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Homoepitaxial growth of Ti–Si–C MAX-phase thin films on bulk Ti3SiC2 substrates
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Journal of Crystal Growth 304 (2007) 264–269 www.elsevier.com/locate/jcrysgro
Homoepitaxial growth of Ti–Si–C MAX-phase thin films on bulk Ti3SiC2 substrates P. Eklunda,, A. Murugaiahb,1, J. Emmerlicha,2, Zs. Cziga`nyc, J. Frodeliusa, M.W. Barsoumb, H. Ho¨gberga, L. Hultmana a
Thin Film Physics Division, Department of Physics, Chemistry, and Biology, IFM, Linko¨ping University, SE-581 83 Linko¨ping, Sweden b Department of Materials Science and Engineering, Drexel University, Philadelphia, PA 19104, USA c Research Institute for Technical Physics and Materials Science, Hungarian Academy of Sciences, P.O. Box 49, H-1525 Budapest, Hungary Received 22 December 2006; received in revised form 5 February 2007; accepted 20 February 2007 Communicated by D.W. Shaw Available online 7 March 2007
Abstract Ti3SiC2 films were grown on polycrystalline Ti3SiC2 bulk substrates using DC magnetron sputtering. The crystallographic orientation of the film grains is shown to be determined by the respective substrate-grain orientation through homoepitaxial MAX-phase growth. For a film composition close to Ti:Si:C ¼ 3:1:2, the films predominantly consist of MAX phases, both Ti3SiC2 and the metastable Ti4SiC3. Lower Si content resulted in growth of TiC with Ti3SiC2 as a minority phase. Thus, MAX-phase heterostructures with preferred crystallographic relationships can also be realized. r 2007 Elsevier B.V. All rights reserved. PACS: 81.05.Je; 81.07.b; 81.15.Kk; 61.10.Nz; 68.55.a Keywords: A1. Scanning electron microscopy; A1. Transmission electron microscopy; A1. X-ray diffraction; A3. Physical vapor deposition processes; B1. Carbides; B1. Nanomaterials
1. Introduction The fascinating properties of Ti3SiC2, combining typical ceramic and metallic attributes, has triggered intense research on this phase as well as the other members of the family of layered ternary carbides and nitrides known as the Mn+1AXn phases (n ¼ 1–3) [1]. Here, M is an early transition metal (e.g., Ti, Zr), A is an A-group element (e.g., Si, Al), and X is carbon and/or nitrogen. The majority of the 50 known MAX phases was originally synthesized and crystallographically determined by Nowotny et al. [2,3], who found that the MAX phases have a hexagonal crystal structure closely related to the interstitial Corresponding author.
E-mail address: [email protected] (P. Eklund). Present address: Momentive Performance Materials, 22557 West Lunn Road, Strongsville, OH 44149, USA. 2 Present address: Department of Materials Chemistry, RWTH Aachen, Kopernikusstrasse 16, D-52074 Aachen, Germany. 1
0022-0248/$ - see front matter r 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2007.02.014
transition metal carbides and nitrides (MX) [4]. The insertion of A-element layers into the MX skeleton results in highly anisotropic materials, with typical c/a ratios of 5–8. However, relatively little attention was paid to the consequences of this crystal structure to the macroscopic properties until the mid-1990s, when interest in the MAX phases was resurrected [5]. Progress in synthesis techniques for bulk Ti3SiC2 [6,7], at present permitting more than 99% phase-pure material [8,9] has been crucial for further exploration of their properties. Thin-film deposition using magnetron sputtering has been demonstrated as a method for single-phase epitaxial growth of Ti3SiC2 [10,11] with impact on characterization of, e.g., electrical properties [11,12], thermal stability [13,14], and mechanical properties [12,15]. Additionally, magnetron sputtering has been employed for growth of other M3AX2 phases, including Ti3GeC2 and Ti3AlC2 [16,17], as well as M2AX phases such as Ti2GeC [16], Ti2AlC [18,19], Cr2AlC [20,21], and Ti2AlN [22,23]. Most
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thin-film growth studies on MAX phases have focused on using single-crystal substrates, predominantly Al2O3(0 0 0 1) and MgO(1 1 1). For heteroepitaxial films, epitaxial relationships such as Al2O3(0 0 0 1)//TiC(1 1 1)// Ti3SiC2(0 0 0 1) (out-of-plane) have been demonstrated [10,11]. The TiC here is either a deliberately inserted seed layer, in order to facilitate epitaxial MAX-phase growth, or an incubation layer that forms during initial deposition. It is worth noting that similar epitaxial relations (without TiC interlayers) were recently demonstrated for Al2O3 formed on bulk Ti3AlC2 subjected to high-temperature oxidation [24]. A further important characteristic of the magnetronsputtering method is that it operates at relatively lowsubstrate temperatures (o900 1C) and at conditions far from thermodynamic equilibrium, and consequently permits synthesis of metastable phases such as Ti4SiC3 and Ti4GeC3 [10,16]. The objective of the present study is to explore the possibility of depositing MAX-phase thin films on MAXphase substrates, with the aim of improved understanding of the nucleation and growth mechanisms. We demonstrate growth of Ti3SiC2 thin films on polycrystalline Ti3SiC2 bulk substrates, as well as growth of the metastable phase Ti4SiC3.
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Cross-sectional samples for transmission electron microscopy (TEM) were mechanically polished to a thickness of 50 mm and then dimpled from both sides to 20 mm at the
Fig. 1. X-ray diffractograms from (a) a polycrystalline bulk Ti3SiC2 substrate, and Ti–Si–C films deposited with Si currents of (b) 26 mA, (c) 31 mA, and (d) 36 mA.
2. Experimental details Ti–Si–C MAX-phase thin films were synthesized by unbalanced (type II) DC magnetron sputtering, using an ultra-high vacuum (UHV) system previously described elsewhere [11]. All experiments were performed at a constant pressure of 5.3 103 mbar, using Ar (99.9999%) as discharge gas. High-purity targets of Ti and graphite (diameter 75 mm) and Si (diameter 25 mm) were used. The magnetrons were operated in currentregulation mode, using fixed target currents of 300 mA for Ti and C, known to result in a Ti:C ratio of 3:2, corresponding to the composition in Ti3SiC2 [11]. To investigate the effect of different Si contents, the Si target current was varied in the 26–36 mA range. The growth conditions were investigated at a substrate temperature of 900 1C on substrates of cut and polished sintered bulk Ti3SiC2 [5]. Prior to the depositions, the substrates were ultrasonically degreased in a three-step process by trichloroethylene, acetone and isopropanol, followed by an in situ heat treatment at 900 1C for 1 h (in vacuum). The deposition time was 5 h, corresponding to a nominal film thickness of 1 mm. The structural properties of the deposited films were investigated by grazing-incidence X-ray diffraction (GI-XRD) using a Philips diffractometer with a parallel-beam detector and a Cu Ka X-ray tube. For all GI-XRD measurements, the incidence angle of the primary X-ray beam was 31. Scanning electron microscopy (SEM) secondary-electron images were obtained in a LEO 1550 SEM operated at an acceleration voltage of 20 kV.
Fig. 2. SEM images of (a) polished polycrystalline bulk Ti3SiC2 substrate and (b) Ti–Si–C film deposited with 31 mA Si current. Inset: saw-tooth appearance from {1 1 2¯ 0} facets.
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center. The specimens were subsequently thinned to electron transparency using a Baltec RES010 ion miller operated at 10 kV using Ar+ ions, with a final polishing step at 3 kV. Bright-field TEM and high-resolution TEM (HREM) images were obtained in a Tecnai G2 20 U-Twin 200 kV FEGTEM. 3. Results and discussion Fig. 1 shows the X-ray diffractograms from (a) an uncoated Ti3SiC2 substrate, and (b)–(d) Ti–Si–C films deposited with Si currents of 26, 31, and 36 mA, respectively. In Fig. 1(a), diffraction peaks originating from different Ti3SiC2 orientations can be observed (0 0 0 ‘ and 1 0 1¯ ‘ peaks), i.e., the material is polycrystalline. The 1 1 1 and 2 0 0 peaks from TiC can also be seen. The majority of these peaks are present in Figs. 1(b)–(d) as well; however, there are a number of important differences. For the sample deposited at 26 mA Si current (Fig. 1(b)), the relative intensities of the 1 0 1¯ ‘ and 0 0 0 ‘ peaks are comparable to the substrate peaks, with the exceptions of the 0 0 0 8 peak (40.91), which is weaker, and the 1 0 1¯ 2 peak (35.21), which is stronger. The most striking difference, however, is apparent for the sample deposited at 31 mA Si current (Fig. 1(c)). In addition to the peaks originating from Ti3SiC2 (and TiC), the diffractogram in Fig. 1(c) contains a relatively strong peak at the 2y angle 7.51, and weaker peaks at 15.51, 231, and 321. These four peaks originate from (0 0 0 ‘) planes of the metastable thinfilm phase Ti4SiC3 [25], specifically the (0 0 0 2) to (0 0 0 8) planes. Additionally, the peak intensities of the Ti3SiC2 peaks are different from the ones in the diffractogram from the substrate. The Ti3SiC2 0 0 0 ‘ peaks are of two to four times higher intensity relative to the 1 0 1¯ 4 peak (39.61), and the 1 0 1¯ 1 (34.11) and 1 0 1¯ 5 peaks (42.51) are even stronger. The appearance of Ti4SiC3 peaks and the large differences in Ti3SiC2 peak intensities compared to the substrate show that MAX-phase thin-film growth has occurred, both of Ti4SiC3 and Ti3SiC2. For the sample deposited at higher Si flux (36 mA Si current, Fig. 1(d)), it can be observed that the Ti3SiC2 0 0 0 ‘ peaks are instead very weak; the 1 0 1¯ ‘-peak intensities, on the other hand, are relatively close to those in Fig. 1(a). In the following, we investigate the nucleation and growth mechanisms of Ti3SiC2 and Ti4SiC3 films on Ti3SiC2 bulk substrates. Fig. 2(a) shows an SEM image of an uncoated Ti3SiC2 substrate, with a flat, polished surface, and grain sizes of the order of 10 mm. Fig. 2(b) shows an SEM image of the film deposited at 31 mA Si current. It can be seen that the granular structure is retained for the as-deposited film; different growth textures were obtained, however, as represented by the different grains labeled I–V in Fig. 2(b). The surface of grain I is flat due to [0 0 0 2]-oriented growth, as explained below. For the grains labeled II–IV, 0002 terraces with growth steps are evident, with decreasing observed spacing in the order II–IV. A similar growth mode was found [11] for growth of
Ti3SiC2(0 0 0 1) on on-axis Al2O3(0 0 0 1) substrates. In that case, however, the steps were provided by threading screw dislocations. The growth steps in the present films are presumably normal to [0 0 0 ‘] with a ½1 0 1¯ ‘ surface normal. Also, since the cut grains of the present substrates provide geometrically determined steps, there is no need for screw dislocations to drive the growth of Ti3SiC2 in the present case. Thus, MAX-phase growth with tilted basal planes occurs. The tilt angle of the basal planes in the film with respect to the substrate surface corresponds to the orientation of the underlying substrate grain, with grain IV having the largest basal-plane tilt angle. In contrast, grain V exhibits surface islands with intersecting growth steps indicative of tilt angles far away from the [0 0 0 1] zone axis.
Fig. 3. TEM overview images of (a) [0 0 0 ‘]-oriented and (b) tilted MAXphase basal planes in Ti–Si–C film grown at 31 mA Si current.
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The growth steps exhibit a saw-tooth appearance (magnified inset in Fig. 2(b)) from {1 1 2¯ 0} facets. The surface pores and holes in the substrate are retained throughout the film. Figs. 3(a) and (b) show two typical TEM overview images of the film deposited using a Si current of 31 mA. Fig. 3(a) shows a grain where c-axis-oriented growth of Ti3SiC2 has occurred, with the basal planes parallel to the substrate surface. Together with the observations above regarding Fig. 2, this demonstrates that homoepitaxial growth of Ti3SiC2 is possible. Fig. 3(b) shows tilted MAX-phase basal planes with an orientation corresponding to the substrate-grain orientation; compare Fig. 2(b). Given their anisotropy, MAX phases tend to grow fastest on f1 1 2¯ 0g and f1 1¯ 0 0g planes [26]. Thus, a distinct morphology with large facets of {0 0 0 1} low-energy planes is formed. The Ti3SiC2{0 0 0 1} planes are epitaxially related to the TiC{1 1 1} planes, and the corresponding interfaces are coherent [27]. This follows from the isomorphism of the TiC(1 1 1) and the Ti3SiC2(0 0 0 1)
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low-energy surfaces and the corresponding strong tendency for basal-plane-oriented growth of Ti3SiC2 on TiC{1 1 1} planes. The growth orientation of the MAX-phase film is therefore governed by the orientation of the substrate grain, upon which it is grown, i.e., we observe tilted basalplane growth following the MAX-phase substrate-grain orientation. Growth voids originating from self-shadowing of film surface facets during deposition are also seen in Fig. 3(b). In particular, voids form when tilted basal planes grow into each other. Similar observations have been made for Ti3SiC2 films grown on MgO(1 0 0) substrates [11]. Thus, we deduce that Ti3SiC2 grown on TiC will be oriented in such a way that the epitaxial relation TiC(1 1 1)// Ti3SiC2(0 0 0 1) is maintained. The present findings, together with related observations for Ti2AlN [23] imply that similar crystallographic-orientation relationships should be common for other MAX phases with respect to their MX counterparts as well as isostructural materials (e.g., MgO).
Fig. 4. HREM images of the sample grown at 31 mA Si current of (a) Ti3SiC2 film on substrate TiC grain, (b) magnification of interface between TiC and intersecting Ti3SiC2, (c) an area further away from the substrate showing intersection between Ti3SiC2, TiC, and Ti4SiC3, and (d) magnified area showing intergrowth of Ti3SiC2, TiC, and Ti4SiC3.
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Fig. 4(a) shows a HREM image from a Ti3SiC2 film on a substrate TiC grain, where Ti3SiC2 has nucleated and grown in two different orientations. Fig. 4(b) is a magnified image of the interface between TiC and two intersecting Ti3SiC2 grains. On TiC, Ti3SiC2 has the possibility to grow in equivalent TiC/1 1 1S directions, which is why growth of tilted basal planes in different directions is observed. Fig. 4(c) shows a HREM image from the same sample further away from the substrate. Fig. 4(d) shows a magnified section of Fig. 4(c), where intergrowth of the phases Ti3SiC2, TiC, and Ti4SiC3 can be observed. In this area, two tilted basal planes of Ti3SiC2 met without void formation (lower half of Fig. 4(c)), instead TiC formed. After a brief occurrence of TiC, Ti4SiC3 nucleated. Figs. 5(a) and (b) show SEM images from Ti–Si–C films deposited with Si currents of (a) 26 mA and (b) 36 mA. For the low-current conditions, MAX phases with their characteristic basal planes do not generally appear. Instead, growth of TiC is preferred (cf. Fig. 6). In Fig. 5(b), the film with the higher Si content shows a number of grains where characteristic tilted MAX-phase basal planes can be seen, however, to a much lesser extent than the film deposited at 31 mA.
Fig. 6 shows a typical HREM image from the film deposited using a Si current of 26 mA. The volume fraction of TiC is much larger than in the films with higher Si content; Ti3SiC2 is a minority phase. It is thus clear that the Si content in the growth flux has a strong influence on the nucleation and growth of the different phases. As discussed above, the 31 mA Si film predominantly consists of MAX phases, both Ti3SiC2 and Ti4SiC3, with TiC present as a minority phase. This indicates that the Si content is slightly lower than the ideal Ti:Si:C ratio of 3:1:2. As for the highSi-current conditions, based on the phase diagram and previous thin-film growth studies [11], it might be expected that a higher Si content (i.e., 36 mA Si current) would lead to formation of Ti3SiC2 together with the Si-rich Nowotny phase Ti5Si3Cx. The latter phase was, however, not observed in the present samples. This may be because of the substrate temperature of 900 1C, which is considerably
Fig. 5. SEM images of Ti–Si–C films deposited with Si currents of (a) 26 mA and (b) 36 mA.
Fig. 6. HREM image of a Ti–Si–C film deposited with Si current of 26 mA, showing Ti3SiC2 as a minority phase besides TiC.
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higher than expected for formation of Ti5Si3Cx. Note, however, that there is significant overlap between XRD peaks of Ti3SiC2 and Ti5Si3Cx. It is nevertheless clear that the large amount of Si is a limiting factor for growth of Ti3SiC2, as seen in Fig. 5(b). We have thus demonstrated a growth window (with respect to the Si flux) in which thinfilm Ti3SiC2 can be grown homoepitaxially on bulk Ti3SiC2. 4. Conclusions Ti3SiC2 films can be grown on polycrystalline Ti3SiC2 bulk substrates using DC magnetron sputtering. The crystallographic orientation of the film grains is determined by the respective substrate-grain orientation, i.e., homoepitaxial MAX-phase growth. Low Si content results in growth of TiC with Ti3SiC2 as a minority phase. For higher-Si content, the films predominantly consist of MAX phases, both Ti3SiC2 and the metastable Ti4SiC3. Excess Si is a limiting factor for Ti3SiC2 growth. We have thus demonstrated a growth window (with respect to the Si flux) in which thin-film Ti3SiC2 can be grown homoepitaxially on bulk Ti3SiC2. Acknowledgments We acknowledge the Swedish Agency for Innovation Systems (VINNOVA), the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the National Science Foundation (DMR-0503711), ABB Ltd., and Kanthal AB for support. J. E., Zs. C., and L. H. also acknowledge support from the EU Commission Integrated Project FOREMOST NMP3-CT-2005-515840. References [1] M.W. Barsoum, Prog. Solid State Chem. 28 (2000) 201. [2] H. Nowotny, Prog. Solid State Chem. 2 (1970) 27. [3] H. Wolfsgruber, H. Nowotny, F. Benesovsky, Monatsh. Chem. 98 (1967) 2403. [4] W. Jeitschko, H. Nowotny, Monatsh. Chem. 98 (1967) 329.
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[5] M.W. Barsoum, T. El-Raghy, J. Am. Ceram. Soc. 79 (1996) 1953. [6] S.B. Li, J.X. Xie, L.I. Zhang, L.F. Zheng, Mater. Sci. Technol. 19 (2003) 1442. [7] V. Gauthier, B. Cochepin, S. Dubois, D. Vrel, J. Am. Ceram. Soc. 89 (2006) 2899. [8] S. Yang, Z.M. Sun, H. Hashimoto, T. Abe, J. Eur. Ceram. Soc. 23 (2003) 3147. [9] J.-O. Zhu, B.-C. Mei, X.W. Xu, J. Liu, Scripta Mater. 49 (2003) 693. [10] J.-P. Palmquist, U. Jansson, T. Seppa¨nen, P.O.A˚. Persson, J. Birch, L. Hultman, P. Isberg, Appl. Phys. Lett. 81 (2002) 835. [11] J. Emmerlich, J.-P. Palmquist, H. Ho¨gberg, J.M. Molina-Aldareguia, Zs. Cziga´ny, Sz. Sasva´ri, P.O.A˚. Persson, U. Jansson, L. Hultman, J. Appl. Phys. 96 (2004) 4817. [12] J. Emmerlich, Ph. D. Thesis, Linko¨ping University, Linko¨ping studies in science and technology no. 1024, ISBN:91-85523-64-X. [13] P. Eklund, C. Virojanadara, J. Emmerlich, L.I. Johansson, H. Ho¨gberg, L. Hultman, Phys. Rev. B 74 (2006) 045417. [14] J. Emmerlich, D. Music, P. Eklund, O. Wilhelmsson, U. Jansson, J.M. Schneider, H. Ho¨gberg, L. Hultman, Acta Mater. 55 (2007) 1479. [15] J.M. Molina-Aldareguia, J. Emmerlich, J.-P. Palmquist, U. Jansson, L. Hultman, Scripta Mater. 49 (2003) 155. [16] H. Ho¨gberg, P. Eklund, J. Emmerlich, J. Birch, L. Hultman, J. Mater. Res. 20 (2005) 779. [17] O. Wilhelmsson, J.-P. Palmquist, T. Nyberg, U. Jansson, Appl. Phys. Lett. 85 (2004) 1066. [18] C. Walter, C. Martinez, T. El-Raghy, J.M. Schneider, Steel Res. Int. 76 (2005) 225. [19] O. Wilhelmsson, J.-P. Palmquist, E. Lewin, J. Emmerlich, P. Eklund, P.O.A˚. Persson, H. Ho¨gberg, S. Li, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, J. Crystal Growth 291 (2006) 290. [20] R. Martens, Z. Sun, D. Music, J.M. Schneider, Adv. Eng. Mater. 6 (2004) 903. [21] C. Walter, D.P. Sigumonrong, T. El-Raghy, J.M. Schneider, Thin Solid Films 515 (2006) 389. [22] T. Joelsson, A. Ho¨rling, J. Birch, L. Hultman, Appl. Phys. Lett. 86 (2005) 111913. [23] M. Beckers, N. Schell, R.M.S. Martins, A. Mu¨cklich, W. Mo¨ller, L. Hultman, J. Appl. Phys. 99 (2006) 34902. [24] Z.J. Lin, M.J. Zhuo, T.C. Zhou, M.S. Li, J.Y. Wang, Scripta Mater. 54 (2006) 1815. [25] J.-P. Palmquist, S. Li, P.O.A˚. Persson, J. Emmerlich, O. Wilhelmsson, H. Ho¨gberg, M.I. Katsnelson, B. Johansson, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, Phys. Rev. B 70 (2004) 165401. [26] T. El-Raghy, M.W. Barsoum, J. Am. Ceram. Soc. 82 (1999) 2849. [27] Z.J. Lin, M.J. Zhuo, Y.C. Zhou, M.S. Li, J.Y. Wang, Scripta Mater. 55 (2006) 445.
Journal of Crystal Growth 304 (2007) 264–269 www.elsevier.com/locate/jcrysgro
Homoepitaxial growth of Ti–Si–C MAX-phase thin films on bulk Ti3SiC2 substrates P. Eklunda,, A. Murugaiahb,1, J. Emmerlicha,2, Zs. Cziga`nyc, J. Frodeliusa, M.W. Barsoumb, H. Ho¨gberga, L. Hultmana a
Thin Film Physics Division, Department of Physics, Chemistry, and Biology, IFM, Linko¨ping University, SE-581 83 Linko¨ping, Sweden b Department of Materials Science and Engineering, Drexel University, Philadelphia, PA 19104, USA c Research Institute for Technical Physics and Materials Science, Hungarian Academy of Sciences, P.O. Box 49, H-1525 Budapest, Hungary Received 22 December 2006; received in revised form 5 February 2007; accepted 20 February 2007 Communicated by D.W. Shaw Available online 7 March 2007
Abstract Ti3SiC2 films were grown on polycrystalline Ti3SiC2 bulk substrates using DC magnetron sputtering. The crystallographic orientation of the film grains is shown to be determined by the respective substrate-grain orientation through homoepitaxial MAX-phase growth. For a film composition close to Ti:Si:C ¼ 3:1:2, the films predominantly consist of MAX phases, both Ti3SiC2 and the metastable Ti4SiC3. Lower Si content resulted in growth of TiC with Ti3SiC2 as a minority phase. Thus, MAX-phase heterostructures with preferred crystallographic relationships can also be realized. r 2007 Elsevier B.V. All rights reserved. PACS: 81.05.Je; 81.07.b; 81.15.Kk; 61.10.Nz; 68.55.a Keywords: A1. Scanning electron microscopy; A1. Transmission electron microscopy; A1. X-ray diffraction; A3. Physical vapor deposition processes; B1. Carbides; B1. Nanomaterials
1. Introduction The fascinating properties of Ti3SiC2, combining typical ceramic and metallic attributes, has triggered intense research on this phase as well as the other members of the family of layered ternary carbides and nitrides known as the Mn+1AXn phases (n ¼ 1–3) [1]. Here, M is an early transition metal (e.g., Ti, Zr), A is an A-group element (e.g., Si, Al), and X is carbon and/or nitrogen. The majority of the 50 known MAX phases was originally synthesized and crystallographically determined by Nowotny et al. [2,3], who found that the MAX phases have a hexagonal crystal structure closely related to the interstitial Corresponding author.
E-mail address: [email protected] (P. Eklund). Present address: Momentive Performance Materials, 22557 West Lunn Road, Strongsville, OH 44149, USA. 2 Present address: Department of Materials Chemistry, RWTH Aachen, Kopernikusstrasse 16, D-52074 Aachen, Germany. 1
0022-0248/$ - see front matter r 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2007.02.014
transition metal carbides and nitrides (MX) [4]. The insertion of A-element layers into the MX skeleton results in highly anisotropic materials, with typical c/a ratios of 5–8. However, relatively little attention was paid to the consequences of this crystal structure to the macroscopic properties until the mid-1990s, when interest in the MAX phases was resurrected [5]. Progress in synthesis techniques for bulk Ti3SiC2 [6,7], at present permitting more than 99% phase-pure material [8,9] has been crucial for further exploration of their properties. Thin-film deposition using magnetron sputtering has been demonstrated as a method for single-phase epitaxial growth of Ti3SiC2 [10,11] with impact on characterization of, e.g., electrical properties [11,12], thermal stability [13,14], and mechanical properties [12,15]. Additionally, magnetron sputtering has been employed for growth of other M3AX2 phases, including Ti3GeC2 and Ti3AlC2 [16,17], as well as M2AX phases such as Ti2GeC [16], Ti2AlC [18,19], Cr2AlC [20,21], and Ti2AlN [22,23]. Most
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thin-film growth studies on MAX phases have focused on using single-crystal substrates, predominantly Al2O3(0 0 0 1) and MgO(1 1 1). For heteroepitaxial films, epitaxial relationships such as Al2O3(0 0 0 1)//TiC(1 1 1)// Ti3SiC2(0 0 0 1) (out-of-plane) have been demonstrated [10,11]. The TiC here is either a deliberately inserted seed layer, in order to facilitate epitaxial MAX-phase growth, or an incubation layer that forms during initial deposition. It is worth noting that similar epitaxial relations (without TiC interlayers) were recently demonstrated for Al2O3 formed on bulk Ti3AlC2 subjected to high-temperature oxidation [24]. A further important characteristic of the magnetronsputtering method is that it operates at relatively lowsubstrate temperatures (o900 1C) and at conditions far from thermodynamic equilibrium, and consequently permits synthesis of metastable phases such as Ti4SiC3 and Ti4GeC3 [10,16]. The objective of the present study is to explore the possibility of depositing MAX-phase thin films on MAXphase substrates, with the aim of improved understanding of the nucleation and growth mechanisms. We demonstrate growth of Ti3SiC2 thin films on polycrystalline Ti3SiC2 bulk substrates, as well as growth of the metastable phase Ti4SiC3.
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Cross-sectional samples for transmission electron microscopy (TEM) were mechanically polished to a thickness of 50 mm and then dimpled from both sides to 20 mm at the
Fig. 1. X-ray diffractograms from (a) a polycrystalline bulk Ti3SiC2 substrate, and Ti–Si–C films deposited with Si currents of (b) 26 mA, (c) 31 mA, and (d) 36 mA.
2. Experimental details Ti–Si–C MAX-phase thin films were synthesized by unbalanced (type II) DC magnetron sputtering, using an ultra-high vacuum (UHV) system previously described elsewhere [11]. All experiments were performed at a constant pressure of 5.3 103 mbar, using Ar (99.9999%) as discharge gas. High-purity targets of Ti and graphite (diameter 75 mm) and Si (diameter 25 mm) were used. The magnetrons were operated in currentregulation mode, using fixed target currents of 300 mA for Ti and C, known to result in a Ti:C ratio of 3:2, corresponding to the composition in Ti3SiC2 [11]. To investigate the effect of different Si contents, the Si target current was varied in the 26–36 mA range. The growth conditions were investigated at a substrate temperature of 900 1C on substrates of cut and polished sintered bulk Ti3SiC2 [5]. Prior to the depositions, the substrates were ultrasonically degreased in a three-step process by trichloroethylene, acetone and isopropanol, followed by an in situ heat treatment at 900 1C for 1 h (in vacuum). The deposition time was 5 h, corresponding to a nominal film thickness of 1 mm. The structural properties of the deposited films were investigated by grazing-incidence X-ray diffraction (GI-XRD) using a Philips diffractometer with a parallel-beam detector and a Cu Ka X-ray tube. For all GI-XRD measurements, the incidence angle of the primary X-ray beam was 31. Scanning electron microscopy (SEM) secondary-electron images were obtained in a LEO 1550 SEM operated at an acceleration voltage of 20 kV.
Fig. 2. SEM images of (a) polished polycrystalline bulk Ti3SiC2 substrate and (b) Ti–Si–C film deposited with 31 mA Si current. Inset: saw-tooth appearance from {1 1 2¯ 0} facets.
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center. The specimens were subsequently thinned to electron transparency using a Baltec RES010 ion miller operated at 10 kV using Ar+ ions, with a final polishing step at 3 kV. Bright-field TEM and high-resolution TEM (HREM) images were obtained in a Tecnai G2 20 U-Twin 200 kV FEGTEM. 3. Results and discussion Fig. 1 shows the X-ray diffractograms from (a) an uncoated Ti3SiC2 substrate, and (b)–(d) Ti–Si–C films deposited with Si currents of 26, 31, and 36 mA, respectively. In Fig. 1(a), diffraction peaks originating from different Ti3SiC2 orientations can be observed (0 0 0 ‘ and 1 0 1¯ ‘ peaks), i.e., the material is polycrystalline. The 1 1 1 and 2 0 0 peaks from TiC can also be seen. The majority of these peaks are present in Figs. 1(b)–(d) as well; however, there are a number of important differences. For the sample deposited at 26 mA Si current (Fig. 1(b)), the relative intensities of the 1 0 1¯ ‘ and 0 0 0 ‘ peaks are comparable to the substrate peaks, with the exceptions of the 0 0 0 8 peak (40.91), which is weaker, and the 1 0 1¯ 2 peak (35.21), which is stronger. The most striking difference, however, is apparent for the sample deposited at 31 mA Si current (Fig. 1(c)). In addition to the peaks originating from Ti3SiC2 (and TiC), the diffractogram in Fig. 1(c) contains a relatively strong peak at the 2y angle 7.51, and weaker peaks at 15.51, 231, and 321. These four peaks originate from (0 0 0 ‘) planes of the metastable thinfilm phase Ti4SiC3 [25], specifically the (0 0 0 2) to (0 0 0 8) planes. Additionally, the peak intensities of the Ti3SiC2 peaks are different from the ones in the diffractogram from the substrate. The Ti3SiC2 0 0 0 ‘ peaks are of two to four times higher intensity relative to the 1 0 1¯ 4 peak (39.61), and the 1 0 1¯ 1 (34.11) and 1 0 1¯ 5 peaks (42.51) are even stronger. The appearance of Ti4SiC3 peaks and the large differences in Ti3SiC2 peak intensities compared to the substrate show that MAX-phase thin-film growth has occurred, both of Ti4SiC3 and Ti3SiC2. For the sample deposited at higher Si flux (36 mA Si current, Fig. 1(d)), it can be observed that the Ti3SiC2 0 0 0 ‘ peaks are instead very weak; the 1 0 1¯ ‘-peak intensities, on the other hand, are relatively close to those in Fig. 1(a). In the following, we investigate the nucleation and growth mechanisms of Ti3SiC2 and Ti4SiC3 films on Ti3SiC2 bulk substrates. Fig. 2(a) shows an SEM image of an uncoated Ti3SiC2 substrate, with a flat, polished surface, and grain sizes of the order of 10 mm. Fig. 2(b) shows an SEM image of the film deposited at 31 mA Si current. It can be seen that the granular structure is retained for the as-deposited film; different growth textures were obtained, however, as represented by the different grains labeled I–V in Fig. 2(b). The surface of grain I is flat due to [0 0 0 2]-oriented growth, as explained below. For the grains labeled II–IV, 0002 terraces with growth steps are evident, with decreasing observed spacing in the order II–IV. A similar growth mode was found [11] for growth of
Ti3SiC2(0 0 0 1) on on-axis Al2O3(0 0 0 1) substrates. In that case, however, the steps were provided by threading screw dislocations. The growth steps in the present films are presumably normal to [0 0 0 ‘] with a ½1 0 1¯ ‘ surface normal. Also, since the cut grains of the present substrates provide geometrically determined steps, there is no need for screw dislocations to drive the growth of Ti3SiC2 in the present case. Thus, MAX-phase growth with tilted basal planes occurs. The tilt angle of the basal planes in the film with respect to the substrate surface corresponds to the orientation of the underlying substrate grain, with grain IV having the largest basal-plane tilt angle. In contrast, grain V exhibits surface islands with intersecting growth steps indicative of tilt angles far away from the [0 0 0 1] zone axis.
Fig. 3. TEM overview images of (a) [0 0 0 ‘]-oriented and (b) tilted MAXphase basal planes in Ti–Si–C film grown at 31 mA Si current.
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The growth steps exhibit a saw-tooth appearance (magnified inset in Fig. 2(b)) from {1 1 2¯ 0} facets. The surface pores and holes in the substrate are retained throughout the film. Figs. 3(a) and (b) show two typical TEM overview images of the film deposited using a Si current of 31 mA. Fig. 3(a) shows a grain where c-axis-oriented growth of Ti3SiC2 has occurred, with the basal planes parallel to the substrate surface. Together with the observations above regarding Fig. 2, this demonstrates that homoepitaxial growth of Ti3SiC2 is possible. Fig. 3(b) shows tilted MAX-phase basal planes with an orientation corresponding to the substrate-grain orientation; compare Fig. 2(b). Given their anisotropy, MAX phases tend to grow fastest on f1 1 2¯ 0g and f1 1¯ 0 0g planes [26]. Thus, a distinct morphology with large facets of {0 0 0 1} low-energy planes is formed. The Ti3SiC2{0 0 0 1} planes are epitaxially related to the TiC{1 1 1} planes, and the corresponding interfaces are coherent [27]. This follows from the isomorphism of the TiC(1 1 1) and the Ti3SiC2(0 0 0 1)
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low-energy surfaces and the corresponding strong tendency for basal-plane-oriented growth of Ti3SiC2 on TiC{1 1 1} planes. The growth orientation of the MAX-phase film is therefore governed by the orientation of the substrate grain, upon which it is grown, i.e., we observe tilted basalplane growth following the MAX-phase substrate-grain orientation. Growth voids originating from self-shadowing of film surface facets during deposition are also seen in Fig. 3(b). In particular, voids form when tilted basal planes grow into each other. Similar observations have been made for Ti3SiC2 films grown on MgO(1 0 0) substrates [11]. Thus, we deduce that Ti3SiC2 grown on TiC will be oriented in such a way that the epitaxial relation TiC(1 1 1)// Ti3SiC2(0 0 0 1) is maintained. The present findings, together with related observations for Ti2AlN [23] imply that similar crystallographic-orientation relationships should be common for other MAX phases with respect to their MX counterparts as well as isostructural materials (e.g., MgO).
Fig. 4. HREM images of the sample grown at 31 mA Si current of (a) Ti3SiC2 film on substrate TiC grain, (b) magnification of interface between TiC and intersecting Ti3SiC2, (c) an area further away from the substrate showing intersection between Ti3SiC2, TiC, and Ti4SiC3, and (d) magnified area showing intergrowth of Ti3SiC2, TiC, and Ti4SiC3.
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Fig. 4(a) shows a HREM image from a Ti3SiC2 film on a substrate TiC grain, where Ti3SiC2 has nucleated and grown in two different orientations. Fig. 4(b) is a magnified image of the interface between TiC and two intersecting Ti3SiC2 grains. On TiC, Ti3SiC2 has the possibility to grow in equivalent TiC/1 1 1S directions, which is why growth of tilted basal planes in different directions is observed. Fig. 4(c) shows a HREM image from the same sample further away from the substrate. Fig. 4(d) shows a magnified section of Fig. 4(c), where intergrowth of the phases Ti3SiC2, TiC, and Ti4SiC3 can be observed. In this area, two tilted basal planes of Ti3SiC2 met without void formation (lower half of Fig. 4(c)), instead TiC formed. After a brief occurrence of TiC, Ti4SiC3 nucleated. Figs. 5(a) and (b) show SEM images from Ti–Si–C films deposited with Si currents of (a) 26 mA and (b) 36 mA. For the low-current conditions, MAX phases with their characteristic basal planes do not generally appear. Instead, growth of TiC is preferred (cf. Fig. 6). In Fig. 5(b), the film with the higher Si content shows a number of grains where characteristic tilted MAX-phase basal planes can be seen, however, to a much lesser extent than the film deposited at 31 mA.
Fig. 6 shows a typical HREM image from the film deposited using a Si current of 26 mA. The volume fraction of TiC is much larger than in the films with higher Si content; Ti3SiC2 is a minority phase. It is thus clear that the Si content in the growth flux has a strong influence on the nucleation and growth of the different phases. As discussed above, the 31 mA Si film predominantly consists of MAX phases, both Ti3SiC2 and Ti4SiC3, with TiC present as a minority phase. This indicates that the Si content is slightly lower than the ideal Ti:Si:C ratio of 3:1:2. As for the highSi-current conditions, based on the phase diagram and previous thin-film growth studies [11], it might be expected that a higher Si content (i.e., 36 mA Si current) would lead to formation of Ti3SiC2 together with the Si-rich Nowotny phase Ti5Si3Cx. The latter phase was, however, not observed in the present samples. This may be because of the substrate temperature of 900 1C, which is considerably
Fig. 5. SEM images of Ti–Si–C films deposited with Si currents of (a) 26 mA and (b) 36 mA.
Fig. 6. HREM image of a Ti–Si–C film deposited with Si current of 26 mA, showing Ti3SiC2 as a minority phase besides TiC.
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higher than expected for formation of Ti5Si3Cx. Note, however, that there is significant overlap between XRD peaks of Ti3SiC2 and Ti5Si3Cx. It is nevertheless clear that the large amount of Si is a limiting factor for growth of Ti3SiC2, as seen in Fig. 5(b). We have thus demonstrated a growth window (with respect to the Si flux) in which thinfilm Ti3SiC2 can be grown homoepitaxially on bulk Ti3SiC2. 4. Conclusions Ti3SiC2 films can be grown on polycrystalline Ti3SiC2 bulk substrates using DC magnetron sputtering. The crystallographic orientation of the film grains is determined by the respective substrate-grain orientation, i.e., homoepitaxial MAX-phase growth. Low Si content results in growth of TiC with Ti3SiC2 as a minority phase. For higher-Si content, the films predominantly consist of MAX phases, both Ti3SiC2 and the metastable Ti4SiC3. Excess Si is a limiting factor for Ti3SiC2 growth. We have thus demonstrated a growth window (with respect to the Si flux) in which thin-film Ti3SiC2 can be grown homoepitaxially on bulk Ti3SiC2. Acknowledgments We acknowledge the Swedish Agency for Innovation Systems (VINNOVA), the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the National Science Foundation (DMR-0503711), ABB Ltd., and Kanthal AB for support. J. E., Zs. C., and L. H. also acknowledge support from the EU Commission Integrated Project FOREMOST NMP3-CT-2005-515840. References [1] M.W. Barsoum, Prog. Solid State Chem. 28 (2000) 201. [2] H. Nowotny, Prog. Solid State Chem. 2 (1970) 27. [3] H. Wolfsgruber, H. Nowotny, F. Benesovsky, Monatsh. Chem. 98 (1967) 2403. [4] W. Jeitschko, H. Nowotny, Monatsh. Chem. 98 (1967) 329.
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[5] M.W. Barsoum, T. El-Raghy, J. Am. Ceram. Soc. 79 (1996) 1953. [6] S.B. Li, J.X. Xie, L.I. Zhang, L.F. Zheng, Mater. Sci. Technol. 19 (2003) 1442. [7] V. Gauthier, B. Cochepin, S. Dubois, D. Vrel, J. Am. Ceram. Soc. 89 (2006) 2899. [8] S. Yang, Z.M. Sun, H. Hashimoto, T. Abe, J. Eur. Ceram. Soc. 23 (2003) 3147. [9] J.-O. Zhu, B.-C. Mei, X.W. Xu, J. Liu, Scripta Mater. 49 (2003) 693. [10] J.-P. Palmquist, U. Jansson, T. Seppa¨nen, P.O.A˚. Persson, J. Birch, L. Hultman, P. Isberg, Appl. Phys. Lett. 81 (2002) 835. [11] J. Emmerlich, J.-P. Palmquist, H. Ho¨gberg, J.M. Molina-Aldareguia, Zs. Cziga´ny, Sz. Sasva´ri, P.O.A˚. Persson, U. Jansson, L. Hultman, J. Appl. Phys. 96 (2004) 4817. [12] J. Emmerlich, Ph. D. Thesis, Linko¨ping University, Linko¨ping studies in science and technology no. 1024, ISBN:91-85523-64-X. [13] P. Eklund, C. Virojanadara, J. Emmerlich, L.I. Johansson, H. Ho¨gberg, L. Hultman, Phys. Rev. B 74 (2006) 045417. [14] J. Emmerlich, D. Music, P. Eklund, O. Wilhelmsson, U. Jansson, J.M. Schneider, H. Ho¨gberg, L. Hultman, Acta Mater. 55 (2007) 1479. [15] J.M. Molina-Aldareguia, J. Emmerlich, J.-P. Palmquist, U. Jansson, L. Hultman, Scripta Mater. 49 (2003) 155. [16] H. Ho¨gberg, P. Eklund, J. Emmerlich, J. Birch, L. Hultman, J. Mater. Res. 20 (2005) 779. [17] O. Wilhelmsson, J.-P. Palmquist, T. Nyberg, U. Jansson, Appl. Phys. Lett. 85 (2004) 1066. [18] C. Walter, C. Martinez, T. El-Raghy, J.M. Schneider, Steel Res. Int. 76 (2005) 225. [19] O. Wilhelmsson, J.-P. Palmquist, E. Lewin, J. Emmerlich, P. Eklund, P.O.A˚. Persson, H. Ho¨gberg, S. Li, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, J. Crystal Growth 291 (2006) 290. [20] R. Martens, Z. Sun, D. Music, J.M. Schneider, Adv. Eng. Mater. 6 (2004) 903. [21] C. Walter, D.P. Sigumonrong, T. El-Raghy, J.M. Schneider, Thin Solid Films 515 (2006) 389. [22] T. Joelsson, A. Ho¨rling, J. Birch, L. Hultman, Appl. Phys. Lett. 86 (2005) 111913. [23] M. Beckers, N. Schell, R.M.S. Martins, A. Mu¨cklich, W. Mo¨ller, L. Hultman, J. Appl. Phys. 99 (2006) 34902. [24] Z.J. Lin, M.J. Zhuo, T.C. Zhou, M.S. Li, J.Y. Wang, Scripta Mater. 54 (2006) 1815. [25] J.-P. Palmquist, S. Li, P.O.A˚. Persson, J. Emmerlich, O. Wilhelmsson, H. Ho¨gberg, M.I. Katsnelson, B. Johansson, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, Phys. Rev. B 70 (2004) 165401. [26] T. El-Raghy, M.W. Barsoum, J. Am. Ceram. Soc. 82 (1999) 2849. [27] Z.J. Lin, M.J. Zhuo, Y.C. Zhou, M.S. Li, J.Y. Wang, Scripta Mater. 55 (2006) 445.