- Email: [email protected]
Annealing effect on microstructure and mechanical properties of amorphous Al–C–N films
Applied Surface Science 257 (2011) 6079–6084
Contents lists available at ScienceDirect
Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Annealing effect on microstructure and mechanical properties of amorphous Al–C–N films Q.H. Luo a,b , Y.H. Lu a,∗ , Y.Z. Lou b a b
National Center for Materials Service Safety, University of Science and Technology Beijing, Beijing 100083, China Beijing Institute of Aeronautical Materials, Beijing 100095, China
a r t i c l e
i n f o
Article history: Received 2 January 2011 Received in revised form 28 January 2011 Accepted 29 January 2011 Available online 21 February 2011 Keywords: Al–C–N films Annealing Microstructure Mechanical properties
a b s t r a c t Al–C–N thin films with different Al contents were deposited on Si (1 0 0) substrates by closed-field unbalanced reactive magnetron sputtering in the mixture of argon and nitrogen gases. These films were subsequently vacuum-annealed at 700 ◦ C and 1000 ◦ C, respectively. The microstructures of as-deposited and annealed films were characterized by X-ray diffraction (XRD) and high-resolution transmission electron microscopy (HRTEM); while the hardness and elastic modulus values were measured by nanoindention method. The results indicated that the microstructure of both as-deposited and annealed Al–C–N films strongly depended on Al content. For thin films at low Al content, film delamination rather than crystallization occurred after the sample was annealed at 1000 ◦ C. For thin films at high Al content, annealing led to the formation of AlN nanocrystallites, which produced nanocomposites of AlN embedded into amorphous matrices. Both the density and size of AlN nanocrystallites were found to decrease with increasing depth from the film surface. With increasing of annealing temperature, both hardness and elastic modulus values were decreased; this trend was decreased at high Al content. Annealing did not change elastic recovery property of Al–C–N thin films. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The amorphous C–N films (a-CNx ) have received a great deal of attention since Liu and Cohen [1] predicted that hardness of -C3 N4 might exceed that of diamond [2–4]. Addition of one or more third elements, B, Si, Al, etc., in C–N thin films can improve its thermal stability and reduce its residual stress, which meets special properties [5–8]. For example, addition of Al element formed h-AlN/a-CN nano-composite structure, which effectively increased the nitrogen content, as a result enhanced film hardness. On the other hand, Al–C–N film has a potential large band gap, high heat conductivity coefficient, high hardness, good thermal and chemical stability, as well as unique performance in the semiconductor, optoelectronics, and piezoelectric, which result in large application potential in optoelectronic devices, microelectronic devices and micro-friction domain for Al–C–N films. Al–C–N film can be prepared by many methods such as reactive magnetron sputtering [9–12], implantation of carbon and nitrogen ions into aluminum block and so on [13]. In most case, Al–C–N films prepared by reactive magnetron sputtering are amorphous or low degree of crystallization [9–12]. It is well known that the carbon-related films usually have a low thermal
∗ Corresponding author. E-mail address: lu [email protected] (Y.H. Lu). 0169-4332/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2011.01.140
stability, which is essential to be noticed in practical application. Microstructure and composition of Al–C–N films are believed to be important controlling factors for such a property. Therefore, the study on structure and properties of annealed Al–C–N films is necessary. This study is based on our earlier research on structure, electrical conductivity and Hall effect of amorphous Al–C–N thin films [14], from which two typical thin films were vacuum annealed. The as-deposited and annealed thin films were characterized ex situ in terms of nanostructure by X-ray diffraction and highresolution transmission electron microscopy, their hardness and elastic modulus by nano-indenter, in order to investigate evolution of nanostructure and changes in mechanical properties with annealing temperature.
2. Experimental procedure In this study, a series of Al–C–N films with different Al contents were deposited on Si (1 0 0) substrates by medium frequency magnetron sputtering method using one aluminum and one graphite targets with a high-purity (99.999%). The background pressure was pumped down to ≤2.7 × 10−4 Pa. The thin films were deposited at a pulsed biased voltage of −80 V (250 kHz) at a working pressure of ∼0.26 Pa under a mixture gas of high purity argon and high-purity nitrogen. During deposition, a substrate rotation speed of 10 rpm
6080
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
Fig. 1. The change in C and N content with Al content in presently deposited Al–C–N films.
and a target–substrate distance of 10 cm were applied; while the flow rates of argon and nitrogen were controlled to 15 and 9 sccm, respectively. Al–C–N thin films with different Al contents were deposited by way of adjusting electric current of Al target under a fixed graphite target current of 6.0 A. All thin films were deposited at 150 ◦ C for 40 min, and the film thickness was about 650 nm by examination using cross-sectional SEM (field emission JSM-6335F). The substrate pre-sputtering before Al–C–N film deposition was same with that in the literature [14]. A CABROLITE high temperature vacuum furnace equipped with a mechanical/diffusion oil pumping system was used to conduct annealing process. After inserting the specimens, the quartz tube was sealed and pumped down to 0.018 mPa. Annealing treatment was carried out at two different temperatures, 700 and 1000 ◦ C, for 5 h. The structures of as-deposited and annealed Al–C–N films were analyzed by X-ray diffraction (XRD) using a Bragg–Brentano diffractometer (D/Max-RB, Cu target, voltage 40 kV, current 150 mA) in –2 configuration. Their nanostructures were examined by highresolution transmission electron microscopy (HRTEM, JEOL 2010) operated at 200 kV. Film chemical composition and phase configuration were determined by X-ray photoelectron spectroscopy (XPS, Thermofisher, voltage 4 kV, current 4 A). Before XPS examination, the film surface was cleaned using methanol, and subsequently etched down to about 40 nm using 1 keV Ar ion to remove the surface oxide. The hardness and elastic modulus values of thin films were measured by a Nanoindent II nanoindentation instrument by a single loading–unloading mode until a maximum indentation depth of 65 nm was reached. Here, a four-sided diamond pyramid tip was used. The calibration of the hardness value was frequently checked by measuring on a standard fused silica sample. In order to minimize the substrate effect, the maximum indentation depth was kept less than 10% of the total film thickness. Six separate measurements were taken for each sample in order to get a mean value. 3. Results and discussion Fig. 1 shows changes of C and N content with increase of Al content in the presently deposited thin films. The element concentrations were measured by XPS. It is found that with increase of Al content N content monotonically increases at the cost of the C content. It should be pointed out that there was some oxygen contamination present in all samples due to the environment but it was mainly on the surface. After sputtering down to ∼5–10 nm from the surface, the oxygen content was approximately 2–5 at.% and neglected.
Fig. 2. High resolution spectra of N 1s XPS in Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films, as well as the corresponding peak fitting of Al3 C84 N13 film.
In order to characterize the bonding structure of these films, typical high-resolution core-level XPS analyses were carried out. Fig. 2 shows high resolution XPS N 1s core-level spectra of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 , respectively. To precisely characterize their bonds, these XPS spectra were fitted using XPSPEAK41 software based on the Lorentzian–Gaussian function. As an example the spectra peak fitting of Al3 C84 N13 is also shown in Fig. 2. From Fig. 2 it is found that two nitrogen-related bonds appear in N 1s spectra of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films. They are located at binding energies 396.5 and 397.7 eV, and were assigned to N–Al and N–C, respectively. The fitted areas of XPS N 1s peak in Al3 C84 N13 , Al13 C71 N16 , and Al32 C43 N25 films are presented in Table 1. It is found that, with increase of Al content, N–Al bond is significantly enhanced, and instead N–C bonding is greatly decreased. It indicates that Al–N has a much larger affinity than N–C, which results in formation of Al–N rather than N–C. Fig. 3(a) and (b) shows XRD patterns of Al3 C84 N13 and Al32 C43 N25 thin films in as-deposited and after annealed states, respectively. It is found from Fig. 3(a) that besides (1 0 0) characteristic peak of single crystal silicon substrate there is no peak in Al3 C84 N13 thin film in as-deposited state, which indicates that the as-deposited Al3 C84 N13 thin film exhibits amorphous. In fact, all presently deposited Al–C–N thin films are amorphous. After Al3 C84 N13 was vacuum annealed at 700 and 1000 ◦ C, respectively, no obvious change in their XRD spectra is found, which indicates that annealing not higher than 1000 ◦ C did not make Al3 C84 N13 crystallize. From Fig. 3(b), it is found that Al32 C43 N25 thin film remains its amorphous structure till 700 ◦ C; however, additional (1 0 0) peak of AlN appears in the XRD patterns after it was annealed at 1000 ◦ C, which indicates crystallization took place after 1000 ◦ C annealing, resulting in the formation of hexagonal structural AlN crystal in Al32 C43 N25 . It is apparent that the nanostructure of annealed Al–C–N thin films is strongly dependent of the film composition. At low Al content there is no enough Al–N to concentrate through diffusion to produce AlN nanocrystal at 1000 ◦ C, as a result there is no AlN peak Table 1 The fitted area of XPS N 1s peak in Al3 C84 N13 , Al13 C71 N16 , and Al32 C43 N25 films. Peak area (× 104 )
Al3 C84 N13 Al13 C71 N16 Al32 C43 N25
N 1s N–Al
N–C
5.8 6.5 9.9
2.7 2.0 1.1
Uncertainties of the data were estimated to be within 10%.
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
6081
Fig. 3. XRD patterns of as-deposited and annealed (a) Al3 C84 N13 and (b) Al32 C43 N25 .
in the XRD pattern shown in Fig. 3(a). Only when the Al concentration reaches a certain level, enough Al–N concentration in the local region is built at high temperature, producing nanocrystal AlN, as a result AlN (1 0 0) diffraction peak appears in the XRD pattern, as shown in Fig. 3(b). By the Scherrer formula average size of nanocrystallites after such a treatment is estimated to be about 9 nm. Fig. 4(a) and (b) shows cross-section TEM and high, resolution TEM images of as-deposited Al3 C84 N13 thin film, respectively, together with their Fourier Transform (FFT) spectra in Fig. 4(c) and (d). From Fig. 4(a) there is no cavity in the film and the film–substrate interface, which indicates a dense film and a good cohesion between film–substrate interface. From Fig. 4(b) it is found that Al3 C84 N13 thin film is amorphous, which is consistent with its XRD result. Compared with the HRTEM lattice in region A, a strained-layer with a thickness of about 1.5 nm is found in region B in Si substrate adjacent to the film/substrate interface. Compared with that in Fig. 4(c), the streaking (spot splitting) in the FFT spectrum in Fig. 4(d) reveals the presence of a large strain (lattice aberration) in the Si substrate, which is due to strain deformation because of residual stress in Fig. 4(a). Its deformation direction is normal to the stretching of the diffraction spots. Fig. 5 is cross-section TEM and HRTEM images of Al3 C84 N13 thin film annealed at 1000 ◦ C. From the TEM bright field image (Fig. 5(a)) it is found that there are four clear layers, the boundaries of which are almost parallel to the film/substrate interface. Fig. 5(b) shows HRTEM image across the layer boundary A in Fig. 5(a). It is found that there are two amorphous layers across the boundary, which is confirmed by the FFT spectrum (see the inset). It seems that there are two different alignments in adjacent amorphous layers. Fig. 5(c) is a HRTEM image showing the film–substrate interface structure. It is found that there is no any strain in the Si substrate,
Fig. 4. Cross-section TEM image of as-deposited Al3 C84 N13 thin film: (a) bright field image; (b) high-resolution image close to film/substrate interface; (c and d) FFT spectra of A and B regions, respectively in (b).
6082
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084 Table 2 Chemical composition at different depths perpendicular to the surface of Al3 C84 N13 after annealed at 1000 ◦ C. Compositions
P1
P2
P3
P4
C N O Al
87.4 1.7 8.0 2.9
87.3 2.0 7.8 2.9
86.4 3.1 8.0 2.5
86.9 3.3 7.4 2.4
film after 1000 ◦ C annealing, which is distinguished from that in Al3 C84 N13 . Fig. 6(b) and (c) shows HRTEM images of positions A and B in Fig. 6(a), respectively. From Fig. 6(b) the nanocrystallites with a diameter size of about 5 nm were found to be embedded into amorphous matrices in the film interior. In nano-grains single {1 0 0}
Fig. 5. Cross-section TEM image of Al3 C84 N13 after annealed at 1000 ◦ C: (a) bright field image; (b and c) high-resolution image from A and B regions, respectively in (a). The FFT spectra of selected parts in (b) and (c) are inserted.
which is evidenced by the FFT spectrum obtained from region A in the interface. It indicates that annealing effectively released the residual stress in the film/substrate interface. These results suggest that, after 1000 ◦ C annealing, layer delamination takes place in Al3 C84 N13 thin film although such a layer delamination does not change the amorphous structure, which is believed to be related to short-range order in each layer [15,16]. In order to further characterize nanostructure of the annealed film, TEM–EDS analyses on P1, P2, P3 and P4 were carried out. Their results are listed in Table 2. It is found that the N content in annealed Al3 C84 N13 varies in the range of 1.7–3.3 at.%, much lower than that in as-deposited state (13 at.%). The decrease in N content is believed to be contributed to diffusion in annealing process. Fig. 6 is cross-section TEM image of Al32 C43 N25 film after annealed at 1000 ◦ C. From TEM bright field image (see Fig. 6(a)) it is apparent that no obvious layer delamination is found in Al32 C43 N25
Fig. 6. Cross-section TEM image of Al32 C43 N25 thin films after annealed at 1000 ◦ C: (a) bright field image; (b and c) high-resolution images from A and B regions in (a). The FFT spectra of both (b) and (c) are inserted.
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
6083
Table 3 Chemical composition at different depths perpendicular to the surface of Al32 C43 N25 after annealed at 1000 ◦ C. Compositions (at.%)
P1
P2
P3
P4
C N O Al
32.4 16.7 7.9 43.0
36.4 15.5 6.9 41.2
41.4 14.3 6.4 37.9
41.8 13.9 6.1 38.2
lattice fringes with a lattice space of ∼3.0 A˚ are observed. This value is consistent with that shown in JCPDS data for AlN (3.05 A˚ for (1 0 0) lattice space). It suggests that 1000 ◦ C annealing makes Al32 C43 N25 crystallize, which produces nanocomposite structure of AlN nanocrystals embedded in amorphous matrices. This process is related to fast concentration of Al–N bond at high Al content in Al–C–N film. On the other hand, as shown in Fig. 6(c), lower density and smaller sized nanocrystallites are found in the vicinity close to the film/substrate surface, which indicates that there is no uniformity in structure along the film depth after annealing. With increase of the depth from the film surface both the size and density of nanocrystallites decreased. Besides, from Fig. 6(c) it is found that there is no any strain in the Si substrate close to the film/substrate interface, which indicates that annealing can effectively release the residual stress in the film/substrate interface. In order to further characterize nanostructure of annealed Al32 C43 N25 thin films, TEM–EDS analyses on P1, P2, P3 and P4 (Fig. 6(a)) were carried out. The results are listed in Table 3. It is found that N content of Al32 C43 N25 film after 1000 ◦ C-annealing varies in the range of 13.9–16.7 at.%, lower than that of as-deposited one. It slightly decreases with increase of film depth. Fig. 7 shows the dependence of hardness on indentation depth in Al3 C84 N13 thin films in as-deposited state and after annealed at 700 and 1000 ◦ C, respectively. It is found that after impression into film about 30 nm the hardness values of both as-deposited and annealed films almost do not change with the indentation depth, which indicates the present hardness measurement has avoided substrate effect. Fig. 8(a) and (b) shows nanohardness and elastic modulus values of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 in as-deposited state and after annealed at different temperatures, respectively. From Fig. 8(a) it is found that as-deposited Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 have nano-hardness values of 14.4, 12.8 and 9.4 GPa, respectively. It is apparent that with increase of Al content the nanohardness value is decreased, which is contributed to decrease of C–N bonding with Al content. On the other hand, when the
Fig. 7. The dependence of nano-hardness value on indentation depth in asdeposited and annealed Al3 C84 N13 thin films.
Fig. 8. The dependences of (a) hardness (b) elastic modulus values on annealing temperature in Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films.
annealing temperature increases nano-hardness values of the three films gradually decrease, which is mostly contributed to relaxation of the residual stress and decrease of nitrogen content (see Tables 2 and 3) during annealing. It is presumed that decrease of nitrogen content decreased the amount of C–N bonding, which resulted in decrease of hardness. It is worthy of noticing that less decrease in hardness value with the annealing temperature takes place for high Al contented thin films. A similar phenomenon in elastic modulus value is found in Fig. 8(b). Fig. 9 presents the load–displacement curve of Al3 C84 N13 in asdeposited sate and after annealed at 700 and 1000 ◦ C, respectively. It is found that for the present measurement the maximum inden-
Fig. 9. The dependence of load on indentation depth in as-deposited and annealed Al3 C84 N13 thin films.
6084
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
tation depth is about 65 nm, about one tenth of the film thickness. It is well known that the elastic recovery rate ER is calculated as follows: ER =
dmax − dres dmax
where dmax is the maximum indentation depth and dres is the residual indentation depth after unloading [17,18]. It is calculated that the elastic recovery rates of all three samples are about 69%, which reveals that annealing does not change the elastic recovery rate of Al–C–N film. 4. Conclusions In this study, three Al–C–N films with different Al contents were deposited on Si (1 0 0) substrates by closed-field unbalanced reactive magnetron sputtering. The phase configuration, microstructure and mechanical properties of three films in asdeposited and annealed states were subsequently studied. The main results are listed as follows: (1) The phase configuration and microstructure in both asdeposited and annealed Al–C–N thin films strongly depended on Al content. Usually, for Al–C–N films with low Al content delamination rather than crystallization occurred after the films were annealed at 1000 ◦ C. By comparison, for thin films at high Al content, annealing led to formation of AlN nanocrystallites, which produced nanocomposites of AlN embedded into amorphous matrices. (2) With increase of annealing temperature, both hardness and elastic modulus values of Al–C–N thin films were decreased, however less decrease was found for Al–C–N thin films at high
Al content. On the other hand, annealing does not change the elastic recovery rate of Al–C–N films. Acknowledgement The work described in this paper was supported by The National Nature Research Fund of China (No. 50871018). References [1] A.Y. Liu, M.L. Cohen, Science 245 (1989) 841. [2] M. Bai, K. Kato, N. Umehara, Y. Miyake, J. Xu, H. Tokisue, Thin Solid Films 376 (2000) 170. [3] X.C. Wang, Z.Q. Li, P. Wu, E.Y. Jiang, H.L. Bai, Appl. Surf. Sci. 253 (2006) 2087. [4] R. Prioli, S.I. Zanette, A.O. Caride, M.M. Lacerda, F.L. Freire Jr., Diamond Relat. Mater. 8 (1999) 993. [5] S. Ma, B. Xu, G. Wu, Y. Wang, F. Ma, D. Ma, K. Xu, T. Bell, Surf. Coat. Technol. 202 (2008) 5379. [6] D. Watanabe, H. Aoki, R. Moriyama, M.K. Mazumder, C. Kimura, T. Sugino, Diamond Relat. Mater. 17 (2008) 669. [7] Y.L. Yang, D. Zhang, W. Yan, Y.R. Zheng, Opt. Laser Eng. 48 (2010) 119. [8] J.L. Lin, J.J. Moore, B. Mishra, M. Pinkas, W.D. Sproul, Acta Mater. 58 (2010) 1554. [9] N. Jiang, S. Xu, K.N. Ostrikov, E.L. Tsakadze, J.D. Long, J.W. Chai, Z.L. Tsakadze, Int. J. Mod. Phys. B 16 (2002) 1132. [10] A.L. Ji, Y. Du, L.B. Ma, Z.X. Cao, J. Cryst. Growth 279 (2005) 420. [11] A.L. Ji, L.B. Ma, C. Liu, C.R. Li, Z.X. Cao, Diamond Relat. Mater. 14 (2005) 1348. [12] A.L. Ji, L.B. Ma, C. Liu, P. Zheng, C.R. Li, Z.X. Cao, Appl. Phys. Lett. 86 (2005) 021918. [13] V.V. Uglov, N.N. Cherenda, V.V. Khodasevich, V.A. Sokol, I.I. Abramov, A.L. Danilyuk, A. Wenzel, J. Gerlach, B. Rauschenbach, Nucl. Instrum. Methods Phys. Res. B 147 (1999) 332. [14] Q.H. Luo, D.L. Yu, Y.H. Lu, Chin. J. Vac. Sci. Technol. 30 (2010) 138. [15] M.J. Zhou, S.F. Wong, C.W. Ong, Q. Li, Thin Solid Films 516 (2007) 336. [16] D.G. McCulloch, X.L. Xiao, J.L. Peng, P.C.T. Ha, D.R. McKenzie, M.M.M. Bilek, S.P. Lau, D. Sheeja, B.K. Tay, Surf. Coat. Technol. 198 (2005) 217. [17] W.Y. Ni, Y.T. Cheng, D.S. Grummon, Appl. Phys. Lett. 80 (2002) 3311. [18] Y.H. Lu, Y.G. Shen, J.P. Wang, Z.F. Zhou, K.Y. Li, Surf. Coat. Technol. 201 (2007) 7368.
Contents lists available at ScienceDirect
Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Annealing effect on microstructure and mechanical properties of amorphous Al–C–N films Q.H. Luo a,b , Y.H. Lu a,∗ , Y.Z. Lou b a b
National Center for Materials Service Safety, University of Science and Technology Beijing, Beijing 100083, China Beijing Institute of Aeronautical Materials, Beijing 100095, China
a r t i c l e
i n f o
Article history: Received 2 January 2011 Received in revised form 28 January 2011 Accepted 29 January 2011 Available online 21 February 2011 Keywords: Al–C–N films Annealing Microstructure Mechanical properties
a b s t r a c t Al–C–N thin films with different Al contents were deposited on Si (1 0 0) substrates by closed-field unbalanced reactive magnetron sputtering in the mixture of argon and nitrogen gases. These films were subsequently vacuum-annealed at 700 ◦ C and 1000 ◦ C, respectively. The microstructures of as-deposited and annealed films were characterized by X-ray diffraction (XRD) and high-resolution transmission electron microscopy (HRTEM); while the hardness and elastic modulus values were measured by nanoindention method. The results indicated that the microstructure of both as-deposited and annealed Al–C–N films strongly depended on Al content. For thin films at low Al content, film delamination rather than crystallization occurred after the sample was annealed at 1000 ◦ C. For thin films at high Al content, annealing led to the formation of AlN nanocrystallites, which produced nanocomposites of AlN embedded into amorphous matrices. Both the density and size of AlN nanocrystallites were found to decrease with increasing depth from the film surface. With increasing of annealing temperature, both hardness and elastic modulus values were decreased; this trend was decreased at high Al content. Annealing did not change elastic recovery property of Al–C–N thin films. © 2011 Elsevier B.V. All rights reserved.
1. Introduction The amorphous C–N films (a-CNx ) have received a great deal of attention since Liu and Cohen [1] predicted that hardness of -C3 N4 might exceed that of diamond [2–4]. Addition of one or more third elements, B, Si, Al, etc., in C–N thin films can improve its thermal stability and reduce its residual stress, which meets special properties [5–8]. For example, addition of Al element formed h-AlN/a-CN nano-composite structure, which effectively increased the nitrogen content, as a result enhanced film hardness. On the other hand, Al–C–N film has a potential large band gap, high heat conductivity coefficient, high hardness, good thermal and chemical stability, as well as unique performance in the semiconductor, optoelectronics, and piezoelectric, which result in large application potential in optoelectronic devices, microelectronic devices and micro-friction domain for Al–C–N films. Al–C–N film can be prepared by many methods such as reactive magnetron sputtering [9–12], implantation of carbon and nitrogen ions into aluminum block and so on [13]. In most case, Al–C–N films prepared by reactive magnetron sputtering are amorphous or low degree of crystallization [9–12]. It is well known that the carbon-related films usually have a low thermal
∗ Corresponding author. E-mail address: lu [email protected] (Y.H. Lu). 0169-4332/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2011.01.140
stability, which is essential to be noticed in practical application. Microstructure and composition of Al–C–N films are believed to be important controlling factors for such a property. Therefore, the study on structure and properties of annealed Al–C–N films is necessary. This study is based on our earlier research on structure, electrical conductivity and Hall effect of amorphous Al–C–N thin films [14], from which two typical thin films were vacuum annealed. The as-deposited and annealed thin films were characterized ex situ in terms of nanostructure by X-ray diffraction and highresolution transmission electron microscopy, their hardness and elastic modulus by nano-indenter, in order to investigate evolution of nanostructure and changes in mechanical properties with annealing temperature.
2. Experimental procedure In this study, a series of Al–C–N films with different Al contents were deposited on Si (1 0 0) substrates by medium frequency magnetron sputtering method using one aluminum and one graphite targets with a high-purity (99.999%). The background pressure was pumped down to ≤2.7 × 10−4 Pa. The thin films were deposited at a pulsed biased voltage of −80 V (250 kHz) at a working pressure of ∼0.26 Pa under a mixture gas of high purity argon and high-purity nitrogen. During deposition, a substrate rotation speed of 10 rpm
6080
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
Fig. 1. The change in C and N content with Al content in presently deposited Al–C–N films.
and a target–substrate distance of 10 cm were applied; while the flow rates of argon and nitrogen were controlled to 15 and 9 sccm, respectively. Al–C–N thin films with different Al contents were deposited by way of adjusting electric current of Al target under a fixed graphite target current of 6.0 A. All thin films were deposited at 150 ◦ C for 40 min, and the film thickness was about 650 nm by examination using cross-sectional SEM (field emission JSM-6335F). The substrate pre-sputtering before Al–C–N film deposition was same with that in the literature [14]. A CABROLITE high temperature vacuum furnace equipped with a mechanical/diffusion oil pumping system was used to conduct annealing process. After inserting the specimens, the quartz tube was sealed and pumped down to 0.018 mPa. Annealing treatment was carried out at two different temperatures, 700 and 1000 ◦ C, for 5 h. The structures of as-deposited and annealed Al–C–N films were analyzed by X-ray diffraction (XRD) using a Bragg–Brentano diffractometer (D/Max-RB, Cu target, voltage 40 kV, current 150 mA) in –2 configuration. Their nanostructures were examined by highresolution transmission electron microscopy (HRTEM, JEOL 2010) operated at 200 kV. Film chemical composition and phase configuration were determined by X-ray photoelectron spectroscopy (XPS, Thermofisher, voltage 4 kV, current 4 A). Before XPS examination, the film surface was cleaned using methanol, and subsequently etched down to about 40 nm using 1 keV Ar ion to remove the surface oxide. The hardness and elastic modulus values of thin films were measured by a Nanoindent II nanoindentation instrument by a single loading–unloading mode until a maximum indentation depth of 65 nm was reached. Here, a four-sided diamond pyramid tip was used. The calibration of the hardness value was frequently checked by measuring on a standard fused silica sample. In order to minimize the substrate effect, the maximum indentation depth was kept less than 10% of the total film thickness. Six separate measurements were taken for each sample in order to get a mean value. 3. Results and discussion Fig. 1 shows changes of C and N content with increase of Al content in the presently deposited thin films. The element concentrations were measured by XPS. It is found that with increase of Al content N content monotonically increases at the cost of the C content. It should be pointed out that there was some oxygen contamination present in all samples due to the environment but it was mainly on the surface. After sputtering down to ∼5–10 nm from the surface, the oxygen content was approximately 2–5 at.% and neglected.
Fig. 2. High resolution spectra of N 1s XPS in Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films, as well as the corresponding peak fitting of Al3 C84 N13 film.
In order to characterize the bonding structure of these films, typical high-resolution core-level XPS analyses were carried out. Fig. 2 shows high resolution XPS N 1s core-level spectra of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 , respectively. To precisely characterize their bonds, these XPS spectra were fitted using XPSPEAK41 software based on the Lorentzian–Gaussian function. As an example the spectra peak fitting of Al3 C84 N13 is also shown in Fig. 2. From Fig. 2 it is found that two nitrogen-related bonds appear in N 1s spectra of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films. They are located at binding energies 396.5 and 397.7 eV, and were assigned to N–Al and N–C, respectively. The fitted areas of XPS N 1s peak in Al3 C84 N13 , Al13 C71 N16 , and Al32 C43 N25 films are presented in Table 1. It is found that, with increase of Al content, N–Al bond is significantly enhanced, and instead N–C bonding is greatly decreased. It indicates that Al–N has a much larger affinity than N–C, which results in formation of Al–N rather than N–C. Fig. 3(a) and (b) shows XRD patterns of Al3 C84 N13 and Al32 C43 N25 thin films in as-deposited and after annealed states, respectively. It is found from Fig. 3(a) that besides (1 0 0) characteristic peak of single crystal silicon substrate there is no peak in Al3 C84 N13 thin film in as-deposited state, which indicates that the as-deposited Al3 C84 N13 thin film exhibits amorphous. In fact, all presently deposited Al–C–N thin films are amorphous. After Al3 C84 N13 was vacuum annealed at 700 and 1000 ◦ C, respectively, no obvious change in their XRD spectra is found, which indicates that annealing not higher than 1000 ◦ C did not make Al3 C84 N13 crystallize. From Fig. 3(b), it is found that Al32 C43 N25 thin film remains its amorphous structure till 700 ◦ C; however, additional (1 0 0) peak of AlN appears in the XRD patterns after it was annealed at 1000 ◦ C, which indicates crystallization took place after 1000 ◦ C annealing, resulting in the formation of hexagonal structural AlN crystal in Al32 C43 N25 . It is apparent that the nanostructure of annealed Al–C–N thin films is strongly dependent of the film composition. At low Al content there is no enough Al–N to concentrate through diffusion to produce AlN nanocrystal at 1000 ◦ C, as a result there is no AlN peak Table 1 The fitted area of XPS N 1s peak in Al3 C84 N13 , Al13 C71 N16 , and Al32 C43 N25 films. Peak area (× 104 )
Al3 C84 N13 Al13 C71 N16 Al32 C43 N25
N 1s N–Al
N–C
5.8 6.5 9.9
2.7 2.0 1.1
Uncertainties of the data were estimated to be within 10%.
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
6081
Fig. 3. XRD patterns of as-deposited and annealed (a) Al3 C84 N13 and (b) Al32 C43 N25 .
in the XRD pattern shown in Fig. 3(a). Only when the Al concentration reaches a certain level, enough Al–N concentration in the local region is built at high temperature, producing nanocrystal AlN, as a result AlN (1 0 0) diffraction peak appears in the XRD pattern, as shown in Fig. 3(b). By the Scherrer formula average size of nanocrystallites after such a treatment is estimated to be about 9 nm. Fig. 4(a) and (b) shows cross-section TEM and high, resolution TEM images of as-deposited Al3 C84 N13 thin film, respectively, together with their Fourier Transform (FFT) spectra in Fig. 4(c) and (d). From Fig. 4(a) there is no cavity in the film and the film–substrate interface, which indicates a dense film and a good cohesion between film–substrate interface. From Fig. 4(b) it is found that Al3 C84 N13 thin film is amorphous, which is consistent with its XRD result. Compared with the HRTEM lattice in region A, a strained-layer with a thickness of about 1.5 nm is found in region B in Si substrate adjacent to the film/substrate interface. Compared with that in Fig. 4(c), the streaking (spot splitting) in the FFT spectrum in Fig. 4(d) reveals the presence of a large strain (lattice aberration) in the Si substrate, which is due to strain deformation because of residual stress in Fig. 4(a). Its deformation direction is normal to the stretching of the diffraction spots. Fig. 5 is cross-section TEM and HRTEM images of Al3 C84 N13 thin film annealed at 1000 ◦ C. From the TEM bright field image (Fig. 5(a)) it is found that there are four clear layers, the boundaries of which are almost parallel to the film/substrate interface. Fig. 5(b) shows HRTEM image across the layer boundary A in Fig. 5(a). It is found that there are two amorphous layers across the boundary, which is confirmed by the FFT spectrum (see the inset). It seems that there are two different alignments in adjacent amorphous layers. Fig. 5(c) is a HRTEM image showing the film–substrate interface structure. It is found that there is no any strain in the Si substrate,
Fig. 4. Cross-section TEM image of as-deposited Al3 C84 N13 thin film: (a) bright field image; (b) high-resolution image close to film/substrate interface; (c and d) FFT spectra of A and B regions, respectively in (b).
6082
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084 Table 2 Chemical composition at different depths perpendicular to the surface of Al3 C84 N13 after annealed at 1000 ◦ C. Compositions
P1
P2
P3
P4
C N O Al
87.4 1.7 8.0 2.9
87.3 2.0 7.8 2.9
86.4 3.1 8.0 2.5
86.9 3.3 7.4 2.4
film after 1000 ◦ C annealing, which is distinguished from that in Al3 C84 N13 . Fig. 6(b) and (c) shows HRTEM images of positions A and B in Fig. 6(a), respectively. From Fig. 6(b) the nanocrystallites with a diameter size of about 5 nm were found to be embedded into amorphous matrices in the film interior. In nano-grains single {1 0 0}
Fig. 5. Cross-section TEM image of Al3 C84 N13 after annealed at 1000 ◦ C: (a) bright field image; (b and c) high-resolution image from A and B regions, respectively in (a). The FFT spectra of selected parts in (b) and (c) are inserted.
which is evidenced by the FFT spectrum obtained from region A in the interface. It indicates that annealing effectively released the residual stress in the film/substrate interface. These results suggest that, after 1000 ◦ C annealing, layer delamination takes place in Al3 C84 N13 thin film although such a layer delamination does not change the amorphous structure, which is believed to be related to short-range order in each layer [15,16]. In order to further characterize nanostructure of the annealed film, TEM–EDS analyses on P1, P2, P3 and P4 were carried out. Their results are listed in Table 2. It is found that the N content in annealed Al3 C84 N13 varies in the range of 1.7–3.3 at.%, much lower than that in as-deposited state (13 at.%). The decrease in N content is believed to be contributed to diffusion in annealing process. Fig. 6 is cross-section TEM image of Al32 C43 N25 film after annealed at 1000 ◦ C. From TEM bright field image (see Fig. 6(a)) it is apparent that no obvious layer delamination is found in Al32 C43 N25
Fig. 6. Cross-section TEM image of Al32 C43 N25 thin films after annealed at 1000 ◦ C: (a) bright field image; (b and c) high-resolution images from A and B regions in (a). The FFT spectra of both (b) and (c) are inserted.
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
6083
Table 3 Chemical composition at different depths perpendicular to the surface of Al32 C43 N25 after annealed at 1000 ◦ C. Compositions (at.%)
P1
P2
P3
P4
C N O Al
32.4 16.7 7.9 43.0
36.4 15.5 6.9 41.2
41.4 14.3 6.4 37.9
41.8 13.9 6.1 38.2
lattice fringes with a lattice space of ∼3.0 A˚ are observed. This value is consistent with that shown in JCPDS data for AlN (3.05 A˚ for (1 0 0) lattice space). It suggests that 1000 ◦ C annealing makes Al32 C43 N25 crystallize, which produces nanocomposite structure of AlN nanocrystals embedded in amorphous matrices. This process is related to fast concentration of Al–N bond at high Al content in Al–C–N film. On the other hand, as shown in Fig. 6(c), lower density and smaller sized nanocrystallites are found in the vicinity close to the film/substrate surface, which indicates that there is no uniformity in structure along the film depth after annealing. With increase of the depth from the film surface both the size and density of nanocrystallites decreased. Besides, from Fig. 6(c) it is found that there is no any strain in the Si substrate close to the film/substrate interface, which indicates that annealing can effectively release the residual stress in the film/substrate interface. In order to further characterize nanostructure of annealed Al32 C43 N25 thin films, TEM–EDS analyses on P1, P2, P3 and P4 (Fig. 6(a)) were carried out. The results are listed in Table 3. It is found that N content of Al32 C43 N25 film after 1000 ◦ C-annealing varies in the range of 13.9–16.7 at.%, lower than that of as-deposited one. It slightly decreases with increase of film depth. Fig. 7 shows the dependence of hardness on indentation depth in Al3 C84 N13 thin films in as-deposited state and after annealed at 700 and 1000 ◦ C, respectively. It is found that after impression into film about 30 nm the hardness values of both as-deposited and annealed films almost do not change with the indentation depth, which indicates the present hardness measurement has avoided substrate effect. Fig. 8(a) and (b) shows nanohardness and elastic modulus values of Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 in as-deposited state and after annealed at different temperatures, respectively. From Fig. 8(a) it is found that as-deposited Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 have nano-hardness values of 14.4, 12.8 and 9.4 GPa, respectively. It is apparent that with increase of Al content the nanohardness value is decreased, which is contributed to decrease of C–N bonding with Al content. On the other hand, when the
Fig. 7. The dependence of nano-hardness value on indentation depth in asdeposited and annealed Al3 C84 N13 thin films.
Fig. 8. The dependences of (a) hardness (b) elastic modulus values on annealing temperature in Al3 C84 N13 , Al13 C71 N16 and Al32 C43 N25 thin films.
annealing temperature increases nano-hardness values of the three films gradually decrease, which is mostly contributed to relaxation of the residual stress and decrease of nitrogen content (see Tables 2 and 3) during annealing. It is presumed that decrease of nitrogen content decreased the amount of C–N bonding, which resulted in decrease of hardness. It is worthy of noticing that less decrease in hardness value with the annealing temperature takes place for high Al contented thin films. A similar phenomenon in elastic modulus value is found in Fig. 8(b). Fig. 9 presents the load–displacement curve of Al3 C84 N13 in asdeposited sate and after annealed at 700 and 1000 ◦ C, respectively. It is found that for the present measurement the maximum inden-
Fig. 9. The dependence of load on indentation depth in as-deposited and annealed Al3 C84 N13 thin films.
6084
Q.H. Luo et al. / Applied Surface Science 257 (2011) 6079–6084
tation depth is about 65 nm, about one tenth of the film thickness. It is well known that the elastic recovery rate ER is calculated as follows: ER =
dmax − dres dmax
where dmax is the maximum indentation depth and dres is the residual indentation depth after unloading [17,18]. It is calculated that the elastic recovery rates of all three samples are about 69%, which reveals that annealing does not change the elastic recovery rate of Al–C–N film. 4. Conclusions In this study, three Al–C–N films with different Al contents were deposited on Si (1 0 0) substrates by closed-field unbalanced reactive magnetron sputtering. The phase configuration, microstructure and mechanical properties of three films in asdeposited and annealed states were subsequently studied. The main results are listed as follows: (1) The phase configuration and microstructure in both asdeposited and annealed Al–C–N thin films strongly depended on Al content. Usually, for Al–C–N films with low Al content delamination rather than crystallization occurred after the films were annealed at 1000 ◦ C. By comparison, for thin films at high Al content, annealing led to formation of AlN nanocrystallites, which produced nanocomposites of AlN embedded into amorphous matrices. (2) With increase of annealing temperature, both hardness and elastic modulus values of Al–C–N thin films were decreased, however less decrease was found for Al–C–N thin films at high
Al content. On the other hand, annealing does not change the elastic recovery rate of Al–C–N films. Acknowledgement The work described in this paper was supported by The National Nature Research Fund of China (No. 50871018). References [1] A.Y. Liu, M.L. Cohen, Science 245 (1989) 841. [2] M. Bai, K. Kato, N. Umehara, Y. Miyake, J. Xu, H. Tokisue, Thin Solid Films 376 (2000) 170. [3] X.C. Wang, Z.Q. Li, P. Wu, E.Y. Jiang, H.L. Bai, Appl. Surf. Sci. 253 (2006) 2087. [4] R. Prioli, S.I. Zanette, A.O. Caride, M.M. Lacerda, F.L. Freire Jr., Diamond Relat. Mater. 8 (1999) 993. [5] S. Ma, B. Xu, G. Wu, Y. Wang, F. Ma, D. Ma, K. Xu, T. Bell, Surf. Coat. Technol. 202 (2008) 5379. [6] D. Watanabe, H. Aoki, R. Moriyama, M.K. Mazumder, C. Kimura, T. Sugino, Diamond Relat. Mater. 17 (2008) 669. [7] Y.L. Yang, D. Zhang, W. Yan, Y.R. Zheng, Opt. Laser Eng. 48 (2010) 119. [8] J.L. Lin, J.J. Moore, B. Mishra, M. Pinkas, W.D. Sproul, Acta Mater. 58 (2010) 1554. [9] N. Jiang, S. Xu, K.N. Ostrikov, E.L. Tsakadze, J.D. Long, J.W. Chai, Z.L. Tsakadze, Int. J. Mod. Phys. B 16 (2002) 1132. [10] A.L. Ji, Y. Du, L.B. Ma, Z.X. Cao, J. Cryst. Growth 279 (2005) 420. [11] A.L. Ji, L.B. Ma, C. Liu, C.R. Li, Z.X. Cao, Diamond Relat. Mater. 14 (2005) 1348. [12] A.L. Ji, L.B. Ma, C. Liu, P. Zheng, C.R. Li, Z.X. Cao, Appl. Phys. Lett. 86 (2005) 021918. [13] V.V. Uglov, N.N. Cherenda, V.V. Khodasevich, V.A. Sokol, I.I. Abramov, A.L. Danilyuk, A. Wenzel, J. Gerlach, B. Rauschenbach, Nucl. Instrum. Methods Phys. Res. B 147 (1999) 332. [14] Q.H. Luo, D.L. Yu, Y.H. Lu, Chin. J. Vac. Sci. Technol. 30 (2010) 138. [15] M.J. Zhou, S.F. Wong, C.W. Ong, Q. Li, Thin Solid Films 516 (2007) 336. [16] D.G. McCulloch, X.L. Xiao, J.L. Peng, P.C.T. Ha, D.R. McKenzie, M.M.M. Bilek, S.P. Lau, D. Sheeja, B.K. Tay, Surf. Coat. Technol. 198 (2005) 217. [17] W.Y. Ni, Y.T. Cheng, D.S. Grummon, Appl. Phys. Lett. 80 (2002) 3311. [18] Y.H. Lu, Y.G. Shen, J.P. Wang, Z.F. Zhou, K.Y. Li, Surf. Coat. Technol. 201 (2007) 7368.