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Si interface chemistry and epitaxial growth mode
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Acta Materialia 57 (2009) 232–236 www.elsevier.com/locate/actamat
NiSi2/Si interface chemistry and epitaxial growth mode S.B. Mi a,*, C.L. Jia a, Q.T. Zhao b, S. Mantl b, K. Urban a a
Institute of Solid State Research and Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons (ER-C), Research Centre Juelich, D-52425 Juelich, Germany b Institute of Bio- and Nanosystems (IBN1), and Center of Nanoelectronic Systems for Information Technology (CNI), Forschungszentrum Ju¨lich GmbH, D-52425 Ju¨lich, Germany Received 9 May 2008; received in revised form 27 August 2008; accepted 2 September 2008 Available online 25 September 2008
Abstract Epitaxial NiSi2 thin films are formed by annealing of Ni on sulfur-implanted silicon (1 0 0). The atomic structure and chemistry of the NiSi2/Si interface are investigated by aberration-corrected transmission electron microscopy. The interface is atomically sharp and runs mainly along the (1 0 0) plane. {1 1 1} segments of interface are also observed as minor facets. The atomic structure of the (1 0 0) and (1 1 1) interface has been determined. Interfacial dislocations with Burgers vectors a/4<1 1 1> and a/2<1 1 0> are observed near {1 1 1} facets. In particular, these dislocations have extra half atomic planes in the Si substrate. This configuration of dislocation does not agree with the sign of the lattice mismatch between bulk NiSi2 and Si. This novel phenomenon is understood by the fact that a high concentration of sulfur in the interface area leads to an expansion of the NiSi2 lattice and thus inverts the sign of the lattice mismatch. It is suggested that the change of the strain status, in addition to the doping effect of S, also plays a role in the tunable Schottky barrier height in this system. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Thin films; High-resolution electron microscopy; Silicides; Interface
1. Introduction Many attempts have been made to produce epitaxial thin films of transition metal silicides on silicon substrate for both fundamental studies on electrical and structural properties of interface, and the application as Schottky and ohmic contacts in silicon-based devices [1–3]. The electronic properties in metal/Si junctions strongly depend on the structural features of the interface, such as the coordination of metal atoms, reconstruction and defects at the interface. In recent years one has learned how to modify the electronic properties by adjusting film growth parameters as well as employing chemical additives. It was found that the Schottky barrier height (SBH) of NiSi on Si(1 0 0) can be controlled by the segregation of sulfur at the NiSi/Si(1 0 0) interfaces [4,5]. The concentration of S at the NiSi/Si interface increases
*
Corresponding author. Tel.: +49 2461 612413; fax: +49 2461 616444. E-mail address: [email protected] (S.B. Mi).
with the S+ implant dose, and correspondingly decreases the SBH gradually on n-Si(1 0 0). For nanometer scale device applications, it is desired to obtain an atomically sharp metal/Si interface and be able to control the Schottky barrier height of the metal/Si contacts. Among the metal/Si candidates, NiSi2 has a cubic structure (CaF2-type) with a lattice parameter (a = 0.541 nm) [6] very close to that of Si (a = 0.54309 nm) [7]. Due to the small lattice mismatch of about 0.4% between NiSi2 and Si, an atomically sharp NiSi2/Si interface has been reported [8–14]. Experimental and theoretical studies have been dedicated to the NiSi2/ Si system [8–17]. Nevertheless, the atomic and electronic structures and properties of the interface are not yet fully understood. Recently, NiSi2/Si junctions were obtained by annealing of Ni on S+-implanted silicon (1 0 0) and the segregation of sulfur (S) atoms at the NiSi2/Si(1 0 0) interface was observed [18]. The effect of the S segregation on the SBH of the NiSi/Si contacts [4,5] brings additional interest to research into the atomic details of the interface in the NiSi2/S+-implanted Si systems. In the present work, we study the atomic structure and chemistry of the interface
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.09.002
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in the NiSi2/S+-implanted Si junction by aberration-corrected transmission electron microscopy. Our aim is to study the effect on various structural features of the interface due to the S-additives. 2. Experimental In preparation of the samples, n-type Si(1 0 0) with a resistivity of 2.5 8.5 Xcm was implanted with S+ at an energy of 50 keV to a dose of 5 1014 S+ cm 2 at an incidence angle of 7°. Subsequently, a 20 nm thick Ni layer was deposited on the S+ as-implanted sample after a standard cleaning and an HF dip to remove the native oxide. The NiSi2 layer was formed by rapid thermal annealing at 950 °C for 30 s in a 90% N2 + 10% H2 atmosphere. A detailed account of the epitaxial NiSi2 growth has been reported elsewhere [18]. In order to determine the interface structure, cross-sectional specimens were prepared by a standard procedure for high-resolution transmission electron microscopy (HRTEM). The interface structure of NiSi2/Si was investigated in a microscope FEI Titan 80-300 equipped with a spherical aberration corrector for the objective lens, operated at 300 kV. HRTEM images were recorded using negative CS imaging (NCSI) technique [19]. 3. Results and discussion Fig. 1 shows a cross-sectional low-magnification image of the NiSi2/Si interface viewed along the <0 1 1> zone axis of Si. NiSi2 thin films grow on Si in a cubic-on-cubic relation. The interface looks very sharp, with a precise change of image contrast. The interface runs along the (1 0 0) plane, as indicated by a horizontal arrow. Small {1 1 1} facets exist at the interface, denoted by vertical arrows. Fig. 2a and b shows high-resolution images of the interface between NiSi2 and Si. The images were recorded along the <0 1 1> zone axis under the NCSI conditions, which led to a negative phase contrast, with atom columns appearing bright against a dark background. Image simulations were performed to confirm the conditions using the MacTempas
Fig. 2. HRTEM images of the interface variants A (a) and B (b) viewed along a <0 1 1> zone axis. Structure models of the interfaces are superimposed in the images. The unit cells of NiSi2 and Si are indicated by white rectangular boxes. The simulated images are inserted in a brokenline box in (a) (thickness = 6.4 nm, spherical abberation (Cs) = 15 lm, defocus = 4.4 nm and crystal tilt = 6 mrad) and (b) (thickness = 6.8 nm, Cs = 15 lm, defocus = 4.2 nm and crystal tilt = 8 mrad).
Fig. 1. A cross-sectional low-magnification TEM image of NiSi2/Si viewed along the <0 1 1> zone axis. The (1 0 0) interface is indicated by a horizontal arrow. {1 1 1} facets are indicated by vertical arrows.
software package [20]. A slight difference in image contrast can be seen, which is due to the local crystal tilt from the <0 1 1> zone axis. The projected dumbbell configuration of silicon along the <0 1 1> direction is clearly visible and extends up to the interface. NiSi2 has a CaF2 structure and the projected lattice images in all <0 1 1> directions are identical. The columns of Ni atoms (white circles) and Si atoms (black circles) are identified by checking the image contrast. The NiSi2 film terminates at the interface with a Ni atomic plane in both Fig. 2a and b. However, the terminating plane of Si substrate looks different. In Fig. 2a, the terminating plane shows a half dumbbell character, while in Fig. 2b it shows a full dumbbell character. In addition, the atomic columns of Si in the terminating plane
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locate in the middle position below two Ni columns in the terminating plane of the NiSi2 film in Fig. 2a. In Fig. 2b the atomic columns of Si in the terminating plane locate directly below the Ni columns in the last plane of the NiSi2 film. The reconstructed feature can be recognized in the terminating plane of Si: small and large spacings between the Si columns occur alternately, denoted by S and L. The simulated image that optimally matches with the experimental image is inserted in Fig. 2a and b, respectively. The structural features of Fig. 2a and b agree with the interface structure mode of 5-coordinated silicon facing 7-coordinated nickel at the interface, as reported in Ref. [8] for NiSi2 films prepared by molecular beam allotaxy. According to the crystal structure of Si, the two interfaces in Fig. 2a and b are indeed identical. The observed difference in images of Fig. 2a and b is due to the different viewing direction, i.e. the interface structure being projected in two orthogonal <0 1 1> directions. Nevertheless, since the structure of silicon does not have fourfold symmetry, the appearance of the two orthogonal projections of the interface in the same viewing direction must signify a lattice defect. Therefore, it is necessary to distinguish the two projections of the interface in a given <1 1 0> direction. In the following, these two projections are named as 90° variants. The projection of the interface in Fig. 2a is denoted as variant A and that in Fig. 2b as variant B. The two 90° interface variants are usually separated by a {1 1 1} facet with lattice defects. Fig. 3a shows an image of a {1 1 1} facet which separates two 90° variants A and B. The interface between NiSi2 films and Si is marked by a dotted line. An interfacial dislocation was observed, and is denoted by an arrow. Performing a Burgers circuit around the dislocation, a closure failure defining a Burgers vector a/4 <1 1 1> (a is the lattice constant of Si) was obtained. The related lattice distortion can be seen extending into NiSi2 thin films with a few nanometers. Fig. 3b shows an image of another {1 1 1} facet between the two 90° variants A and B. Two {1 1 1} layers of Si dumbbells directly on the right side of the facet show a twin relation with the Si substrate. The {1 1 1} twin boundaries are indicated by an arrow. Performing a Burgers circuit around the {1 1 1} facet, a closure failure with a displacement vector a/4<1 1 1> was obtained, indicating the existence of a dislocation here. A similar twin structure was reported in the CoSi2/Si system [21]. However, in our investigation we find that the extra half atomic plane of the dislocation stays in the Si substrate side, as shown in Fig. 3b, while the extra half plane was reported in the CoSi2 side in the CoSi2/Si system. In Fig. 3c, a similar interfacial twin structure is displayed at a {1 1 1} facet. However, the (1 0 0) interface parts beside the facet are only variant A. An interfacial dislocation was observed. The Burgers vector of the dislocation measured to be a/2<1 1 0> is different from that of the dislocation shown in Fig. 3a and b. However, the extra atomic planes again locate in the Si substrate side.
Fig. 3. (a and b) HRTEM images of the {1 1 1} facet between two (1 0 0) interface variants, A and B. (c) A {1 1 1} facet between two segments of variant A. The dotted line traces the interface. {1 1 1} twin boundaries in Si are seen, as indicated by an arrow in (b) and (c). Burgers circuits show the projected Burgers vectors of the dislocations.
Two types of {1 1 1} interface of NiSi2/Si have been observed, as displayed in Fig. 3. In Fig. 3a, NiSi2 has the same crystallographic orientation as the Si substrate (Atype geometry). In Fig. 3b and c, due to the two-layer twin
S.B. Mi et al. / Acta Materialia 57 (2009) 232–236
structure in Si, the NiSi2 film shares the same <1 1 1> axis as Si but is rotated by 180° about the axis (B-type geometry). Several structure modes have been proposed for the NiSi2/Si (1 1 1) interface, in which different coordinates of the interfacial Ni atoms – 5, 7 and 8 – are included [10,13,16,17]. Fig. 4a and b shows the NiSi2/Si (1 1 1) interface with A-type and B-type geometry, respectively. Fig. 4b is a magnification of the {1 1 1} facet in Fig. 3c, in which the twin boundary is indicated by an arrow. At the NiSi2/Si (1 1 1) interface marked by a dotted line in Fig. 4a and b, both A-type and B-type geometry show that NiSi2 terminates at the interface with a Si plane. With this configuration of the termination, the Ni atoms in the last Ni plane of
Fig. 4. HRTEM image of the NiSi2/Si (111) interface with A-type geometry (a) and B-type geometry (b). The dotted line traces the interface. The simulated images are inserted in the left side in (a) (thickness = 6 nm, Cs = 15 lm and defocus = 4.6 nm) and in a broken-line box in (b) (thickness = 6.4 nm, Cs = 15 lm, defocus = 4.2 nm and crystal tilt = 15 mrad). The structure mode of the NiSi2/Si (111) interface is superimposed. The {111} twin boundaries in Si are indicated by an arrow in (b).
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the NiSi2 film are seven coordinated, differing from the eight coordinated nickel in the bulk NiSi2. As shown by the structure model overlaying the interface area in Fig. 4a and b, the white and black circles represent the Ni and the Si atomic columns, respectively. The calculated image that optimally fits the experimental image is inserted in Fig. 4a and b, respectively. In Fig. 4a, the spacing of adjacent Si layers across the interface is about 0.239 nm, which is very close to the Si bond length in Si lattice of 0.235 nm in <1 1 1> directions. The value of the spacing is calibrated using the lattice parameter of Si (a = 0.54309 nm). It should be noted that the NiSi2/Si interface with seven coordinated Ni atoms has the lowest energy based on theoretical calculations [16,17]. In our NiSi2/Si system, the interface adopts mainly the (1 0 0) plane with {1 1 1} facets. The structure of the (1 0 0) interface exhibits two 90° variants, which are usually separated by a {1 1 1} facet accompanying a lattice defect. The displacement vector of the defect accommodates the 90° rotation, which is not a symmetry element for the silicon structure. The important finding in our study is that the interfacial dislocations have extra half atomic planes locating in the Si substrate side. If the lattice constant of pure NiSi2 and silicon is considered, this configuration of the dislocations is expected to enhance strain due to the lattice mismatch in the system. Since the lattice constant of bulk NiSi2 is about 0.4% smaller than that of Si, an energy increase in the system is expected, due to the enhancement of mismatch strain, which would produce a driving force to move the dislocation away from the interface. These expectations based on the lattice parameters of pure NiSi2 and silicon are contradictory to our experimental observations. Therefore, it is reasonable to relate the dislocation configuration to the S-additives. Segregation of S in the interface area may be responsible for the dislocation arrangement. One can imagine that S atoms as substitution and interstitial defects in the NiSi2 lattice will increase the lattice parameters, depending on the concentration. Above a critical concentration, the lattice constant of the NiSi2 becomes larger than that of Si, leading to a lattice mismatch in the opposite direction. In this circumstance, the dislocation configuration as observed in the NiSi2/S+-implanted Si system will accommodate the lattice mismatch. In order to test the effect of the S-additives on the dislocation configuration, we performed an investigation of the microstructure and interface structure of NiSi2 grown on Si (1 0 0) without S+implantation. It was found that the interface structure of NiSi2/Si without S is the same as that of NiSi2/Si with S. The important difference is that the interfacial dislocations show the extra half atomic planes in the NiSi2 film side. Now, it can be concluded that the different configuration of dislocation can be related to the segregation of S at the interface. We note that the different dislocation configuration implies a different strain status at the interface. This means that the lattice strain at the Scho-
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ttky and ohmic contacts is tunable by controlling the concentration of S. In interpreting the change in the SBH, the changed strain configuration should be considered in addition to the electrical effect of S atoms. 4. Conclusion By means of aberration-corrected transmission electron microscopy, atomic structure and defect configuration at the NiSi2/S+-implanted Si interface have been investigated, with the following findings. At the (1 0 0) interface, a 2 1 reconstructed interface has been determined with the atom arrangement of five coordinated silicon facing seven coordinated nickel. At {1 1 1} facets, two types of (1 1 1) interface exist, with NiSi2 being terminated on an atomic plane of silicon and seven coordinated nickel atoms in the last nickel plane. Interfacial dislocations with Burgers vectors of a/4<1 1 1> and a/2<1 1 0> are observed close to NiSi2/ S+-implanted Si {1 1 1} facets. These dislocations show a different character from those found for the interface in the systems of NiSi2/Si without S, implying a change in the strain status in the interface area. This finding provides a new structural feature, which should be considered when interpreting the change of SBH observed in this system.
skopie mit Elektronen) of the initiative and networking fund of the Helmholtz Assosciation. References [1] [2] [3] [4] [5] [6] [7]
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
Acknowledgements S.B. Mi gratefully acknowledges funding by project vhfz-001 (Deutsches Zentrum fu¨r Mikroskopie und Spektro-
[19] [20] [21]
Tung RT, Gibson JM, Poate JM. Phys Rev Lett 1983;50:429. von Ka¨nel H. Mater Sci Rep 1992;8:193. Schwarz C, von Ka¨nel H. J Appl Phys 1996;79:8798. Zhao QT, Breuer U, Rije E, Lenk St, Mant S. Appl Phys Lett 2005;86:062108. Zhao QT, Breuer U, Lenk St, Mant S. J Appl Phys 2007;102:023522. Hansen M. Constitution of binary alloys. New York: McGraw-Hill; 1958. p. 1040. Villars P, Calvert LD. Pearson’s handbook of crystallographic data for intermetallic phases, vol. 4. Meterial Park, OH: ASM International; 1991. Falke U, Bleloch A, Falke M, Teichert S. Phys Rev Lett 2004;92:116103. Cherns D, Hetherington CJD, Humphreys CJ. Philos Mag A 1983;49:165. Cherns D, Anstis GR, Hutchison JL, Spence JCH. Philos Mag A 1982;46:849. Khang Y, Kahng SJ, Mang KM, Jeon D, Lee JH, Kim YN, et al. J Vacuum Sci Technol B 1994;12:2094. Tung RT, Levi AFJ, Sullivan JP, Schrey F. Phys Rev Lett 1991;66:72. Chen WJ, Chen FR. Philos Mag A 1993;68:605. Teichert S, Falke M, Giesler H, Beddies G, Hinneberg HJ. Thin Solid Films 1998;336:222. Yu BD, Miyamoto Y, Sugino O, Sakai A, Sasaki T, Ohno T. J Vacuum Sci Technol B 2001;19:1180. Hamann DR. Phys Rev Lett 1988;60:313. Van den Hoek PJ, Ravenek W, Baerends EJ. Phys Rev Lett 1988;60:1743. Zhao QT, Mi SB, Jia CL, Breuer U, Lenk S, Mantl S. J. Mater Res, submitted for publication.. Jia CL, Lentzen M, Urban K. Science 2003;299:870. O’Keefe MA, Kilaas R. Scan Microsc Suppl 1988;2:225. Falke M, Falke U, Bleloch A, Teichert S, Beddies G, Hinneberg HJ. Appl Phys Lett 2005;86:203103.
Acta Materialia 57 (2009) 232–236 www.elsevier.com/locate/actamat
NiSi2/Si interface chemistry and epitaxial growth mode S.B. Mi a,*, C.L. Jia a, Q.T. Zhao b, S. Mantl b, K. Urban a a
Institute of Solid State Research and Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons (ER-C), Research Centre Juelich, D-52425 Juelich, Germany b Institute of Bio- and Nanosystems (IBN1), and Center of Nanoelectronic Systems for Information Technology (CNI), Forschungszentrum Ju¨lich GmbH, D-52425 Ju¨lich, Germany Received 9 May 2008; received in revised form 27 August 2008; accepted 2 September 2008 Available online 25 September 2008
Abstract Epitaxial NiSi2 thin films are formed by annealing of Ni on sulfur-implanted silicon (1 0 0). The atomic structure and chemistry of the NiSi2/Si interface are investigated by aberration-corrected transmission electron microscopy. The interface is atomically sharp and runs mainly along the (1 0 0) plane. {1 1 1} segments of interface are also observed as minor facets. The atomic structure of the (1 0 0) and (1 1 1) interface has been determined. Interfacial dislocations with Burgers vectors a/4<1 1 1> and a/2<1 1 0> are observed near {1 1 1} facets. In particular, these dislocations have extra half atomic planes in the Si substrate. This configuration of dislocation does not agree with the sign of the lattice mismatch between bulk NiSi2 and Si. This novel phenomenon is understood by the fact that a high concentration of sulfur in the interface area leads to an expansion of the NiSi2 lattice and thus inverts the sign of the lattice mismatch. It is suggested that the change of the strain status, in addition to the doping effect of S, also plays a role in the tunable Schottky barrier height in this system. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Thin films; High-resolution electron microscopy; Silicides; Interface
1. Introduction Many attempts have been made to produce epitaxial thin films of transition metal silicides on silicon substrate for both fundamental studies on electrical and structural properties of interface, and the application as Schottky and ohmic contacts in silicon-based devices [1–3]. The electronic properties in metal/Si junctions strongly depend on the structural features of the interface, such as the coordination of metal atoms, reconstruction and defects at the interface. In recent years one has learned how to modify the electronic properties by adjusting film growth parameters as well as employing chemical additives. It was found that the Schottky barrier height (SBH) of NiSi on Si(1 0 0) can be controlled by the segregation of sulfur at the NiSi/Si(1 0 0) interfaces [4,5]. The concentration of S at the NiSi/Si interface increases
*
Corresponding author. Tel.: +49 2461 612413; fax: +49 2461 616444. E-mail address: [email protected] (S.B. Mi).
with the S+ implant dose, and correspondingly decreases the SBH gradually on n-Si(1 0 0). For nanometer scale device applications, it is desired to obtain an atomically sharp metal/Si interface and be able to control the Schottky barrier height of the metal/Si contacts. Among the metal/Si candidates, NiSi2 has a cubic structure (CaF2-type) with a lattice parameter (a = 0.541 nm) [6] very close to that of Si (a = 0.54309 nm) [7]. Due to the small lattice mismatch of about 0.4% between NiSi2 and Si, an atomically sharp NiSi2/Si interface has been reported [8–14]. Experimental and theoretical studies have been dedicated to the NiSi2/ Si system [8–17]. Nevertheless, the atomic and electronic structures and properties of the interface are not yet fully understood. Recently, NiSi2/Si junctions were obtained by annealing of Ni on S+-implanted silicon (1 0 0) and the segregation of sulfur (S) atoms at the NiSi2/Si(1 0 0) interface was observed [18]. The effect of the S segregation on the SBH of the NiSi/Si contacts [4,5] brings additional interest to research into the atomic details of the interface in the NiSi2/S+-implanted Si systems. In the present work, we study the atomic structure and chemistry of the interface
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.09.002
S.B. Mi et al. / Acta Materialia 57 (2009) 232–236
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in the NiSi2/S+-implanted Si junction by aberration-corrected transmission electron microscopy. Our aim is to study the effect on various structural features of the interface due to the S-additives. 2. Experimental In preparation of the samples, n-type Si(1 0 0) with a resistivity of 2.5 8.5 Xcm was implanted with S+ at an energy of 50 keV to a dose of 5 1014 S+ cm 2 at an incidence angle of 7°. Subsequently, a 20 nm thick Ni layer was deposited on the S+ as-implanted sample after a standard cleaning and an HF dip to remove the native oxide. The NiSi2 layer was formed by rapid thermal annealing at 950 °C for 30 s in a 90% N2 + 10% H2 atmosphere. A detailed account of the epitaxial NiSi2 growth has been reported elsewhere [18]. In order to determine the interface structure, cross-sectional specimens were prepared by a standard procedure for high-resolution transmission electron microscopy (HRTEM). The interface structure of NiSi2/Si was investigated in a microscope FEI Titan 80-300 equipped with a spherical aberration corrector for the objective lens, operated at 300 kV. HRTEM images were recorded using negative CS imaging (NCSI) technique [19]. 3. Results and discussion Fig. 1 shows a cross-sectional low-magnification image of the NiSi2/Si interface viewed along the <0 1 1> zone axis of Si. NiSi2 thin films grow on Si in a cubic-on-cubic relation. The interface looks very sharp, with a precise change of image contrast. The interface runs along the (1 0 0) plane, as indicated by a horizontal arrow. Small {1 1 1} facets exist at the interface, denoted by vertical arrows. Fig. 2a and b shows high-resolution images of the interface between NiSi2 and Si. The images were recorded along the <0 1 1> zone axis under the NCSI conditions, which led to a negative phase contrast, with atom columns appearing bright against a dark background. Image simulations were performed to confirm the conditions using the MacTempas
Fig. 2. HRTEM images of the interface variants A (a) and B (b) viewed along a <0 1 1> zone axis. Structure models of the interfaces are superimposed in the images. The unit cells of NiSi2 and Si are indicated by white rectangular boxes. The simulated images are inserted in a brokenline box in (a) (thickness = 6.4 nm, spherical abberation (Cs) = 15 lm, defocus = 4.4 nm and crystal tilt = 6 mrad) and (b) (thickness = 6.8 nm, Cs = 15 lm, defocus = 4.2 nm and crystal tilt = 8 mrad).
Fig. 1. A cross-sectional low-magnification TEM image of NiSi2/Si viewed along the <0 1 1> zone axis. The (1 0 0) interface is indicated by a horizontal arrow. {1 1 1} facets are indicated by vertical arrows.
software package [20]. A slight difference in image contrast can be seen, which is due to the local crystal tilt from the <0 1 1> zone axis. The projected dumbbell configuration of silicon along the <0 1 1> direction is clearly visible and extends up to the interface. NiSi2 has a CaF2 structure and the projected lattice images in all <0 1 1> directions are identical. The columns of Ni atoms (white circles) and Si atoms (black circles) are identified by checking the image contrast. The NiSi2 film terminates at the interface with a Ni atomic plane in both Fig. 2a and b. However, the terminating plane of Si substrate looks different. In Fig. 2a, the terminating plane shows a half dumbbell character, while in Fig. 2b it shows a full dumbbell character. In addition, the atomic columns of Si in the terminating plane
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locate in the middle position below two Ni columns in the terminating plane of the NiSi2 film in Fig. 2a. In Fig. 2b the atomic columns of Si in the terminating plane locate directly below the Ni columns in the last plane of the NiSi2 film. The reconstructed feature can be recognized in the terminating plane of Si: small and large spacings between the Si columns occur alternately, denoted by S and L. The simulated image that optimally matches with the experimental image is inserted in Fig. 2a and b, respectively. The structural features of Fig. 2a and b agree with the interface structure mode of 5-coordinated silicon facing 7-coordinated nickel at the interface, as reported in Ref. [8] for NiSi2 films prepared by molecular beam allotaxy. According to the crystal structure of Si, the two interfaces in Fig. 2a and b are indeed identical. The observed difference in images of Fig. 2a and b is due to the different viewing direction, i.e. the interface structure being projected in two orthogonal <0 1 1> directions. Nevertheless, since the structure of silicon does not have fourfold symmetry, the appearance of the two orthogonal projections of the interface in the same viewing direction must signify a lattice defect. Therefore, it is necessary to distinguish the two projections of the interface in a given <1 1 0> direction. In the following, these two projections are named as 90° variants. The projection of the interface in Fig. 2a is denoted as variant A and that in Fig. 2b as variant B. The two 90° interface variants are usually separated by a {1 1 1} facet with lattice defects. Fig. 3a shows an image of a {1 1 1} facet which separates two 90° variants A and B. The interface between NiSi2 films and Si is marked by a dotted line. An interfacial dislocation was observed, and is denoted by an arrow. Performing a Burgers circuit around the dislocation, a closure failure defining a Burgers vector a/4 <1 1 1> (a is the lattice constant of Si) was obtained. The related lattice distortion can be seen extending into NiSi2 thin films with a few nanometers. Fig. 3b shows an image of another {1 1 1} facet between the two 90° variants A and B. Two {1 1 1} layers of Si dumbbells directly on the right side of the facet show a twin relation with the Si substrate. The {1 1 1} twin boundaries are indicated by an arrow. Performing a Burgers circuit around the {1 1 1} facet, a closure failure with a displacement vector a/4<1 1 1> was obtained, indicating the existence of a dislocation here. A similar twin structure was reported in the CoSi2/Si system [21]. However, in our investigation we find that the extra half atomic plane of the dislocation stays in the Si substrate side, as shown in Fig. 3b, while the extra half plane was reported in the CoSi2 side in the CoSi2/Si system. In Fig. 3c, a similar interfacial twin structure is displayed at a {1 1 1} facet. However, the (1 0 0) interface parts beside the facet are only variant A. An interfacial dislocation was observed. The Burgers vector of the dislocation measured to be a/2<1 1 0> is different from that of the dislocation shown in Fig. 3a and b. However, the extra atomic planes again locate in the Si substrate side.
Fig. 3. (a and b) HRTEM images of the {1 1 1} facet between two (1 0 0) interface variants, A and B. (c) A {1 1 1} facet between two segments of variant A. The dotted line traces the interface. {1 1 1} twin boundaries in Si are seen, as indicated by an arrow in (b) and (c). Burgers circuits show the projected Burgers vectors of the dislocations.
Two types of {1 1 1} interface of NiSi2/Si have been observed, as displayed in Fig. 3. In Fig. 3a, NiSi2 has the same crystallographic orientation as the Si substrate (Atype geometry). In Fig. 3b and c, due to the two-layer twin
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structure in Si, the NiSi2 film shares the same <1 1 1> axis as Si but is rotated by 180° about the axis (B-type geometry). Several structure modes have been proposed for the NiSi2/Si (1 1 1) interface, in which different coordinates of the interfacial Ni atoms – 5, 7 and 8 – are included [10,13,16,17]. Fig. 4a and b shows the NiSi2/Si (1 1 1) interface with A-type and B-type geometry, respectively. Fig. 4b is a magnification of the {1 1 1} facet in Fig. 3c, in which the twin boundary is indicated by an arrow. At the NiSi2/Si (1 1 1) interface marked by a dotted line in Fig. 4a and b, both A-type and B-type geometry show that NiSi2 terminates at the interface with a Si plane. With this configuration of the termination, the Ni atoms in the last Ni plane of
Fig. 4. HRTEM image of the NiSi2/Si (111) interface with A-type geometry (a) and B-type geometry (b). The dotted line traces the interface. The simulated images are inserted in the left side in (a) (thickness = 6 nm, Cs = 15 lm and defocus = 4.6 nm) and in a broken-line box in (b) (thickness = 6.4 nm, Cs = 15 lm, defocus = 4.2 nm and crystal tilt = 15 mrad). The structure mode of the NiSi2/Si (111) interface is superimposed. The {111} twin boundaries in Si are indicated by an arrow in (b).
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the NiSi2 film are seven coordinated, differing from the eight coordinated nickel in the bulk NiSi2. As shown by the structure model overlaying the interface area in Fig. 4a and b, the white and black circles represent the Ni and the Si atomic columns, respectively. The calculated image that optimally fits the experimental image is inserted in Fig. 4a and b, respectively. In Fig. 4a, the spacing of adjacent Si layers across the interface is about 0.239 nm, which is very close to the Si bond length in Si lattice of 0.235 nm in <1 1 1> directions. The value of the spacing is calibrated using the lattice parameter of Si (a = 0.54309 nm). It should be noted that the NiSi2/Si interface with seven coordinated Ni atoms has the lowest energy based on theoretical calculations [16,17]. In our NiSi2/Si system, the interface adopts mainly the (1 0 0) plane with {1 1 1} facets. The structure of the (1 0 0) interface exhibits two 90° variants, which are usually separated by a {1 1 1} facet accompanying a lattice defect. The displacement vector of the defect accommodates the 90° rotation, which is not a symmetry element for the silicon structure. The important finding in our study is that the interfacial dislocations have extra half atomic planes locating in the Si substrate side. If the lattice constant of pure NiSi2 and silicon is considered, this configuration of the dislocations is expected to enhance strain due to the lattice mismatch in the system. Since the lattice constant of bulk NiSi2 is about 0.4% smaller than that of Si, an energy increase in the system is expected, due to the enhancement of mismatch strain, which would produce a driving force to move the dislocation away from the interface. These expectations based on the lattice parameters of pure NiSi2 and silicon are contradictory to our experimental observations. Therefore, it is reasonable to relate the dislocation configuration to the S-additives. Segregation of S in the interface area may be responsible for the dislocation arrangement. One can imagine that S atoms as substitution and interstitial defects in the NiSi2 lattice will increase the lattice parameters, depending on the concentration. Above a critical concentration, the lattice constant of the NiSi2 becomes larger than that of Si, leading to a lattice mismatch in the opposite direction. In this circumstance, the dislocation configuration as observed in the NiSi2/S+-implanted Si system will accommodate the lattice mismatch. In order to test the effect of the S-additives on the dislocation configuration, we performed an investigation of the microstructure and interface structure of NiSi2 grown on Si (1 0 0) without S+implantation. It was found that the interface structure of NiSi2/Si without S is the same as that of NiSi2/Si with S. The important difference is that the interfacial dislocations show the extra half atomic planes in the NiSi2 film side. Now, it can be concluded that the different configuration of dislocation can be related to the segregation of S at the interface. We note that the different dislocation configuration implies a different strain status at the interface. This means that the lattice strain at the Scho-
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ttky and ohmic contacts is tunable by controlling the concentration of S. In interpreting the change in the SBH, the changed strain configuration should be considered in addition to the electrical effect of S atoms. 4. Conclusion By means of aberration-corrected transmission electron microscopy, atomic structure and defect configuration at the NiSi2/S+-implanted Si interface have been investigated, with the following findings. At the (1 0 0) interface, a 2 1 reconstructed interface has been determined with the atom arrangement of five coordinated silicon facing seven coordinated nickel. At {1 1 1} facets, two types of (1 1 1) interface exist, with NiSi2 being terminated on an atomic plane of silicon and seven coordinated nickel atoms in the last nickel plane. Interfacial dislocations with Burgers vectors of a/4<1 1 1> and a/2<1 1 0> are observed close to NiSi2/ S+-implanted Si {1 1 1} facets. These dislocations show a different character from those found for the interface in the systems of NiSi2/Si without S, implying a change in the strain status in the interface area. This finding provides a new structural feature, which should be considered when interpreting the change of SBH observed in this system.
skopie mit Elektronen) of the initiative and networking fund of the Helmholtz Assosciation. References [1] [2] [3] [4] [5] [6] [7]
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Acknowledgements S.B. Mi gratefully acknowledges funding by project vhfz-001 (Deutsches Zentrum fu¨r Mikroskopie und Spektro-
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