REVISTA MEXICANA DE FÍSICA S REVISTA MEXICANA DE FÍSICA S PÁGINAS 1-74 PÁGINAS 1-74 ISSN IX Abril 2011 VOLUMEN 57 NUMERO 2

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1 REVISTA MEXICANA DE FÍSICA S VOL. 57 NUM. 2 ABRIL 2011 PÁGINAS 1-74 PÁGINAS 1-74 Abril 2011 ISSN IX REVISTA MEXICANA DE FÍSICA S VOLUMEN 57 NUMERO 2 ABRIL 2011 PÁGINAS 1-74 VOLUMEN 57 NUMERO 2 ABRIL 2011 PÁGINAS 1-74 CODEN: RMFXFAT Ejemplar $ REVISTA MEXICANA DE FÍSICA S Abril

2 Director: Fransisco Ramos Gómez Facultad de Ciencias, UNAM, México Consejeros Eméritos Marcos Moshinsky Instituto de Física, UNAM, México Leopoldo García-Colín Universidad Autónoma Metropolitana Iztapalapa, México Manuel Peimbert Instituto de Astronomía UNAM, México Fernando Alba Instituto de Física, UNAM, México Consejo Editorial Materia Condensada: Carlos Balseiro Centro Atómico de Bariloche, Argentina Alipio G. Calles Facultad de Ciencias, UNAM, México Manuel Cardona Institute Max Planck, Stuttgart, Alemania Robert Cava University of Princeton, USA Roberto Escudero Instituto de Investigaciones en Materiales, UNAM, México Francisco Jaque Universidad Autónoma de Madrid, España Harold Kroto Florida State University Física Atómica y Molecular: Gerardo Delgado-Barrio Consejo Superior de Investigación Científica, España James McGuire Tulane University, USA Alfred Schlachter Advanced Light Source, LBL Berkeley California, USA Física Nuclear: Alejandro Frank Instituto de Ciencias Nucleares, UNAM, México Arturo Menchaca Instituto de Física, UNAM, México Andrés Sandoval GSI, Alemania & CERN, Suiza Termodinámica y Física Estadística: Eugenio E. Vogel Universidad de la Frontera, Chile Ivan L heureux University of Ottawa, Canada Víctor Romero Instituto de Física, UNAM, México REVISTA MEXICANA DE FÍSICA Óptica: Alejandro Cornejo Instituto Nacional de Astrofísica, Óptica y Electrónica, México Eugenio Méndez CICESE, México Jumpei Tsujiuchi Institute of Technology, Tokio, Japón Fernando Mendoza Centro de Investigaciones en Óptica, México Gravitación y Física Matemática: Octavio Obregón Instituto de Física, Universidad de Guanajuato, México Fernando Quevedo University of Cambridge, Inglaterra Instrumentación: Victor Castaño Centro de Física Aplicada y Tecnología Avanzada, UNAM, México Daniele Finotello Kent State University, USA Partículas Elementales y Teoría de Campo: Gerardo Herrera Corral CINVESTAV, IPN, México Física Médica: María Ester Brandan Instituto de Física, UNAM, México Secretaria: María Magdalena López Reynoso Sociedad Mexicana de Física Revisión de Estilo: José Luis Alvarez García Facultad de Ciencias, UNAM Juan Pablo Flores del Villar Sociedad Mexicana de Física Edición Técnica: Raúl A. Espejel Morales Facultad de Ciencias, UNAM Asistentes Técnicos: Efraín R. Garrido Román Sociedad Mexicana de Física Paris M. Sánchez Carreón Sociedad Mexicana de Física La Revista Mexicana de Física es una publicación bimestral de la Sociedad Mexicana de Física, A.C., Apartado Postal , Coyoacán, México, D.F., MÉXICO. Se publica con el patrocinio de: Instituto Nacional de Astrofísica Óptica y Eléctronica, Puebla, Instituto Potosino de Investigación Científica y Tecnológica, San Luis Potosí, y de la UNAM: Rectoría, Coordinación de la Investigación Científica, Instituto de Astronomía, Centro de Ciencias de la Materia Condensada, Instituto de Ciencias Nucleares, Instituto de Investigaciones en Materiales, Instituto de Física, Facultad de Ciencias e Instituto de Matemáticas. Indizada en: Actualidad Iberoamericana, Astron. & Astrophys. Abstr., Bull. Signal., Chem. Abstr., Curr. Cont., Curr. Math. Pub., Curr. Pap. Phys., Electr. & Electron. Abstr., INIS Atomind., Math. Sci., LatIndex, Math. Rev., Nucl. Sci. Abstr., PERIODICA, Phys. Abstr., Phys. Ber., Res. Alert, Sci. Abstr., Sci. Cit. Ind., y SciSearch. Incluida en el Indice de Revistas Mexicanas de Investigación Científica y Tecnológica del Consejo Nacional de Ciencia y Tecnología (CONACyT). Las instrucciones para autores aparecen publicadas en el número 6 de cada volúmen. El costo de la suscripción anual es de $ pesos para la República Mexicana, $130 USD para América Central y del Norte y $160 USD para el resto del mundo. Precio del ejemplar $170.00

3 Revista Mexicana de Física S ISSN IX Publicación de la Sociedad Mexicana de Física, A.C. Apartado postal , Coyoacán, México, D.F. Director: Fransisco Ramos Gómez Oficinas: 2 piso, Departamento de Física, Facultad de Ciencias, Ciudad Universitaria, México, D.F. Tel.: (55) ; FAX: (55) Se autoriza la reproducción parcial o total de su contenido citando la fuente: Rev. Mex. Fis S. Los artículos firmados son responsabilidad de los autores. Certificado de licitud número 3110 y de contenido número 2775 otorgado por la Comisión Calificadora de Publicaciones y Revistas Ilustradas de la Secretaría de Gobernación. Reserva del título número de la Dirección General de Derechos de Autor. Publicación periódica: Registro número , características , otorgado por la oficina del Servicio Postal Mexicano. El volumen 57, número 2, abril de 2011, se terminó de imprimir en abril de 2011; se tiraron 100 ejemplares impresos y 100 electrónicos. Impresión: Impresiones Integradas del Sur, S.A. de C.V., Amatl No. 20, Col. Santo Domingo, Delegación Coyoacán, México, D.F., Tel.: Diseño de portada: Arte Gráfico, Sur 71 No. 501, Col. Justo Sierra, México, D.F. Impreso en México Printed in Mexico

4 i VI International Topical Meeting on Nanostructured Materials and Nanotechnology, Nanotech 2009 San Carlos, Nuevo Guaymas September 17-19, 2010 Editor: MARCELINO BARBOZA FLORES BEATRIZ DEL CARMEN CASTANEDA MEDINA ALVARO POSADA AMARILLAS RAFAEL GARCIA GUTIERREZ ELDER DE LA ROSA CRUZ

5 ii ORGANIZING COMMITTEE Rafael García Gutiérrez Alvaro Posada Amarillas Elder de la Rosa Cruz Marcelino Barboza Flores

6 iii PREFACIO El Sexto Encuentro Internacional sobre Materiales Nanoestructurados y Nanotecnología, NANOTECH 2009, es un congreso internacional que se ha organizado en la República Mexicana desde el año La primera reunión se llevó a cabo en las instalaciones del Centro de Investigación en Óptica, en León, Guanajuato. En el 2005 se organizó en el Centro de Ciencias de la Materia Condensada-UNAM en Ensenada. En el 2006 se llevó a cabo en la ciudad de Puebla en el Instituto de Física-BUAP. Monterrey fue la sede de la organización de la Conferencia del 2007, en las instalaciones de la Universidad Autónoma de Nuevo León. En el 2008 el congreso se celebró en Ciudad Universitaria-UNAM en México D.F. En esta ocasión la reunión fue organizada del 17 al 19 de septiembre del 2009 por el Departamento de Investigación en Física de la Universidad de Sonora en San Carlos Nuevo Guaymas. El congreso fue el escenario para la presentación de 2 cursos cortos, 2 mesas redondas para tratar los temas de las aplicaciones de la nanotecnología en la industria, 15 exposiciones orales y 88 carteles, además de 10 conferencias magistrales presentadas por empresarios sonorenses y científicos de renombre mundial. El objetivo principal del NANOTECH 2009 fue el de proporcionar un foro para que científicos, ingenieros y empresarios busquen solución a problemas científicos que conduzcan a aplicaciones prácticas. Los tópicos que se trataron en dicho evento comprendieron desde ciencia básica hasta aplicaciones y técnicas de comercialización de alta tecnología. Algunos de los principales temas discutidos aquí fueron los nanotubos de carbono, nanomateriales magnéticos y nanoestructuras metálicas (plasmones), celdas solares y de combustible, nanofósforos incluyendo óxidos, nitruros, tierras raras, y orgánicos; nanomedicina, nanocristales lineales y no lineales y cristales fotónicos. El comité organizador agradece profundamente el apoyo financiero otorgado por la Universidad de Sonora, la Dirección Adjunta de Desarrollo Científico y Académico, CONACYT (México), Red Temática de Nanociencias y Nanotecnología, y las empresas, Rubio Pharma y Asociados y RD Research & Technology.

7 REVISTA MEXICANA DE FÍSICA S 57 (2) 1 6 ABRIL 2011 Catalytic activity of MoS 2 nanotubes in the hydrodesulphurization reaction of dibenzothiophene F. Leonard-Deepak a,b, R.Pérez-Hernández b,c, J. Cruz-Reyes d, S. Fuentes e, and M.J. Yacaman,b a International Iberian Nanotechnology Laboratory, Avda Mestre Jose Veiga, Braga 4715, Portugal. b Department of Physics and Astronomy, One UTSA Circle, The University of Texas at San Antonio, Texas, 78249, USA, c Instituto Nacional de Investigaciones Nucleares, Carr. México-Toluca S/N La Marquesa, Ocoyoacac, Edo. de México 52750, México. d Facultad de Ciencias Químicas e Ingeniería, Universidad Autónoma de Baja California, Tijuana, B.C., México. e Centro de Nanociencias y Nanotecnología de la Universidad Nacional Autónoma de México, Km. 107 Carretera Tijuana-Ensenada, Apartado Postal, 356, Ensenada, B.C., 22800, México. Recibido el 20 de noviembre de 2009; aceptado el 18 de enero de 2010 In the need for developing better fuels and as a consequence better hydrodesulphurization catalysts (HDS), new generations of catalysts are necessary to reduce substantially the sulfur content in diesel and gasoline fuels. HDS are catalytic processes that involve Mo or W- based catalysts, often doped with other transition metals. We synthesized MoS 2 nanotubes by reacting MoO 3 with thiourea and used them as catalysts for the hydrodesulfurization of dibenzothiophene in a batch reactor. X-ray diffraction, scanning electron microscopy, and transmission electron microscopy techniques were used to characterize their morphology and structure. The results indicated the hexagonal crystalline structure of MoS 2 and large yields of the MoS 2 nanotubes with unusual square or rhomboid faceted shapes. The catalytic behavior of the MoS 2 nanotube catalysts showed that the direct desulfurization pathway prevailed over the hydrogenation (HYD) pathway. This finding was attributed to the low rim/edge sites ratio, induced by the size and morphology of the nanotubes showing large flat area which is responsible for the biphenyl (BP) selectivity. Keywords: Hydrodesulfurization; selectivity; dibenzothiophene (DBT); molybdenum sulfide (MoS 2); nanotubes; TEM. En la necesidad de desarrollar mejores combustibles y como consecuencia mejores catalizadores para la hidrodesulfuracion (HDS), nuevas generaciones de catalizadores son necesarios para reducir sustancialmente el contenido de azufre en los combustibles diesel y gasolina. HDS es un proceso catalítico que involucra catalizadores basados en Mo y W, a menudo dopados con otros metales de transición. Se sintetizaron nanotubos de MoS 2 reaccionando MoO 3 con thiourea. Los nanotubos se utilizaron como catalizadores para la hidrodesulturacion de dibenzotiofeno en un reactor discontinuo (batch reactor). Las técnicas de difracción de rayos X, microscopía electrónica de barrido y de transmisión fueron utilizadas para caracterizar la morfología y la estructura de los catalizadores. Los resultados mostraron la estructura cristalina hexagonal del MoS 2 y grandes rendimientos de nanotubos de MoS 2 con formas facetadas cuadradas o romboidales inusuales. El comportamiento catalítico de los nanotubos de MoS 2 demostró que la vía de desulfuración directa prevaleció sobre la vía de hidrogenación (HYD). Este resultado se atribuyó a la baja relación diámetro/superficie (rim/edge), inducida por el tamaño y morfología de los nanotubos, mostrando un área grande y plana, que es la responsable de la selectividad del bifenilo (BP). Descriptores: hidrodesulfuracion; selectividad; dibenzotiofeno; sulfuro de molibdeno (MoS 2); MET. PACS: Hc, cp, Lp, Og 1. Introduction Elimination of sulfur from petroleum feedstocks is necessary in order to meet the severe restrictions on the sulfur concentrations in fuels [1,2]. The hydrodesulfurization (HDS) of polyaromatic sulfur compounds or deep HDS is especially difficult for the case of heavy oils containing high concentration of sulfur (2 3 wt %). Catalysts based on molybdenum sulfide are widely used in oil refineries for the HDS, hydrodenitrogenation (HDN) and hydrodeoxygenation (HDO) reactions of petroleum-derived feedstocks [3-5]. Due to the stringent environmental legislation that set the sulfur level at 15 ppm, new catalysts with significantly improved catalytic performance must be developed. Sulfur compounds that are known to remain in fuels such as diesel at sulfur levels below 500 ppm include dibenzothiophene (DBT) and alkyl-substituted DBT s such as 4,6- dimethyldibenzothiophene (4,6-DMDBT) [6,7]. The HDS generally proceeds through two pathways: a hydrogenation (HYD) pathway involving aromatic ring hydrogenation and a hydrogenolysis pathway via direct C S bond cleavage, also called the direct desulfurization (DDS) pathway [8]. The contribution of both pathways defining the selectivity depends on the catalyst type. The HDS of DBT in CoMo catalysts occurs predominantly via the DDS pathway yielding HYD/DDS ratios from 0.3 to 0.5. However, for the HDS of 4,6-DMDBT [5], due to the steric hindrance of the methyl groups it is necessary for the hydrogenation of at least one aromatic ring before the elimination of sulfur. In

8 2 F.L. DEEPAK, R.P. HERNÁNDEZ, J. CRUZ-REYES, S. FUENTES, AND M.J. YACAMAN that case, new catalysts with higher specific hydrogenolysis activity and/or higher hydrogenation activity are required. The addition of acid functionality through the use of zeolite [9-11] or amorphous alumina-silicate supports [11,12] to the standard promoted molybdenum sulfide-based catalysts led to noticeable enhancement in the HDS of alkylsubstituted DBT enabling the dealkylation and isomerization of the alkyl substituents, thereby transforming the refractory components into more reactive species. Acidic supports have also improved the catalytic performance of the catalyst particles by increasing their electron deficient character, resulting in greater sulfur resistance and intrinsic activity [13-15]. However, support acidity is also associated with catalyst deactivation by coke formation [16], a phenomenon that led to numerous efforts to fine-tune the effects of the support acidity [17-21]. MoS 2 nanoparticles can have different morphologies depending on the preparation conditions (nanotubes, nanorods, onion-like nanoparticles, 2D nanoparticles, etc). All the morphologies are derived from its layer structure in which atoms within a layer are bound by strong covalent forces, while individual layers are held together by van der Waals interactions. The stacking sequence of the layers can lead to the formation of either a hexagonal polymorph with two layers in the unit cell (2H), rhombohedral with three layers (3R) or trigonal with one layer (1T). Nanotubes of the transition metal dichalcogenides (ex: MoS 2, WS 2 ) have attracted considerable attention in recent years [22-26]. One of the first methods of synthesis of MoS 2 nanotubes was developed by Feldman, et al. [27]. This method involved the gas-phase reaction of MoO 3 x and H 2 S at 850 C in a reducing atmosphere. Nath, et al. [28] used thermal decomposition of ammonium thiomolybdate at higher temperatures, which resulted in the formation of MoS 2 nanotubes. Li, et al. [29] have developed an atmospheric pressure chemical vapor deposition (APCVD) route for the synthesis of MoS 2 nanostructures. These nanostructures, including three-dimensional nanoflowers (NF), were obtained by the reaction of chlorides (MoCl 5 ) and sulfur, under controlled conditions. The measured surface area and field emission of these nanostructures showed them to be promising candidates as catalysts. In all the methods of synthesis outlined so far, the reducingsulphidizing agents included H 2 S (or a mixture of H 2 and H 2 S) and S powder. In general, the methods of synthesis of MoS 2 nanotubes obtained them in low yields ( 20 %) as well as by long tedious procedures (ex: two-stage synthesis). The most important application of MoS 2 is as catalyst for the HDS of fuels, typically, they are evaluated in model test reactions as the HDS of thiophene, dibenzothiophene and 4,6 DMDBT [30-34]. In order to scale the use of MoS 2 nanotubes in catalysis or other applications it is important to devise new synthetic routes to obtain them in large yields. The present work proposes a simple one step synthetic process, using thiourea and MoO 3 as the starting materials to produce large quantities of MoS 2 nanotubes. The resulting nanotubes have unusual faceted shapes (square or rhomboid) which are reported here for the first time. Thiourea has not been employed previously as a sulphur source for making MoS 2 nanotubes; it generates a suitable reducing-sulphidizing environment in-situ, eliminating the use of a separate reducing agent. 2. Experimental methods and characterization 2.1. Synthesis Synthesis of the MoS 2 nanotubes was carried out as follows. About 0.6 g of MoO 3 (mp = 795 C) and 1.0 g of thiourea (CSN 2 H 4,mp = C) were placed in an alumina boat (ratio of Mo:S was kept at 1:2.5 to ensure an excess of the sulphur source). The boat was placed in an alumina tube at the heating zone of a horizontal furnace. Before the reaction the system was flushed with N 2 for 1/2 hr to remove any traces of oxygen. The tube was then heated to 1000 C in flowing N 2 (flow rate = 200 cc/min) [35]. Previously cleaned silicon substrates were placed at regular intervals in the outlet region of the alumina tube to collect the product as a deposit during the course of the reaction. The reaction was carried out for 1 hr, after which it was gradually cooled down to room temperature in flowing N 2. At the end of the reaction the resulting grey colored powder was collected from the alumina boat and the silicon substrate (nanotubes) for further analysis Characterization X-ray diffraction (XRD) powder patterns were recorded in a Siemens D-5000 diffratometer, using Cu Kα (λ= nm) radiation. Scanning electron microscopy (SEM) was performed in a FEG Hitachi S-5500 ultra high resolution electron microscope (0.4 nm at 30 kv) with a BF/DF Duo-STEM detector. Transmission electron microscopy (TEM) and selected area electron diffraction (SAED) were performed using a Tecnai 20 TEM equipped with a Schottky-type field emission gun, ultra-high resolution pole piece (Cs=0.5 mm), and a scanning transmission electron microscope (STEM) unit with high angle annular dark field (HAADF) detector operating at 200 kv Catalytic experiments The HDS of DBT was tested in a 300 ml high pressure Parr reactor by placing 4.4 g DBT, 100 ml of decalin and the calculated amount of precursor needed to produce 0.68 g of catalyst. The reactor was purged of residual air, pressurized with H 2 to 3.1 MPa (450 psi) and then heated to the reaction temperature of 623 K in about 10 min. A stirring rate of 600 rpm was used. The advance of the reaction was monitored by gas chromatography with a HP 6890 gas chromatograph, using samples taken every 20 min during the first hour, then every 30 min for the next four hours. Reduction of sample volume due to sampling was 5% of total volume. The identity of the reaction products was confirmed by mass spectrometry Rev. Mex. Fís. 57 (2) (2011) 1 6

9 CATALYTIC ACTIVITY OF MOS 2 NANOTUBES IN THE HYDRODESULPHURIZATION REACTION OF DIBENZOTHIOPHENE 3 3. Results and discussion FIGURE 1. Powder XRD patterns of the MoS 2 nanotubes. Red lines-mos 2 and blue lines-moo 2 Figure 1 shows the XRD patterns of the MoS 2 nanotubes synthesized using thiourea as the S source and MoO 3 as the Mo source. The XRD patterns are in good agreement with those reported for the hexagonal crystalline structure of MoS 2 (JCPDS ). The principal diffraction peak of the MoS 2 nanotubes appeared at 2θ=14.397, corresponding to the (002) planes, which are a measure of crystal growth in the c direction; similar to the growth of 1D ZnO nanorods [36]. However, a small quantity of monoclinic MoO 2 (JCPDS ) was also identified. This finding showed that it is possible to use this method to obtain MoS 2 nanotubes with high crystallinity and purity. Figure 2 shows SEM micrographs of the MoS 2 nanotubes obtained by the reaction of MoO 3 and thiourea. The large yield and the hollow empty core of the nanotubes are evident in the micrograph in Fig. 2a. A closer look at the nanotubes, Fig. 2b, reveals the unusual faceted shape of the tubes (square or rhomboid). To our knowledge, this is the first time that MoS 2 nanotubes with such unusual faceted shapes have been observed. The hollow empty core of these structures is another outstanding feature seen in Fig. 2a. The nanotubes measure between nm in diameter and extend up to several microns in length. FIGURE 3. (a) BF-TEM image of a facetted MoS 2 nanotube. A close up image of the tip is shown in the inset. b) STEM-HAADF image of the nanotube. c) STEM-HAADF image used for EDX analysis and drift correction. d) Point EDX analysis performed at the center of the nanotubes. e) Line scan analysis carried out on the line from Fig. c. FIGURE 2. (a) SEM micrographs of the as-obtained MoS 2 nanotubes. (b) Close-up view of the nanotubes showing the faceted morphology (square or rhomboid shapes) of the nanotubes. with a HP 6890 GC-MS, using a HP-5 MS capillary column (30 m 0.25 mm 0.25 µm). Catalytic activity was expressed in terms of % conversion of DBT vs reaction time, and from these data, the reaction rates were calculated for the MoS 2 nanotubes. The mean standard deviation for catalytic measurements was about 2.5% FIGURE 4. (a) HRTEM image of the edge of a MoS 2 nanotube. A close up of the image and the FFT are shown in the inset. The spacing of 0.63 nm (002 planes) is distinctly seen. (b) The internal part of the nanotube showing the oxide core (MoO 2). The lattice spacing in this case is 0.24 nm, which corresponds to the (101) planes of MoO 2. Rev. Mex. Fís. 57 (2) (2011) 1 6

10 4 F.L. DEEPAK, R.P. HERNÁNDEZ, J. CRUZ-REYES, S. FUENTES, AND M.J. YACAMAN FIGURE 5. Results of activity and product selectivity of the HDS of DBT for the MoS 2 nanotubes catalyst. Fig. 3b (STEM-HAADF image) a higher contrast is observed on the central part of the nanotube due to one face which is on top of the other. This is also confirmed by the EDX line analysis because the number of counts obtained in that part was much higher due to the higher thickness of the material at that point where the two faces were being analyzed. The EDX drift-corrected spectrum profile shows the characteristic and distinct Mo and S lines (Fig. 3d). The line scan in the EDX analysis reveals the presence of Mo (K,L) and S (K) lines and a small proportion of oxygen (O-K line) in the nanotubes. According to line scan carried out on the nanotube, Mo and S seem to be homogenously distributed (Fig. 3e). High resolution Transmission electron microscopy (HRTEM) analysis performed on the nanotube reveals that indeed the MoS 2 nanotube (Fig. 4a) exhibited a different phase at the core consisting of the oxide (Fig. 4b). The d- spacing obtained on the walls (shell) was characteristic of MoS 2, with a distance between layers about 0.63 nm, corresponding of the (002) planes. The HRTEM and the Fast Fourier Transform (FFT) (Fig. 4a inset) found very good crystallinity in the nanotubes which are oriented perpendicularly to the c axis. The inner part of the nanotube revealed a different distance between lattice fringes (Fig. 4b), the values obtained were about 0.24 nm which can be attributed to MoO 2 (101) that still remained in the material without being completely sulphided [38]. This is in agreement with the XRD pattern (Fig. 1) Catalytic activity FIGURE 6. Pseudo-first order reaction over the MoS 2 nanotube catalyst. The value of kinetic parameter k is mol/g s. High spatial resolution Energy Dispersive X-ray Analysis (SEM-EDAX) and elemental mapping of individual nanotubes was done to verify the presence of Mo and S. The elemental map clearly reveals the presence of Mo and S in the nanotubes. This is also confirmed by EDAX, which reveals the presence of the characteristic and distinct Mo (K,L) and S(K) lines. The Mo:S ratio of the nanotubes is found to be close to 1:2, according to EDAX analysis [35]. The compositional analysis of the respective elements has been carried out from the integration of the respective peaks of Mo and S in the EDAX spectrum. Although the peaks of the S(K) line from the Mo(L) line are too close to be clearly distinguishable a comparison with the standard sample of MoS 2 (purchased from Aldrich) can be used to resolve the composition between the two elements. Thus qualitatively the presence of Mo is resolved by the presence of Mo(K) line in the EDX spectrum and the quantitative analysis has been carried out by comparison of the compositions with a standard sample of MoS 2. Figure 3a shows a low magnification TEM micrograph of a MoS 2 nanotube. The nanotubes are facetted and empty wherein the faces are folded onto them to form the tube. In The catalytic activity of the MoS 2 nanotubes has been investigated for the HDS of DBT at 623 K under hydrogen pressure of 3.1 MPa. The five hour reaction time allowed for a better kinetic analysis of the pathway reaction following the evolution of products. The main HDS products detected from DBT over the MoS 2 nanotubes catalyst are: biphenyl (BP), obtained through the DDS pathway and tetrahydrodibenzothiophene (THDBT) and phenylcyclohexane (PCH) obtained through the HYD pathway. These products indicate a reaction scheme in agreement with prior reports, as shown in Scheme 1 [39,40]. Since both pathways are parallel [5], the ratio between HYD and DDS can be approximated in terms of experimental selectivity by Eq. (1) [40]. The selectivity calculated was 0.66 indicating that the DDS pathway is dominant over the HYD pathway. PCH + [THDBT] Selectivity = HYD/DDS = (1) [BP] Results of activity and product selectivity of the HDS of DBT for the MoS 2 nanotubes catalyst are displayed in Fig. 5. The MoS 2 nanotube catalyst showed values of 19 % conversion of DBT after 5 hours of reaction, which is in agreement with previous reports [41]. The HDS reaction of DBT using the MoS 2 nanotube catalyst was found to follow a pseudo-first order reaction mechanism (Fig. 6). The rate constant calculated from the optimum fitting process of the present catalyst was mol/g s. Some HDS catalysts require an ac- Rev. Mex. Fís. 57 (2) (2011) 1 6

11 CATALYTIC ACTIVITY OF MOS 2 NANOTUBES IN THE HYDRODESULPHURIZATION REACTION OF DIBENZOTHIOPHENE 5 SCHEME 1. Reaction network for the HDS of DBT. tivation period, where the activity increases with time-onstream as the catalyst is sulfided. (e.g., sulfided and/or reduced). In our case the catalyst showed good stability since the beginning of the reaction. Dungey et al. [42], observed an initial period of instability in the reaction rate, attributed to the fact that their materials were not pretreated. The main reaction products observed in this study were biphenyl (BP) and THDBT which are the primary products of DDS and HYD reactions, respectively. Phenylcyclohexane (PCH) is a secondary product resulting from C S bond cleavage of THDBT, an intermediate product formed by hydrogenation of one of the aromatic rings of DBT. There is a debate over relating the HDS catalytic activity of molybdenum sulfide-like crystal structures to their edge and/or basal plane stacking [43-45]. However, some studies proposed that the activity of molybdenum sulfide was localized at the edges and not on the flat basal planes [45]. It has been proposed [45], that hydrogenation is carried out in the rim-sites (end of the tube) due to the presence of active sites with high unsaturation (usually three missing sulfur atoms); meanwhile HDS is done on edge-sites of low unsaturation (usually one missing sulfur atom). With this background of catalytic activity and electron microscopy results, we propose that the low rim/edges sites ratio is responsible for the high BP selectivity because nanotubes have a larger flat area than rim sites (Scheme 2). In addition, the sulfur vacancies play a critical role on the selectivity because there are more vacancies in the rim sites than in the edge sites [45]. 4. Conclusions SCHEME 2. Distinction between rim or edge sites for stacked or unstacked M o S 2 particles. A HDS catalyst containing MoS 2 nanotubes was prepared by in-situ reaction of MoO 3 with thiourea. Large yields of the MoS 2 -nanotubes with an unusual faceted shape (square or rhomboid) and high internal surface area was obtained. The selectivity HYD/DDS ratio of the MoS 2 -nanotubes catalyst was 0.66; in this case the direct desulfurization pathway (DDS) was dominant over hydrogenation (HYD). This finding is attributed primarily to the size and morphology of the nanotubes showing low rim/edge ratio, due to the larger presence of flat surfaces than rim areas, responsible for the higher BP selectivity. Indeed, sulfur vacancies cannot be discarded to play an important role on selectivity as the sulfur insaturation of sites which is also related with the position of atoms. Acknowledgments We thank M. Del Valle for reviewing the manuscript. SCHEME 3. Distinction between rim or edge sites for M os 2 nanotubes Rev. Mex. Fís. 57 (2) (2011) 1 6

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13 REVISTA MEXICANA DE FÍSICA S 57 (2) 7 9 ABRIL 2011 Structural and optical characterization of In x Ga 1 x N nano-structured grown by chemical vapor deposition A. Ramos-Carrazco and E. Chaikina Centro de Investigación Científica y de Educación Superior de Ensenada, Ensenada, Baja California, CP 22860, México. O.E. Contreras Centro de Nanociencias y Nanotecnología, Universidad Nacional Autónoma de México, Ensenada, Baja California, CP 22860, México. M. Barboza-Flores and R. Garcia Centro de Investigación en Física Universidad de Sonora, Hermosillo, Sonora, México. rgarcia@cifus.uson.mx Recibido el 24 de noviembre de 2009; aceptado el 15 de enero de 2010 Nitrides of group III have generated important applications in optoelectronic devices. Principally InGaN is a novel alloy for the development of solid-state lighting and photovoltaic systems, since it is possible to control its bandgap from 3.4 ev to 0.7 ev by simply varying the indium concentration. However during the growth of InGaN inherent defects are obtained in the material, degrading its optical properties. In this work, the effect of the indium concentration is studied. The results of the optical and structural characterization of a series of In x Ga 1 x N films (0 x 0.3) deposited by chemical vapor deposition (CVD) are reported. Keywords: InGaN; semiconductor; luminescence and optoelectronics. Los nitruros del grupo III han generado aplicaciones importantes en los dispositivos optoelectronicos. Principalmente el InGaN ha mostrado ser una aleación novedosa para el desarrollo de la iluminación de estado sólido y sistemas fotovoltaicos, ya que es posible controlar el ancho de su banda prohibida desde 3.4 ev a 0.7 ev con solo variar la concentración de indio. Sin embargo durante el crecimiento de las películas de InGaN aparecen defectos en el material debido a las diferencias ente los átomos indio y galio. En este trabajo se estudia el efecto de la concentración de indio en las propiedades del InGaN. Se reportan los resultados de las caracterizaciones ópticas y estructurales de las películas de In xga 1 xn (0 x 0.3) depositadas por vapores químicos (CVD). Descriptores: InGaN; semiconductor; luminiscencia. PACS: Hk; Rx; Eq; Vv; Hk; b 1. Introduction In the development field of new materials, the compound semiconductors continue being an area of great interest and rapid expansion [1]. The ternary semiconductor InGaN is an important alloy for the development of lighting emitting devices, photovoltaic systems and power electronic, due to the capacity to control the band gap (E g ), which varies according to the indium concentration in a range of energies from 0.7 ev (InN) to 3.4 ev (GaN) [2]. Recently some attempts to grow high-quality low-cost InGaN have been done. One of the techniques that more likely fulfill the requirements is the chemical vapor deposition (CVD). This technique has reduced the cost of the synthesis maintaining an acceptable level in the optoelectronic properties of InGaN. However, the inherent mismatch between the lattice parameters of the substrate (sapphire, SiC, AsGa, Si, LiGaO) [3,4] and the InGaN phase, plus the indium incorporation (0 x 1) limits the growth of the material and degrade the optical and electronic InGaN properties. [5,6,8] In this work spectroscopy UV-VIS and photoluminescence (PL) have been used to study the optical properties of InGaN films grown by CVD [9]. Scanning electron microscopy and X-ray diffraction were used to characterize the morphology and structure of the InGaN films. 2. Experimental The synthesis of In x Ga 1 x N multilayer films with an indium composition of 0 x 0.3 deposited on sapphire at temperature of 900 C were grown by CVD. These films use the layers of aluminium nitride (AlN) and gallium nitride (GaN) as buffer and nucleation layer, respectively. The Fig. 1 shows the schematic diagram of the multilayer structure. The absorption measurements were made by two different techniques: transmission and diffuse reflectance. The absorption spectra were obtained with an AVANTES spectrometer (AvaSpec 256) in the wavelength range from 180 nm to 1100 nm. The diffuse reflectance was carried out in a UV-visible spectrometer Cary 300. All measurements of absorption were realized at room temperature. The PL measurements were obtained using two different light sources. The first, using a He-Cd laser (74 Series omnichrome - λ=325 nm). The luminescence of the sample is collimated through a spectrometer (SPECTRAPRO 500i) where the

14 8 A. RAMOS-CARRAZCO, E. CHAIKINA, O.E. CONTRERAS, M. BARBOZA-FLORES, AND R. GARCIA FIGURE 1. Multilayer structure of InGaN films grown on sapphire by chemical vapor deposition. FIGURE 3. Photoluminescence emission of InGaN films obtained by excitation He-Cd laser (λ = 325 nm). FIGURE 2. Absorption coefficients of the In xga 1 xn films obtained by diffuse reflectance. Excitation source: tungsten and deuterium lamps (λ = 190 nm to 850 nm). signal is quantified. In the second PL measurement, a UVvisible spectrometer, Hitachi Digilab F4500, with xenon lamp as excitation source was utilized. The XRD characterization was carried out in a powder diffractometer (Philips X pert). The surface of the InGaN films was studied in a SEM Jeol Results and discussion The absorption results obtained by diffuse reflectance (Fig. 2) were very different in comparison with the transmission measurements. The attenuation zone (including tails) varies in a region of energies from 2 ev to 3.3 ev (620 nm to 375 nm), which are near to the values of energies band gap expected in the In x Ga 1 x N films according to the Vegard s law. The origin of these absorption tails are attributed to the deformation of the crystalline lattice and to the existence of defects such as oxygen impurities and gallium/nitrogen vacancies [10]. Figure 3 shows the PL spectra of the In x Ga 1 x N films. The samples with indium composition smaller than FIGURE 4. X-ray diffractograms of (a) the In xga 1 xn (0 x 30) films from 33 to 37 2θ, the (0002) plane is marked and, (b) traces of other phases (impurities) present in the In x Ga 1 x N films grown by CVD in this work. 20 atomic percent (x<0.20) showed peaks with a FWHM of 500 mev whereas samples with higher indium composition (x>0.20) presented a broader peaks with a FWHM of 1 ev. Therefore, as well as the composition is increased in the In x Ga 1 x N phase the band gap energy is modified, showing a red-shift of the PL peak and also broader luminescence in the high indium samples. This behavior has its origin in the deformation of the In x Ga 1 x N lattice (stress due indium incorporation and the formation of a wide range of different In x Ga 1 x N crystals) and the existence of defects (oxygen impurities and gallium/nitrogen vacancies). Furthermore, in some parts of the spectra some modulations were observed due to the interference effect (Fabry-Perot) caused by internal reflections within the multilayer In x Ga 1 x N films [11]. Figure 4a shows the XRD results of the In x Ga 1 x N films. These diffractograms showed a hexagonal crystalline phase (wurzite) for the films. In x Ga 1 x N (0002) and GaN (0002) planes are marked. The In x Ga 1 x N crystalline phase was correlated with GaN phase located in the 2θ (34.56 ) position for the crystallographic plane (0002) according to the ICDD crystallographic letters [12]. In Fig. 4b is shown traces of Rev. Mex. Fís. 57 (2) (2011) 7 9

15 STRUCTURAL AND OPTICAL CHARACTERIZATION OF In x Ga 1 x N NANO-STRUCTURED GROWN BY CHEMICAL... 9 some impurities that appear in the In x Ga 1 x N films, indium oxide (In 2 O 3 ), indium nitride (InN) and indium metallic (clusters). The indium oxide can be related with the emission in the 550 nm region (emission by an indirect transition of 2.09 ev reported by Novkovski) [7]. Figure 5 shows SEM images of the InGaN films. The surface morphology of the films does not follow a pattern of growth that has a relation with the indium composition. The growth mode of the In x Ga 1 x N films appears to be the Volmer-Weber type. This growth mode is characterized by island formation due to nucleation crystals in diverse crystallographic directions. In this case, the crystals are the structures of columnar type which self-ensemble to form In x Ga 1 x N islands. 4. Conclusions FIGURE 5. Images of the InGaN films surface obtained in the SEM. Amplification of 750 X and a scale of 40 µm. A series of InGaN films deposited by CVD were characterized. It was found that the In x Ga 1 x N films with indium composition, x 0.20 present absorption and emission spectra that follow the Vegard s law. In x Ga 1 x N with higher content of indium (x 0.20) showed a broad PL emission (FHWM 1 ev) and large tails of absorption. In addition an extrinsic emission in the region of 570 nm ( 2.17 ev) was observed in this films. XRD showed the presence (traces) of undesirable phases such as In 2 O 3, InN and metallic indium in the films. SEM analysis found the formation of In x Ga 1 x N islands that affect the smoothness of the film surface. Acknowledgments The author grateful acknowledge the use of the facilities at the CICESE, CIFUS, ASU and CNyN. This research was partially supported by the project PAPIIT-UNAM IN Thanks to the support granted by CONACYT during my studies in CICESE. 1. Ariza C. H, Rev. Acad. Colomb. Cienc., , (2003), pp S. Strite and H. Morkoc, American Vaccum Society, B10 4, (1992), pp S. L. Hwang, K. S. Jang, K. H. Kim, H. S. Jeon, H. S. Ahn, M. Yang, Phys. Stat. Sol., 4 1, (2007), pp M.A. Sánchez García, J.L. Pau, F. Naranjo, A. Jiménez, S. Fernández, J. Ristic, F. Calle, E. Calleja y E. Muñoz, Material Science and Engineering B, 93 1, (2002), pp H. J. Chang, C. H. Chen and Y. F. Chen, T.Y. Lin, L. C. Chen, K.H. Chen and Z. H. Lan, Applied physics letters, 86 2, id(021911), (2005), pp Feng Shih Wei, Tang Tsung-Yi, Lu Yen-Cheng, Liu Shin-Jiun, Lin En-Chiang, Yang C.C., Ma Kung-Jen, Ching-Hsing, Chen L.C., Kim K.H., Lin J. Y., Jiang H.X., Journal of Applied Physics, 95 10, (2004), pp Novkovski, N., Tanusevski, A., Origin of the optical absorption of In2O3 thin films in the visible range, Semiconductor Science and Technology, 23 (9), id. (095012), 1-4 pp. 8. Michael A. Reshchikov and Hadis Morkoc, Journal of Applied Physics, 97, , (2005), pp M. U. González, J. A. Sánchez-Gil, Y. González and L. González, E. R. Méndez, American Vaccum Society, B18 4, (2000), pp O. Vigil y R. Zabala, Revista de Física Cubana, 72, (1987), pp C. Hums, T. Finger, T. Hempel, J. Christen and A. Dadgar, A. Hoffman, A. Krost, Journal of Applied Physics, 101, , (2007), pp ICDD crystallographic letters: In ( ), In 2O 3 ( ), InN ( ), Al 2 O 3 ( ), GaN ( ). Rev. Mex. Fís. 57 (2) (2011) 7 9

16 REVISTA MEXICANA DE FÍSICA S 57 (2) ABRIL 2011 Synthesis and characterization of In-doped ZnO nano-powders produced by combustion synthesis R. Garcia a, R. Nuñez-Gonzalez b, D. Berman-Mendoza a, M. Barboza-Flores a, and R. Rangel c a Departamento de Investigación en Física Universidad de Sonora, Hermosillo, Sonora, México, rgarcia@cifus.uson.mx b Departamento de Matemáticas Universidad de Sonora, Hermosillo, Sonora, México. c División de estudios de posgrado, Facultad de Ingeniería Química, UMSNH, Edificio V-1, Ciudad Universitaria, Morelia, Michoacán, México. Recibido el 7 de enero de 2010; aceptado el 18 de enero de 2010 Indium-doped ZnO powder was performed by a solution combustion technique using metal nitrates as oxidizer agents and carbohydrazide as fuel. The powders synthesized by this method are spongy clusters consisting of platelet-shaped nanocrystals with a wurtzite structure and narrow particle size distribution. Photoluminescence studies reveal that the powders emit high intensity luminescence. Defect-related green-yellow luminescence was found to be dependent upon the level of indium doping. Keywords: Combustion synthesis; luminescence; ZnO; semiconducting II-VI materials. Se sintetizó ZnO impurificado con indio usando la técnica síntesis por combustión partiendo de los nitratos como agentes oxidantes y carbohidrazina como combustible. Los polvos sintetizados por este método están formados por aglomerados compuestos de nano-cristales con una estructura tipo wurtzita y con una distribución de partícula uniforme. Estudios de fotoluminiscencia mostraron que los polvos emiten una luminiscencia de gran intensidad. Se encontró que la luminiscencia amarilla-verdosa que emiten estos polvos está relacionada con la concentración de indio en el ZnO. Descriptores: Síntesis por combustión, luminiscencia; ZnO; materiales semiconductores II-VI. PACS: m; w; Et; Dz; Wx; Ka 1. Introduction ZnO has attracted much attention towards applications in electronic and optoelectronic devices, such as UV photodetectors, solar cells, light emitting diodes and diode lasers [1,2]. Normally n-type dopants for ZnO are the III group elements such as indium [3,4], aluminum [5] and gallium [6]; while silver [7] and lithium [8] have been used for p-type doping. Indium doping is known to cause a red-shift in the band gap [3], while aluminum doping causes a blue shift, which increases with doping concentration [9,10]. In this work, a one-step synthesis method by combustion has been used to produce In-doped ZnO powder, using the nitrates of the metals as oxidizer agents and carbohydrazide as fuel. The effect of doping concentration on the structure and luminescence of ZnO has been investigated by x-ray diffraction and photoluminescence. 2. Experimental Procedure Undoped and indium-doped ZnO powders were prepared by combustion synthesis, using zinc nitrate hexahydrate (Zn (NO 3 ) 2 6H 2 O), de-ionized (DI) water as the solvent and carbohydrazide (CH 6 N 4 O) as fuel. Indium nitrate pentahydrate (In(NO 3 ) 3 5H 2 O) was added into the solution as a doping source with the molar concentration of 0.1%, 0.5%, 1%, and 5%, respectively. The solution was thoroughly stirred and homogenized in a beaker, and then it was transferred to a preheated furnace at 500 C. Combustion occurred after few minutes in the furnace and ZnO powder are formed in the beaker. The powders showed white and yellow color depending on indium concentration. 3. Results and discussion Figure 1 shows SEM images of ZnO samples. The powders present a sponge-like appearance within homogeneous sized grains. Doping with indium has no significant effect on the powder morphology. XRD spectra of undoped and In-doped ZnO are shown in Fig. 2. The effect of indium doping on the ZnO lattice structure is studied by monitoring the diffraction peak position and its FWHM. The main diffraction peaks can be related to the hexagonal wurtzite structure. The Bragg equation and Scherrer s formula were used to determine the lattice parameter and the grain diameter d, shown in Fig. 3. A slight shift to lower diffraction angles, lower peak intensity, and peak broadening are observed with increasing Indoping concentration. The slight shift in peak position can be related to the substitution of Zn 2+ ions with In 3+ ions as the difference between the ionic radii of In 3+ and Zn 2+ is very small (0.076 nm and nm respectively) [11]. The expansion of the lattice can be observed only at higher doping concentration (> 5 at.%). The crystalline quality diminishes with the introduction of indium, as seen in the broadening of diffraction peaks related to the presence of smaller grains. The optical properties of undoped and In-doped ZnO were

17 SYNTHESIS AND CHARACTERIZATION OF IN-DOPED ZNO NANO-POWDERS PRODUCED BY COMBUSTION SYNTHESIS 11 FIGURE 2. XRD spectra of (a) annealed and (b) as-grown undoped ZnO powders, (c) annealed and (d) as-grown 1% In-doped ZnO powders, and (e) annealed and (f) as-grown 5% In-doped ZnO powders. FIGURE 3. Calculated lattice parameter and grain size in ZnO powders with different indium doping concentrations. FIGURE 1. Secondary electron images of (a) undoped ZnO, (b) 1% In-doped ZnO, and (c) 5% In-doped ZnO. The scale is the same for the 3 images. characterized by PL spectroscopy at room temperature; the results are shown in Fig. 4. For the undoped ZnO powder there are two dominant emission bands: one is in the ultraviolet (UV) region with the emission peak at 388 nm corresponding to near-band-edge emission; and the other is a broad peak in the green-yellow region centered at 520 nm. FIGURE 4. Room temperature PL spectra of (a) undoped, (b) 0.1% In-doped, (c) 0.5% In-doped, (d) 1% In-doped, and (e) 5% Indoped ZnO powders. Rev. Mex. Fís. 57 (2) (2011) 10 12

18 12 R. GARCIA, R. NUÑEZ-GONZALEZ, D. BERMAN-MENDOZA, M. BARBOZA-FLORES, AND R. RANGEL With indium doping in various concentrations, the near-bandedge emission has the same energy, but its intensity is significantly reduced in the 5 at. % doped sample. The other emission peak in the green-yellow region undergoes a red shift to 580 nm and quenches gradually with indium concentration. It has been previously reported that indium doping leads to blue shift and broadening of the UV emission peak [4,12]. In our study, the introduction of indium into the ZnO lattice with concentrations less than 1% does not affect the luminescence intensity, and does not produce a noticeable blue shift in the UV emission line. This indicates that the indium as dopant is not involved in the near band edge transition. The broad green emission centered at 520 nm in undoped ZnO has been attributed to oxygen vacancies (V + 0 ) [10,13,14]. The 0.1% indium introduced into the ZnO shifts the green luminescence towards 580 nm in the yellow region with a considerable reduction in the intensity. This is due to In-doping introduces negatively-charged oxygen interstitials (O i ), which help to maintain charge equilibrium and contribute to the yellow luminescence [10,14,15]. When indium concentration increases, the yellow luminescence decreases instead of the expected increase. This suggests that at higher concentration levels, more indium atoms take up the lattice or interstitial sites in the ZnO lattice, which has no contribution to radiative recombination and only expands the lattice parameter and deteriorates the material quality. This also explains the suppression of both UV and green-yellow band emission at higher doping concentrations. The red shift of the green emission from 540 nm in undoped to 550 nm in In-doped ZnO is due to the formation of In 3+ - V + 0 complexes [14]. Furthermore, it can be found that the intensity of the green emission at 550 nm is reduced with increasing indium doping concentration, as illu. Janotti and Van de Walle [16] have presented in a model for the formation energy of oxygen vacancies in ZnO, which establishes a relationship between green emission intensity and the indium doping concentration. 4. Conclusions Homogeneous undoped and indium-doped ZnO nano-sized powders with a hexagonal wurtzite structure have been produced by combustion synthesis. It is observed that indium doping has no significant effect on the UV emission from ZnO and only influences the green-yellow luminescence. This may be due to In 3+ ions inducing the generation of oxygen interstitials to retain the charge neutrality, an event that causes a deep level emission shift from green to yellow. Also, it was found that the formation of In 3+ - V + 0 complexes induces a red shift of green emission in In-doped ZnO. Acknowledgements The authors gratefully acknowledge the use of facilities within the University of Sonora. This research has been partially supported by CONACyT. 1. A. Osinsky et al., Appl. Phys. Lett. 85 (2004) M. Law, L.E. Greene, J.C. Johnson, R. Saykally, and P.D. Yang, Nature Mater. 4 (2005) K.J. Kim and Y.P. Park, Appl. Phys. Lett. 78 (2001) J. Jie et al., Chem. Phys. Lett. 387 (2007) S.Y. Kuo et al., J. Crys. Growth 287 (2006) J.D. Ye et al., J. Crys. Growth 283 (2005) L. Guan, B.X. Lin, W.Y. Zhang, S. Zhong, and Z. Fu, Appl. Phys. Lett. 88 (2006) O. Lopatituk, L. Chemyak, O. Osinsky, and J.Q. Xie, Appl. Phys. Lett. 87 (2005) H.P. He et al., Appl. Phys. Lett. 90 (2007) M.S. Wang et al., Mater. Lett. 61 (2007) R.D. Shannon, Acta Crystallogr. A32 (1976) S.Y. Bae, H.C. Choi, C.W. Na, and J. Park, Appl. Phys. Lett. 86 (2005) K. Vanheusden, W.L. Warren, C.H. Seager, D.R. Tallant, and J.A. Voigt, J. Appl. Phys. 79 (1996) X.L. Wu, G.G. Siu, C.L. Fu, and H.C. Ong, Appl. Phys. Lett. 78 (2001) M. Liu, A.K. Kitai, and P. Mascher, J. Lumin. 54 (1992) A. Janotti and C.G. Van de Walle, Appl. Phys. Lett. 87 (2005) Rev. Mex. Fís. 57 (2) (2011) 10 12

19 REVISTA MEXICANA DE FÍSICA S 57 (2) ABRIL 2011 Photoconductivity studies of gold nanoparticles supported on amorphous and crystalline TiO 2 matrix prepared by sol-gel method G. Valverde-Aguilar, J.A. García-Macedo, and V. Renteria-Tapia, Departamento de Estado Sólido, Instituto de Física, Universidad Nacional Autónoma de México, Apartado Postal México, D.F., 04510, México, Tel. (5255) ; Fax (5255) valverde@fisica.unam.mx M. Aguilar-Franco Departamento de Física Química. Instituto de Física, Universidad Nacional Autónoma de México, Apartado Postal México, D.F., 04510, México. Recibido el 7 de diciembre de 2009; aceptado el 13 de julio de 2010 Gold metallic nanoparticles embedded in amorphous and crystalline TiO 2 matrix as powders and films were synthesized by the sol gel process at room temperature. The TiO 2 matrix was synthesized by using tetrabutyl orthotitanate as the inorganic precursor. The films were spin-coated on glass wafers. The samples were annealed at at 100 C for 30 minutes and sintered at 520 C for 1 hour to generated anatase and rutile phases. The film shows a light blue colour. The amorphous film exhibits an absorption band at 568 nm. The crystalline film exhibit two absorption peaks located at around 402 (from TiO 2 matrix) and 651 nm is due to the surface plasmon resonance of the gold nanoparticles. The films were studied using X-ray diffraction, infrared spectroscopy, scanning electron microscopy, high resolution transmission electronic microscopy and UV-Vis absorption spectroscopy. Photoconductivity studies were performed on amorphous and crystalline TiO 2 /Au films. The experimental data were fitted with straight lines at darkness and under illumination at 515 nm and 645 nm. This indicates an ohmic behavior. Transport parameters were calculated. Keywords: Titania; gold nanoparticles; sol-gel; photoconductivity; Gans theory; refractive index. Nanopartículas metálicas de oro insertadas en una matriz de TiO 2 (amorfa y cristalina) fueron sintetizadas en forma de polvos y películas por el método sol-gel a temperatura ambiente. La matriz de TiO 2fue sintetizada usando el tetrabutil ortotitanato como precursor inorgánico. Las películas fueron depositadas por spin-coating sobre substratos de vidrio. Las muestras fueron recocidas a 100 C por 30 minutos y sinterizadas a 520 C por 1 hora para generar las fases cristalinas anatasa y rutilo. Estas películas cristalinas muestran un color azul, y su absorción está en 645 nm, la cual es debido a su plasmón de resonancia. Las películas fueron caracterizadas por difracción de rayos X, espectroscopia infrarroja, microscopia de barrido y de alta resolución. Los estudios de fotoconductividad fueron realizados en las muestras amorfas y cristalinas de TiO 2 /Au. Los datos experimentales obtenidos en la oscuridad y bajo iluminación a 515 nm y 645 nm fueron ajustados por mínimos cuadrados. Esto indica un comportamiento óhmico. Los parámetros de transporte fueron calculados. Descriptores: Titanio; Nanopartículas metálicas de oro; sol-gel; películas delgadas; fotoconductividad; teoría de Gans; índice de refracción. PACS: r; r 1. Introduction Titanium dioxide (TiO 2 ) is a non-toxic material. TiO 2 thin films exhibit high stability in aqueous solutions and no photocorrosion under band gap illumination and special surface properties. TiO 2 thin films are already widely used in the study of the photocatalysis and photoelectrocatalysis of organic pollutants [1,2]. Photoelectrocatalytic system has received a great deal of attention due to drastically enhanced quantum efficiency [3]. By applying small bias, recombination of generated electron hole pairs is retarded. TiO 2 is the subject of intensive research, especially with regard to its end uses in solar cells, chemical sensors, photoelectrochemical cells, photocatalysis and electronic devices [4,5]. Due to its wide-ranging chemical and physical properties (electrical conductivity, photosensitivity, and aqueous environments) TiO 2 has a large variety of potential applications. As a wide band gap semiconductor, TiO 2 shows a diverse heterogeneity of crystalline phases, whereby it is possible to find it in anatase, rutile or brookite form [6]. TiO 2 are almost impossible to measure in great detail in powder form, due to the difficulty in manipulating grain sizes in the range of 1 50 nm [7]. Furthermore, measurements carried out on powder represent only an average value for many grains oriented in all possible directions. This difficulty in working with powder samples, together with the ongoing search for new applications, has compelled many researchers to work with TiO 2 thin films instead. In the present work, we described the synthesis, characterization and photoconductivity behaviour of amorphous and crystalline TiO 2 films doped with gold nanoparticles (NP s). The films were produced by the sol gel process at room temperature by using the spin-coating method and deposited on glass wafers. The samples were sintered at 520 C for 1 hour. The obtained films were studied by X-ray diffraction (XRD), optical absorption (OA), infrared spectroscopy (IR), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) studies. Photoconductivity studies were performed on these films. Transport parameters were calculated.

20 14 G. VALVERDE-AGUILAR, J.A. GARCÍA-MACEDO, V. RENTERIA-TAPIA, AND M. AGUILAR-FRANCO 2. Experimental Glass substrates were cleaned in boiling acidic solution of sulphuric acid-h 2 O 2 (4:1) under vigorous stirring for 30 minutes. They were then placed in deionized water and boiled for 30 minutes, rinsed three times with deionized water and stored in deionized water at room temperature. Preparation of TiO 2 solution. All reagents were Aldrich grade. The precursor solutions for TiO 2 films were prepared by the following method. Tetrabutyl orthotitanate and diethanolamine (NH(C 2 H 4 OH) 2 ) which prevent the precipitation of oxides and stabilize the solutions were dissolved in ethanol. After stirring vigorously for 2 h at room temperature, a mixed solution of deionized water and ethanol was added dropwise slowly to the above solution with a pipette under stirring. Finally, Tetraethyleneglycol (TEG) was added to the above solution. This solution is stirred vigorously to obtain a uniform sol. The resultant alkoxide solution was kept standing at room temperature to perform hydrolysis reaction for 2h, resulting in the TiO 2 sol. Preparation of Au stock solution M of Hydrogen Tetrachloroaurate(III) hydrate (HAuCl 4 aq) was dissolved in a mixture of deionized water and ethanol. It was stirred for 5 minutes. The Au stock solution was added to 20 ml of TiO 2 solution. This final solution was stirred for 17 hours at room temperature to obtain a purple colour. The final chemical composition of this solution was Ti(OC 4 H 9 ) 4 : NH(C 2 H 4 OH) 2 : C 2 H 5 OH : DI H 2 O : TEG: nitric acid: HauCl 4 = 1:1: 14.1:1:1.028:0.136: The TiO 2 with gold NP s solution has a ph = 6.0. The TiO 2 films were deposited by the spincoating technique. The precursor solution was placed on the glass wafers ( cm 2 ) using a dropper and spun at a rate of 3000 rpm for 20 s. After coating, the film was dried at 100 C for 30 min in a muffle oven and sintered at 520 C for 1 h in a muffle oven in order to remove organic components. The procedure was repeated two times to achieve the film thickness with two layers. The crystalline films show a light blue colour. UV-vis absorption spectra were obtained on a Thermo Spectronic Genesys 2 spectrophotometer with an accuracy of ±1 nm over the wavelength range of nm. The structure of the final films was characterized by XRD patterns. These patterns were recorded on a Bruker AXS D8 Advance diffractometer using Ni-filtered CuKα radiation. A step-scanning mode with a step of 0.02 in the range from 1.5 to 60 in 2θ and an integration time of 2 s was used. IR spectra were obtained from a KBr pellet using a Bruker Tensor 27 FT-IR spectrometer. Pellets were made from a finely ground mixture of the sample and KBr at a ratio of KBr:sample = 1: The thickness of the films was measured using a SEM microscopy Model STEREOSCAN at 20 kv. FIGURE 1. X-ray diffraction pattern at high angle of the amorphous and crystalline TiO 2 films with gold NP s. For photoconductivity studies [8] silver electrodes were painted on the sample. It was maintained in a 10 5 Torr vacuum cryostat at room temperature in order to avoid humidity. For photocurrent measurements, the films were illuminated with light from an Oriel Xe lamp passed through a 0.25 m Spex monochromator. Currents were measured with a 642 Keithley electrometer connected in series with the voltage power supply. The applied electrostatic field E was parallel to the film. Light intensity was measured at the sample position with a Spectra Physics 404 power meter. 3. Results and discussion 3.1. X-ray diffraction patterns The X-ray diffraction patterns of the amorphous and crystalline TiO 2 films with gold NP s is presented in Fig. 1. From amorphous film, its spectrum reveals the presence of gold NP s by the diffraction peaks located at 2θ = 38.24, 44.39, and which can be indexed as (111), (200), (220) and (311) respectively. The position of the diffraction peaks is in good agreement with those given in ASTM data card (# ). The crystalline film sintered at 520 C for 1 hour exhibits very good crystallization that corresponds to anatase and rutile phases. The anatase phase was identified by the diffraction peaks located at 2θ = 25.33, 47.97, 54.00, and which can be indexed as (101), (200), (105), (211) and (204) respectively. The rutile phase was identified by the diffraction peaks located at 2θ = 27.47, and which can be indexed as (110), (101) and (111) respectively. The position of the diffraction peaks in the film is in good agreement with those given in ASTM data card (# ) for anatase and ASTM data card (# ) for rutile. The presence of gold NP s was detected by the same diffraction peaks identified in the amorphous film. The average crystalline size (D) was calculated from Scherrer s formula [9] by using the diffraction peak (101) for anatase phase and the peak (110) for rutile phase: with λ= m. D = 0.9λ B cos θ (1) Rev. Mex. Fís. 57 (2) (2011) 13 18