ARTICLE pubs.acs.org/JPCC
Complex Defect Structure in the Core of Sn-Doped In2O3 Nanorods and Its Relationship with a Dislocation-Driven Growth Mechanism D. Maestre,*,† D. H€aussler,‡ A. Cremades,† W. J€ager,‡ and J. Piqueras† † ‡
Departamento de Física de Materiales, Facultad de Ciencias Físicas, Universidad Complutense de Madrid, 28040 Madrid, Spain Mikrostrukturanalytik, Christian-Albrechts Universit€at zu Kiel, D-24143, Kiel, Germany ABSTRACT: Sn-doped In2O3 arrow-shaped nanorods have been grown by a catalyst-free evaporation deposition method. Transmission electron microscopy (TEM) investigations reveal a complex defect structure consisting of nanoprecipitates, dislocation loops, and voids, in the core of the nanorods, extending all along the growth axis. The nanoprecipitates are Sn or Sn-rich and appear sometimes associated with small voids. The voids appear aligned in rows or discontinuous empty nanochannels along the nanorod core. TEM images depict the presence of surface ripples perpendicular to the growth direction, with typical distances of less than 10 nm. The relationship of the observed defect structure with a dislocation-driven growth of the nanorods is discussed.
’ INTRODUCTION Sn-doped In2O3 (ITO) is an extensively investigated transparent conducting oxide material with relevant technological applications in many fields, such as optoelectronic devices, solar cells, field emitters, or gas sensors.1 4 Recently, the synthesis and characterization of ITO nano- and microstructures in forms of nanowires, nanorods, nanopyramids, or nanoarrows have attracted increasing attention, as the exploitation of the phenomena appearing at the nanoscale could spread its applications as nanoelectronic and optoelectronic devices.5 7 Hence, much work is underway to gain new insights into understanding the processes involved in the growth and the doping of such nanostructures, in order to bridge the gap between the fundamental crystal growth theories and the technological applications of nanomaterials. Diverse fabrication processes have been reported for the synthesis of ITO 1D nanostructures, such as thermal evaporation, electrodeposition, or chemical synthesis.8 10 This study reports results of the structural characterization of Sn-doped In2O3 nanorods and nanoarrows that were fabricated via a catalyst-free vapor solid method avoiding the presence of catalyst or external substrates. The nature of the extended defects in the nanoarrows, such as nanoprecipitates and dislocations, and their influence on the growth mechanism have been investigated by TEM. The synthesis procedure used for the growth of these nanostructures and their chemical and optical characterization have been previously reported.11 The nanostructures have been characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM), conventional and high-resolution transmission electron microscopy (TEM), selected area electron diffraction (SAED), energy-dispersive X-ray spectroscopy (EDS) in TEM, and high-angle annular dark-field scanning-TEM (HAADF-STEM). r 2011 American Chemical Society
’ EXPERIMENTAL METHODS Sn-doped In2O3 nanorods were synthesized via a catalyst-free thermal treatment, which enables the growth of different semiconducting oxides nanostructures without using templates or external substrates.12 14 In this study, a mixture of high-purity InN (99.99%) and SnO2 (99.9%) powders, with an InN/SnO2 weight ratio of 10:1, has been used as starting material. The initial powders were first mixed and milled in a centrifugal ball mill and then pressed into disk-shaped samples. The pellets were annealed in a furnace using a two-step treatment under a controlled argon flow. The annealing treatment consists of an initial step at 350 C for 3 h, which favors the nitrogen desorption from InN, followed by a second step at 650 C for 20 h, which leads to the growth of the Sn-doped In2O3 nanostructures on the surface of the pellet. The morphology of the as-grown structures was studied by secondary electron mode in a Leica 440 SEM. The crystalline structure of the samples was investigated by XRD measurements carried out in a Philips diffractometer working at 45 kV and 40 mA, using the Cu Kα raditation. For the TEM studies, the nanostructures were carefully removed from the pellet surface and dispersed by ultrasonic treatment in highpurity isopropanol. A few drops of the solution containing the nanostructures were then deposited onto TEM copper grids coated with a holey carbon film. The TEM investigation was performed with a Philips CM30 TEM and an FEI-Tecnai F30 G2 TEM, both operated at 300 kV. The microscopic structure of the nanorods was investigated by diffraction contrast TEM imaging conditions, as well as by selected area electron diffraction (SAED). Spatially resolved energy-dispersive X-ray spectroscopy Received: May 17, 2011 Revised: July 11, 2011 Published: August 15, 2011 18083
dx.doi.org/10.1021/jp204579u | J. Phys. Chem. C 2011, 115, 18083–18087
The Journal of Physical Chemistry C
Figure 1. (a) SEM and (b) TEM images showing the nanorods and nanoarrows grown on the Sn-doped In2O3 sample treated at 650 C.
Figure 2. HRTEM image acquired on the nanoarrow, in the lower inset, showing interplanar distances of 0.506 nm, which correspond to the (200) distances in cubic In2O3. The corresponding FFT pattern is shown in the upper inset.
(EDS) and high-angle annular dark-field (HAADF) STEM were used to investigate the local elemental composition of the nanostructures. For the X-ray microanalyses, the EDAX DX-4 system, in combination with an UTW-Si (Li)-detector attached to the FEI-Tecnai F30 G2 TEM, was used.
’ RESULTS AND DISCUSSION Elongated nano- and microstructures, mainly rods and arrows, grow on the surface of the pellet after annealing at 650 C, as shown in Figure 1a. The temperature required for the thermal growth of In2O3/Sn one-dimensional nanostructures usually ranges from 900 to 1100 C. In fact, Li et al. stated that there should be not enough energy to crystallize indium oxide nanowires below 750 C.15 In our case, the use of InN as the precursor enables us to grow the Sn-doped as well as undoped In2O3 nanostructures at a lower temperature. Sn-doped In2O3 nanostructures obtained by this process generally show a longer and thinner morphology, as compared with the undoped ones grown by the same method,16 and present arrow-, beak-, and elbow-like shapes, whereas in the case of undoped In2O3, most of the structures show an arrow-like morphology.16 Our EDX measurements show that the Sn content averaged
ARTICLE
Figure 3. (a) TEM image of a nanoarrow showing a high defective region at the core extending up to the apex (upper inset). A TEM image of a nondoped In2O3 arrow-like nanorod showing a nanopipe at the core is included in the lower inset. (b) Bright-field TEM image corresponding to the central region of the nanoarrow shown in (a). The contrast features can be interpreted as nanoprecipitates, voids, and dislocation loops.
across larger areas in the structures is about 2.6 at. %. The presence of Sn in the precursor materials may result in significant variations in growth parameters, such as local vapor pressures, supersaturation ratios, surface energies, or growth rates along specific crystallographic directions. All of these variations will govern the final morphology of the nanostructures. It has been reported that the presence of Sn in the precursor powders leads to varying morphologies also for other doped oxides, such as GeO2 and ZnO.17,18 Our TEM investigations confirm the high aspect ratio of many of these doped nanoarrows, which have typical lengths up to some micrometers and constant diameters of tens of nanometers, as shown in Figure 1b. The absence of tin droplets at the end of the nanoarrows indicates that the growth mechanism occurs by a vapor solid growth mechanism. Figure 2 shows a HRTEM image of the tip region of a single-crystalline nanoarrow (lower inset). The lattice distance between fringes is 0.506 nm, which corresponds to the (200) interplanar distances in cubic In2O3. The fast Fourier transform (FFT) pattern, as that displayed in Figure 2 (upper inset), corresponds to the In2O3 bixbyite cubic structure in the [001] zone axis. The growth direction of the nanoarrows determined from the TEM observations is Æ100æ. TEM bright-field diffraction contrast imaging performed on Sn-doped In2O3 nanoarrows with typical diameters between 100 and 300 nm, and some micrometers length, shows the presence of a high defective region in the core, as shown in Figure 3a. In the case of the nondoped In2O3 nanorods, a perfect nanopipe is formed in the core (lower inset in Figure 3a), instead of the highly defective region observed in the Sn-doped In2O3 case. This region, which appears with a dark contrast in the bright-field TEM images, extends along the nanoarrows, up to the apex (upper inset in Figure 3a). Figure 3b, at higher magnification, reveals that the core region contains different defects, such as nanoprecipitates, voids, and dislocation loops. The nanoprecipitates appear well distributed along the core region, with sizes ranging from 2 to 7 nm in diameter. Their nature was analyzed by local nanoprobe EDX-STEM measurements. The local element composition in regions with nanoprecipitates shows the presence of only In, Sn, and O in the probed areas (as shown in the EDX spectrum in Figure 4a). The spectral lines originate from the nanoprecipitates and the host material surrounding them. Comparisons 18084
dx.doi.org/10.1021/jp204579u |J. Phys. Chem. C 2011, 115, 18083–18087
The Journal of Physical Chemistry C
ARTICLE
Figure 4. (a) Local EDX spectrum acquired on the core region of a nanoarrow with nanoprecipitates. A local EDX spectrum acquired on a region without nanoprecipitates has been included as an inset. (b) HAADF image corresponding to the region of the nanoarrow marked on the general view in the upper inset.
with measurements taken from areas without precipitates (inset in Figure 4a) indicate that the particles are Sn or Sn-rich precipitates. Furthermore, high-angle annular dark-field (HAADF) investigations in STEM have been carried out, as they allow for locally element monitoring by exploiting the high-angle electronscattering behavior. This technique is also known as Z-contrast imaging. HAADF images (Figure 4b) show bright spots with a distribution and sizes as those of the small rounded nanoparticles observed in the bright-field TEM images along the nanoarrow core. This high HAADF contrast confirms the presence of higher Z elements in the core region, in accordance with the results from the EDX measurements, and indicates that the nanoprecipitates consist of Sn or Sn-rich material. The dislocation loops observed in the core region of the nanoarrows (Figure 3b) may have a large influence on their physical properties. To achieve a better understanding of the nature of these observed defects, two-beam condition imaging investigations have been performed. The contrast of the dislocation loops is visible under two-beam conditions with the (220) diffraction spot near the [001] zone axis pattern, whereas the loops become almost invisible when the two-beam image is performed, exciting the orthogonal family of spots, as the noncollinear (220) and (022) (Figure 5). On the basis of the invisibility criterion, the direction of the Burgers vector associated with the loops has been estimated to be along the [111] direction. This result is compatible b = a/2 observed in bcc compounds, such as In2O3. The formation of the loops during the nanorod growth could lead to the relaxation of the local strain induced in the core region during the growth. A possible source of local lattice strain is the formation of precipitates in the rod core, which would explain the presence of precipitates and loops in the same regions. Ovidko et al.19 have found in a theoretical analysis of composite nanowires that dislocation loops may form from misfitting inclusions at smaller misfit values than straight dislocations and that the formation of rows of loops is energetically favorable. The presence of nanoprecipitates in the center of the rods, as those observed in this study, may lead to lattice strains and formation of loops relaxing this strain. In the core regions of a number of nanoarrows, rows of voids (Figure 6a), or discontinuous empty channels (Figure 6c,d),
Figure 5. Dislocation loops observed in the core region of a nanoarrow, imaged by different g vectors.
aligned along the growth axis are observed. In some regions, the voids are associated with the presence of the nanoprecipitates, as observed in Figure 6b. The presence of both Sn-rich nanoprecipitates and voids homogeneously distributed along the central part of the nanoarrows indicates diffusion phenomena occurring during the growth process. These diffusion processes could be related to an inhomogeneous Sn dopant distribution, which leads to the formation of voids possibly via the Kirkendall effect.20 However, some other phenomena should be considered in order to clarify the formation of voids, as the fluctuations in the growth environment due to the presence of Sn.21 Some similar phenomena have been reported for Sn-doped ZnO, where an outward Sn diffusion occurs toward the rod surface22 and groups of voids are formed in the rod core. The presence of aligned voids and dislocations loops in the core region of the nanoarrows could be explained in the frame of a dislocation-driven growth mechanism. Recently, some research has focused on the role played by a screw dislocation in the 18085
dx.doi.org/10.1021/jp204579u |J. Phys. Chem. C 2011, 115, 18083–18087
The Journal of Physical Chemistry C
ARTICLE
Figure 7. TEM images showing parallel features normal to the nanoarrow growth direction.
Figure 6. TEM images of the core region of different nanoarrows, where groups of voids (a), voids associated with nanoprecipitates (b), or empty channels (c, d) can be observed.
anisotropic crystal growth of nanowires and other elongated nanostructures.23,24 Indeed, under growth conditions without applying a catalyst and for low supersaturations, as those used in our case, the 1D growth can be explained attending to a dislocation-driven growth mechanism where a screw dislocation extends along the core of the nanowires, providing self-perpetuating steps, as proposed in the fundamental crystal growth theories of Sears.25 Bierman et al.23 observed screw dislocations in PbS nanowires, analyzing their nature by using diffraction contrast TEM. However, as dislocations may be mobile and unstable in small volumes, usually some other signs associated with screw dislocations should be also taken under consideration as supporting evidence for the dislocation-driven growth. One of these is the presence of a nanopipe extending along the core of nanowires and nanorods. It was demonstrated by Frank26 that screw dislocations with a large Burgers vector accumulate high stress in the dislocation core, and if the Burgers vector exceeds a critical value, the creation of a pipe by removing the highly strained material around the dislocation core is energetically favorable. Indeed we have observed in a previous work for undoped In2O3 nanorods27 perfect regular nanopipes with a constant diameter of 18.5 nm extending along the core. For the formation of such nanopipes, we suggested a dislocation-driven growth mechanism where the nanopipe represents the empty dislocation core. The channels and rows of voids in the core of Sn-doped ZnO nanowires reported in ref 22 could be also be related to this mechanism. The nanotube growth driven by screw dislocations in ZnO has been recently demonstrated21 with agreement between experimental growth kinetics and those predicted from fundamental crystal growth theories. The same authors have shown the screw-dislocation-driven growth of ZnO nanowires seeded by dislocations in substrates in refs 28 and 29, where a more general discussion on the role of screw dislocation in nanowire and nanotube formation has been reported. Contrary to the case of undoped In2O3 nanoarrows that contain a perfect nanopipe, the core region of the doped In2O3 arrows
investigated in this work has a complex structure of voids, nanoprecipitates, and dislocation loops that we attribute to the presence of the Sn dopant. The formation of the nanopipe to relieve the dislocation strain energy can be locally influenced by the presence of dopants around the dislocation and the formation of nanoprecipitates and loops, leading to the observed discontinuous nanopipe. The occurrence of micropipes containing aligned rows of cavities has been also reported for SiC films.30 A striking feature observed on some nanoarrows is an array of faint dark lines perpendicular to the growth axis, observed in brightfield TEM images (Figure 7a). The spacing between these lines ranges from 2.5 to 10 nm. Images of the rod edges (Figure 7b,c) show that the lines correspond to the presence of steps. There is no observation of a direct relation between the stepped structure with the features, in particular nanopipes, observed in the core of the nanoarrows. HRTEM studies of micropipes in bulk SiC of ref 30 have shown that the wall of a screw-dislocation-related pipe can contain steps and that this induces a disturbed lattice up to 10 nm away from the surface of the pipe. A similar effect could lead to the formation of the surface steps for the nanorods observed here.
’ CONCLUSIONS Sn-doped In2O3 arrow-shaped nanorods have been grown by a catalyst-free evaporation deposition method. The nanoarrows have been characterized by different TEM methods. These TEM investigations reveal a complex defect structure in the core of the nanoarrows, including Sn, or Sn-rich, nanoprecipitates, small dislocation loops, as well as rows of voids or discontinuous empty nanochannels parallel to the growth axis. Also, arrays of nanosteps, perpendicular to the growth axis, have been observed in a number of the rods. Comparison with previously investigated nanoarrows of undoped In2O3 that contain a straight perfect nanopipe in the core, leads to suggest that the core structure reported in the present work is due to the presence of a dopant, which prevents the formation of a regular nanopipe all along the nanorod and favors the appearance of rows of voids. The results support the possibility of a dislocation-driven mechanism for the growth of the nanorods and nanoarrows. ’ AUTHOR INFORMATION Corresponding Author
*E-mail: davidmaestre@fis.ucm.es. 18086
dx.doi.org/10.1021/jp204579u |J. Phys. Chem. C 2011, 115, 18083–18087
The Journal of Physical Chemistry C
ARTICLE
’ ACKNOWLEDGMENT This work was supported by MICINN (MAT2009-07882, CSD2009-00013) and UCM-BSCH (Group 910146). ’ REFERENCES (1) Hamberg, I.; Granqvist, C. G. J. Appl. Phys. 1986, 60, R123. (2) Zhu, F.; Zhang, K.; Guenther, E.; Jin, C. S. Thin Solid Films 2000, 363, 314. (3) Sawada, M.; Higuchi, M.; Kondo, S.; Saka, H. Jpn. J. Appl. Phys. 2001, 40, 3332. (4) Ozasa, K.; Ye, T.; Aoyagi, Y. Thin Solid Films 1994, 246, 58. (5) Wan, Q.; Feng, P.; Wang, T. H. Appl. Phys. Lett. 2006, 89, 123102. (6) Berengue, O. M.; Chiquito, A. J.; Pozzi, L. P.; Lanfredi, A. J. C.; Leite, E. R. Nanotechnology 2009, 20, 245706. (7) Wan, Q.; Dattoli, E. N.; Fung, W. Y.; Guo, W.; Chen, Y.; Pan, X.; Lu, W. Nano Lett. 2006, 6, 2909. (8) Xue, X. Y.; Chen, Y. J.; Liu, Y. G.; Shi, S. L.; Wang, Y. G.; Wang, T. H. Appl. Phys. Lett. 2006, 88, 201907. (9) Yu, D.; Wang, D.; Yu, W.; Qian, Y. Mater. Lett. 2004, 58, 84. (10) Lee, C. H.; Lee, S. W.; Lee, S. S. Nanoscale Res. Lett. 2010, 5, 1128. (11) Maestre, D.; Cremades, A.; Gregoratti, L.; Piqueras, J. J. Phys. Chem. C 2010, 114, 3411. (12) Maestre, D.; Martínez de Velasco, I.; Cremades, A.; Amati, M.; Piqueras, J. J. Phys. Chem. C 2010, 114, 11748. (13) Díaz-Guerra, C.; Vila, M.; Piqueras, J. Appl. Phys. Lett. 2010, 96, 193105. (14) Aleman, B.; Hidalgo, P.; Fernandez, P.; Piqueras, J. J. Phys. D: Appl. Phys. 2009, 42, 225101. (15) Li, S. Y.; Lee, C. Y.; Lin, P.; Tseng, T. Y. Nanotechnology 2005, 16, 451. (16) Magdas, D. A.; Cremades, A.; Piqueras, J. Appl. Phys. Lett. 2006, 88, 113107. (17) Hidalgo, P.; Liberti, E.; Rodríguez-Lazcano, Y.; Mendez, B.; Piqueras, J. J. Phys. Chem. C 2009, 113, 17200. (18) Ortega, Y.; Fernandez, P.; Piqueras, J. J. Cryst. Growth 2009, 311, 3231. (19) Ovidko, I. A.; Sheinerman, A. G. Philos. Mag. 2004, 84, 2103. (20) Fan, H. J.; Knez, M.; Scholz, R.; Nielsch, K.; Pippel, E.; Hesse, D.; Zacharias, M.; G€osele, U. Nat. Mater. 2006, 5, 627. (21) Morin, S. A.; Bierman, M. J.; Tong, J.; Jin, S. Science 2010, 328, 476. (22) Ortega, Y.; Dieker, Ch.; J€ager, W.; Piqueras, J.; Fernandez, P. Nanotechnology 2010, 21, 225604. (23) Bierman, M. J.; Lau, Y. K. A.; Kvit, A. V.; Schmitt, A. L.; Jin, S. Science 2008, 320, 1060. (24) Lau, Y. K. A.; Chernak, D. J.; Bierman, M. J.; Jin, S. J. Am. Chem. Soc. 2009, 131, 16461. (25) Sears, G. W. Acta Metall. 1955, 3, 367. (26) Frank, F. C. Acta Crystallogr. 1951, 4, 497. (27) Maestre, D.; H€aussler, D.; Cremades, A.; J€ager, W.; Piqueras, J. Cryst. Growth Des. 2011, 11, 1117. (28) Morin, S. A.; Jin, S. Nano Lett. 2010, 10, 3459. (29) Jin, S.; Bierman, M. J.; Morin, S. A. J. Phys. Chem. Lett. 2010, 1, 1472. (30) Heindl, J.; Strunk, H. P. Phys. Status Solidi B 1996, 193, K1.
18087
dx.doi.org/10.1021/jp204579u |J. Phys. Chem. C 2011, 115, 18083–18087
