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Leopoldo Garc´ıa-Col´ın Universidad Aut´ onoma Metropolitana – Iztapalapa, M´exico

Manuel Peimbert Instituto de Astronom´ıa UNAM, M´exico

Fernando Alba Instituto de F´ısica, UNAM, M´exico

Consejo Editorial Materia Condensada: Carlos Balseiro Centro At´ omico de Bariloche, Argentina

Alipio G. Calles Facultad de Ciencias, UNAM, M´exico

Manuel Cardona Institute Max Planck, Stuttgart, Alemania

Robert Cava University of Princeton, USA

Roberto Escudero Instituto de Investigaciones en Materiales, UNAM, M´exico

Francisco Jaque Universidad Aut´ onoma de Madrid, Espa˜ na

Harold Kroto Florida State University

F´ısica At´ omica y Molecular: Gerardo Delgado-Barrio Consejo Superior de Investigaci´ on Cient´ıfica, Espa˜ na

James McGuire Tulane University, USA

Alfred Schlachter Advanced Light Source, LBL Berkeley California, USA

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Arturo Menchaca Instituto de F´ısica, UNAM, M´exico

Andr´es Sandoval

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´ Instituto Nacional de Astrof´ısica, Optica y Electr´ onica, M´exico

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Jumpei Tsujiuchi Institute of Technology, Tokio, Jap´ on

Fernando Mendoza

´ Centro de Investigaciones en Optica, M´exico

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´xico–Printed in Mexico Impreso en Me

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VI International Topical Meeting on Nanostructured Materials and Nanotechnology, Nanotech 2009

San Carlos, Nuevo Guaymas September 17-19, 2010

Editor: M ARCELINO BARBOZA F LORES B EATRIZ DEL C ARMEN C ASTANEDA M EDINA A LVARO P OSADA A MARILLAS R AFAEL G ARCIA G UTIERREZ E LDER DE LA ROSA C RUZ

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ORGANIZING COMMITTEE Rafael Garc´ıa Guti´errez Alvaro Posada Amarillas Elder de la Rosa Cruz Marcelino Barboza Flores

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P REFACIO El Sexto Encuentro Internacional sobre Materiales Nanoestructurados y Nanotecnolog´ıa, NANOTECH 2009, es un congreso internacional que se ha organizado en la Rep´ublica Mexicana desde el a˜no 2004. La primera reuni´on se llev´o a cabo en ´ las instalaciones del Centro de Investigaci´on en Optica, en Le´on, Guanajuato. En el 2005 se organiz´o en el Centro de Ciencias de la Materia Condensada-UNAM en Ensenada. En el 2006 se llev´o a cabo en la ciudad de Puebla en el Instituto de F´ısica-BUAP. Monterrey fue la sede de la organizaci´on de la Conferencia del 2007, en las instalaciones de la Universidad Aut´onoma de Nuevo Le´on. En el 2008 el congreso se celebr´o en Ciudad Universitaria-UNAM en M´exico D.F. En esta ocasi´on la reuni´on fue organizada del 17 al 19 de septiembre del 2009 por el Departamento de Investigaci´on en F´ısica de la Universidad de Sonora en San Carlos Nuevo Guaymas. El congreso fue el escenario para la presentaci´on de 2 cursos cortos, 2 mesas redondas para tratar los temas de las aplicaciones de la nanotecnolog´ıa en la industria, 15 exposiciones orales y 88 carteles, adem´as de 10 conferencias magistrales presentadas por empresarios sonorenses y cient´ıficos de renombre mundial. El objetivo principal del NANOTECH 2009 fue el de proporcionar un foro para que cient´ıficos, ingenieros y empresarios busquen soluci´on a problemas cient´ıficos que conduzcan a aplicaciones pr´acticas. Los t´opicos que se trataron en dicho evento comprendieron desde ciencia b´asica hasta aplicaciones y t´ecnicas de comercializaci´on de alta tecnolog´ıa. Algunos de los principales temas discutidos aqu´ı fueron los nanotubos de carbono, nanomateriales magn´eticos y nanoestructuras met´alicas (plasmones), celdas solares y de combustible, nanof´osforos incluyendo o´ xidos, nitruros, tierras raras, y org´anicos; nanomedicina, nanocristales lineales y no lineales y cristales fot´onicos. El comit´e organizador agradece profundamente el apoyo financiero otorgado por la Universidad de Sonora, la Direcci´on Adjunta de Desarrollo Cient´ıfico y Acad´emico, CONACYT (M´exico), Red Tem´atica de Nanociencias y Nanotecnolog´ıa, y las empresas, Rubio Pharma y Asociados y RD Research & Technology.

REVISTA MEXICANA DE F´ISICA S 57 (2) 1–6

ABRIL 2011

Catalytic activity of MoS2 nanotubes in the hydrodesulphurization reaction of dibenzothiophene F. Leonard-Deepaka,b , R.P´erez-Hern´andezb,c , J. Cruz-Reyesd , S. Fuentese , 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. M´exico-Toluca S/N La Marquesa, Ocoyoacac, Edo. de M´exico 52750, M´exico. d Facultad de Ciencias Qu´ımicas e Ingenier´ıa, Universidad Aut´onoma de Baja California, Tijuana, B.C., M´exico. e Centro de Nanociencias y Nanotecnolog´ıa de la Universidad Nacional Aut´onoma de M´exico, Km. 107 Carretera Tijuana-Ensenada, Apartado Postal, 356, Ensenada, B.C., 22800, M´exico. 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 Wbased catalysts, often doped with other transition metals. We synthesized MoS2 nanotubes by reacting MoO3 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 MoS2 and large yields of the MoS2 nanotubes with unusual square or rhomboid faceted shapes. The catalytic behavior of the MoS2 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 (MoS2 ); 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 catal´ıtico que involucra catalizadores basados en Mo y W, a menudo dopados con otros metales de transicio´ n. Se sintetizaron nanotubos de MoS2 reaccionando MoO3 con thiourea. Los nanotubos se utilizaron como catalizadores para la hidrodesulturacion de dibenzotiofeno en un reactor discontinuo (batch reactor). Las t´ecnicas de difracci´on de rayos X, microscop´ıa electr´onica de barrido y de transmisi´on fueron utilizadas para caracterizar la morfolog´ıa y la estructura de los catalizadores. Los resultados mostraron la estructura cristalina hexagonal del MoS2 y grandes rendimientos de nanotubos de MoS2 con formas facetadas cuadradas o romboidales inusuales. El comportamiento catal´ıtico de los nanotubos de MoS2 demostr´o que la v´ıa de desulfuraci´on directa prevaleci´o sobre la v´ıa de hidrogenaci´on (HYD). Este resultado se atribuy´o a la baja relaci´on di´ametro/superficie (rim/edge), inducida por el tama˜no y morfolog´ıa de los nanotubos, mostrando un a´ rea grande y plana, que es la responsable de la selectividad del bifenilo (BP). Descriptores: hidrodesulfuracion; selectividad; dibenzotiofeno; sulfuro de molibdeno (MoS2 ); MET. PACS: 81.16.Hc, 61.05.cp, 68.37.Lp, 68.37.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 4,6dimethyldibenzothiophene (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

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´ F.L. DEEPAK, R.P. HERNANDEZ, 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]. MoS2 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: MoS2 , WS2 ) have attracted considerable attention in recent years [22-26]. One of the first methods of synthesis of MoS2 nanotubes was developed by Feldman, et al. [27]. This method involved the gas-phase reaction of MoO3−x and H2 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 MoS2 nanotubes. Li, et al. [29] have developed an atmospheric pressure chemical vapor deposition (APCVD) route for the synthesis of MoS2 nanostructures. These nanostructures, including three-dimensional nanoflowers (NF), were obtained by the reaction of chlorides (MoCl5 ) 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 H2 S (or a mixture of H2 and H2 S) and S powder. In general, the methods of synthesis of MoS2 nanotubes obtained them in low yields (∼ 20 %) as well as by long tedious procedures (ex: two-stage synthesis). The most important application of MoS2 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 MoS2 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 MoO3 as the starting materials to produce large quantities of MoS2 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 MoS2 nanotubes; it generates a suitable reducing-sulphidizing environment in-situ, eliminating the use of a separate reducing agent.

2. 2.1.

Experimental methods and characterization Synthesis

Synthesis of the MoS2 nanotubes was carried out as follows. About 0.6 g of MoO3 (mp = 795◦ C) and 1.0 g of thiourea (CSN2 H4 ,mp = 170-176◦ 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 N2 for 1/2 hr to remove any traces of oxygen. The tube was then heated to 1000◦ C in flowing N2 (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 N2 . 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. 2.2.

Characterization

X-ray diffraction (XRD) powder patterns were recorded in a Siemens D-5000 diffratometer, using Cu Kα (λ=0.15406 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. 2.3.

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 H2 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. F´ıs. 57 (2) (2011) 1–6

CATALYTIC ACTIVITY OF MOS2 NANOTUBES IN THE HYDRODESULPHURIZATION REACTION OF DIBENZOTHIOPHENE

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F IGURE 1. Powder XRD patterns of the MoS2 nanotubes. Red lines-MoS2 and blue lines-MoO2

3

Results and discussion

Figure 1 shows the XRD patterns of the MoS2 nanotubes synthesized using thiourea as the S source and MoO3 as the Mo source. The XRD patterns are in good agreement with those reported for the hexagonal crystalline structure of MoS2 (JCPDS 03-066-0160). The principal diffraction peak of the MoS2 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 MoO2 (JCPDS-01-076-1807) was also identified. This finding showed that it is possible to use this method to obtain MoS2 nanotubes with high crystallinity and purity. Figure 2 shows SEM micrographs of the MoS2 nanotubes obtained by the reaction of MoO3 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 MoS2 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 200-800 nm in diameter and extend up to several microns in length.

F IGURE 3. (a) BF-TEM image of a facetted MoS2 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.

F IGURE 2. (a) SEM micrographs of the as-obtained MoS2 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 MoS2 nanotubes. The mean standard deviation for catalytic measurements was about 2.5%

F IGURE 4. (a) HRTEM image of the edge of a MoS2 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 (MoO2 ). The lattice spacing in this case is 0.24 nm, which corresponds to the (101) planes of MoO2 .

Rev. Mex. F´ıs. 57 (2) (2011) 1–6

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F IGURE 5. Results of activity and product selectivity of the HDS of DBT for the MoS2 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 MoS2 nanotube (Fig. 4a) exhibited a different phase at the core consisting of the oxide (Fig. 4b). The dspacing obtained on the walls (shell) was characteristic of MoS2 , 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 MoO2 (101) that still remained in the material without being completely sulphided [38]. This is in agreement with the XRD pattern (Fig. 1). 3.1.

F IGURE 6. Pseudo-first order reaction over the MoS2 nanotube catalyst. The value of kinetic parameter k is 7.49×10−7 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 MoS2 (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 MoS2 . Figure 3a shows a low magnification TEM micrograph of a MoS2 nanotube. The nanotubes are facetted and empty wherein the faces are folded onto them to form the tube. In

Catalytic activity

The catalytic activity of the MoS2 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 MoS2 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 MoS2 nanotubes catalyst are displayed in Fig. 5. The MoS2 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 MoS2 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 7.49×10−7 mol/g s. Some HDS catalysts require an ac-

Rev. Mex. F´ıs. 57 (2) (2011) 1–6

CATALYTIC ACTIVITY OF MOS2 NANOTUBES IN THE HYDRODESULPHURIZATION REACTION OF DIBENZOTHIOPHENE

5

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].

S CHEME 1. Reaction network for the HDS of DBT.

4.

S CHEME 2. Distinction between “rim” or “edge” sites for stacked or unstacked Mo S2 particles.

Conclusions

A HDS catalyst containing MoS2 nanotubes was prepared by in-situ reaction of MoO3 with thiourea. Large yields of the MoS2 -nanotubes with an unusual faceted shape (square or rhomboid) and high internal surface area was obtained. The selectivity HYD/DDS ratio of the MoS2 -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.

S CHEME 3. Distinction between “rim” or “edge” sites for Mo S2 nanotubes

Rev. Mex. F´ıs. 57 (2) (2011) 1–6

6

´ F.L. DEEPAK, R.P. HERNANDEZ, J. CRUZ-REYES, S. FUENTES, AND M.J. YACAMAN

∗. e-mail: [email protected]

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Structural and optical characterization of Inx Ga1−x N nano-structured grown by chemical vapor deposition A. Ramos-Carrazco and E. Chaikina Centro de Investigaci´on Cient´ıfica y de Educaci´on Superior de Ensenada, Ensenada, Baja California, CP 22860, M´exico. O.E. Contreras Centro de Nanociencias y Nanotecnolog´ıa, Universidad Nacional Aut´onoma de M´exico, Ensenada, Baja California, CP 22860, M´exico. M. Barboza-Flores and R. Garcia∗ Centro de Investigaci´on en F´ısica Universidad de Sonora, Hermosillo, Sonora, 83190 M´exico. ∗ e-mail: [email protected] 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 Inx Ga1−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 aleaci´on novedosa para el desarrollo de la iluminaci´on de estado s´olido 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 concentraci´on de indio. Sin embargo durante el crecimiento de las pel´ıculas de InGaN aparecen defectos en el material debido a las diferencias ente los a´ tomos indio y galio. En este trabajo se estudia el efecto de la concentraci´on de indio en las propiedades del InGaN. Se reportan los resultados de las caracterizaciones o´ pticas y estructurales de las pel´ıculas de Inx Ga1−x N (0 ≤ x ≤ 0.3) depositadas por vapores qu´ımicos (CVD). Descriptores: InGaN; semiconductor; luminiscencia. PACS: 61.46.Hk; 61.82.Rx; 71.55.Eq; 61.72.Vv; 78.60.Hk; 78.60. 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 (Eg ), 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 Inx Ga1−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

8

A. RAMOS-CARRAZCO, E. CHAIKINA, O.E. CONTRERAS, M. BARBOZA-FLORES, AND R. GARCIA

F IGURE 1. Multilayer structure of InGaN films grown on sapphire by chemical vapor deposition. F IGURE 3. Photoluminescence emission of InGaN films obtained by excitation He-Cd laser (λ = 325 nm).

F IGURE 2. Absorption coefficients of the Inx Ga1−x N 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 5300.

3. 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 Inx Ga1−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 Inx Ga1−x N films. The samples with indium composition smaller than

F IGURE 4. X-ray diffractograms of (a) the Inx Ga1−x N (0≤ x ≤30) films from 33 ˚ to 37 ˚ 2θ, the (0002) plane is marked and, (b) traces of other phases (impurities) present in the Inx Ga1−x N films grown by CVD in this work.

20 atomic percent (x0.20) presented a broader peaks with a FWHM of ∼1 eV. Therefore, as well as the composition is increased in the Inx Ga1−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 Inx Ga1−x N lattice (stress due indium incorporation and the formation of a wide range of different Inx Ga1−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 Inx Ga1−x N films [11]. Figure 4a shows the XRD results of the Inx Ga1−x N films. These diffractograms showed a hexagonal crystalline phase (wurzite) for the films. Inx Ga1−x N (0002) and GaN (0002) planes are marked. The Inx Ga1−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. F´ıs. 57 (2) (2011) 7–9

STRUCTURAL AND OPTICAL CHARACTERIZATION OF Inx Ga1−x N NANO-STRUCTURED GROWN BY CHEMICAL. . .

9

graphic directions. In this case, the crystals are the structures of columnar type which self-ensemble to form Inx Ga1−x N islands.

4.

F IGURE 5. Images of the InGaN films surface obtained in the SEM. Amplification of 750 X and a scale of 40 µm.

some impurities that appear in the Inx Ga1−x N films, indium oxide (In2 O3 ), 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 Inx Ga1−x N films appears to be the Volmer-Weber type. This growth mode is characterized by island formation due to nucleation crystals in diverse crystallo-

1. Ariza C. H, Rev. Acad. Colomb. Cienc., 27 104, (2003), pp. 357-369. 2. S. Strite and H. Morkoc, American Vaccum Society, B10 4, (1992), pp. 1237-1266. 3. S. L. Hwang, K. S. Jang, K. H. Kim, H. S. Jeon, H. S. Ahn, M. Yang, Phys. Stat. Sol., 4 1, (2007), pp. 125-128. 4. M.A. S´anchez Garc´ıa, J.L. Pau, F. Naranjo, A. Jim´enez, S. Fern´andez, J. Ristic, F. Calle, E. Calleja y E. Mu˜noz, Material Science and Engineering B, 93 1, (2002), pp. 189-196. 5. 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. 1-3. 6. 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. 5388-5396.

Conclusions

A series of InGaN films deposited by CVD were characterized. It was found that the Inx Ga1−x N films with indium composition, x ≤ 0.20 present absorption and emission spectra that follow the Vegard’s law. Inx Ga1−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 In2 O3 , InN and metallic indium in the films. SEM analysis found the formation of Inx Ga1−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 IN101509. Thanks to the support granted by CONACYT during my studies in CICESE.

7. Novkovski, N., Tanusevski, A., 2008. 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, 061301, (2005), pp 1-95. 9. M. U. Gonz´alez, J. A. S´anchez-Gil, Y. Gonz´alez and L. Gonz´alez, E. R. M´endez, American Vaccum Society, B18 4, (2000), pp. 1980-1990. 10. O. Vigil y R. Zabala, Revista de F´ısica Cubana, 72, (1987), pp. 67-76. 11. C. Hums, T. Finger, T. Hempel, J. Christen and A. Dadgar, A. Hoffman, A. Krost, Journal of Applied Physics, 101, 033113-1, (2007), pp. 1-4. 12. ICDD crystallographic letters: In (00-005-0642), In2 O3 (00006-0416), InN (00-050-1239), Al2 O3 (00-010-0173), GaN (00-050-0792).

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ABRIL 2011

Synthesis and characterization of In-doped ZnO nano-powders produced by combustion synthesis R. Garciaa , R. Nu˜nez-Gonzalezb , D. Berman-Mendozaa , M. Barboza-Floresa , and R. Rangelc a Departamento de Investigaci´on en F´ısica Universidad de Sonora, Hermosillo, Sonora, 83190 M´exico, e-mail: [email protected] b Departamento de Matem´aticas Universidad de Sonora, Hermosillo, Sonora, 83190 M´exico. c Divisi´on de estudios de posgrado, Facultad de Ingenier´ıa Qu´ımica, UMSNH, Edificio V-1, Ciudad Universitaria, Morelia, Michoac´an, M´exico. 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 sintetiz´o ZnO impurificado con indio usando la t´ecnica s´ıntesis por combusti´on partiendo de los nitratos como agentes oxidantes y carbohidrazina como combustible. Los polvos sintetizados por este m´etodo est´an formados por aglomerados compuestos de nano-cristales con una estructura tipo wurtzita y con una distribuci´on de part´ıcula uniforme. Estudios de fotoluminiscencia mostraron que los polvos emiten una luminiscencia de gran intensidad. Se encontr´o que la luminiscencia amarilla-verdosa que emiten estos polvos est´a relacionada con la concentraci´on de indio en el ZnO. Descriptores: S´ıntesis por combusti´on, luminiscencia; ZnO; materiales semiconductores II-VI. PACS: 78.55.-m; 61.46.+w; 778.55.Et; 81.05.Dz; 81.07.Wx; 81.20.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 (NO3 )2 ·6H2 O), de-ionized (DI) water as the solvent and carbohydrazide (CH6 N4 O) as fuel. Indium nitrate pentahydrate (In(NO3 )3 ·5H2 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 Zn2+ ions with In3+ ions as the difference between the ionic radii of In3+ and Zn2+ is very small (0.076 nm and 0.074 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

SYNTHESIS AND CHARACTERIZATION OF IN-DOPED ZNO NANO-POWDERS PRODUCED BY COMBUSTION SYNTHESIS

11

F IGURE 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.

F IGURE 3. Calculated lattice parameter and grain size in ZnO powders with different indium doping concentrations.

F IGURE 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.

F IGURE 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.

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˜ R. GARCIA, R. NUNEZ-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 In3+ - V+ 0 com-

plexes [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 In3+ 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 In3+ - 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.

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14. X.L. Wu, G.G. Siu, C.L. Fu, and H.C. Ong, Appl. Phys. Lett. 78 (2001) 2285. 15. M. Liu, A.K. Kitai, and P. Mascher, J. Lumin. 54 (1992) 35. 16. A. Janotti and C.G. Van de Walle, Appl. Phys. Lett. 87 (2005) 122102.

Rev. Mex. F´ıs. 57 (2) (2011) 10–12

REVISTA MEXICANA DE F´ISICA S 57 (2) 13–18

ABRIL 2011

Photoconductivity studies of gold nanoparticles supported on amorphous and crystalline TiO2 matrix prepared by sol-gel method G. Valverde-Aguilar, J.A. Garc´ıa-Macedo, and V. Renteria-Tapia, Departamento de Estado S´olido, Instituto de F´ısica, Universidad Nacional Aut´onoma de M´exico, Apartado Postal 20-364 M´exico, D.F., 04510, M´exico, Tel. (5255) 56225103; Fax (5255) 56161535 E-mail: [email protected] M. Aguilar-Franco Departamento de F´ısica Qu´ımica. Instituto de F´ısica, Universidad Nacional Aut´onoma de M´exico, Apartado Postal 20-364 M´exico, D.F., 04510, M´exico. Recibido el 7 de diciembre de 2009; aceptado el 13 de julio de 2010 Gold metallic nanoparticles embedded in amorphous and crystalline TiO2 matrix as powders and films were synthesized by the sol–gel process at room temperature. The TiO2 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 TiO2 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 TiO2 /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. Nanopart´ıculas met´alicas de oro insertadas en una matriz de TiO2 (amorfa y cristalina) fueron sintetizadas en forma de polvos y pel´ıculas por el m´etodo sol-gel a temperatura ambiente. La matriz de TiO2 fue sintetizada usando el tetrabutil ortotitanato como precursor inorg´anico. Las pel´ıculas 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 pel´ıculas cristalinas muestran un color azul, y su absorci´on est´a en 645 nm, la cual es debido a su plasm´on de resonancia. Las pel´ıculas fueron caracterizadas por difracci´on de rayos X, espectroscopia infrarroja, microscopia de barrido y de alta resoluci´on. Los estudios de fotoconductividad fueron realizados en las muestras amorfas y cristalinas de TiO2 /Au. Los datos experimentales obtenidos en la oscuridad y bajo iluminaci´on a 515 nm y 645 nm fueron ajustados por m´ınimos cuadrados. Esto indica un comportamiento o´ hmico. Los par´ametros de transporte fueron calculados. Descriptores: Titanio; Nanopart´ıculas met´alicas de oro; sol-gel; pel´ıculas delgadas; fotoconductividad; teor´ıa de Gans; ´ındice de refracci´on. PACS: 72.80.-r; 73.61.-r

1.

Introduction

Titanium dioxide (TiO2 ) is a non-toxic material. TiO2 thin films exhibit high stability in aqueous solutions and no photocorrosion under band gap illumination and special surface properties. TiO2 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. TiO2 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) TiO2 has a large variety of potential applications. As a wide band gap semiconductor, TiO2 shows a diverse heterogeneity of crystalline phases, whereby it is possible to find it in anatase, rutile or brookite form [6].

TiO2 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 TiO2 thin films instead. In the present work, we described the synthesis, characterization and photoconductivity behaviour of amorphous and crystalline TiO2 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.

14

G. VALVERDE-AGUILAR, J.A. GARC´IA-MACEDO, V. RENTERIA-TAPIA, AND M. AGUILAR-FRANCO

2. Experimental Glass substrates were cleaned in boiling acidic solution of sulphuric acid-H2 O2 (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 TiO2 solution. All reagents were Aldrich grade. The precursor solutions for TiO2 films were prepared by the following method. Tetrabutyl orthotitanate and diethanolamine (NH(C2 H4 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 TiO2 sol. Preparation of Au stock solution. 0.03 M of Hydrogen Tetrachloroaurate(III) hydrate (HAuCl4 ·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 TiO2 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(OC4 H9 )4 : NH(C2 H4 OH)2 : C2 H5 OH : DI H2 O : TEG: nitric acid: HauCl4 = 1:1: 14.1:1:1.028:0.136:0.024. The TiO2 with gold NP’s solution has a pH = 6.0. The TiO2 films were deposited by the spincoating technique. The precursor solution was placed on the glass wafers (2.5×2.5 cm2 ) 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 300-900 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 = 1:0.019. The thickness of the films was measured using a SEM microscopy Model STEREOSCAN at 20 kV.

F IGURE 1. X-ray diffraction pattern at high angle of the amorphous and crystalline TiO2 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. 3.1.

Results and discussion X-ray diffraction patterns

The X-ray diffraction patterns of the amorphous and crystalline TiO2 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, 64.62 and 77.60 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 (#04-0784). 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, 55.16 and 62.71 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, 36.14 and 41.32 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 (#21-1272) for anatase and ASTM data card (#21-1276) 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: D= with λ=1.54056×10−10 m.

Rev. Mex. F´ıs. 57 (2) (2011) 13–18

0.9λ B cos θ

(1)

PHOTOCONDUCTIVITY STUDIES OF GOLD NANOPARTICLES SUPPORTED ON AMORPHOUS AND CRYSTALLINE TiO2 . . .

15

TABLE I. Summary of nanoscopic characteristics of amorphous and crystalline TiO2 /Au films. Phase

B

Radian

D

Crystal phase

(nm)

(wt%)

Anatase (101)

0.44◦

0.00768

18.5

59.7±4

Rutile (110)

0.31◦

0.00543

26.3

37.4±3

Au

-

-

-

2.9±4

F IGURE 3. Experimental optical absorption spectrum (black dotted line) of the crystalline TiO2 /Au film. The calculated optical absorption spectrum (grey solid line) obtained by Gans theory.

F IGURE 2. Absorption spectra of the amorphous (black solid line) and crystalline (grey solid line) TiO2 film with gold NP’s.

The percentage of anatase, rutile and gold phases was calculated by means of a Rietvield refinement. These calculations are shown in Table I. 3.2.

Optical absorption

Figure 2 shows the optical absorption spectra of the amorphous and crystalline TiO2 /Au films taken at room temperature in the range of 300-900 nm. The spectrum of the film calcined at 450◦ C for 15 min shows an absorption band A located at 402 (3.08 eV) corresponding to the TiO2 matrix, and a second band B located 651 nm (1.93 eV) corresponding to the surface plasmon resonance (SPR) of the gold NP’s. The spectrum of the amorphous film shows a peak shoulder C at 568 nm (2.68 eV) which is the SPR band of spherical Au nanoparticles [10,11]. To clarify the XRD and optical absorption experimental results, the formation mechanism of Au nanoparticles is discussed below. It is known that the photolysis of HAuCl4 to the Au atom, Au0 , is a multiphoton event [12,13], and it proceeds by irradiation. Therefore, for amorphous TiO2 /Au film, the Au nucleation was slow and random because the HAuCl4 ions were reduced by daylight (containing a little UV light) and this mostly happened after the gelation. The nuclei were thus distributed randomly within the TiO2 skeleton and consequently led to the growth of the Au particles that were inhomogeneous, and their size distribution very wide. Literature [8,14] reports an absorption peak for surface plasmon resonance (SPR) of gold nanoparticles around 500550 nm. A red-shift in the maximum in absorbance towards larger wavelength (from 568 to 651 nm) with respect to the amorphous TiO2 film is evident as well as a broadening of the peak absorption width compared to the amorphous film.

F IGURE 4. Cross-sectional SEM image of (a) amorphous and (b) crystalline TiO2 films with gold NP’s.

The dependence of this shift on the embedding medium indicates the high sensitivity of surface plasmon band to clustermatrix interface properties. This fact is originated to the increase in the diameter of Au nanoparticles and an increment of the refractive index of TiO2 matrix with increasing the heat-treatment temperature [15,16]. It is well known that the refractive index of TiO2 films is related to the crystal phase (anatase or rutile), the crystalline size and the densities of the films [17]. For these reasons, the optical absorption spectrum (Fig. 2) was fitted very well using Gans theory [18] with a local refractive index nlocal = 2.6 (Fig. 3). This index has a value close to the refractive index reported for the anatase phase (nanatase = 2.54) [19]. This is consistent with the fact we have anatase phase in a proportion of 59.7 wt% according to the X-ray diffraction pattern

Rev. Mex. F´ıs. 57 (2) (2011) 13–18

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G. VALVERDE-AGUILAR, J.A. GARC´IA-MACEDO, V. RENTERIA-TAPIA, AND M. AGUILAR-FRANCO

F IGURE 6. Size-distribution histograms obtained by HRTEM analysis of gold metallic NP’s.

Figure 5 shows the HRTEM image of the crystalline TiO2 /Au film. Figure 5a shows gold NP’s which were identified as brilliant particles. Figure 4a shows a superposition of these populations. The reflection (101) corresponds to the anatase phase; and the reflection (111) corresponds to gold nanoparticle. The diffraction patterns (in the insert of the figure) show these reflections. From HRTEM studies taking into account a population of gold NP’s, the corresponding size-distribution histograms were obtained (Fig. 6). The distributions from the major axis A and minor length axis B and their respective standard deviations are A = 9.8 ± 7.8 nm (Fig. 5 a), B = 6.6 ± 3.9 nm. 3.4. F IGURE 5. (a) HRTEM image of the crystalline TiO2 /Au film exhibits several gold NP’s. (b) The reflections correspond to anatase nanocrystals and gold metallic nanoparticles were identified with white arrows. The inset shows the diffraction pattern showing these reflections.

(Fig. 3), while the rutile phase (nrutile = 2.75) [20] has a proportion of 37.4 wt%.

Photoconductivity studies

Usually [8] Ohm’s law under light illumination is given by →

645

The thickness of the films was measured by SEM technique. Figure 4 shows the SEM image for amorphous and crystalline TiO2 films with gold NP’s. The thickness and the standard deviation for both kinds of films were calculated. The average thickness for amorphous and crystalline TiO2 /Au films is equal to 7.0 ± 1.2 µm and 3.8 ± 1.1 µm, respectively.

→

(2)

TABLE II. Linear fittings of amorphous and crystalline TiO2 films. λ (nm)

3.3. SEM and HRTEM measurements

→

J = J ph + (σd + σph ) E

515 Darkness

Rev. Mex. F´ıs. 57 (2) (2011) 13–18

TiO2 / Au film Crystalline

A1

J0 −7

1.40×10−3

−10

3.14×10

Amorphous

5.36×10

2.89×10−6

Crystalline

3.64×10−7

1.07×10−3

Amorphous

4.97×10−10

2.27×10−6

−7

Crystalline

3.73×10

Amorphous

3.32×10−10

6.66×10−4 1.89×10−6

PHOTOCONDUCTIVITY STUDIES OF GOLD NANOPARTICLES SUPPORTED ON AMORPHOUS AND CRYSTALLINE TiO2 . . .

17

TABLE III. φl0 and φµτ parameters of amorphous and crystalline TiO2 /Au films. λ (nm) 515

Parameters φl0 (cm) 2

φµτ (cm /V) 645

φl0 (cm) 2

φµτ (cm /V)

Amophous

Crystalline

TiO2 / Au

TiO2 / Au

−6

1.41×10−3

5.42×10

−10

3.42×10−8

1.91×10−6

9.71×10−4

1.23×10

−10

3.91×10

7.87×10−8

With the Eq. (3), by measuring I, the dark conductivity and the conductivity under illumination at 645 and 355 nm, and fitting the experimental data by the least squares method, as it is shown in Fig. 7, the photoconductive (φµτ ) and photovoltaic (φl0 ) parameters were obtained by using the next expressions. ¡ ¢ hc φl0 = J0i − Jod eαλI hc φµτ = (A1i − A1d ) eα λI

F IGURE 7. Experimental data of current density vs. electric field spectra from (a) amorphous and (b) crystalline TiO2 /Au films. Linear fits correspond to the dotted lines. →

where J ph is the photovoltaic current density, and σph is the photoconductivity. When the current densities are assumed → to be parallel to the electric field E Eq. (2) becomes into the next one: µ ¶ qφl0 αI qφµτ αI J= + σd + E (3) hν hν with φ as the quantum yield of charge carrier photogeneration, l0 as the charge carrier mean free path, α as the sample absorption coefficient, I as the light intensity at the frequency ν of illumination, h as the Planck’s constant, and τ gs the charge carriers mean lifetime. The first term is the photovoltaic transport effect, the second one is the dark conductivity σd =en0 µ, and the third one is the photoconductivity itself. Eq. (3) can be written as: J = A1 E + J0

(4)

From the absorption spectrum of crystalline film (Fig. 2), the illumination wavelength for photoconductivity studies were chosen: 645 nm that corresponds to the maximum absorption band and 515 nm were there is no absorption. Photoconductivity results of amorphous and crystalline TiO2 films with gold NP’s are shown in Fig. 7. Current density as function of electric applied field on the film was plotted. The experimental data were fitted by least-squares with straight lines at darkness and under illumination. This indicates an ohmic behaviour. The linear fits are shown in Table II. For both kinds of TiO2 /Au films, when the illumination wavelength decreases the J0 value decreases. For crystalline film, when the illumination wavelength decreases, the slopeA1 increases. It indicates a strong photoconductive behavior in these films.

the subscripts i= illumination and d= darkness. Table III contains the φl0 and φµτ values. φl0 and φµτ parameter values are bigger for crystalline films than those from amorphous ones. This indicates a strong photoconductive effect in the crystalline TiO2 /Au films.

4.

Conclusions

High optical quality crystalline TiO2 films with gold NP’s were obtained by sol-gel process. XRD measurements reveal the presence of the anatase and rutile phases, which were produced after sintering treatment of 520◦ C for 1h. The anatase phase has a bigger proportion (59.75 wt%) than the rutile phase (37.4 wt%). The optical absorption spectrum was fitted very good using Gans theory by using a local refractive index nlocal = 2.6. This index is related to the major crystal phase, anatase. The experimental data J vs E were fitted by straight lines corresponding to an ohmic behaviour. Crystalline TiO2 /Au films exhibit a strong photoconductive effect. Anatase phase leads a better conduct on the electron/hole pair than the amorphous phase.

Acknowledgments The authors acknowledge the financial supports of CONACYT 79781, REdNyN, PUNTA and PAPIIT IN107510. The authors are thankful to Luis Rend´on (HRTEM), Roberto Hern´andez-Reyes (SEM) and Diego Quiterio (preparation of the samples for SEM studies) for technical assistance. GVA is grateful for CONACYT support.

Rev. Mex. F´ıs. 57 (2) (2011) 13–18

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G. VALVERDE-AGUILAR, J.A. GARC´IA-MACEDO, V. RENTERIA-TAPIA, AND M. AGUILAR-FRANCO

1. R. Suarez, P.K. Nair, and P.V. Kamat, Langmuir 14 (1998) 3236.

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3. D.W. Kim et al., International Journal of Hydrogen Energy 32 (2007) 3137. 4. C. Graziani Garcia, N-Y. Murakami R. Argazzi, and C.-A. Bignozzi, J. Photochem. Photobiol. A: Chem. 115 (1998) 239. 5. S. Due˜nas et al., Semicond. Sci. Technol. 20 (2005) 1044. 6. L. Hu, T. Yoko, H. Kosuka, and S. Sakka, Thin Solid Films 219 (1992) 18. 7. H. Gerischer, and A. Heller, Electrochem. Soc. 139 (1992) 113.

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Rev. Mex. F´ıs. 57 (2) (2011) 13–18

REVISTA MEXICANA DE F´ISICA S 57 (2) 19–21

ABRIL 2011

Integration and electrical characterization of organic thin film transistor for an active matrix of oleds G. Guti´errez-Herediaa,c,∗ , L.A. Gonz´aleza,c,∗∗ , A. Avenda˜noc , D. Bermana , H.N. Alshareefc,d , B.E. Gnadec , and M. Quevedo-L´opezb,c,∗∗∗ a Centro de Investigaci´on en F´ısica, Universidad de Sonora, Hermosillo, Sonora 83190 M´exico, ∗ e-mail: [email protected]; ∗∗ e-mail: [email protected] b Departamento de Pol´ımeros y Materiales, Universidad de Sonora, Hermosillo, Sonora 83190 M´exico, ∗∗∗ e-mail: [email protected] c Department of Materials Science & Engineering, University of Texas at Dallas, Richardson, Texas 75080-3021 USA. d Department of Materials Science, King Abdullah University of Science and Technology, Jeddah, Saudi Arabia. Recibido el 7 de enero de 2010; aceptado el 15 de enero de 2010 We present a novel integration process of all organic Thin Film Transistor (TFTs) and its electrical characterization. The test circuit is designed to drive an active matrix of organic light emitting diode (AMOLED). The process is performed in both plastic (Polyethylene naphthalene, PEN) and glass substrates. The basic circuit is formed by two pentacene-based transistors and a capacitor. All of these devices use parylene as dielectric. As a result of the electrical characterization, we show that this circuit can deliver up to 40 µA. This current level is appropriate if we consider that the minimum required current to obtain 200 cd/m2 from a typical OLED is of 10 µA. Keywords: Flexible electronic; OTFT; Parylene; Pentacene; AMOLED. Presentamos una nueva forma de integraci´on y caracterizaci´on el´ectrica de TFTs completamente org´anico. El circuito de prueba est´a dise˜nado para controlar una matriz activa de diodos emisores de luz org´anicos. El proceso es desarrollado tanto en substratos de pl´astico (Polyethylene naphthalene, PEN) como en substratos de vidrio. El circuito b´asico consiste de 2 transistores de pentaceno y un capacitor. En estos dispositivos empleamos perileno como diel´ectrico. Como resultado del proceso de caracterizaci´on el´ectrica, reportamos que nuestro circuito puede proveer a cada OLED hasta 40 µA. Este nivel de corriente es apropiado considerando que la corriente m´ınima requerida para lograr 200 cd/m2 en un OLED es de 10 µA. Descriptores: Electr´onica flexible; OTFT; Perileno; Pentaceno; AMOLED. PACS: 72.80.Le, 81.05.Fb, 84.30.-r, 85.30.Tv, 85.40.- e

1.

Introduction

TFTs, formed by nanometric films of semiconductor and dielectric materials, are the most used devices to drive the current supplied to pixels of panel displays due to its good electrical performance and high scale of integration. Currently, most of the driver circuits are formed by amorphous silicon based TFTs. Alternatively, recent investigations on novel materials have shown that electronic devices can be fabricated on flexible substrates. Therefore, flexible electronics is developed for applications such as large area sensors, flexible solar cells and displays [1-3]. Low cost process, large area implementation, and potentially low power consumption are some of the attractive properties of flexible electronics. However, the performance of the devices integrated into a flexible substrate is dependent on the compatibility of the materials process during the fabrication. In order to achieve a good performance on the integration of several devices, fabrication processes are realized with organic and inorganic materials. Here, we show the development of all-organic circuits with materials and processes compatible with an OLED, that

F IGURE 1. Flexible substrate with the all organic active matrix displays, transistor and capacitor.

´ ´ ´ ˜ D. BERMAN, H.N. ALSHAREEF, B.E. GNADE, AND M. QUEVEDO-LOPEZ 20 G. GUTIERREZ-HEREDIA, L.A. GONZALEZ, A. AVENDANO,

F IGURE 4. Selection OTFT electrical characteristics. (a) IDS − VGS and (b) IDS − VDS .

F IGURE 2. Active pixel circuit diagram. Transistor and capacitor all organic based to drive an OLED.

F IGURE 5. Driver OTFT electrical characteristics. (a) IDS − VGS and (b) IDS − VDS .

it was previously reported in Ref. 4, and a flexible PEN substrate. Each circuit has the capability to supply the required current to turn a matrix of OLEDs on. All the processes were carried out at temperatures lower than 80◦ C. The circuit is based on a conventional 2 OTFTs (pentacene) configuration, with a capacitor to store the charge for the driver OTFT. The OTFTs are fabricated using a process previously reported [5-7]. In Fig. 1 we illustrate a PEN substrate containing the fabricated circuits.

2. Fabrication process

F IGURE 3. Driver OTFT cross-section with respectively layers. Selection OTFT are not shown.

The basic pixel circuit design used for the active matrix is shown in Fig. 2. A pulse is sent to the selection OTFT to active the pixel. The driver OTFT uses the voltage provided by the data line to supply the current to turns the OLED on. Charge stored in the capacitor allows a stable current supply to the OLED while the selection OTFT is off. This allows the circuit to be active. The OLED used in this work requires 1 mA/cm2 to provide a brightness of about 200 cd/m2 [4]. This corresponds to a current of about 10 µA for each individual OLED. Each OLED has a dimension of 1×1 mm. From the above calculations we estimated the width (W ) and length (L) of both the driver and selection transistors as well as the required capacitor to fulfill the time response required to get an active pixel. The W/L dimensions of the transistors are 1500/5 µm and 350/5 µm for the driver and selection OTFT, respectively. Figure 2 shows the integration process for the active pixel. We start with either glass or PEN substrates covered with 150 nm of ITO. The ITO is then patterned and wet etched to define the OLED anode contact as shown in the Fig. 3a. Next, 10 nm of chrome (Cr) and 100 nm of gold

Rev. Mex. F´ıs. 57 (2) (2011) 19–21

INTEGRATION AND ELECTRICAL CHARACTERIZATION OF ORGANIC THIN FILM TRANSISTOR FOR AN ACTIVE. . .

(Au) are deposited and patterned to define the bottom capacitor contact and the OTFT gate (Fig 3b). Figure 3c shows the gate dielectric deposition of 150 nm of parylene using chemical vapor deposition (CVD). The dielectric is then patterned and etched using Reactive Ion Etching (RIE). 100 nm of Au are deposited by e-beam for the source-drain contacts and then patterned and etched as show Fig. 3d. 150 nm of organic semiconductor, pentacene is thermally deposited followed by 300 nm of parylene as encapsulation. Both films are then patterned and etched using RIE (Fig. 3e). Another 300 nm film of parylene is deposited to encapsulate all the devices and lines before the OLED deposition as shown in Fig. 3f. This layer is then patterned and etched by RIE to open the anode contact (ITO). Finally, the OLED is deposited as described in the process reported in Ref. 4.

ent gate voltages (VGS ). As it was expected, good transistor characteristics are observed. Figure 5a and 5b show the IDS −VGS and the IDS −VDS (curve family), respectively for the driver OTFT which shows enough current to supply the OLED. At VGS of -10 V the driver OTFT provides about 10 µA, current required to achieve a brigthness of 200 cd/m2 . From these results we extracted the average hole mobility (µsat ) of 0.04 cm2 /V-s and a threshold voltage (VT ) of -2V by fitting the linear region of experimental data in these plots according to the following equation [8]. µ 1/2

ID =

4. 3.

Results and discussions

Previously reported pentacene based TFTs show that their electrical behaviour is that of p-type devices. This is also observed with the resulting TFTs. The electrical response of the fully integrated devices is shown in Fig. 4 and 5. Figure 4a shows the typical IDS − VGS for the selection TFT. Figure 4b shows the IDS − VDS curve family, where drain– source voltage (VDS ) is swept from 0 V to -20 V for differ-

Ci W µsat 2L

¶1/2 (VG − VT ).

We demonstrated an integration process for an AMOLED compatible with flexible substrates. Both driver and selection transistors showed excellent characteristics. All processes are carried out at temperatures lower than 80◦ C. The operation and performance analyses of the circuit show that this integration can deliver more than 4 times (40 µA) the required current (10 µA) to drive an OLED at a brightness of 200 cd/m2 .

5. S. Gowrisanker et al., Organic Electronics 10 (2009) 1024.

2. S.C. Lim et al., Materials Science and Engineering B 121 (2005) 211.

6. S Gowrisanker et al., Organic Electronics 10 (2009) 1217.

4. U. S. Bhansali et al., Appl. Phys. Lett. 94 (2009) 203501.

(1)

Conclusions

1. L. Zhou et al., IEEE Electron Device Letter 26 (2005) 640.

3. S. Yujuan, Z. Yi, C. Xinfa, and L. Shiyong, IEEE Transactions On Electron Devices 50 (2003) 1137.

21

7. S. Gowrisanker et al., Electrochemical and Solid-State Letters 12 (2009) H50. 8. A. Ortiz-Conde et al., Microelectronics Reliability 42 (2002) 583.

Rev. Mex. F´ıs. 57 (2) (2011) 19–21

REVISTA MEXICANA DE F´ISICA S 57 (2) 22–25

ABRIL 2011

Theoretical study of the electronic band gap in β-SiC nanowires A. Trejoa , M. Calvinoa , A. E. Ramosb , E. Carvajala , and M. Cruz-Irissona a Instituto Polit´ecnico Nacional, ESIME-Culhuacan, Av. Santa Ana 1000,M´exico, 04430, D.F., M´exico, e-mail: [email protected] b Universidad Nacional Aut´onoma de M´exico, Instituto de Investigaciones en Materiales, Apartado Postal 70-360, M´exico,04510, D.F., M´exico, e-mail: [email protected] Recibido el 7 de diciembre de 2009; aceptado el 14 de julio de 2010 The structure and electronic properties of β-SiC nanowires in the directions of growth [111] and [001] are carried out by means of density functional theory (DFT) based on the generalized gradient approximation (GGA). The dangling bonds of the surface atoms in the quantum wires are passivated using hydrogen atoms. The calculations show that both nanowires exhibit a direct energy band gap at center of Brillouin zone. The electronic band structure and band gaps show a significant dependence on the diameter, orientation and surface passivation. Keywords: Density functional theory; nanowires; silicon carbide. La estructura y las propiedades electr´onicas de nanoalambres de β-SiC crecidos en las direcciones [111] y [001] son calculadas a trav´es de la teor´ıa del funcional de la densidad (DFT) basada en la aproximaci´on de gradiente generalizado (GGA). Los enlaces rotos de los a´ tomos de la superficie en los alambres cu´anticos son pasivados usando a´ tomos de hidr´ogeno. Los resultados muestran que ambos tipos de nanoalambres presentan una brecha de energ´ıa directa en el centro de la zona de Brilllouin. La estructura de bandas electr´onica y la brecha de energ´ıa muestran una significativa dependencia del di´ametro, orientaci´on y pasivaci´on de la superficie. Descriptores: Teor´ıa del funcional de la densidad; nanoalambres; carburo de silicio. PACS: 71.15.Mb; 73.21.Hb; 62.23.Hj

1. Introduction The study of low-dimensional quantum structures has attracted great attention recently in the field of semiconductors research [1–3]. Nanowires are one of the most common one-dimensional (1-D) structures and many kinds of materials can be synthetized into nanowires structures. They present remarkable different properties and applications from their corresponding bulk forms [4]. An example, of these 1-D systems are SiC nanowires (NWs), due to their wide band gap with high electron mobility, SiCNWs would be favorable for applications in high temperature, high power, and high frequency nanoscale devices [5, 6]. In recent years SiC have been intensively studied for their potential applications in electronic devices and sensors [7]. In this work, we study the hydrogen-passivated β-SiC NWs oriented along both [001] and [111] directions [Figs. 1(a) and 1(b), respectively] using the density functional theory (DFT) based on the pseudopotential plane-wave approach with the supercell technique. The generalized gradient approximation (GGA) exchange-correlation functional used is a revised version of Perdew, Burke, and Enzerhof (RPBE) [8]. We are focusing on the electronic structure and energy gap and their dependence on wire diameter and orientation. Also, the total and partial density of states (DOS) as well as the total electron density are calculated.

2. Calculation procedure As we have mentioned above, our calculations were performed in the framework of DFT-GGA utilizing the RPBE

exchange and correlation functional. The core electrons are described using ultrasoft Vanderbilt pseudopotentials [9] within the CASTEP code [10, 11], as implemented in the Materials Studio software suite. The kinetic energy cutoff for the plane-wave basis set is 280 eV. The Brillouin zone has been sampled with a highly converged set of k points, using grids up to (1×1×6) points according to the Monkhorst Pack scheme [12], the initial bond lengths of Si-H and C-H are 0.147 nm and 0.107 nm, respectively. Nanowires are then placed in a cubic simulation cell with periodic boundary conditions. The size of the simulation cell is chosen so that

F IGURE 1. Side and top view of relaxed structural models grown in the directions (a) [001] and (b) [111] SiCNWs passivated with H (small spheres). C and Si are represented by dark and light gray spheres.

THEORETICAL STUDY OF THE ELECTRONIC BAND GAP IN β-SiC NANOWIRES

23

F IGURE 2. Energy bands for the SiCNWs grown in the directions [001] and [111], respectively. The calculated gaps are 5.04 and ˚ respectively. The maxi4.3 eV for the diameters 3.261 and 3.55 A, mum value of the valence band energy, at the Γ point, was taken as the (zero) reference.

F IGURE 4. Band gap energy of hydrogenated SiCNWs as a function of diameter. (a) [001] SiCNW (solid diamonds) and (b) [111] SiCNW (solid hexagons). The dash line is a guide for the eyes.

the distance between the cluster and its replica (due to the ˚ Under this periodic boundary conditions) is more than 12 A. consideration, the interactions between the nanowires and their replicas are negligible. Finally all the wires are relaxed to minimize the total energy using the conjugate gradient algorithm [13]. The local minimum is achieved when all resid˚ It is ual forces acting on the atoms are less than 0.03 eV/A. known that DFT systematically underestimates the semiconductor band gap energy [14]. A scissors operator of 0.97 eV has been considered, which is corresponding to the difference between experimental (2.34 eV) [15] and our calculated value (1.37 eV) for the energy bandgap of bulk crystalline β-SiC.

3.

F IGURE 3. Calculated density of states (DOS, solid line) and partial density of states (PDOS) projected onto s (dotted line) and p (dashed line) orbitals. (a) Bulk β-SiC, (b) [001], and (c) [111] SiCNWs.

Results

We present here the electronic band structure for βSiC nanowires (SiCNWs) oriented along the [001] and [111] directions. The atomic positions of all atoms were fully relaxed using the first principles methods described above. Calculation of electronic properties performed in one-dimensional Brillouin zone along the wires axis. Figures 2(a) and 2(b) shows the band dispersion along the

Rev. Mex. F´ıs. 57 (2) (2011) 22–25

24

A. TREJO, M. CALVINO, A. E. RAMOS, E. CARVAJAL, AND M. CRUZ-IRISSON

F IGURE 5. (a) Longitudinal and cross sectional view of [001] relaxed SiCNWs showing the total electron density, the isosurface ˚ −3 . (b) Electron charge density in a (110) value used was 0.2 A plane through only C atoms of bulk, showing the ropelike structure resulting from the relaxation.

nanowire direction for SiCNWs with similar width d. Notice that in both cases the SiCNWs have a direct band gap. This is true for all NWs widths considered here. It is worth noting that dangling bond-like states do not appear within the energy band gap region for SiCNWs in the two directions. This is an indication of hydrogen passivation of the surface dangling bonds that provides a smooth termination of the orbitals. Figure 3 shows the calculated total density of states (DOS) and partial density of states (PDOS) corresponding to electronic band structure of Fig. 2. An analysis of the orbital contributions shows that the eigenvalue near the valenceband maximum are pure carbon atom p states. Notice that in all cases, (a)-(c), the conduction band edge near the Brillouin zone center is primarily formed by p states, the contribution of s states being negligible. As expected from quantum size effects, we observed that the absolute value of the conduction band minimum increases in energy as the thickness of the wire decreases. This effect of quantum confinement of electron is observed all widths

F IGURE 6. (a) Longitudinal and cross sectional view of [111] relaxed SiCNWs showing the total electron density. The isosurface ˚ −3 . (b) Electron charge density in a (110) value used was 0.2 A plane through only C atoms of bulk, showing the ropelike structure resulting from the relaxation.

we studied. This leads to an increase in the electronic energy band gap (Eg) with decreasing NWs diameter. For NWs with similar diameter but different orientation we observed, in Fig. 4, that Eg is greater for wires along [001] and lower for wires along [111] directions. Besides this dependence of Eg on the diameter d and the orientation of growth axis, other ways for modifying the electronic properties would be important for applications. The orientation anisotropy in Eg reduces with the nanowire width and is expected to disappear for very thick wires when the Eg approaches that of the bulk material. Figures 5 and 6 show the isosurface (a) and the contour map (b) for the total electron density of the two kinds of nanowires [001] and [111], respectively. The electron densities cover regions in the [001] nanowire only for carbon atoms in transversal planes (observed from the [110] perspective) of growth direction. The substantial difference with [111] nanowire is the orientation of these planes, which is diagonal (observed from the same perspective) to the growth direction. In the contour maps, the lighter regions

Rev. Mex. F´ıs. 57 (2) (2011) 22–25

THEORETICAL STUDY OF THE ELECTRONIC BAND GAP IN β-SiC NANOWIRES

25

are associated with higher field values. In the isosurfaces, the darker values correspond to more populated regions. This information could be corroborated observing the contour maps for both wires. As was expected the greater magnitude values are associated with the nearest regions of the more electronegative nucleus (carbon atoms). Consistently, partial density of states [Fig. 3 (a)-(c)], for the superior edge of valence band has a primordial p character contribution, which are related with carbon atoms shown in the electron density maps. Similarly, the partial density of states for the lower edge of conduction band has a p character, we suppose that these are related with unoccupied p orbitals of Si atoms.

and [111] and their electronic band structure as a function of diameter. These properties are strongly influenced by quantum confinement. Direct fundamental band gaps are found at Gamma point for both wires, which enlarge as diameter shrinks. It is also found that [001] wires have overall a larger gap than [111] wires. SiC nanowires with direct band gap are promising candidates for optoelectronics applications such as light emitting devices and photodetectors. The wave length of the emitted or detected light can be tuned through the choice of the NWs width.

4.

This work was partially supported by SIP-IPN 20090652 and 25231-F from CONACyT. The computing facilities of DGSCA-UNAM are fully acknowledged

Conclusions

In summary we have studied, in the framework of density functional theory within the generalized gradient approximation, the structures of β-SiC nanowires oriented in the [001]

Acknowledgments

9. D. Vanderbilt, Phys. Rev. B 41 (1990) 7892.

1. R. Rurali, Phys. Rev. B 71 (2005) 205405. 2. X.H. Peng et al., J. Appl. Phys. 102 (2007) 024304. 3. E. Bekaroglu, M. Topsakal, S. Cahangirov, and S. Ciraci, Phys. Rev. B 81 (2010) 075433. 4. F. Fabbri et al., Mat. Sci. Semicond. Process. 11 (2008) 179. 5. B.K. Agrawal, A. Pathak, and S. Agrawal, J. Phys. Soc. Jap. 78 (2009) 034721. 6. Z. Wang et al., J. Phys. Chem. C 113 (2009) 12731. 7. H. Kohmo et al., Nanoscale 1 (2009) 344. 8. B. Hammer, L.B. Hansen and J.K. Norskov, Phys. Rev. B 59 (1999) 7413.

10. S.J. Clark et al., Zeitschrift fuer Kristallographie 220 (2005) 567. 11. Accelrys Inc., CASTEP Users Guide (San diego, Accelrys Inc. 2001). 12. H.J. Monkhorst, and J.D. Pack, Phys. Rev. B 13 (1976) 5188. 13. B.G. Pfrommer, M. Cˆot´e, S.G. Louie, and M.L. Cohen, J. Comp. Phys. 131 (1997) 233. 14. H. Ehrenreich and F. Spaepen, Solid State Physics eds. Advances in Research Applications 54 (Academic Press, USA, 2000). p. 179. 15. W. M. Zhou et al., Appl. Surf. Sci. 253 (2006) 2056.

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REVISTA MEXICANA DE F´ISICA S 57 (2) 26–29

ABRIL 2011

Dose dependent shift of the TL glow peak in a silicon rich oxide (SRO) film T.M. Pitersa M. Aceves-Mijaresb , D. Berman-Mendozaa , L.R. Berriel-Valdosb , and J.A. Luna-L´opezc a Centro de Investigaci´on en F´ısica, Universidad de Sonora, Apartado Postal 5-088, Hermosillo, Sonora 83190, M´exico, e-mail: [email protected]; [email protected] b ´ Instituto Nacional de Astrof´ısica, Optica y Electr´onica, Apartado Postal 51, Puebla, Puebla, 72000, Mexico, e-mail: [email protected] c El Centro de Investigaci´on en Dispositivos Semiconductores, Benem´erita Universidad Aut´onoma de Puebla, Apartado postal 1651, Puebla, Pue. 72000, M´exico. Recibido el 7 de febrero de 2010; aceptado el 18 de enero de 2010 Thermoluminescence (TL) properties of UV irradiated silicon oxide films with silicon nano particles were investigated. The TL glow curve exhibits two symmetric glow peaks, one centered at about 120◦ C and the other one centered at around 240◦ C. The position of the peak maximum of the 120◦ C TL peak appears to shift to higher temperatures with increasing radiation dose while the high temperature peak shifts to lower temperature. The shift to lower temperature with increasing radiation dose of the 240◦ C peak is typical for a second order kinetics glow peak. The shift of the 120◦ C peak to higher temperature however is peculiar and is explained in this work as an effect of the multiple phase (silicon nano particles embedded in silicon oxide) nature of the film. Keywords: Thermoluminescence; silicon; SRO; LPCVD. En este trabajo se investigaron las propiedades de Termoluminiscencia (TL) de pel´ıculas de oxido de silicio con exceso de silicio irradiadas con UV. La curva de emisi´on termoluminiscente exhibe dos picos sim´etricos, uno centrado en 120◦ C y la otra centrada en alrededor de 240◦ C. La posici´on del m´aximo pico de TL en 120◦ C parece correrse hacia altas temperaturas con el incremento en la dosis de radiaci´on, mientras el pico que se encuentra a altas temperaturas se desplaza hacia temperaturas bajas. El corrimiento hacia temperaturas bajas con el incremento de la dosis de radiaci´on del pico en 240◦ C es t´ıpico para un pico emisi´on cin´etico de segundo orden. El corrimiento del pico de emisi´on en 120◦ C hacia temperaturas altas sin embargo es peculiar y en este trabajo se explica como un efecto de la naturaleza multifases (nanopart´ıculas de Silicio embebidas en el oxido de silicio) de la pel´ıcula. Descriptores: Termoluminiscencia; silicio; SRO; LPCVD. PACS: 73.61.Cw, 74.25.Gz, 73.63.Kv, 78.60.Kn, 78.67.Bf, 81.15.Gh.

1. Introduction Optical and electrical properties of silicon-rich silica have been extensively investigated during the last decades due to the potential use of these materials in optoelectronic devices [1,2]. Many techniques have been reported to obtain SRO, including: co sputtering [3,4], RF glow discharge of SiH4 -O [5,6], Oxidation, ion-beam-assisted electron beam deposition [7,8], and CVD (Chemical Vapor Deposition) and combinations of these techniques with Si implantation [9-10]. Depending on the fabrication method and conditions SRO presents strong visible photo luminescence emission [11]. Emissions have been reported around 350 nm, 410 nm, 560 nm and 750 nm. The 350, 410 and 560 nm emissions have been ascribed to Si-O related species and oxygen vacancy related defects while the 750 nm emission is associated with some form of quantum confinement effect of the silicon clusters [12]. Evidence for the association of the 750 nm luminescence is the red shift and the decrease of intensity of this luminescence with increasing size of the silicon clusters [13-15]. In this study we investigated the UV induced thermoluminescence (TL) behavior of an annealed SRO film that contained principally 750 nm emission peak, in its photoluminescence spectrum and some emission

in the 300-475 nm range. The low temperature (120◦ C) TL peak shows a peculiar shift to higher temperature when the intensity of this peak increases. It is shown that this behavior can be explained by the multiple phase (silicon nano particles embedded in silicon oxide) nature of the film.

2.

Experimental details

The samples used in this study were cut from a 550 nm thick SRO film deposited on an N type Si wafer with a resistivity in the range of 3-5 Ωcm. The film was prepared by the LPCVD (Low Pressure Chemical Vapor Deposition) method in a hot wall reactor at 700◦ C, using a mixture of N2 O and SiH4 reactor gases. The ratio of the flow rates of the reactor gases, R0 =N2 O/SiH4 was 20 which resulted in a silicon excess of 8%. After the preparation the samples were submitted to a heat treatment at 1100◦ C in a N2 atmosphere for densification. After this treatment the samples contain nano particles of Si. The TL measurements were performed in an automated TL/OSL reader, model TL/OSL System TL-DA15, fabricated by RISO, National Laboratory, Denmark. The TL reader was adapted for the possibility to perform UV-Vis illuminations. The light source comprises a 450 W Xenon

DOSE DEPENDENT SHIFT OF THE TL GLOW PEAK IN A SILICON RICH OXIDE (SRO) FILM

F IGURE 1. TL glow curves of the SRO film sample nr A42 after 300 s irradiation with 230, 260, 290 and 350 nm light and a 600 s storage.

F IGURE 2. Raw data for the low temperature glow peak of sample B with different intensities (left graph). Gaussian shaped peaks were fitted to the data and from these fittings the peak positions T and intensities H were determined. In the middle graph are three examples a, b and c corresponding to the data a, b and c of the left graph. The determined intensities and positions (temperatures) are plotted in the right graph (scattered data points) together with a fitting of a proposed model for the shift (solid line).

lamp operated at 300 W, a home made shutter that could be gradually opened and a monochromator model GM252 (KRATOS). The intensity of the light at the sample position was about 3 mW/cm2 for 300 nm and shutter completely opened. This was determined using a pyroelectric radiometer system model 7080 purchased from Oriel. A homemade computer program controlled the monochromator, the shutter and the TL/OSL system. All experiments in the TL-reader were performed in a N2 environment.

3.

Experimental results and analyses

To determine the TL glow curve, the sample was irradiated for 5 (sample A42) or 10 min (sample B) with UV light of wavelengths between 200 and 400 nm and additionally stored for 10 min to clean “fading sensitive glow peaks”. Figures 1 shows the results of the TL measurements for different irradiation wavelengths for the sample A42. It is seen that the shape of all glow curves is very similar except for a small

27

variation in the peak temperatures. These variations appear to be related to the intensity rather than the wavelength as was verified by TL measurements after irradiations with 350 nm light of different intensities (partially opened shutter). These measurements indicate that the first TL peak shifts to a higher temperature and the second peak to a lower temperature with increasing intensity. The shift of the second peak of the TL glow curve to lower temperatures with increasing intensity (or dose) is typical for a non-first order glow peak. This effect has been well documented [16] and is related to the dependence of the recombination probability with the concentration of recombination centres. On the other hand the behaviour of the first TL glow peak, ie the shift to higher temperature with increasing intensity (or dose), has as far as we know never been described before. Here this behaviour is explained as a consequence of the presence of nano particles. The 120◦ C peak is treated as a distribution of first order peaks with their maxima at slightly different temperatures. The temperature at which the distribution has its maximum is considered as the average peak position of the individual peaks of the distribution weighted by their intensities. For the intensity of the peak, the peak height was taken. The position and peak height were determined by fitting a Gaussian function to the peak excluding the region where the peak overlaps with the second peak and taking into account the background signal. Figure 2 shows the procedure for sample B. To explain the temperature shift of the peak we propose a model in which it is assumed that during the irradiation stage the radiation defects, say electrons and holes, are generated at the Si nano particles. Some of the electrons escape from the nano particles and get trapped at TL traps in the bulk material of the film preferentially close to the nano particle. The holes get trapped at not yet specified luminescence centers. Actually, for explaining the temperature shift of the TL peak, the details about how the holes reach the luminescence centers and the assignation of these centers is not important and here it is simply assumed that the hole is left behind at the nano particle. Further it is assumed that the TL peaks of the distribution corresponding to the electron traps closest to the

F IGURE 3. Energy level diagram of the model for the trap filing. At zero doses the defect ?, generated at the nano particle np by UV light, is trapped at trap 1, which is the trap closest to the nano particle np (a). As the dose increases, the subsequent generated defects at the nano particle are trapped at trap 2, 3 and 4 which are at increasing distance from the nano particle (b), (c), (d). The energy levels of the traps increase with the approximation of the trap to the nano particle due to crystal strain caused by the presence of the nano particle.

Rev. Mex. F´ıs. 57 (2) (2011) 26–29

´ T.M. PITERS, M. ACEVES-MIJARES, D. BERMAN-MENDOZA, L.R. BERRIEL-VALDOS, AND J.A. LUNA-LOPEZ

28

nano particle have the lowest peak temperature (or lowest activation energy). This may be thought of as an effect of crystal strain caused by the presence of the nano particle. Figure 4 shows an energy level diagram and the transfer processes during the irradiation stage corresponding to this model. The readout stage of the TL process is assumed to be similar to the usually assumed, i.e. during the heating the electrons are thermally released from the traps and transfer through the conduction band to the trapped holes where, upon recombination, a photon is emitted. Within this model we make the following three assumptions: (1) The dependence of the individual peak temperatures of traps on the distance (r) between nano particle and trap is: T (r) = T∞ − (T∞ − T0 ) exp (−kr r)

kq = kr ·

4 πV ρq ρT γ 3

¶−1/3

and µ q=r·

4 πV ρq ρT γ 3

¶1/3

where ρq is the density of nano particles, ρT the density of traps, V is the volume of the sample and γ is a ‘set up’ parameter that relates the peak height with the number of filled traps. The solution for (2) is T (H) = cT0 + (1 − c)T∞

(3)

where c is:

(2) Secondly we assume that the traps are distributed homogeneously over space. This implies that the number of available traps at a distant r from a nano particle increases proportional to r2 (3) Finely we assume that the electrons generated during irradiation at a nano particle are trapped at nearest available traps. Thus when N electrons are generated at the nano particle, they are all trapped within a circumference R around the nano particle corresponding to a volume with N traps while al traps outside the circumference remain empty. Note that according to the first and second assumption, R is proportional to 3√ N. Since the intensity √ H is proportional to N, R is also proportional to 3 H. Using these three assumptions and additionally assuming that the nano particle is very small compared to the sphere of filled traps around the nano particle (so that we may integrate from the center of the nano particle) the average peak temperature can be expressed as: 3 H

µ

(1)

where kr is the spatial alteration coefficient for the peak temperature, T∞ is the peak temperature of traps far away from the nano particle and T0 is the peak temperature of traps closest to the nano particle

T (H) =

with

√ 3 ZH

(T∞ − (T∞ − T0 ) exp (−kq q)) q 2 dq (2) 0

6 − 3(Q2 + 2Q + 2) exp(−Q) Q3 √ 3 Q = kq H c=

The function T (H) depends thus only on three parameters: T∞ , T0 and kq . Fitting this function to the scattered data points of the right graph of Fig. 3 leads to the solid line in this graph with parameter values T0 = 352 K (79◦ C) and T∞ =400 K (127◦ C). The parameter kq has the value 0.5185.

4.

Conclusion

We have shown that the temperature shift as function of dose of the UV induced low temperature glow peak of SRO could be an effect of the confinement of the defect creation sites (at silicon nano particles). Good fitting results were obtained for a simple model based on the generation of radiation defects at the nano particles and subsequent trapping at traps in the bulk, in which it is assumed that (1) the traps are homogeneously distributed, (2) the traps closest to the nano particles are filled first and (3) the peak temperature of the traps obey: T = T∞ − (T∞ − T0 )exp(−kr · r) where r is the distance between traps and nano particles.

Acknowledgement We are grateful to the laboratory of microelectronics of INAOE, and especially to Mauro Landa y Pablo Alarc´on for the preparation of the SRO Films.

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˜ de un catalizador Au/TiO2 /SiO2 en la reacci´on de Estudio del desempeno oxidaci´on de CO J.A. Garc´ıa-Macedo∗ , R. Arreola-S´anchez, M.A. R´ıos-Enr´ıquez, V.M. Renter´ıa-Tapia, y G. Valverde-Aguilar Departamento de Estado S´olido, Instituto de F´ısica, Universidad Nacional Aut´onoma de M´exico, Circuito Exterior Cd. Universitaria Coyoac´an. M´exico, D.F., 04510, M´exico, ∗ Tel. (5255) 56225103; Fax (5255) 56161535, e-mail: [email protected] Recibido el 6 de febrero de 2010; aceptado el 3 de agosto de 2010 El presente trabajo reporta el estudio del desempe˜no de un catalizador con baja carga de oro depositado sobre un material compuesto de di´oxido de titanio-di´oxido de silicio en funci´on de par´ametros como la relaci´on TiO2 :SiO2 , la carga de oro, el m´etodo de deposici´on de oro, etc., y su efecto sobre la actividad catal´ıtica medida como conversi´on de CO. Con base en los resultados obtenidos se logr´o mejorar la actividad del catalizador consiguiendo oxidaci´on de CO a temperatura ambiente y la estabilidad en los materiales despu´es de 4 corridas de conversi´on de CO en un per´ıodo de 5 d´ıas. Descriptores: Catalizadores de oro; oxidaci´on de CO; alta estabilidad; materiales compuestos TiO2 /SiO2 . In this work the performance, measured as CO conversion, of a catalyst with low gold loading deposited on titanium dioxide: silicon dioxide composite was studied as function of several parameters like: TiO2 :SiO2 ratio, gold loading, gold deposition method, etc. Based on the results obtained from the experiments an improvement of the catalytic activity of the material was achieved. CO oxidation at room temperature was reached and high long catalyst stability after 4 catalytic runs during 5 days was observed. Keywords: Gold catalyst; CO oxidation; high long catalyst stability; TiO2 /SiO2 composites. PACS: 82.65.+r; 81.16.HC

1. Introducci´on A pesar de que desde los a˜nos 70’s se encuentran en la literatura estudios de catalizadores de oro soportados sobre o´ xidos [1], s´olo hasta que Haruta y col. reportaron que peque˜nas part´ıculas de oro (< 5 nm) altamente dispersas sobre o´ xidos met´alicos son muy activas en la oxidaci´on de mon´oxido de carbono (CO) e H2 a bajas temperaturas [2] se despert´o un gran inter´es en el desarrollo de catalizadores de oro. Sin embargo, la implementaci´on de estos catalizadores en aplicaciones pr´acticas ha sido lenta hasta ahora, esto debido en parte a la gran dependencia del desempe˜no de los catalizadores de las rutas y condiciones de s´ıntesis durante la preparaci´on de ellos [3]. Adem´as, las nanopart´ıculas de oro tienden a aglomerarse, disminuyendo su actividad catal´ıtica conforme crece su di´ametro [4]. Debido a la gran a´ rea superficial y alta estabilidad t´ermica de la s´ılice (SiO2 ) esta ha sido un material muy atractivo como soporte para catalizadores, especialmente para usos industriales [5]. Sin embargo, materiales de Au/SiO2 obtenidos mediante rutas de s´ıntesis similares a las empleadas para obtener catalizadores de Au/TiO2 exhiben una actividad catal´ıtica pobre, debido a que las part´ıculas de oro sobre materiales inertes como el SiO2 suelen ser grandes debido a la acidez de la s´ılice [6]. El mejor desempe˜no catal´ıtico del oro sobre o´ xidos reducibles, como el TiO2 o CeO2 se atribuye a la estabilizaci´on de las nanopart´ıculas por las interacciones m´as fuertes metal-soporte [7]. Por otro lado, el TiO2 tiene un a´ rea superficial mucho menor que la s´ılice, pero pertenece a la clasificaci´on de o´ xidos reducibles y semiconductores tipo n. Estudios recientes reportan el efec-

to positivo de TiO2 altamente disperso sobre s´ılice respecto a la estabilizaci´on de las part´ıculas soportadas de oro [8]. En general, los estudios se realizan con altas cargas de oro, alrededor del 1 % al 3 %. Investigaciones a cargas menores son muy escasas.

2.

Experimental

2.1. 2.1.1.

Preparaci´on de los catalizadores Preparaci´on del soporte de di´oxido de titanio depositado sobre s´ılice (TiO2 /SiO2 )

El dep´osito de TiO2 sobre SiO2 se realiz´o mediante procesos sol-gel. Para lograr el sol, a 5 mL de una suspensi´on coloiR dal comercial, (Ludox°AS-40; 40 % en agua, con di´ametro de part´ıcula de 24 nm, a´ rea superficial de 135 m2 /g) o suspensi´on de s´ılice pirog´enica (Sigma-Aldrich con di´ametro de part´ıcula de 7 nm, a´ rea superficial de 300 m2 /g) se agregaron lentamente las cantidades correspondientes de una disoluci´on al 20 % en isopropanol anhidro (Sigma-Aldrich) de tetra-isoprop´oxido de titanio (98 % Sigma-Aldrich) bajo fuerte agitaci´on para lograr materiales con relaciones de 11 y 22 % en peso de TiO2 . Al sol resultante se le permiti´o la gelaci´on por 3 semanas. El material resultante se sec´o a 200◦ C por 12 h. 2.1.2.

Deposici´on de oro sobre el soporte

Previo a la deposici´on de Au sobre los materiales se prepar´o una disoluci´on de HAuCl4 (Sigma-Aldrich, 99.99 %)

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separ´o el s´olido por centrifugaci´on y se sec´o a 150◦ C por 8 h y se almacen´o sin cuidados especiales. 2.2.

F IGURA 1. Efecto de la temperatura de activaci´on de la muestra sobre la actividad catal´ıtica, para muestras activadas a 300 y 600◦ C en atm´osfera de hidr´ogeno.

F IGURA 2. Efecto del soporte de la muestra sobre la actividad R catal´ıtica, para TiO2 depositado sobre s´ılice Ludox° y s´ılice pirog´enica.

con una concentraci´on de 1×10−4 M, de NaOH 0.1 M y de urea 1.2 M. A) Para el m´etodo de deposici´on con NaOH (DP-NaOH) el valor del pH de la disoluci´on de HAuCl4 fue ajustado a 9.0 a˜nadiendo lentamente peque˜nas cantidades de la disoluci´on NaOH. Acto seguido se adicion´o el soporte de TiO2 /SiO2 en una proporci´on de 1 g de soporte por 50 mL de disoluci´on. La suspensi´on resultante se agit´o vigorosamente por 2 h a temperatura ambiente, para despu´es separar el s´olido por centrifugaci´on. La adici´on del soporte a la disoluci´on disminuy´o el valor de su pH, por lo que fue necesario a˜nadir, posteriormente, peque˜nas cantidades de disoluci´on de NaOH para reajustar el valor a 9.0. El material resultante se sec´o a 150◦ C por 8 h. El material fr´ıo se almacen´o en viales sin cuidados especiales [3]. B) Para la deposici´on empleando urea (DP-Urea) al´ıcuotas de las disoluciones de HAuCl4 y la de urea se mezclaron para obtener una diluci´on con una concentraci´on de 0.42 M de urea y la cantidad necesaria de Au en disoluci´on. A esta mezcla se le agreg´o, bajo fuerte agitaci´on, el soporte en una proporci´on de 1 g por 100 mL de disoluci´on. Despu´es de 24 h de agitaci´on se

Actividad catal´ıtica

La determinaciones de la disminuci´on de la concentraci´on de CO se realiz´o en un microreactor de lecho catal´ıtico de cuarzo (∅DI =0.9 mm) a presi´on atmosf´erica (585 mmHg) conectado en l´ınea a un cromat´ografo de gases Perkin-Elmer Clarus 500 equipado con un detector de ionizaci´on de flama y una columna empacada. 40 mg de catalizador con una carga de 0.01 % en peso fueron colocados en el reactor y este fue alimentado con una mezcla de concentraci´on certificada (Praxair) al 1 % de CO y 1 % de O2 en balance de nitr´ogeno (v:v), con un flujo de 20 mL/min. Con el objeto de reducir el Au3+ del ion [AuCl4 ]− a Au met´alico (Au0 ) las muestras del catalizador Au/TiO2 /SiO2 se sometieron in situ a una corriente de hidr´ogeno con un flujo de 60 mL/min a 300◦ C y 600◦ C por 2 horas. El porcentaje de conversi´on en funci´on de la temperatura de reacci´on de determin´o seg´un la Ec. (1): µ ¶ A0 %Convco = 100 ∗ 1 − (1) AT Donde: % ConvCO = porcentaje de conversi´on de CO a la ´ temperatura de reacci´on T A0 = Area cromatogr´afica del CO en ausencia de catalizador correspondiente a la concentraci´on ´ inicial AT = Area cromatogr´afica del CO a la temperatura T en presencia de catalizador. 2.3.

Microscop´ıa electr´onica

Microfotograf´ıas de alta resoluci´on fueron tomadas con un microscopio JEOL JEM-2010F FasTEM con un voltaje de aceleraci´on de 200 KV y una resoluci´on punto apunto de 0.19 nm.

3.

Resultados y discusi´on

Actividad catal´ıtica Efecto de la temperatura de activaci´on Muestras de catalizador con cargas del 11 % de TiO2 y 0.01 % de Au fueron activadas en presencia de hidr´ogeno a 300 y 600◦ C por 2 h. Las muestras que fueron activadas a 300◦ C durante la primera corrida de oxidaci´on de CO presentaron conversiones comparativamente menores que las muestras activadas a 600◦ C. Sin embargo, al emplear las muestras activadas a 300◦ C la curva de conversi´on de CO se desplaz´o hacia temperaturas menores logrando conversiones mayores en comparaci´on con los resultados de la segunda corrida con la muestra activada a 600◦ C, como se observa en la Fig. 1. Cuando el catalizador es activado a 300◦ C, la cantidad de hidr´oxido de oro convertido a oro met´alico es comparativamente menor que en el catalizador activado a 600◦ C,

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tos resultados obtenidos en estas muestras, los subsecuentes materiales fueron activados a 300◦ C en presencia de H2 . Efecto del soporte en la actividad

F IGURA 3. Efecto de la concentraci´on de TiO2 (15 y 22 %) en muestras Au/TiO2 /SiO2 sobre la actividad catal´ıtica (s´ılice R Ludox° ).

permaneciendo oro no reducido remanente sobre la superficie del soporte, como es reportado en la literatura [10]. Por otra parte, Haruta reporta que hidr´oxidos de oro pueden reaccionar con el ox´ıgeno para producir o´ xidos de oro que son inestables, los cuales pueden ser reducidos f´acilmente a Au0 en presencia CO para formar CO2 [11], respaldando esta hip´otesis. Lo anterior podr´ıa explicar la mayor actividad durante el primer ciclo de oxidaci´on y, parcialmente el mejor desempe˜no del catalizador activado 300◦ C durante la segunda corrida. Por otra parte, la muestra que fue activada a 600◦ C no presenta mejor´ıa significativa, durante la segunda corrida de oxidaci´on de CO, en su actividad catal´ıtica. Lo anterior puede ser debido a que a mayores temperaturas se logra reducir una mayor cantidad de oro proveniente del precursor, pero al ser expuesta a una temperatura tan alta las part´ıculas de oro tienden a aglomerarse y aumentando as´ı su tama˜no durante el proceso de calentamiento (activaci´on) [9,11]. Si bien, al ser activadas a 300◦ C el proceso de reducci´on del precursor de oro a oro met´alico es significativamente m´as lenta, los resultados mostrados en la Fig. 1, muestran una mejora en el desempe˜no catal´ıtico del material en funci´on del tiempo, lo cual siguiere que una reducci´on a temperaturas menores puede resultar en un mejor desempe˜no del catalizador. Durante la tercera corrida, los resultados obtenidos con el material activado a 300◦ C exhiben una ligera mejor´ıa en la respuesta, presentando un mayor porcentaje de oxidaci´on (Fig. 1), mientras que la muestra activada a 600◦ C no present´o ninguna mejor´ıa significativa. A pesar de que empleando temperaturas menores para la activaci´on de los catalizadores, la reducci´on de oro parece ser m´as lenta y que requiere de un tratamiento en presencia de CO y O2 para lograr su mejor desempe˜no catal´ıtico, se disminuye la posibilidad del crecimiento de las part´ıculas reduciendo la p´erdida de actividad del material. Estos resultados son congruentes con la explicaci´on anterior y con lo reportado por Haruta [11] y Zanella et al. [4,9] de que las part´ıculas met´alicas de oro tienden a crecer en tama˜no al aumentar la temperatura. De la Fig.1 se observa que ninguno de los materiales present´o p´erdida de la actividad catal´ıtica despu´es de tres corridas. Con base en es-

Los resultados de la actividad catal´ıtica de muestras con las mismas proporciones de TiO2 (11 %) y una carga de 0.01 % de oro soportadas sobre dos diferentes s´ılices se muestran en la Fig. 2, en la cual se observa una gran diferencia en la respuesta de ambas muestras. El resultado es coherente con los reportados por Moreau y Bond [12] quienes observaron que al aumentar el a´ rea superficial del soporte por arriba de 200 m2 /g disminuye la actividad catal´ıtica, logrando mayor actividad cuando emplearon materiales con a´ reas superficiales entre 30 y 100 m2 /g. Efecto de la concentraci´on de TiO2 De los resultados arriba discutidos se observa que el tipo de s´ılice empleada tiene un papel importante en la actividad catal´ıtica del material final. Con base en lo anterior y con el objeto de estudiar el efecto de la cantidad de di´oxido de titanio en el material sobre la actividad catal´ıtica, se prepaR como soporte con raron muestras empleado s´ılice Ludox° dos proporciones diferentes de TiO2 (15 y 22 %). Los resultados del catalizador con una carga del 22 % de di´oxido de titanio muestran actividad catal´ıtica a menores temperaturas, alrededor de 50◦ C por debajo, respecto al material con 15 % de TiO2 (Fig. 3). Venezia et al. reportan que catalizadores de oro sobre TiO2 /SiO2 con concentraciones menores al 5 % en peso de TiO2 tienen mayor actividad comparativamente a la de catalizadores con contenidos mayores de TiO2 . Ellos calcinan los materiales compuestos, con el fin de eliminar el disolvente y material org´anico, antes de realizar la deposici´on de oro. Discuten que al aumentar la concentraci´on de TiO2 sobre el soporte de s´ılice, e´ ste tiende a aglomerarse y formar cristales de anatasa durante el proceso de calcinaci´on a altas temperaturas, originando un cambio estructural sobre la superficie del soporte [13]. Tai et al., por su parte, explican que estos cambios estructurales sobre la superficie de los materiales compuestos de TiO2 /SiO2 hacen que la movilidad de las part´ıculas met´alicas de oro sea mayor, aumentando la posibilidad de coalescencia; provocando as´ı, el aumento en el tama˜no de part´ıculas de oro y la subsecuente disminuci´on de la actividad catal´ıtica [15]. As´ı mismo, Venezia et al. concluyen que, con bajos contenidos de TiO2 , menores al 5 > %, e´ ste se encuentra altamente disperso, permaneciendo de esta manera con una estructura amorfa [13]. Estas observaciones [13,14] son coherentes con los resultados de los estudios de espectroscopia de IR y de difracci´on de Rx realizados por Renter´ıa et al., donde se observa que el TiO2 no forma fases cristalina por debajo de 600◦ C [15]. As´ı mismo, Renter´ıa observa que grupos alc´oxido no hidrolizados permanecen presentes en el material incluso despu´es de haber sido tratados t´ermicamente por 6h a 500◦ C [15]. Se puede decir, con base a los resultados obtenidos por Renter´ıa et al., que el TiO2 se

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electr´onica de transmisi´on (TEM) de alta resoluci´on, en ella se ve claramente como una part´ıcula de oro se deposit´o sobre un peque˜no casquete de TiO2 que creci´o sobre una part´ıcula R. de s´ılice Ludox° Est´a bien documentado que el tama˜no de las part´ıculas met´alicas es de suma importancia en la actividad catal´ıtica de los catalizadores de oro [11] y e´ ste aumenta al someter el catalizador a altas temperaturas o con su empleo en la oxidaci´on de CO [4,9,11]. La muestra con una concentraci´on del 22 % de TiO2 y una carga de 0.01 % de oro fue sometida a 2 corridas adicionales en un periodo de 5 d´ıas. Se observa en la Fig. 5 un desplazamiento hacia temperaturas m´as altas en la F IGURA 4. Efecto de la carga de Au (0.01 y 0.002 %) en muestras R de Au/TiO2 /SiO2 sobre la actividad catal´ıtica (s´ılice Ludox°).

F IGURA 5. Efecto del n´umero de corridas sobre la actividad caR tal´ıtica (s´ılice Ludox°).

encuentra con una estructura amorfa incluso en los materiales con contenidos mayores al 20 %, debido a que las muestras de TiO2 /SiO2 no fueron sometidos a temperaturas mayores a 150◦ C antes de ser probadas como catalizadores. Se ha reportado en diversas investigaciones que la actividad catal´ıtica depende de la carga de oro en los catalizadores [11]. Muestras de TiO2 /SiO2 al 22 % de TiO2 fueran depositadas con dos diferentes cargas de Au empleando el m´etodo DP-urea [8], los materiales obtenidos fueron empleados durante la oxidaci´on de CO y como se muestra en la Fig. 4. Estos dos materiales ya presentan actividad a temperatura ambiente. En esta figura es claro que al disminuir la carga de oro en un 80 % la actividad catal´ıtica mejora, logrando una conversi´on de CO del 50 % durante la segunda corrida aproximadamente a los 138◦ C en comparaci´on con el material con una carga de oro de 0.01 % que alcanz´o la misma conversi´on alrededor de los 255◦ C, es decir, 117◦ C por encima de la anterior. La disminuci´on en la actividad catal´ıtica est´a relacionada con el aumento en el tama˜no de las part´ıculas de oro sobre el catalizador, al aumentar la carga de oro, aunado a que el oro se deposita preferentemente sobre el di´oxido de titanio y no sobre la s´ılice [13], aumentando la posibilidad de crecimiento de las part´ıculas. La Fig. 6a muestra una micrograf´ıa de contraste Z, realizada mediante microsc´opica

F IGURA 6. a) b)Micrograf´ıa de contraste Z de muestras de nanopart´ıculas de oro depositadas sobre TiO2 /SiO2 ; concentraci´on de oro de 0.01 %.

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F IGURA 7. a) Micrograf´ıa de una nano part´ıcula de oro tomada por TEM. b) Espectro de dispersi´on de energ´ıa (EDS) realizado en el TEM sobre la nanopart´ıcula de oro.

tercera y cuarta corrida, de 49 y 58◦ C al 50 % de conversi´on, respectivamente; es decir, una disminuci´on en la actividad catal´ıtica debido, presumiblemente, al crecimiento de las part´ıculas de oro. Comotti et al., prepararon catalizadores de Au/TiO2 partiendo de una suspensi´on coloidal de oro empleando alcohol polivin´ılico o glucosa monohidratada como agentes protectores de los coloides. Ellos observan que estos catalizadores mejoran su actividad catal´ıtica despu´es de varias corridas de oxidaci´on de CO y, que a partir de la cuarta corrida se observa una ca´ıda en la actividad catal´ıtica [16]. Estas observaciones coinciden con el comportamiento de nuestro catalizador. Comotti explica que, esta mejor´ıa en la respuesta catal´ıtica se debe a que el catalizador requiere de un tratamiento t´ermico m´as prolongado en presencia de ox´ıgeno, esto con el fin de eliminar los agentes protectores que tienen un efecto inhibitorio en la reacci´on y, para lograr la m´axima activaci´on del catalizador [16]. Por u´ ltimo, reali-

zan tratamientos t´ermicos a diferentes temperaturas y observan que cuando se calientan los catalizadores por arriba de los 400◦ C las p´erdidas en la actividad catal´ıtica son mayores. Los estudios de espectrofotometr´ıa de IR de Renter´ıa et al. [15] revelan la presencia de grupos alc´oxido no hidrolizados que permanecen sobre el sobre el soporte de Au/TiO2 . Estos grupos normalmente son eliminados por calcinaci´on a temperaturas arriba de los 500◦ C, con la consecuente modificaci´on estructural el TiO2 para formar cristales de anatasa [13,14]. Estos grupos alc´oxido persistentes en nuestros materiales pueden actuar como inhibidores de la reacci´on de CO y ser eliminados paulatinamente en presencia de O2 para que la actividad catal´ıtica mejore durante las primeras corridas como se observa en la Fig. 5. El decaimiento claramente marcado de la actividad catal´ıtica, a partir de la tercera corrida, tiene explicaci´on en el hecho de que estos materiales durante las corridas catal´ıticas fueron calentados hasta los 550◦ C en cada una de ellas, lo cual coincide con los resultados reportados por Comotti et al. [16]. Lo anterior supone una explicaci´on alternativa y complementaria del desplazamiento de las cuervas hac´ıa temperaturas m´as bajas. Se plantean realizar an´alisis termogravim´etricos y t´ermico diferencial para obtener datos que expliquen a detalle este comportamiento. Cabe destacar que, los resultados obtenidos en estas dos u´ ltimas corridas son muy similares, las curvas se superponen alrededor del 60 % de conversi´on y la diferencia de temperatura entre la tercera y cuarta corrida al 50 % es de s´olo 9◦ C despu´es de 3 d´ıas, es decir, la p´erdida de actividad es mucho menor, lo cual indica que el material tiende a ser estable. DeR y bido al dep´osito de di´oxido de titanio sobre s´ılice Ludox° a que la carga de oro es muy baja, la dispersi´on de oro en el material no es homog´enea (Fig. 6b). De acuerdo al an´alisis qu´ımico hecho por TEM-EDS, la composici´on de la part´ıcula mostrada en la Fig. 7, contiene oro, titanio, silicio y ox´ıgeno, los cuales corresponden al material Au/TiO2 /SiO2 , las se˜nales de cobre corresponden a la rejilla en la cual est´a montada la muestra.

4. Conclusiones El desempe˜no, medido como actividad catal´ıtica, de materiales compuestos de TiO2 /SiO2 , es sensible a par´ametros como el contenido de TiO2 , contenido de oro, tipo de SiO2 que se emplea, m´etodo de deposici´on de las nanopart´ıculas de oro y la temperatura de activaci´on de los mismos. Con base en los resultados obtenidos durante el estudio, se logr´o mejorar la actividad del catalizador consiguiendo actividad a temperatura ambiente y estabilidad en los materiales durante 5 d´ıas despu´es de 4 corridas de conversi´on de CO.

Agradecimiento Los autores agradecen el apoyo financiero de los proyectos: CONACYT 79781, NSF-CONACYT, PUNTA, PAPIIT IN107510 y al ICyT-DF. M. A. R´ıos-Enr´ıquez y GVA agradecen al ICyT-DF por la Beca posdoctoral.

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REVISTA MEXICANA DE F´ISICA S 57 (2) 36–40

ABRIL 2011

Photocatalytic activity in the visible region of high energy milled TiO2 :N nanopowders L. Rojas-Blanco, F.J. Espinoza-Beltr´an, P.G. Mani-Gonz´alez and R. Ram´ırez-Bon Centro de Investigaci´on y Estudios Avanzados del IPN. Unidad Quer´etaro, apartado postal 1-798, Quer´etaro, Qro., 76001, M´exico, e-mails: smile [email protected]; [email protected]; [email protected]; [email protected] G. Zambrano Thin Films Group, Universidad del Valle, Cali, Colombia, e-mail: [email protected] J. Vel´asquez-Salazar Departamento de F´ısica de la Universidad de Austin Texas, University Station C1600 Austin, Texas, e-mail: [email protected] Recibido el 26 de febrero de 2009; aceptado el 16 de julio de 2010 In this work, TiO2 :N nanopowders were synthesized by high-energy ball milling using commercial titanium dioxide (TiO2 ) in the anatase crystalline phase and urea to introduce nitrogen into the TiO2 lattice in order to enhance their photocatalytic properties in the visible spectral region. Several samples were prepared by milling a mixture of TiO2 -urea powders during 2, 4, 8, 12 and 24 hours and characterized by spectroscopic and analytical techniques. X-ray diffraction (XRD) results showed the coexistence of anatase and the high-pressure srilankite TiO2 crystalline phases in the samples. Scanning electron microscopy (SEM) revealed that the grain size of the powder samples decreases to about 200 nm after 24 h of milling. UV–Vis diffuse reflectance spectroscopy measurement showed a clear red-shift in the onset of light absorption from about 390 to 470 nm as consequence of nitrogen doping in the samples. The photocatalytic activity of the TiO2 :N samples was evaluated by methylene blue degradation under visible light irradiation. We found that all samples have a higher photocatalytic activity than the undoped TiO2 , which can be attributed to the effect of the introduction of N atoms into the TiO2 lattice. X-ray photoelectron spectroscopy (XPS) measurements were performed to confirm the presence of N and determine its chemical bonding in the samples. Keywords: Photocatalysis; titanium oxide; nitrogen doping; srilankite. En este trabajo, nanopolvos de TiO2 :N fueron sintetizados en un molino de bolas de alta energ´ıa usando di´oxido de titanio comercial (TiO2 ) en fase cristalina anatasa y urea para introducir nitr´ogeno en la red de TiO2 con el objetivo de optimizar sus propiedades fotocatal´ıticas en la regi´on espectral visible. Varias muestras fueron preparadas moliendo una mezcla de TiO2 -urea en polvo durante 2, 4, 8, 12 y 24 horas y caracterizadas por t´ecnicas espectrosc´opicas y anal´ıticas. Los resultados de difracci´on de rayos X (DRX) mostraron la coexistencia de las fases cristalinas anatasa y srilankita de alta presi´on de TiO2 . La microscop´ıa electr´onica de barrido (MEB) revel´o que el tama˜no de grano decrece a alrededor de 200 nm despu´es de 24 horas de molienda. Los datos de espectroscopia por reflectancia difusa en el UV-Vis mostraron un corrimiento en el borde de absorci´on de 390 a 470 nm debido a la introducci´on de nitr´ogeno en las muestras. La actividad fotocatal´ıtica de las muestras de TiO2 :N fue evaluada por la degradaci´on del azul de metileno bajo irradiaci´on de luz visible. Encontramos que todas las muestras tienen una actividad fotocatal´ıtica m´as alta que las del TiO2 sin dopar, lo cual puede ser atribuido al efecto de la introducci´on de a´ tomos de N. Se realizaron tambi´en mediciones de espectroscop´ıa de electrones fotoemitidos por rayos X para confirmar la presencia y enlace qu´ımico de los a´ tomos de N en las muestras. Descriptores: Fotocat´alisis; oxido de titanio; dopaje con nitr´ogeno; srilankita. PACS: 82.30.Vy; 82.45.Jn; 81.07.Wx; 82.80.Pv

1. Introduction Heterogeneous photocatalysis of organic compounds on semiconductor surfaces in aqueous media is an efficient method for waste water treatment and its use has been increased in recent years [1]. TiO2 is a low-cost photocatalyst and is the most extensively investigated material in this field. However, its high energy band gap (3.3 eV for anatase phase) limits the photocatalytic process to irradiation wavelengths in the UV region (λ