while the CH 4 flow rate was varied from 0 to 24 sccm. The total working gas pressure was approximately 0.17᎐2.8 Pa for a-C and 0.17᎐0.5 Pa for a-C:H deposition. The discharge power was varied from 300 to 960 W, but was mainly at 960 W. Bias up to y150 V was induced due to a 13.56 MHz power source applied to the substrate carrier. Substrates were not specially heated during deposition process. Thick films on SiŽ111. substrates Žsize 20 = 8 = 0.5 mm. were used for measurement of the microhardness and determination of the intrinsic stress. Thin films Ž30᎐60 nm. were prepared on KCl substrates for investigation by transmission electron diffraction ŽTED.. The film thickness measured by Tencor ALPHA STEP 500 was approximately 1.5 m for the majority of films. The film structure was investigated by Raman spectroscopy, and also by TED. Electrical resistivity of the films on glass substrates was determined using standard four-point probe measurements. The microhardness measurements were carried out on a FISCHERSCOPE H100 apparatus at a growing load of up to 10 mN. Radii of curvature of the substrate before and after coating deposition were measured by observation of Newton’s rings on an optical interferometry system ŽZYGO Mark IV. and the intrinsic stresses were calcu-
1077
lated using Stoney’s classical equation w7x. All radius measurements and presented tendencies were reproducible.
3. Results 3.1. Microhardness, stress and resisti¨ ity Fig. 1 shows a strong dependency of the microhardness on intrinsic stress for all the thick Žapprox. 1.5 m. a-C and a-C:H films obtained at different values of discharge power Ž300᎐900 W., Ar or Arq CH 4 pressure Ž0.17᎐3.3 Pa. and negative substrate bias Žfrom 0 up to y100 V.. Similar dependencies were reported in w8,9x. Both hardness and intrinsic stress are closely related to the film microstructure. The increase in Ar pressure from 0.17 to 2.8 Pa leads to a decrease in the microhardness and intrinsic stress and an increase in resistivity ŽFig. 2.. Such behavior of the film properties could be explained by an increasing amount of absorbed gas on the boundaries of clusters, which leads to their imperfect interconnection. Note that at low Ar pressure, when the mean free path of sputtered C atoms is equal to or larger than the cathode᎐substrate distance L, the energy, not only of the Ar ions, but of every carbon atom arriving at the substrate, is of several eV. This is already sufficient to remove Žre-sputter. the gas impurities weakly bound to the growing film. In such a case, the value of the microhardness can be approximately 25 GPa without any ion bombardment ŽFig. 2.. A decrease in rL
Fig. 1. The microhardness of ; 1.5-m thick a-C and a-C:H films, sputtered under different conditions, depending on the intrinsic stress.
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V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Fig. 2. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films on the Ar pressure at discharge power of 960 W; L, cathode᎐substrate distance; and , mean free path of carbon atoms.
Fig. 3. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films on the O 2 partial pressure at discharge power of 960 W; pAr s 0.17 Pa.
leads to a reduction in the average energy of condensing C atoms and as a result, to a decrease in bonds between adjacent clusters. Similar, but more pronounced changes in microhardness, intrinsic stress and resistivity are observed when even a small amount of oxygen is added to the argon ŽFig. 3.. Raman spectra of the films deposited at pAr s 0.17 and 2.8 Pa and at different oxygen additions to Ar Ž pAr s 0.17 Pa, pO s 0.02 and 0.11 Pa. are shown in 2
Fig. 4. The increase in Ar pressure, and especially in the partial O 2 pressure, leads to an increase in the intensity of both the D- Ž; 1350 cmy1 . and G-band Ž; 1580 cmy1 .. It is particularly visible for the D-band. An increase in the ratio of intensity I D rIG is usually observed at a rise in substrate temperature w10x, and is due to an increase in the size andror number of the sp 2 clusters, according to w11x. Together with the increase in resistivity and decrease in intrinsic stress, the Raman spectra support the aforementioned idea of the
Fig. 4. Raman spectra of a-C films sputtered at discharge power of 960 W.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
Fig. 5. Dependence of the microhardness, intrinsic stress, resistivity and deposition rate of a-C films Ž; 1.5 m. on the negative substrate bias at discharge power of 960 W; pAr s 0.17 Pa.
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against previous cases can mainly be explained by the re-sputtering of weakly bound gas impurities on the boundaries of clusters. This leads to enhancement of the contacts between adjacent clusters and a decrease in the electron scattering on numerous cluster boundaries. In addition, it leads to an increase in the intrinsic stress due to formation of the dense, incoherent boundaries of clusters. An increase in the film hardness can be caused by this growth process and also by a displacement of C atoms by Ar ions into subsurface positions. Fig. 6 shows the dependence of the intrinsic stress and microhardness of a-C:H films prepared without Žand partially with. ion bombardment on CH 4rAr flow ratio. A small amount of CH 4 added to Ar leads to an increase in the intrinsic stress and microhardness of a-C:H films in comparison to a-C ones. This can be caused by a decrease in the short-range order of the film, due to the formation of additional sp 3 bonds and a decrease in cluster size in the amorphous matrix. Further addition of CH 4 in Ar leads to the growth of numerous clusters with hydrogen terminating the carbon bonds at their boundaries, thereby decreasing the number of eventual cross-links in the carbon network. This results in the steep decrease in the film hardness and intrinsic stress. 3.2. Structure
clusters disconnecting due to absorbed gas impurities, which can lead to an increase in the number of sp 2 clusters. A microhardness of approximately 50 GPa was obtained for the 1.5-m thick film deposited at a discharge power of 960 W and substrate bias of y100 V ŽFig. 5.. The same result was achieved by decreasing these values to 300 W and y60 V, respectively. This increase in microhardness and intrinsic stress and simultaneous decrease in resistivity to a very low value
The electron diffraction patterns of all the investigated films are very similar to each other and consist of three visible haloes. The first small-angle halo appears only if the film contains fragments of a layer structure. It corresponds to the diffraction from prime interplanes wanalogous to the diffraction from Ž002. graphite planesx. The prime interplanar distance d and radius r of the first co-ordination sphere for C᎐C bonds de-
Fig. 6. Dependence of the microhardness and intrinsic stress of a-C:H films on the CH 4 rAr flow ratio at discharge power of 960 W; pAr s 0.17 Pa.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
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Table 1 Dependence of the prime interplanar distance Ž d . and the radius Ž r1 . of the first co-ordination sphere for a-C and a-C:H films on the depositions parameters Ar flow rate wsccmx
CH4 flow rate wsccmx
Pressure wPax
Substrate bias wVx
d wnmx
r1 " 0.001 wnmx
13 13 13 13 13 13
᎐ ᎐ ᎐ 1 12 24
0.17 0.17 0.17 0.19 0.32 0.44
y150 y100 Floating Ž; 5. Floating Ž; 5. Floating Ž; 5. Floating Ž; 5.
0.298᎐0.305 0.298᎐0.301 0.326᎐0.337 Not determined 0.389᎐0.405 0.405᎐0.410
0.142 0.142 0.142 0.143 0.143 0.143
termined are given in Table 1. The radius r was calculated by Ehrenfest’s equation w12x: 2 r sin s 1.23 ew in which 2 is scattering angle of the maximum of corresponding Žthird. visible halo and ew is the electron wavelengths. This procedure for all investigated a-C and a-C:H films gives a radius value r s 0.142᎐0.143 nm, corresponding closely to that of graphite. This means that carbon atoms in all our films, including the hardest a-C ones, are bonded to each other mainly through sp 2 bonding. The interplanar distance d can be determined from the Bragg’s formula: 2 dsin s n ew For a-C films obtained without ion bombardment, the interplanar distance d was found to be 0.326᎐0.337 nm Ž0.335 nm is the interlayer spacing in graphite.. Bias Žy100, y150 V. during deposition shifted these values to 0.298᎐0.305 nm ᎏ see Table 1. This con-
siderable decrease in d can be explained by the highly stressed sp 2 structure of the films. Such a structure is known for bulk graphite under high pressure, e.g. w13x. The main source of pressure in our films was intrinsic stress induced by ion bombardment. An intrinsic stress of approximately 10 GPa was obtained for a-C film with a thickness of approximately 50 nm deposited onto SiŽ111. at a substrate bias of y100 V ŽFig. 7.. The CH 4 addition to Ar led to a considerable increase in distance d ŽTable 1.. This agrees with the data for nanostructured graphite prepared by mechanical milling under a hydrogen atmosphere w14x. The main peak corresponding to this interplanar diffraction strongly decreases and shifts to lower angles, indicating an expansion of the graphite interlayer with increasing hydrogen in the samples. Part of the hydrogen probably occupied sites between the planes, as supported by neutron diffraction measurements w14x.
4. Discussion In the case of amorphous carbon film deposited by
Fig. 7. Dependence of the intrinsic stress of a-C films on the film thickness at discharge power of 960 W; pAr s 0.17 Pa; substrate bias y100 V.
V. Kuliko¨ sky et al. r Diamond and Related Materials 10 (2001) 1076᎐1081
magnetron sputtering, it would be possible to reduce or almost avoid the influence of gas impurities on the boundaries between the growing grains or clusters. There are two ways to achieve this: Ž1. by the preferential re-sputtering of gas adsorbed, mainly at the joints of growing film islands, through enhanced ion bombardment; Ž2. through a decrease in Ar pressure during sputtering to the value at which the mean free path of carbon atoms is larger than the cathode᎐substrate distance. In the latter case, every depositing carbon atom has an energy of approximately several eV. This is sufficient for both sputtering the weakly bound gas impurity, and occupying a stable position without significant heating of the surrounding atoms. The structure formed in this way can be very dense with strong interactions between atoms at grain boundaries, even if they occupy incoherent positions. This interaction can be strongly enhanced due to the biaxial compressive stress which accompanies such a process of film growth. The possible sliding along prime planes within clusters is difficult, due to the high disordering within clusters and high compressive stress, which, in turn, can bind these planes, see for example w15x. In any case, if such sliding is possible, it is realized partially during film growth under compressive stress. That is why these films demonstrate high elasticity Ž89᎐94%.. If the substrate temperature is low during film deposition and the recrystallization process is not developed Žfor carbon film, the recrystallization temperature is over 900 K., a nanostructure having a high hardness can be obtained. The experimental results presented support the aforementioned concept, according to the references cited.
5. Conclusions The superhardness of amorphous carbon films is usually explained by the presence of sp 3 bonds in the film. The higher the sp 3rsp 2 ratio is, the higher the film hardness. However, our results Ždata on electron diffraction and the low level of film resistivity. and those reported in w1᎐4x indicate that very hard carbon films with sp 2 C᎐C bonds can be obtained. It has been shown on our 1.5-m thick carbon films Ža-C and a-C:H. that with increasing negative substrate bias andror decreasing Ar, O 2 or CH 4 pressure, the resistivity decreases with a simultaneous increase in microhardness and compressive stress. The a-C films condensed under ion bombardment showed a microhardness of 50 GPa and a very low level of resistivity. Analysis of the electron diffraction patterns of thin films Ž30᎐60 nm. prepared under the same conditions
1081
showed that the radius of the first co-ordination sphere of C atoms for all the films Ža-C and a-C:H. was 0.142᎐0.143 nm, which is in good agreement with the value for sp 2 bonded carbon. The interplanar distance d shifted from approximately 0.326᎐0.337 nm Žfor unbiased a-C films. to 0.298᎐0.305 Žfor those deposited at a substrate bias of y150 V.. This can be caused by strong compression due to the intrinsic stress. All the data indicate an sp 2 bonding configuration of the hard carbon films deposited. The sp 2 hard carbon films presumably consist of cross-linked randomly oriented dense clusters, bonded to each other through incoherent boundaries on which high stresses are accumulated. The fewer the gas impurities on these boundaries are, the stronger is the bonding between the clusters. In this way, a continuous rigid network is formed. The high compressive stress can result in strong bonding between the cluster boundaries, on which a small concentration of sp 3 sites may be occurring.
Acknowledgements This work has been supported by the Concerned European Action on Tribology COST 516 under contract no. 516.50 and by grant GA CR 202r00r1592. References w1x G.A.J. Amaratunga, M. Chhowalla, C.J. Kiely, I. Alexandrou, R. Aharonov, R.M. Devenish, Nature 383 Ž1996. 321᎐323. w2x D. Camino, A.H.S. Jones, D. Mercs, D.G. Teer, Vacuum V52 Ž1999. 125. w3x R.G. Lacerda, P. Hammer, F. Alvarez, F.C. Marques, F.L. Freire, Jr, Diamond Relat. Mater. 9 Ž2000. 796. w4x R.G. Lacerda, F.C. Marques, Appl. Phys. Lett. 73 Ž5. Ž1998. 617. w5x I. Alexandrou, H.-J. Scheibe, C.J. Kiely, A.J. Papworth, G.A.J. Amaratunga, B. Schultrich, Phys. Rev. B 60 Ž1999. 10903. w6x V.Y. Kulikovsky, F. Fendrych, L. Jastrabik, D. Chvostova, Surf. Coat. Technol. 91 Ž1997. 122. w7x G.G. Stoney, Proc. R. Soc. Lond. Ser. A 82 Ž1909. 172. w8x J.W. Zou, K. Reichelt, K. Schmidt, B. Dischler, J. Appl. Phys. 65 Ž1989. 3914. w9x D.R. McKenzie, Y. Yin, N.A. Marks et al., Diamond Relat. Mater. 3 Ž1994. 353. w10x R.S. Rubino, E.S. Takcuchi, J. Power Sources 81r82 Ž1999. 373. w11x R.O. Dillon, J.A. Woollam, V. Katkanant, Phys.Rev. B 29 Ž1984. 3482. w12x P. Ehrenfest, Prog. Amst. Acad. 17 Ž1915. 1132. w13x R.W. Lynch, H.G. Drickamer, J. Chem. Phys. 44 Ž1966. 181. w14x S. Orimo, G. Majer, T. Fukunaga, A. Zuttel, L. Schlapbach, H. Fujii, Appl. Phys. Lett. 75 Ž1999. 3093. w15x J. Schwan, S. Ulrich, T. Theel et al., J. Appl. Phys. 82 Ž1997. 6025.