Nondestructive Monitoring of Defect Evolution in ... - ACS Publications

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Nondestructive Monitoring of Defect Evolution in Epitaxial CdTe Thin Layers Grown on Si(111) J. M. Oliveira,† A. Malachias,‡ C. A. Ospina,§ and S. O. Ferreira*,† †

Departamento de Física, Universidade Federal de Viçosa, Viçosa, MG, Brazil Departamento de Física, Universidade Federal de Minas Gerais, Belo Horizonte, MG, Brazil § Laboratório Nacional de Nanotecnologia, Centro Nacional de Pesquisa em Energia e Materiais, Campinas, SP, Brazil ‡

ABSTRACT: Reduction of crystalline defect density and size is a major issue in optimization of thin film for industrial scale applications. It is particularly hard to achieve defect minimization due to the lack of statistically relevant techniques where defectinduced strain information is also available. In this work, we use a combination of highresolution transmission electron microscopy and synchrotron X-ray diffraction to investigate the evolution of defects during the growth of highly mismatched cadmium telluride (CdTe) films on silicon(111) as a function of growth temperature and layer thickness. High-resolution transmission electron microscopy was employed to show that the grown layers are epitaxial, following the [111] substrate orientation, and to identify the main defect type (double twins). Detailed reciprocal space maps obtained by synchrotron X-ray diffraction were used to identify the presence of additional intensities near the (113) and (002) CdTe bulk reciprocal space positions caused by these defects. Analysis of these maps allowed to retrieve the strain of the atoms inside defects as well as to evaluate their average size and density. We show that defect density decreases as a function of temperature but increases as a function of layer thickness, while defect size and the elastic energy of the atoms inside each defect increases up to 400 °C.

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INTRODUCTION Thin film epitaxy has been a long-standing requirement for many optoelectronic applications that are nowadays already integrated in many commercial devices. It is crucial to understand in general heteroepitaxial growth whether a given compound film will be able to keep the registry of a host substrate and develop a limited mosaic spread or will become fully incoherent, resulting in polycrystalline grains without any crystallographic orientation. In particular, Si substrates are standard platforms for growth of many semiconductor devices due to its well-known physical properties, high crystalline quality, and low price. It is crucial to determine if semiconductors of interest for applications can grow coherently on Si surfaces even when large lattice mismatches and defects are present. The onset of different defect types and their density, which depend on growth parameters such as coverage and substrate temperature, has to be understood in order to establish possible limitations of a given deposition procedure. Cadmium telluride (CdTe) plays a significant role as one of the most important materials for infrared and gamma detectors and for solar cells in the form of bulk crystals, thin/thick layers, and quantum dots.1−4 Heteroepitaxial growth of CdTe on Si substrate is a key technology for low cost mercury cadmium telluride (MCT) infrared detectors.5 The main interest in using CdTe/Si as substrate for the epitaxial growth of MCT is the trend toward mega-pixel focal-plane arrays and the possibility of integrating them with Si-based readout electronics.6 CdTe/Si heterojunctions are also a promising way to achieve large-area, thick and high-quality crystals for X-ray and gamma ray © 2014 American Chemical Society

detectors for medical imaging applications working up to 140 keV energy range7 since bulk CdTe crystals are expensive, extremely soft, and tend to develop a high density of dislocations.8 Although remarkable progress has been achieved using nominal and misoriented Si(100), Si(111), and Si(211) substrates, the large difference in thermal expansion coefficients and the lattice constant mismatch of approximately 19% usually result in a high density of dislocations and other stacking defects, especially close to the interface.9,10 Previous studies from our group have shown that this system follows the Volmer−Weber growth mode with initial formation of almost perfect CdTe islands and that most defects are formed during the coalescence process as the growth proceeds.11,12 In this work, we investigate the evolution of defects in CdTe thin films grown on Si(111) substrates, aiming to elucidate what is the dominant defect type and to understand its relative density and size evolution as a function of growth time and temperature. By combining high-resolution transmission electron microscopy (HRTEM) and X-ray diffraction (HRXRD) using a synchrotron source, we were able identify the main defect type formed and to retrieve quantitative structural information about it. We show that the use of reciprocal space maps comprising the region where defect scattering is observed allows a direct evaluation of defect size Received: September 24, 2013 Revised: January 9, 2014 Published: January 13, 2014 1968

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and strain providing a suitable combination of nondestructive analysis and statistical relevance.

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EXPERIMENTAL SECTION All samples were grown on silicon(111) substrates by molecular beam epitaxy (MBE) using a homemade growth system. High purity polycrystalline CdTe (99.999%) was evaporated from a single effusion cell maintained at 580 °C. The substrates were degreased, dipped in a 2% HF solution for 2 min and thoroughly rinsed in deionized water, just before introduction in the growth system, to eliminate the native SiO2 layer and produce a hydrogen passivated growth surface. The deposition times varied from about 40 min to more than 12 h at a growth rate of 0.2 Å/s, resulting in film thickness between 50 and 900 nm, as determined by profilometry and/or atomic force microscopy (AFM). Height profiles were measured using a XP1 stylus profiler (AMBIOS). AFM images of the grown samples were obtained using a NTEGRA Prima scanning probe microscope (NT-MDT) operating in the intermittent contact mode. All images shown have a lateral scan range of 2 μm. Reciprocal space maps (RSM) near (113) and (002) bulk CdTe reflection were measured by synchrotron HRXRD at the XRD2 beamline of the Brazilian Synchrotron Light Laboratory (LNLS). The energy of the incident photons was fixed to 8 keV, corresponding to a wavelength of 1.5498 Å. The RSMs were collected using a Pilatus 100k detector (DECTRIS), with 487 × 195 elements and pixel size of 172 × 172 μm2. The detector was placed 1.0 m from the sample, covering a 4.8° scattering angle (2θ range). HRTEM measurements were performed at the Brazilian Nanotechnology National Laboratory using a JEM-3010 URP TEM (JEOL) with LaB6 thermionic electron gun, operating at 300 kV, with lateral spatial resolution of 0.17 nm and spherical and chromatic aberration coefficients of 0.7 and 1.2 mm, respectively. The TEM cross-section specimens were prepared along the Si[101] zone axis by manual and dimpler polishing followed by argon ion milling, with variable energy from 3.5 to 2 keV and low angles between 2 and 7.5°, for final thinning. High-resolution images were acquired using a 1024 × 1024 charge coupled device digital camera, driven by Gatan’s DigitalMicrograph software. Multislice simulations were carried out with the Java Electron Microscopy Software (JEMS),13 using supercell files obtained from a crystal model generated by MEGACELL software.14

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Figure 1. AFM images showing the effect of growth temperature and layer thickness on surface morphology: (a) T = 300 °C; d = 500 nm. (b) T = 400 °C; d = 500 nm. (c) T = 400 °C; d = 50 nm.

supported by a series of growth models, which consider the presence of different kinetic barriers and/or thermodynamically driven growth instabilities.15,16 The effect of the total coverage (film thickness) is evidenced by comparing Figure 1b,c, which show the AFM images of two samples grown at the same temperature (400 °C) but having thickness of 500 and 50 nm, respectively. As expected, roughness increases with growth time. It is worth mentioning that the pyramidal features observed in these images support the idea that these MBE layers are formed by coalescence of 3D islands nucleated on Si surface, as observed for CdTe thin film samples grown by hot wall epitaxy (HWE).11 In order to understand the crystalline registry of the studied films, HRTEM was performed in selected samples. Figure 2 shows a HRTEM image of the CdTe/Si interface for a 500 nm sample grown at 300 °C, demonstrating that the layers grow epitaxialy, as can be seen by the alignment of lattice planes of Si substrate and CdTe layer. A previous study on CdTe quantum dots grown on Si(111) by HWE has already shown that this system follows the Volmer−Weber growth mode, and despite a lattice mismatch of about 19%, almost perfect epitaxial islands (with very reduced mosaic) are formed.12 HRTEM was also used to identify the most frequent defect type observed in the studied samples. The defect, observed in all samples analyzed independently of growth temperature and thickness, is the double twin, which is shown in detail in Figure 3a. Figure 3b shows the fast Fourier transform (FFT) image of

RESULTS AND DISCUSSION

A first approach to the thin film structure can be drawn by performing simple topographic AFM measurements. AFM results reveal that all samples studied here have essentially the same surface morphology, with surface roughness larger than 10 nm. The typical morphology behavior as a function of growth temperature and thickness is shown in Figure 1, where AFM images of 3 samples are shown. Figure 1a,b shows samples with the same thickness (500 nm), grown at 300 and 400 °C, respectively, illustrating the effect of growth temperature. Comparing the z range of the two images, one can directly see that surface roughness increases by a factor of at least 3 when growth temperature increases from 300 to 400 °C. Although contrary to the common sense, which would expect a decrease in roughness with larger adatom diffusion length, such an increase has already been observed for other systems. This behavior of surface roughness dependence with temperature is 1969

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reflection is expected, thus not interfering with defect scattering. For all samples, besides the (113) Bragg peak, one observes an additional peak that can be associated with the presence of defects. Although, at this point, one cannot directly associate these peaks to a specific type of defect, their intensity ratio to the (113) reflection can be used as a measurement of the relative defect density.17 As observed in the lower panels of Figure 4, the reciprocal space maps for all samples reveal intensity streaks near the (002) reciprocal space condition. Such streaks are crystal truncation rods arising from the defect planes, and their reciprocal space width can be associated to the defect size. However, the low incident scattering angles related to this reciprocal space condition do not allow for a complete mapping of these reflections, which would be necessary for defect type and symmetry identification. The RSMs of the upper panels, near the (113) reflection, can be completely mapped out, as shown in Figure 4. In the cases shown here, all maps were measured for 500 nm thick samples, but growth temperatures varying from 300 to 400 °C. In all maps, the CdTe lattice peak is observed at the (113) position, while a peak that can be assigned to the defects is seen near the (0.667, 0.667, 2.667) reciprocal space condition. Several quantitative direct comparisons can be drawn based on the data extracted from the (113) maps. A very simple approach to identify modifications in the defect density is obtained by evaluating the intensity ratio between the CdTe (113) film peak and the defect peak. Since the integrated intensity of both features is directly related to the volume of the lattice in each condition (ordered or defective) one can estimate the relative defect density by considering a fixed normalized ratio of 1 at the 300 °C sample. The data, summarized in Table 1, show that from 300 °C (used here as a reference) to 350 °C, the defect density has been reduced by a factor of ∼200, while from 300 to 400 °C a factor of ∼40 is obtained. Such behavior is expected since a higher growth temperature increases surface adatom mobility leading to the formation of a reduced number of defects as compared to lower temperatures. At 400 °C the defect density slightly increases. We must recall here that the defect formation is by itself dependent on an activation energy that was probably overcome at this higher temperature, where larger defects are observed. Similarly to this evaluation as a function of the temperature, we can find out if the defect density depends on the film thickness. Samples grown at 300 °C with 50 nm CdTe coverage (RSMs not shown here) exhibited a peak ratio that indicates a 40% increase in the defect density as the film thickness reach 500 nm. This behavior could be explained by the results obtained for CdTe layers grown on Si(111) by HWE,12 which show that they are formed by coalescence of 3D islands with a high degree of mosaicity, and an in-plane rotation of up to 30°. The complete coalescence of islands, not reached at 50 nm thickness, is effectively obtained for the 500 nm thick film, thus increasing the defect concentration. The defect sizes can be retrieved by evaluating the reciprocal space widths of the peaks along the main crystallographic directions for the CdTe system. Such results are obtained from scans along the (111) direction, which provide the average defect length, along its longitudinal axis, and scans along the (110) direction, that allow for the extraction of their average width. The main results are also summarized in Table 1. As the temperature increases, the average defect length increases from 14.0 nm at 300 °C to 21.8 nm at 350 °C and 35.5 nm at 400 °C. This behavior is followed by an increase in defect width as

Figure 2. HRTEM image of the CdTe/Si interface, showing the alignment of the film and substrate crystalline lattices near the interface.

Figure 3. (a) Detailed HRTEM image of the double twin defect. (b) Fast Fourier transform of image panel a showing spots related to the defect.

the same region and allows the identification of the diffraction peaks corresponding to the defect. Although this result unambiguously shows the presence of an additional lattice organization, it lacks statistical relevance and does not allow for quantitative estimation of defect size and strain. Detailed structural characterization was addressed by synchrotron X-ray diffraction, the technique of choice to provide averaged information about the deposited layers. Figure 4 shows reciprocal space maps measured in the vicinity of the (113) and (002) CdTe reflections (upper and lower panels, respectively) for three selected samples. In this figure, all axes are shown in reciprocal lattice units (R.L.U.) using as reference the bulk CdTe lattice. These reflections were chosen since they can reveal the symmetry of the defect lattice. A strong CdTe lattice peak is expected at the (113) position, providing a lattice reference, on the other side, at the (002) position, a weak 1970

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Figure 4. Reciprocal space maps in the vicinity of CdTe (113) reflection (upper panels) and (002) reflection (lower panels) for samples grown at 300 °C (left panels), 350 °C (middle panels), and 400 °C (right panels). The regions outside the dashed lines in the lower panel maps were not mapped. All samples are 500 nm thick. All axes are scaled in reciprocal lattice units (R.L.U.) using the bulk CdTe lattice as reference.

Table 1. Summary of the Information Retrieved for the Defects from the Interpretation of Reciprocal Space Map Data in the Vicinity of the CdTe (113) Reflection growth temperature

relative defect density

average defect length, L (nm)

average defect width, w (nm)

longitudinal strain inside defect, εl (%)

transversal strain inside defect, εt (%)

300 °C 350 °C 400 °C

1.000 0.006(1) 0.030(1)

14.0(2) 21.8(3) 35.5(3)

5.9(1) 8.0(2) 10.2(2)

0.14(1) 0.10(1) 0.03(1)

0.65(1) 0.55(1) 0.33(1)

well, ranging from 5.9 nm at 300 °C to 8.0 and 10.2 nm at 350 and 400 °C, respectively. Therefore, although the defect density decreases for higher temperatures, the average defect size becomes larger.17 Since defect formation is a thermally activated phenomenon, as stated before, this latter effect is directly related to the enhancement of atomic diffusion at higher temperatures. Since at higher growth temperatures adatoms impinging the surface exhibit larger mobility they are able to reach the preferential nucleation sites, which are in the vicinity of the defects, leading to the observed defect elongation as well as to the lower defect density. As mentioned before, the HRTEM images were used to provide information of defect type and detailed atomic positions. An image of a double twin defect was reconstructed using the MEGACELL software to build an atomic model, which can serve as input for XRD simulations. Such reconstruction is validated by direct comparison of the experimental image and a simulated one, calculated using the JEMS software. Figure 5 shows a superposition of an experimental HRTEM image, the structural model, and the simulated image (highlighted region). The simulated image was obtained with an objective defocus of 41 nm, 2% of noise, 0.084

Figure 5. Superposition of an experimental HRTEM image, the structural model, and a simulated image.

nm·s−1 of drift, and 0.01 nm of vibration in x and y directions.18 The crystal model, where Cd and Te atoms are in red and yellow colors (respectively), was constructed with a thickness of 22.7 nm. Three different crystalline regions, with the same CdTe structure, were labeled as Block 1, Twin region, and Block 2, with a relative in-plane rotation of 109.7 ± 0.1 between the twin region block and blocks 1 and 2, which have the same orientation. This rotation aligns two (111) reflections of blocks 1971

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defects.20 In particular, the regions limited by dashed lines in Figure 6a are seen in more detail in Figure 6b,c. Since the simulated reciprocal space map considers bulk atomic distances and only geometrical symmetry operations are used to produce the defects, the defect peak near the (113) CdTe position is centered exactly at the (2/3, 2/3, 1/3) position. By comparing its position with the positions measured in the maps of the upper panels of Figure 4, one can directly obtain the strain of the atoms inside the double twin defects. We have compared here the position of (111) and (101) cuts of the measured maps with the positions of the respective cuts in the simulated maps, allowing the evaluation of the longitudinal and transversal strain of the atoms inside the defects. We obtained that atoms inside the defects are subjected to compressive strain in both directions along the temperature range used for the investigated samples. As summarized in Table 1, the strain along the defect axis is always smaller than the strain obtained along the direction transversal to the defect. As the temperature increases a reduction of both quantities is observed. A simple comparison of the relative defect energy can also be drawn among the studied samples. Let us consider that defects are cylindrical and that the elastic energy of the atoms inside the double twin boundaries (Ed) is proportional to their volume (therefore, the square of their width (w) times their length (L)) and proportional to the sum of the square of the strains in each direction.21 In our case Ed ∝ Lw2(2εt2 + εl2), where εt and εl are the transversal and longitudinal strains, respectively. By applying this simple rule and normalizing to the unity the energy of each defect at 300 °C, one finds out that the energy per defect approximately doubles (increases by a factor 2.1) at 350 °C and is almost three times larger at 400 °C (increases by a factor of 2.8). Therefore, in spite of the lower strain, the energy per defect increases significantly due to their size at higher growth temperatures. The energy evaluation in the previous paragraph takes into account only the atoms inside the double twin defects. However, the energy at their boundaries (defect cores) are, in principle, higher or as significant as the strain energy stored inside the twins. A simple approximation to defect energy (Ud) due to the stacking fault of atoms on one side of the defect is given by

1 and 2, with its inverse for the twin region block. To correctly simulate the phase contrast image of the twin boundaries and obtain a good high-resolution simulated image as shown in Figure 5, a separation of 2.776 Å between the blocks was set. The structural model obtained using this procedure was used as input to simulate the X-ray diffraction maps with a matrix of atoms using a Matlab script. In our simulation, the double-twin X-ray scattering is retrieved by the following method.19 Initially the atomic coordinates, considering the defect atoms and atoms surrounding it, are embedded into a large spherical crystal (diameter of about 300 Å) with atoms in the bulk lattice sites. The coordinates of the defect atoms are then restricted to follow the bulk spacing, but rotated/translated to match the TEM coordinates. In order to take into account distinct defect orientations observed in the CdTe film, we averaged the scattering over all equivalent orientations of the simulated defective-crystal. Finally, given the reduced defect size in all directions, the X-ray kinematical approximation is used to calculate the diffracted intensity by summing over the scattering of each atom, as expressed below: 2

I(q) = A 0 ∑ fi e

−σ 2R i 2 /2a 2

e

iqR i

i

where q is the reciprocal space condition in which the scattering intensity is evaluated, A0 is a constant used for normalizing the total scattering intensity, f i is the atomic scattering factor of the atom at position i, σ is an effective fluctuation of the defect/ host crystal interface position, set to 0.05, which allows for broadening the profile, washing out finite size oscillations from the simulated defects, and a is the lattice parameter. The result of a (HHL) map (K = H) ranging from −3 ≤ H ≤ 3 to 0 ≤ L ≤ 6 map is shown in Figure 6a. The simplified

Ud = (μb2 /4π )[(1 − v cos2 β )/(1 − v)] ln(4h/b)

where b is the Burgers vector for the defect type, μ is the CdTe shear modulus, ν is the CdTe Poisson’s ratio, β is the angle between the Burgers vector and the dislocation line (differs from zero if the defect has a screw component), and h is the average distance among defects.22 Since, except for their density, all other parameters are similar for the investigated samples, one expects that h, and therefore Ud, should scale with the cubic root of the defect density. In our case, the relative defect core energy for films grown at each temperature studied here will be 1 at 300 °C, 0.182 at 350 °C, and 0.311 at 400 °C. Therefore, because of its reduced defect density and average energy of the atoms inside the defects, CdTe films grown at 350 °C are subjected to a minimum energy spent on defect formation and can be considered as the higher quality CdTe films of the series investigated here.

Figure 6. (a) Simulated reciprocal space map of the (HHL) plane for CdTe crystals containing double twin defects. (b,c) Detailed views of the regions in the vicinity of the (113) and (002) CdTe reciprocal space positions. Axis are shown in R.L.U. using the bulk CdTe lattice as reference.

model used here directly provides the position and symmetry of the additional Bragg peaks due to the scattering from the defect and, although not reproducing the exact shape of the experimental results, can be used to infer the defect core strain and the average defect size. According to Figure 6, several extra peaks are generated besides the bulk crystal peaks due to the presence of twin

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CONCLUSIONS We have used reciprocal space maps and high-resolution transmission electron microscopy to investigate the effect of 1972

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Based on Epitaxially Grown Thick CdTe on n(+)-Si Substrates. IEEE Electron Device Lett. 2006, 27, 890−892. (8) Ward, D. K.; Zhou, X.; Wong; Doty, F. P.; Zimmerman, J. A. Accuracy of Existing Atomic Potentials for the CdTe Semiconductor Compound. J. Chem. Phys. 2011, 134, 244703−244703−13. (9) Smith, D. J.; Tsen, S. C. Y.; Chandrasekhar, D.; Crozier, P. A.; Rujirawat, S.; Brill, G.; Chen, Y. P.; Sporken, R.; Sivananthan, S. Growth and Characterization of CdTe/Si Heterostructures: Effect of Substrate Orientation. Mater. Sci. Eng., B 2000, 77, 93−100. (10) Benson, J. D.; Jacobs, R. N.; Markunas, J. K.; Jaime-Vasquez, M.; Smith, P. J.; Almeida, L. A.; Martinka, M.; Vilela, M. F.; Lee, U. Structural Analysis of CdTe Hetero-Epitaxy on (211) Si. J. Electron. Mater. 2008, 37, 1231−1236. (11) Ferreira, S. O.; Paiva, E. C.; Fontes, G. N.; Neves, B. R. A. Characterization of CdTe Quantum Dots Grown on Si(111) by Hot Wall Epitaxy. J. Appl. Phys. 2003, 93, 1195−1198. (12) Suela, J.; Ribeiro, I. R. B.; Ferreira, S. O.; Malachias, A.; Fontes, G. N.; Montoro, L. A.; Ramirez, A. J. Evolution of Crystalline Domain Size and Epitaxial Orientation of CdTe/Si(111) Quantum Dots. J. Appl. Phys. 2010, 107, 064305−064305−6. (13) Stadelmann, P. A. EMS: A Software Package for ElectronDiffraction Analysis and HREM Image Simulaton in Materials Science. Ultramicroscopy 1987, 21, 131−145. (14) Stroppa, D. G.; Righetto, R. D.; Montoro, L. A.; Ramirez, A. J. MEGACELL: A Nanocrystal Model Construction Software for HRTEM Multislice Simulation. Ultramicroscopy 2011, 111, 1077− 1082. (15) Elsholz, F.; Schöll, E.; Rosenfeld, A. Control of Surface Roughness in Amorphous Thin-Film Growth. Appl. Phys. Lett. 2004, 84, 4167−4169. (16) Leal, F. F.; Oliveira, T. J.; Ferreira, S. C. Kinetic Modelling of Epitaxial Film Growth with up- and Downward Step Barriers. J. Stat. Mech. 2011, P09018−P09018−15. (17) Nordlund, K.; Metzger, T. H.; Malachias, A.; Capello, L.; Calvo, P.; Claverie, A.; Cristiano, F. Measurement of Si 311 Defect Properties Using X-ray Scattering. J. Appl. Phys. 2005, 98, 073529−073529−7. (18) Bettini, J.; Rodrigues, V.; González, J. C.; Ugarte, D. Real-Time Atomic Resolution Study of Metal Nanowires. Appl. Phys. A: Mater. Sci. Process. 2005, 81, 1513−1518. (19) Salonen, E.; Nordlund, K.; Keinonen, J.; Runeberg, N.; Wu, C. H. Reduced Chemical Sputtering of Carbon by Silicon Doping. J. Appl. Phys. 2002, 92, 2216−2218. (20) Richard, M. I.; Malachias, A.; Rouviere, J. L.; Yoon, T.; Holmstrom, E.; Xie, Y. H.; Favre-Nicolin, V.; Holy, V.; Nordlund, K.; Renaud, G.; et al. Tracking Defect Type and Strain Relaxation in Patterned Ge/Si(001) Islands by X-ray Forbidden Reflection Analysis. Phys. Rev. B 2011, 84, 075314−075314−9. (21) Chou, P. C.; Pagano, N. J. Elasticity: Tensor, Dyadic and Engineering Approaches; Dover Publications: Mineola, NY, 1992. (22) Tsao, J. Y. Materials Fundamentals of Molecular Beam Epitaxy; Academic Press: San Diego, CA, 1992.

growth temperature and growth time on defects in CdTe/ Si(111) thin layers grown by molecular beam epitaxy. HRTEM shows that, despite the very big lattice mismatch of almost 19%, the layers grow epitaxially and that the main defect type is the double twin. A comparison of experimental and simulated HRTEM images is used to validate a structural atomic model, which is applied to simulate the RSMs around reciprocal space positions with nonzero defect-induced X-ray scattering. Such possibility consists in a powerful nondestructive defect characterization tool that can be applied to any material, without any sample preparation requirement. The analysis of the synchrotron X-ray diffraction maps shows unambiguously that at 350 °C both the defect density and the defect energy are minimized with respect to the other growth temperatures. Therefore, the knowledge of defect type and analysis of its structural information from reciprocal space maps results is an invaluable tool to evaluate crystal quality in thin layers, in which defect formation cannot be avoided but must be uttermost reduced for application purposes.

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AUTHOR INFORMATION

Corresponding Author

*(S.O.F.) E-mail: [email protected]. Phone: +55 31 38992015. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work has been supported by “Conselho Nacional de ́ Desenvolvimento Cientifico e Tecnológico (CNPq),” “Cooŕ Superior denaçaõ de Aperfeiçoamento de Pessoal de Nivel (CAPES),” and “Fundaçaõ de Amparo à Pesquisa do Estado de Minas Gerais (FAPEMIG)” through research grants and graduate scholarships. The authors would also like to thank the Brazilian Synchrotron Light Laboratory and the Brazilian Nanotechnology National Laboratory at the Brazilian Center for Research in Energy and Materials for technical support during the X-ray diffraction and electron microscopy work, respectively.

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REFERENCES

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dx.doi.org/10.1021/jp409538p | J. Phys. Chem. C 2014, 118, 1968−1973