Nanocomposites and Nanoporous Materials VIII [1 ed.] 9783038131960, 9783908451488

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Nanocomposites and Nanoporous Materials VIII

Nanocomposites and Nanoporous Materials VIII ISNNM8

Selected, peer reviewed papers from the 8th International Symposium on Nanocomposites and Nanoporous Materials (ISNNM8), February 22-24, 2007, Jeju, Korea

Edited by

Chang Kyu Rhee

TRANS TECH PUBLICATIONS LTD Switzerland • UK • USA

Copyright  2008 Trans Tech Publications Ltd, Switzerland

All rights reserved. No part of the contents of this book may be reproduced or transmitted in any form or by any means without the written permission of the publisher.

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Volume 135 of Solid State Phenomena ISSN 1012-0394 (Pt. B of Diffusion and Defect Data - Solid State Data (ISSN 0377-6883)) Full text available online at http://www.scientific.net

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International Advisory Committee A. Ilyuschenko (Belarussian State Powder Metall. Inst., Belarus) Anatoly Y. Yermakov (IMP, Russia) E. Olevsky (San Diego State Univ., USA) K. Niihara (Nagaoka Univ. Tech., Japan) T. Sekino (Osaka Univ., Japan) Valery N. Charushin (IOS, Russia) V. Ivanov (IEP, Russia) Y. Kotov (IEP, Russia) Z. H. Stachurski (Aus. Nat. Univ., Australia) T. Pieczonka (AGH Univ. of Sci. & Tech., Poland) Jeff W. Eerkens (Univ. Missouri-Columbia, USA)

ISNNM8 Steering Committee Jai Sung Lee (Hanyang Univ., Korea) Whung-Whoe Kim (KAERI, Korea) Sang-Eon Park (Inha Univ., Korea) Suk-Joong L. Kang (KAIST, Korea) Young Soon Kwon (Ulsan Univ., Korea) Sung Soo Kim (KRICT, Korea) Soo-Jin Park (Inha Univ., Korea) Jae Sung Lee (POSTECH, Korea) Chang Kyu Rhee (KAERI, Korea) Jong-San Chang (KRICT, Korea) Yoo Dong Hahn (KIMM, Korea) Hong Pyo Kim (KAERI, Korea)

ISNNM8 Academic Committee Young Keun Jeong (Pusan Nat. Univ., Korea) Seok Kim (KRICT, Korea) Young Do Kim (Hanyang Univ., Korea) Ji Soon Kim (Ulsan Univ., Korea) Sung-Tag Oh (SNUT, Korea) Si-Young Jang (Hankuk Aviation Univ., Korea) Yong-Ho Choa (Hanyang Univ. Korea) Jin-Soo Hwang (KRICT, Korea) Min Ku Lee (KAERI, Korea) Moon-Hee Hong (DATQ, Korea)

Local Organizing Committee Jin Ju Park (KAERI, Korea) Dong Jin Kim (KAERI, Korea) Young Rang Uhm (KAERI, Korea) Jung Gu Lee (KAERI, Korea) Jae Woo Kim (KAERI, Korea) Young Kyu Hwang (KRICT, Korea)

PREFACE The 8th International Symposium on Nanocomposites & Nanoporous Materials 2007, ISNNM8 was held in Jeju, Korea, from 22nd to 24th February 2007. Unlike previous series of nanocomposite materials symposiums, the Symposium is a joint meeting together with nanocomposites, nanoporous materials and environment-friendly material research program. Approximately 200 participants from 5 countries attended the symposium and 164 papers were presented; 12 plenary lectures from Poland, India and Russia and 15 oral presentations as well as 137 poster presentations. After the reviewing process, 36 papers were accepted for publication in this special volume of Solid State Phenomena. The strong international participation and high quality of presentations is an indication of the interest in the field of nanocomposite, nano-catalyst, ultrafine polymers, nano-adsorption, nano-characterization etc., in fundamental research and applied engineering. All main aspects of these materials would be covered, including synthesis, mechanism, microstructures, properties and applications. The symposium also provides the latest research results and a state-of-the-art overview of technology in the exciting and rapidly evolving field of nanomaterials. In the first day of the Symposium, lots of papers in the field of nano structured materials, synthesis of nanocomposites, mechanical milling were mainly presented. In the second day, invited lectures from India, Russia, Poland were actively proceed based on nanocomposite coating, nanocomposite photocatalyst, photodynamic therapy, carbon nanotube etc. In the last day, a lot of oral presentations were made progress on the subject of nano-catalyst, synthesis of nanopowder and environmentally friendly materials. Above all, in the second day of the Symposium, SU-IL PYUN’s Workshop on Progress in Materials and Corrosion was also held. This workshop was dedicated to Professor Su-Il Pyun on the occasion of his retirement from Korea Advanced Institute of Science and Technology (KAIST). He pioneered remarkable advances in corrosion and interfacial electrochemistry at KAIST on the highest intellectual and emotional levels for 31 years. Up to now he has devoted his whole life to the fundamental research works, and recently continues his thorough investigation in the areas of interfacial electrochemistry of carbon and transition metal oxide electrodes based on the fractal theory and corrosion of metals using electrochemical noise analysis. This workshop covered the fundamentals of materials and corrosion, and their applications into practices investigated by himself and his disciples who are actively working at the university, national institute, industry and government in the areas of materials science and engineering, especially, corrosion and interfacial electrochemistry. We would like to express our sincere thanks to all the colleagues which contributed to the success of ISNNM8 2007; the members of the ISNNM Steering Committee, session chairs, invited speakers and all the participants. Furthermore we would like to appreciate the Local Organizing Committee and the invaluable administrative support of High Performance Nanocomposites Program (KAERI), Nano Center for Fine Chemicals Fusion Technology (Inha University), Research Center for Nano Catalysis & Environment-friendly Material Research Center (KRICT) and The Korean Powder Metallurgy Institute, before and during the Symposium. Finally, we would like to acknowledge the financial support of many industrial companies and the following Korean government agencies; Ministry of Commerce, Industry and Energy (MOCIE). Without their support, the Symposium would have been less successful. Jai-Sung Lee Whung-Whoe Kim Sang-Eon Park Jong-San Chang Soo-Jin Park

Table of Contents Committees Preface Catalytic Combustion of Effluents from Methane-Based MCFC Device over Cordierite Supported Pd/La-Al2O3 Catalyst J.M. Lee, Y.K. Hwang, A.S. Mamman, S.M. Lee, D.Y. Hong, K.Y. Ahn and J.S. Chang Low-Temperature Fabrication of Polycrystalline Yttrium Aluminum Garnet Powder via a Mechanochemical Solid Reaction of Nanocrystalline Yttria with Transition Alumina H.G. Jung, Y.H. Cheong, I.D. Han, S.J. Kim and S.G. Kang Formation of Hollow Zinc Oxide by Oxidation and Subsequent Thermal Treatment J.G. Lee, R. Nakamura, D. Tokozakura, H. Nakajima, H. Mori and J.H. Lee Antifungal Effectiveness of Nanosilver Colloid against Rose Powdery Mildew in Greenhouses H.S. Kim, H.S. Kang, G.J. Chu and H.S. Byun Fabrication of Oriented TiO2-Based Nanotube Array Thin Films Y.P. Guo, N.H. Lee, H.J. Oh, C.R. Yoon, C.K. Rhee, K.S. Lee and S.J. Kim Formation of Lanthanum Hydroxide and Oxide via Precipitation S.J. Kim, W.K. Han, S.G. Kang, M.S. Han and Y.H. Cheong Colloidal Crystal Templating of Two-Dimensional Ordered Macroporous SiCN Ceramics I.H. Song, Y.J. Kim, H.D. Kim and D.P. Kim Reflectometry Studies of Mesoporous Silica Thin Films Y.K. Hwang, A.S. Mamman, K.R. Patil, L.K. Kim, J.S. Hwang and J.S. Chang Conductive Property of Carbon-Nanotube Dispersed Nanocomposite Coatings for Steel Y.S. Yang, M.J. Kim, Y.C. Lee and S.T. Noh Preparation of Platinum-Ruthenium Nanoparticles on Graphite Nanofibers S. Kim and S.J. Park Thermal Behaviors and Fracture Toughness of Polyurethane-Dispersed Difunctional Epoxy Resins S.J. Park and J.R. Lee Influence of Multiwalled Carbon Nanotube on Rheological Behavior of Mesophase Pitches Y.S. Lee and S.J. Park Preparation and Characterization of AuNP/Al2O3 with Bimodal Nanoporous Structure Y.H. Kim, J.B. Joo, W.Y. Kim, J.J. Lee and J.H. Yi The Effect of Physicochemical Treatment on Pd Dispersion of Carbon-Supported Pd Catalysts Y.T. Kim, E.D. Park, M. Kang and J.E. Yie The Effect of Si/Al Ratio on Selective Catalytic Reduction of NOx with NH3 over Pt/AlSBA-15 M. Kang, J.H. Park, E.D. Park, J.M. Kim, D.J. Kim and J.E. Yie Effect of Ball-Milling Method on the Formation of ODS Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3 Powders J.H. Ahn, B.H. Park and J. Jang Nano-Sized Yttria Dispersed Ferritic Stainless Steels for SOFC Interconnect Applications J.H. Ahn, H.J. Kim, I.H. Oh, B.H. Park and S.H. Jang Methane Storage on Surface Modified Activated Carbons S.M. Yun, J.W. Kim, H.K. Jin, Y.H. Kim and Y.S. Lee Nickel Decoration on Multi-Walled Carbon Nanotubes Using Multi-Step Impregnation Method S.D. Kim, S.J. Park and Y.S. Lee Preparation and Characterization of Electrospun Carbon Nanofibers with Na2CO3/H3PO4 Activation J.S. Im, S.J. Park and Y.S. Lee Removal of Hexavalent Chromium on Chitosan-Deposited Activated Carbon J.M. Lee, K. Palanivelu and Y.S. Lee

1 7 11 15 19 23 27 31 35 39 43 47 53 57 61 65 69 73 77 81 85

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Oxidation of Sulfur Components in Diesel Fuel with Tert-Butyl Hydroperoxide Using Chromium Containing Catalysts K.E. Jeong, H.J. Chae, U.C. Kim, S.Y. Jeong and W.S. Ahn Development and Application of Irradiation Technology in HANARO K.N. Choo and B.G. Kim Synthesis of Cu(In0.75Al0.25)Se2 Thin Films from Binary Selenides Powder Compacted Targets by Sputtering and Selenization B. Munir, R.A. Wibowo and K.H. Kim A Passive Film Formed on Alloy 600 in High Temperature Aqueous Solution D.J. Kim, H.C. Kwon, S.S. Hwang and H.P. Kim Formation of Nano M2X Particles by a Tempering in High Cr Ferritic/Martensitic Steel S.H. Kim, C.H. Han and W.S. Ryu Effect of Vanadium on Development of Acicular Ferrite Microstructure in Low Carbon Steel H.H. Jin, C.S. Shin and H.C. Lee Nucleation of Intragranular Ferrite on B1-Type Non-Metallic Inclusions H.H. Jin, C.S. Shin, H.C. Lee and W.W. Kim Microstructure and Nano-Indentation Properties of Ion-Irradiated Fe-9wt%Cr Alloy H.H. Jin, C.S. Shin and W.W. Kim Synthesis and Characterization of NiFe2O4 Nanoparticles Synthesized by Levitational Gas Condensation (LGC) Y.R. Uhm, B.S. Han, H.M. Lee and C.K. Rhee Corrosion Behavior of Ceramic Dispersion Strengthened High-Cr Stainless Steel G.J. Lee, S.H. Lee, S.M. Hong, B.S. Han, J.J. Park, M.K. Lee and C.K. Rhee Brazing of Ti Using a Zr-Based Amorphous Filler J.G. Lee, J.K. Lee, M.K. Lee, D.N. Shim and C.K. Rhee The Effect of Ag Diffusion Barrier on the Microstructure of a Titanium Dissimilar Joining M.K. Lee, J.G. Lee, J.K. Lee, J.J. Park, Y.R. Uhm and C.K. Rhee Effect of Magnesium Addition on the Phase Transformation of α-Alumina Prepared from Route of Ammonium Aluminum Carbonate Hydroxide Y.M. Jung and S.W. Kim Reduction Behavior of W and Cu Oxides Powder Mixture S. Lee, J.W. Noh, E.P. Kim and M.H. Hong Comparative Studies of Cellular Permeable Solids as Catalyst Supports O. Smorygo, A. Leonov, Z.R. Ismagilov and C.K. Rhee

89 93 99 103 107 111 115 119 123 127 131 135 139 143 150

Solid State Phenomena Vol. 135 (2008) pp 1-6 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.1

Catalytic Combustion of Effluents from Methane-based MCFC Device over Cordierite Supported Pd/La-Al2O3 Catalyst Jong-Min Lee,aYoung Kyu Hwang,a* Ajit Singh Mamman,a Sang Min Lee,b Do-Young Hong,a Kook-Young Ahnb and Jong-San Chang a* a

Research Center for Nanocatalysts, Korea Research Institute of Chemical

Technology (KRICT),PO Box 107, Daejeon, 305-600, Korea E-mail: [email protected] and [email protected] b

Emission Control Group, KIMM, Daejeon, 305-343, Korea [email protected] and [email protected]

Keywords: methane combustion, decomposition of PdO, light-off temperature, lanthanan-modified Al2O3 washcoat reformation

Abstract. Influence of steam, loading amounts of Pd (viz., 0.7, 2.1, 5.4 wt %) and other process parameters (viz. GHSV and temperature) on methane-based MCFC anodeeffluents over Pd-based La2O3-Al2O3 supported on honeycomb substrate have been investigated. TGA study reveals precise temperature of decomposition of Pd-O in catalytic combustion at high temperature. An increase in methane conversion at elevated temperature is attributable to decomposition of PdO to Pd active phase. The XRD analysis of spent catalysts suggests that the reflections corresponding to PdO gradually become weaker with increase in the Pd loading. The addition of 4% water in the feed decreases the methane conversion due to the reversible reaction of PdO with water adduct resulting in the formation of inactive Pd(OH)2 phase. However, the methane conversion remains unaffected and increases monotonously with 100% water concentration in the feed. Introduction Post-catalytic combustion of effluents from methane/natural gas-based MCFC or gas turbines at elevated temperature is an environmental-and material-driven approach performed with an endeavor to mitigate concomitantly the harmful emissions of nitrogen oxides. Several thermo-stable catalysts such as hexaaluminates and its substituted form [1], perovskites [2] have, recently been reported for such purpose. Ciuparu et al. [3] have recently reviewed the literature related to catalytic combustion of methane over Pd-based catalysts. Prasad et al. [4] have discussed the requisite attributes of combustion catalysts. Several thermo-stable support materials [5] have been proposed for such applications. Support such as monolithic honeycomb provides

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necessary mechanical strength, its shape and size results in the pressure drop across the reactor to the minimum. Pre-washcoated honeycomb monolith is commonly used for such purposes. Trimm has addressed various aspects of the substrate material selection and overall stability of washcoat of combustion catalysts [6]. In the present investigation, the catalytic combustion of feed identical to the typical composition of effluents from MCFC device over La-modified alumina wash coat, supported on monolith honeycomb is presented. The influence of steam, wt % Pd loading and that of various process parameters on the overall combustion activity of catalyst has been studied. Further, TGA/DTA study has been performed to know precisely the decomposition temperature of Pd-O in catalytic combustion at high temperature. The XRD analysis has been carried out to ascertain the state of Pd on cooling for catalysts with different Pd wt % loading Experimental The dip coating of lanthania-modified alumina washcoat and that of Pd component is performed as described below. The washcoat-slurry containing 10 wt% alumina sol / H2O (Alintech) and pre-calculated amount of La(NO3)2 (Aldrich, 99.9) was subjected to aging for 1 h at pH = 2. A segment of cordierite honeycomb (1.1 x 1.1 x 5 cm, 200 cpsi) was dip-coated in the slurry several times to obtain 15 wt % Al2O3 loading. The wash-coated monolith was dried after each dip coating. Prior to dipcoating step, substrate was pretreated with 1M HNO3 at 70oC, washed thoroughly with distilled water and dried in vacuum oven at 80oC for 4h. The precursors of La-Al2O3 wash-coated substrate was decomposed at 550oC for 5 h in static air. Finally, 0.5M PdCl2 / EtOH solution (after aging for 1h at pH=2) was used to deposit 0.7-5.4 wt% Pd on the substrate. Subsequently, the catalysts were calcined at 800 and 1000oC in static air for 5h by increasing the furnace temperature at the rate of 3oC/min. XRD measurements were performed on a Rigaku D/Max-RC X-ray diffractometer. TGA study on 2.1 wt% Pd/La-Al2O3/ honeycomb was carried out using SDT Q600, TA Instruments. The catalytic activity was performed by passing a preheated feed (viz. methane, air CO, H2 and CO2) through pre-mixer over a segment of Pd/Al-La2O3/Cordierite honeycomb (2 x 2 x 5 cc) catalyst mounted in catalyst holder of steel reactor as a function temperatures (200-800oC), GHSVs (36000 and 72000 h-1) and λ (ratio of Air(O2) : CO, H2, CH4: λ). The products of combustion reaction were analyzed with GC equipped with TCD.

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Results and Discussion Catalytic activity data for combustion of feed of identical composition to that of effluents from MCFC over cordierite supported 0.7, 2.1 and 5.4 wt% Pd/La-Al2O3 catalysts under typical reaction are presented in Fig.1

CH4 conversion(%)

100 80

(d)

60

(c) (a)

40

(b)

20 0 0

200

400

600

Temperature ( C)

800

o

Figure 1. CH4 conversion over

cordierite supported (a) 0.7wt% Pd/La-Al2O3, (b)

2.1wt% Pd/La-Al2O3, (c) 5.4wt% Pd/La-Al2O3, (GHSV: 72000, LV : 100 cm/s, λ :3) and (d) 2.1wt% Pd/La-Al2O3 (GHSV: 36000, LV : 50 cm/s, λ:3) It may noted that light-off temperature (defined as temperature required to attain 10 % methane conversion) is reached at 250oC for catalysts with 2.1 and 5.4 wt% Pd loadings where as same is attained at 500oC for a catalyst with 0.7 wt% Pd at space velocity of 72,000 h-1. Methane conversion steadily increases up to 600oC for all catalysts after which it passes through minimum and increases further beyond 730oC. Increase in methane conversion with temperature supports the contention that PdO which is active at lower temperature undergoes decomposition at higher temperature. The drop in activity can be pronounced [7] depending upon the experimental conditions used for catalytic evaluation. DTA study performed on spent catalyst (Fig.2). The appearance of endothermic peak at 820oC which indicates the decomposition of PdO to Pd. The reoxidation of Pd on cooling is indicated by another endothermic at 800oC Forrauto et al.[8] reported decomposition of freshly prepared PdO/Al2O3 in between 800~850oC. The thermal stability of PdO under oxygen atmosphere depends upon the nature of support.

Nanocomposites and Nanoporous Materials VIII

0.0

o

Temperature Difference ( C)

4

-0.5 o

o 800 C 820 C

Cooling -1.0 Heating -1.5

0

200

400

600

Temperature( C)

800

o

Figure 2. DTA study: Decomposition and reformation of PdO on 2.1%Pd/La-Al2O3 Catalyst (spent) XRD analyses of cordierite support and those of spent catalysts with different Pd wt % loadings are presented in Fig.3. It reveals that the reflections for PdO (101) phase becomes strong contrary to those for Pd (111), (200) phases with increase in Pd wt % loading.

Intensity(a.u.)

Pd(111)

(d)

Pd(200)

PdO(101)

(c)

(b) (a) 10

20

30

40

50

60

70

80

2theta/degree Figure 3. XRD pattern of (a) cordierite and (b) 0.7wt% Pd/La-Al2O3, (c) 2.1wt% Pd /La-Al2O3, (d) 5.4wt% Pd/La-Al2O3 The degree of dispersion decreases with increase of wt % loading of Pd, lower the dispersion, the lower the Pd reoxidation extent in the spent catalysts after reaction on cooling. The influence of water addition in the feed on the overall conversion is

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complex event. The effect of 10 and 100 % water concentrations on the combustion activity over 0.7 wt% PdO/La-Al2O3 as a function of reaction temperature is depicted in Fig. 4.

CH4 conversion(%)

100

(b) 80

60

40

(a)

20

0

0

200

400

600

Temperature( C)

800

o

Figure 4. Methane conversion of (a) 100% feeding, (b) 10% feeding The combustion activity decreases in the temperature range between 600 and 760oC beyond which it builds up steadily. It is proposed that the inhibiting effect of water is due to the occurrence of reversible reaction of PdO with water adduct leading to the formation of inactive phase of Pd(OH)2. On this basis it may be postulated that the loss of water from Pd(OH)2 could be rate determining step rather than methane activation However, the inhibiting effect of addition of 100% water as an adduct in the feed is not observed. Conclusions From the above investigation, following conclusions can be drawn:
Pd/La-Al2O3 is highly active and thermo-stable catalyst for combustion of effluents from methane-based MCFC device.
PdO undergoes decomposition to Pd at a temperature higher than 820oC
For high temperature combustion applications, Pd and/or Pd/PdO is found to be catalytically active phase.
The PdO phase in spent catalyst decreases with increase in Pd wt% loadings.
Low concentration of water as an adduct in the feed has a deactivating effect on methane conversion in the temperature range between 500oC and 700oC or higher due to its reversible reaction occurring with PdO thus leading to the formation of Pd(OH)2, a catalytically inactive phase.

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Acknowledgment This work was supported by the Korea Ministry of Commerce, Industry and Energy (MOCIE) through the Korea Energy Management Corporation and Electric Power Industry Technology Evaluation & Planning (TS073-03). The authors thank Mr. Y.-K. Seo for his experimental assistance. A. S. Mamman gratefully acknowledges the financial support by KOFST. References [1] M. Machida, K. Eguchi and H. Arai, J. Catal. 120 (1989) 377. [2] S. Cimino, L. Lisi, R. Pirone, G. Russo and M. Turco, Catal, Today, 59 (2000) 19. [3] D. Ciuparu, M. R. Lyubovsky, E. Altman, L. D. Pfefferele and A Datye, Catal. Rev. 44 (2002) 593. [4] R. Prasad, L. A. Kennedy and E. Ruckenstein, Catal. Rev. Sci. Eng., 26 (1984) 1 [5] D. L. Trimm, Appl. Catal., 7 (1983) 249. [6] D. L. Trimm, Catal. Today, 26 (1995) 231 [7] R. J. Farrauto, M. C. Hobson, T. Kennelly, E. M. Waterman, Appl. Catal., A 81 (1992) 227. [8] R. J. Farrauto, J. K. Lampert, M. C. Hobson, E. M. Waterman, Appl. Catal. B 6 (1995) 263.

Solid State Phenomena Vol. 135 (2008) pp 7-10 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.7

Low-temperature fabrication of polycrystalline yttrium aluminum garnet powder via a mechanochemical solid reaction of nanocrystalline yttria with transition alumina Hyun-Gi Jung1,a, Young-Hun Cheong2,b, In-Dong Han3,c, So-Jin Kim4,d, and Sung-Goon Kang5,e 1,4,5

Division of Advanced Materials Science and Engineering, Hanyang University, Seoul 133-791, South Korea 2

Korea Testing Laboratory, Seoul 152-848, South Korea

3

School of Materials Science and Engineering, Hongik University, Chungnam 330-701, South Korea a

b

c

d

[email protected], [email protected], [email protected], [email protected], e [email protected]

Keywords: Yttrium aluminum garnet, Grinding, Phase transformation, Synthesis

Abstract. Mechanical processing of nanocrystalline Y2O3 and transition alumina (AlOOH) was performed, using a Spex mixer mill under atmospheric conditions, to synthesize yttrium aluminum garnet (YAG, Y3Al5O12) powder at lower temperature. The reaction of nanocrystalline Y2O3 with AlOOH was activated by mechanical energy and use of fine instead of heat energy, leading to the direct formation of pure YAG without second phases at a lower temperature, 800°C, which is significantly lower than that required by the conventional solid-state reaction process. Introduction Yttrium aluminum garnet (YAG, Y3Al5O12) has received much attention because of its superior creep resistance, excellent thermal stability and unique optical properties. Single crystal YAG is widely used as solid-state laser host materials. Polycrystalline YAG has many potential applications as thermal barrier coatings and high-temperature structure materials [1,2,3]. However, fabrication of polycrystalline YAG by a solid-state reaction requires high sintering temperatures (>1600°C), owing to low reactivity between Y2O3 and Al2O3. Heat treatments of mixed oxide powders at temperatures below 1600°C yield intermediate phases such as YAP (yttrium aluminum perovskite, YAlO3), YAM (yttrium aluminum monoclinic,Y4Al2O9) [3,4]. The use of fine, reactive YAG powder with improved sinterability is essential to obtain YAG ceramics with a full density and transparency at low temperatures [1]. In order to synthesize such powders, various chemical methods have been employed including co-precipitation [5], sol-gel process [6], spray pyrolysis [7], and hydrothermal synthesis [8]. Recently, a mechanochemical synthesis, that produces reactive, fine YAG powders by grinding of the constituent oxides, has been used as an alternative to chemical methods [9,10]. Literature concerning the synthesis of YAG powder shows that mechanochemical reaction depends on starting materials. Zang and Saito [9], who used Y2O3, α-Al2O3, Al(OH)3, and AlOOH powders, reported no mechanochemical reaction between Y2O3 and α-Al2O3 or Al(OH)3. Kong and Huang [10] claimed pure YAG powder from Y2O3 and α-Al2O3 by grinding and calcining at 1000°C and attributed the formation of YAG phase to reduced particle size resulting from grinding. In this study, the reactivity of nanocrystalline Y2O3 with AlOOH during grinding and calcination was investigated to obtain further information on the synthesis of YAG ceramics. Experimental Procedure The yttrium and aluminum sources for YAG synthesis were nanocrystalline Y2O3 and AlOOH, respectively. Nanocrystalline Y2O3 was prepared by precipitation of 99.9% Y(NO3)3⋅6H2O (Aldrich)

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[11] and AlOOH was heating Al(OH)3 (Aldrich) powder at 400°C for 2 h [9,12]. Y2O3 and AlOOH powder were mixed using a mortar and acetone at a 5:3 molar ratio of Al to Y. Then, the sample was dried at 80°C for 1 h. The milling process was carried out using a Spex 8000-D mixer mill (USA) consisting of a zirconia vial (50cc) and balls (10 mm in diameter, 3 EA) in air at room temperature up to 24 h. The samples ground for 12 h were calcined at temperatures between 600°C and 1300°C for 2 h. Phases present in the powders ground and calcined were investigated with X-ray powder diffractometry (CuKα radiation, 40 kV, 100 mA, Rigaku). The specific surface area (SSA) of the ground powder was measured using a nitrogen gas adsorption instrument (Micrometrics). The particle size and morphology of the ground powder were examined with a field emission scanning electron microscope (Hitachi, S-4700). Results and Discussion Fig. 1 shows the X-ray diffraction patterns of the samples ground for different intervals of time. As the grinding time increases, Bragg peaks of Y2O3 crystals gradually disappear and a nearly amorphous phase comes to be observed in 12 h. The stablility of the crystal structure of the materials against grinding depend on their inherent nature as well as grinding conditions. Some materials, called soft materials, tend to transform easily into amorphous states with grinding; others, called hard materials, tend to remain in the crystalline state [9]. In this case, nanocrystalline Y2O3 and AlOOH is considered as a soft material. Therefore, it is easy to transform into an amorphous state. Unlike previous work using commercially available Y2O3 and AlOOH [9], which yield YAP and YAG after grinding for 12 h, an intermediate or YAG phase was not observed in this work.

Fig. 1. XRD patterns as a function of grinding time for the mixture of Y2O3 and AlOOH. Fig. 2 shows SEM micrographs of the samples ground for 24 h. As grinding time increased, the sample became agglomerated and consisted of primary particles with sizes of ~50 nm (Fig. 2(b)). The agglomeration is generally observed in dry grinding process, where particles are broken in a fine scale and held together by moisture present or van der Waals attraction. Fig. 3 shows the thermal analysis of the as-milled powder. It was performed with thermogravimetry and differential thermal analysis (TG-DTA 2000, MAC Science, Japan) at a heating rate of 5°C/min up to 1200°C in air. DTA analysis indicates a exothermic peak due to crystallization at 864°C for sample, which corresponds to the formation of YAG (Y3Al5O12), as confirmed by XRD. On the other hand, the weight gradually decreased as the temperature was raised. This would be mainly due to the desorption of the hydroxyl group or moisture originated by raw materials and mechanical processing. The large step in TG trace between 550°C and 850°C would be probably due to crystallization of

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YAG. The reaction was accelerated by mechanochemical processing, to direct crystallization of YAG without YAP and YAM, leaving the OH base in the AlOOH as free water, as follows [9]. 3Y2O3 +10AlOOH → 2Y3Al5O12 +5H2O

(1)

The present of the hydroxyl in the sample and a mchanochemical processing cause the formation of aggregated particles with smooth surfaces, resulting in the low surface area value of the ground sample. The starting sample exhibits a high surface area (98.2 m2/g), because of the fine of Y2O3 and AlOOH, as shown in Fig. 2(a). However, the surface area of the sample decreased rapidly as grinding time increases, reaching a value of 3.4 m2/g for a grinding time of 24 h. These results are consistent with the observations of Zhang et al [9].

Fig. 2. SEM micrographs of a Y2O3 and AlOOH mixture ground for (a) 0 h, (b) 12 h, and (c) 24 h.

Fig. 3. DTA/TG traces of a mechanochemically synthesized powder with a Y:Al ratio of 3:5. It is known that in a solid-state reaction between Y2O3 and Al2O3, Al2O3-rich YAM phase first forms and then transforms into YAP, which continues to react with Al2O3, producing YAG. The reaction can also be accelerated if the particle size of the starting powders is reduced [1]. The same phase formation sequence was expected in this work, but results did not correspond. Phases present in the samples ground for 12 h and calcined at various temperatures for 2 h are shown in Fig. 4. YAG phase in ground sample appears directly with calcinations at 600°C without the formation of any intermediate phases, and than, increasing temperatures to 800°C, a well-crystallized phase of pure YAG was obtained. This means that the sample containing fine Y2O3 and AlOOH has better reactivity for YAG formation as compared to the other yttrium and aluminum sources, because it has the largest surface area and the smallest primary particle size. Also, it is noted that the sample increased internal energy by mechanochemical reaction showed more reactivity than conventional solid state reaction at lower temperatures. The sample ground for 24 h and heated at 800°C for 2 h, despite of long period of grinding times, exhibits YAP because of severe agglomerations, as shown in Fig. 2(c). It is concluded from the results of this work that the type and grinding characteristic of starting material and grinding time are important factors in the mechanochemical synthesis of YAG.

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Fig. 4. XRD patterns of the sample ground for 12 h and calcined at various temperatures for 2 h (● represents YAG). Conclusions YAG powders were synthesized via mechanochemical reaction and calcination using nanocrystalline Y2O3 and AlOOH. Grinding for 12 h caused the reaction to form amorphous phases, which increasing internal energy, corresponds to occur direct formation of YAG at low temperature (800°C) without second phases. However, the sample ground for 24 h, which had the severe agglomeration after grinding, resulted in YAP along with YAG at same heating condition. The sample ground for 12 h and calcined at 1300°C exhibits very small particles with ~200nm size of YAG particles. The type and grinding characteristic of the starting material and grinding time are believed to be important factors in the mechanochemical synthesis of YAG. Acknowledgement This work was supported by the research fund of Hanyang University. (HY-2006-I) References [1] K. M. Kinsman, and J. McKittrick, J. Am. Ceram. Soc., 77 (1994) 2866-72 [2] H. Yagi, J. F. Bisson, K. Ueda, and T. Yanagitani, J. Lumines., 121 (2006) 88-94 [3] N. P. Padture, and P. G. Klemens, J. Am. Ceram. Soc., 80 (1997) 1018-20 [4] A. Ikesue, and I. Furusato, J. Am. Ceram. Soc., 78 (1995) 225-28 [5] J. G. Li, T. Ikegami, J. H. Lee, T. Mori, and Y. Yajima, J. Eur. Ceram. Soc., 20 (2000) 2395-2405 [6] H. M. Wang, M. C. Simmonds, and J. M. Rodenburg, Mater. Chem. Phys., 77 (2002) 802-807 [7] M. Nyman, J. Caruso, and M. J. Hampden-Smith, J. Am. Ceram. Soc., 80 (1997) 1231-38 [8] Y. Hakuta, T. Haganuma, K. Sue, T. Adschiri, and K. Arai., Mater. Res. Bull., 38 (2003) 1257-65 [9] Q. Zhang, and F. Saito, Powder Technol., 129 (2003) 86-91 [10] L. B. Kong, J. Ma, and H. Huang, Mater. Lett., 56 (2002) 344-348 [11] L. Wen, X. Sun, Z. Xiu, S. Chen, and C. T. Tsai., J. Eur. Ceram. Soc., 24 (2004) 2681-2688 [12] P. Alphonse, and M. Courty, Thermochim. Acta, 425 (2005) 75-89

Solid State Phenomena Vol. 135 (2008) pp 11-14 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.11

Formation of Hollow Zinc Oxide by Oxidation and Subsequent Thermal Treatment Jung-Goo Lee1,a, Ryusuke Nakamura2,b, Daisuke Tokozakura2,c, Hideo Nakajima2,d, Hirotaro Mori3,e, and Jong-Hoon Lee1,f 1

Advanced Materials Research Division, Korea Institute of Machinery and Materials, 66 Sangnam-dong, Changwon, Kyungnam 641-010, Korea 2 The Institute of Scientific and Industrial Research, Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan 3 Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, 7-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan a [email protected], [email protected], [email protected], d [email protected], [email protected], [email protected] Keywords: Hollow, Zinc oxide, in-situ TEM, Oxidation, Evaporation,

Abstract. The formation of hollow zinc oxide has been studied by oxidation and subsequent thermal treatment of nanometer-sized zinc particles using in-situ TEM. The zinc particles produced under UHV condition were exposed to air at room temperature for 0.6 ks, which resulted in the formation of oxide layer with thickness of 3 nm. Subsequent heating inside UHV chamber of TEM induced the evaporation of the inner zinc, which resulted in the formation of hollow zinc oxide. The produced hollow zinc oxide had the wurtzite structure. Based upon the vapor pressure of the inner zinc, it seems reasonable to consider that the internal zinc vapor leaks away through the interface between the oxide layer and the amorphous carbon film used as a supporting substrate. Introduction Fabrication of nanometer-sized particles with specific size, composition, and morphology is one of the most important current issues in the field of nanotechnology. This is because one can create new advanced materials with properties which do not appear in bulk materials by controlling these parameters. Especially the shape control of nanometer-sized particles enables us to add new functions them with specific size and composition. Recently, nanometer-sized hollow particles of metal oxides and sulfides were successfully synthesized through a mechanism analogous to the Kirkendall Effect associated with the different mobilities of atoms moving in and out of the particles [1], which stimulated many researchers to engage in fabricating nanometer-sized hollow particles. These efforts made it possible to synthesize nanometer-sized hollow particles of not only pure elements [2-5] but also alloy materials [6]. However, few mechanisms have been proposed to explain hollow structure formation in nanometersized particles [7] although they are of significance to design new types of nanometer-sized hollow particles. Especially, with regard to the design of these nanometer-sized hollow particles, it is also important to control both the size of particles and the thickness of outer shell (or the size of inner void) because they can significantly modify the properties of the particles [8]. The combination of the routine methods for synthesis of nanometer-sized particles and the phenomenon of evaporation can give a chance for us to perform much more delicate shape control of nanometer-sized hollow particles. In the present work, the formation of hollow zinc oxide by the phenomenon of oxidation and subsequent thermal treatment of nanometer-sized zinc particles was studied by in-situ transmission electron microscopy (TEM). Experimental procedures To minimize the effect from artifacts such as surface contamination during the preparation and also to investigate the initial oxidation feature of nanometer-sized zinc particles, the particles were

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prepared in the TEM chamber without exposing the surface of the particles to any undesired atmosphere using specially designed evaporator which was installed in the specimen chamber of Hitachi H-800 type TEM. The details can be seen in previous paper [9]. Using this evaporator, zinc (99.99%, Nilaco Inc) was evaporated from tungsten filament onto the supporting film (amorphous carbon film) kept at ambient temperature and nanometer-sized zinc particles were produced on the film. Before zinc deposition, the amorphous carbon film was baked at 773 K for 0.6 ks to make the surface clean. The particles were then exposed to atmosphere at room temperature (RT) to produce zinc oxide shell, which resulted in core (zinc) – shell (zinc oxide) particles. The particles were again transferred into TEM. Finally the particles were subjected to in-situ annealing experiments in the microscope. Namely, they were heated quite slowly up to 523 K and any evolution in shape and structure was studied by both bright-field images (BFIs) and selected area electron diffraction (SAED) patterns. The TEM was equipped with a turbo-molecular pumping system to achieve a base pressure of around ~2 × 10-5 Pa in the specimen chamber. To study details of produced hollow zinc oxide, high resolution electron microscopy (HREM) observation was also carried out in a Hitachi HF-2000 type HREM. Through the experiments, the operating voltage of the microscopes was 200 kV. The chemical composition of the particles was examined with energy dispersive X-ray spectroscopy (EDS) which was attached to the latter microscopy. Results and discussion At first, in an attempt to confirm whether the oxidation of as-produced zinc particles takes place or not inside TEM, after the preparation of zinc particles the substrate was heated up to 523 K inside TEM. Figure 1 shows a typical formation and sublimation of zinc particles inside TEM. Figure 1 (a) and (a’) show a BFI of as-produced zinc particles on an amorphous carbon film and the corresponding SAED, respectively. The diameter of the zinc particles is in the range of 10-20 nm. The corresponding Debye-Scherrer rings can consistently be indexed as those of bulk zinc with a hexagonal structure. The outward appearance of the zinc particles well reflects the typical crystal habit of the hexagonal structure of zinc although the directions of crystal axis are random. Figure 1 (b) and (b’) show a BFI of the same substrate after heated up to 523 K inside TEM and the corresponding SAED, respectively. All zinc particles disappeared from the substrate by heating as shown in Fig. 1 (b), which was caused by zinc sublimation. It should be noted that there is no trace of the formation of zinc oxide in Fig. 1 (b). This fact indicates that as-produced zinc particles did not undergo oxidation during in-situ heating inside the specimen chamber of TEM whose base pressure was around ~2 × 10-5 Pa. For the next step, the as-produced zinc particles were exposed to air at RT for 0.6 ks to bring about oxidation of the surface of the zinc particles. The particles were then reintroduced into the microscope and slowly heated up to 523 K. Figure 2 shows a typical change in the particles exposed to air after heating inside TEM. Figure 2 (a) and (a’) show a BFI of particles after exposed to air at RT for 0.6 ks and the corresponding SAED, respectively. It is

Figure 1 A series of electron micrographs showing as-produced zinc particles and subsequent evaporation of zinc. (a) as-produced zinc particles on an amorphous carbon film kept at ambient temperature, (b) the same film after heating up to 523 K. (a’) and (b’) the corresponding SAED of (a) and (b), respectively.

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Figure 2 A series of electron micrographs showing morphological change in core (zinc) – shell (zinc oxide) particles by heating inside TEM. (a) zinc particles after exposed to air at RT for 0.6 ks, (b) the same area as (a) after heating up to 523 K inside TEM. (a’) and (b’) SAEDs corresponding to (a) and (b), respectively.

certain that exposing the particles to air resulted in the formation of thin oxide layer. The average thickness of the oxide layer is around 3 nm which can be clearly seen in the inset of Fig. 2 (a). The Debye-Scherrer rings in the SAED (Fig. 2 (a’)) can be consistently indexed as those of pure zinc superimposed with those of zinc oxide (ZnO) having the wurtzite structure which is the thermodynamically stable structure of ZnO at RT [10]. This fact indicates that exposing the particles to air at RT resulted in the formation of 3-nm oxide layer of ZnO surrounding the inner zinc. This thickness of oxide layer is well consistent with our previous results and it grows with temperature [11]. Figure 2 (b) and (b’) show a BFI of the same area as shown in Fig. 2 (a) after heating up to 523 K inside TEM and the corresponding SAED, respectively. The Debye-Scherrer rings in the SAED (Fig. 2 (b’)) can be consistently indexed as only those of ZnO having the wurtzite structure. In addition, it is clear that nano-scale holes were formed at the center of the particles from Fig. 2 (b). The formation of the hollow ZnO particles as shown in Fig. 2 can be attributed to the sublimation of internal zinc due to its high equilibrium vapor pressure which, for example, reaches 1.4 × 10-2 Pa at 523 K [12]. This value seems too low to break the zinc oxide shell but high enough to break the interface between the oxide shell and the amorphous carbon film because the adhesion of ZnO on an amorphous carbon film is weak [13]. Based upon these facts, it seems reasonable to consider that the internal zinc vapor leaks away through the interface between the oxide shell and the amorphous carbon film. In order to identify the atomistic structure and the composition of the hollow ZnO shown in Fig. 2 (b), HREM and EDS analysis was carried out. Figure 3 (a) shows the HREM image of a typical hollow ZnO whose (0001) axis is parallel to the direction of incident electron beam. It should be noted that the lattice fringes corresponding {1010} planes of wurtzite zinc oxide (ZnO) can be seen at not only the outer dark contrast region but also the inner light contrast region. This fact indicates that the particles are not donut-type but faceted dome-type. In Fig. 3 (a), it is also certain that the hollow zinc oxide is a single crystal. The EDS analysis verifies that the particle is composed of zinc and oxygen as shown in Fig. 3 (b). Conclusion In conclusion, we studied the formation of hollow ZnO related to the phenomenon of oxidation and subsequent evaporation using in-situ TEM. Exposing nanometer-sized zinc particles produced under UHV conditions to air in 0.6 ks resulted in the formation of oxide layer with thickness of 3 nm. Subsequent heating inside UHV chamber of TEM up to 523 K induced the evaporation of the inner zinc, which resulted in the formation of hollow ZnO. The produced hollow ZnO had the

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Figure 3 (a) high-magnified electron micrograph of the hollow ZnO particle, (b) EDS spectrum from the particle in (a).

wurtzite structure. The present observation shows a possibility of controlling the morphology such as the size of inner voids and the thickness of oxide layer in hollow oxides by both the oxidation condition and the subsequent evaporation. In addition, the present observation provides a unique method for measuring the adhesive energy between oxides and solid substrates such as an amorphous carbon film. Acknowledgements This work was supported by the Grant-in-Aid for Scientific Research (category S) and 21st century COE program (Towards Creating New Industries Based on Inter-Nanoscience) from the Ministry of Education, Culture, Sports, Science and Technology of Japan.

References [1] Y. Yin, R. M. Rioux, C. K. Erdonmez, S. Hughes, G. A. Somorjai, and A. P. Alivisatos, Science 304 (2004), 711. [2] Q. Liu, H. Liu, M. Han, J. Zhu, Y. Liang, Z. Xu, and Y. Song, Adv. Mater. 17 (2005), 1995. [3] M. Sastry, A. Swami, S. Mandal, and PR. Selvakanman, J. Mater. Chem. 15 (2005), 3161. [4] H. F. Shao, X. F. Qian, and Z. K. Zhu, J. Solid State Chem. 178 (2005), 3522. [5] S. W. Kim, M. Kim, W. Y. Lee, and T. Hyeon, J. Am. Chem. Soc. 124 (2002), 7642. [6] J. Yang, J. Y. Lee, and H. P. Too, J. Phys. Chem. B 109 (2005), 19208. [7] C. M. Wang, D. R. Baer, L. E. Thomas, J. E. Amonette, J. Antony, Y. Qiang, and G. Duscher, J. Appl. Phys. 98 (2005), 094308. [8] D. Goll, A. E. Berkowitz, and H. N. Bertram, Phys. Rev. B 70 (2004), 184432. [9] J. -G. Lee, H. Mori, and H. Yasuda, J. Mater. Res. 20 (2005), 1708. [10] T. B. Massalski et al., in: Binary Alloy Phase Diagrams, American Society for Metals, Metals Park, OH, (1986). [11] R. Nakamura, J. -G. Lee, D. Tokozakura, H. Mori, and H. Nakajima, Mater. Lett. 61 (2007), 1060. [12] D. R. Lide, in: Handbook of Chemistry and Physics, 84th edition, CRC Press, (2004). [13] J. -G. Lee, R. Nakamura, H. Nakajima, and H. Mori, in preparation.

Solid State Phenomena Vol. 135 (2008) pp 15-18 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.15

Antifungal effectiveness of nanosilver colloid against rose powdery mildew in greenhouses Haesic Kim1,a, Hyunsuk Kang 1,b, Gyojin Chu1,c, and Hongsik Byun2,d 1

Cluster Instruments Corporation, #777-97, Geomdan-dong Buk-gu, Daegu, Korea, 702-800 2

Keimyung University, #1000, Sindang-dong, Dalseo-Gu, daegu, Korea, 704-701 d

[email protected]

Keywords: nanosilver, antifungal, powdery mildew, rose

Abstract The antifungal effectiveness against rose powdery mildew using antimicrobial nanosilver colloidal solution was investigated. Double-capsulized nanosilver was prepared by chemical reaction of silver ion with aid of physical method, reducing agent and stabilizers. The average diameter of nanosilver was about 1.5 nm. They were highly stable and very well dispersive in aqueous solution. The Transmission electron microscopy and UV-vis spectrometer were used for measurements of size analysis and their stability, respectively. The nanosilver colloidal solution of concentration of 5000 ppm was diluted in 10 ppm of 500 kg and sprayed at large area of 3306 m2polluted by rose powdery mildew. The white rose powdery mildew fade out above 95 % after 2 days and was not recurred for a week. The antifungal effects were observed by an optical microscope and photographs. 1.

Introduction Powdery mildew, caused by Sphaerotheca pannosa var. rosae, is one of the most

common and widespread fungal diseases of greenhouse and outdoor roses [1, 2]. It mainly appear first on the under surface of young leaves in early summer and the infection spreads to stems, shoots and buds [1]. It often causes leaf distortion, color turning, curling and premature defoliation of leaves. The disease reduces flower production and causes weakening of the plants [3]. Silver have long been known to have strong inhibitory and bactericidal effects as well as a broad spectrum of antimicrobal activities [4]. Silver nanoparticles, which have the high specific surface area and high fraction of surface atoms, will have high antimicrobial activity compared to bulk silver metal. Nanosilver colloid that is a well-dispersed and stabilized silver nanoparticle solution will be more adhesive on bacteria and fungus and so have enhanced antibacterial activity. The purpose of this study was to examine the effectiveness of nanosilver colloid as new fungicide against rose powdery mildew in greenhouses.

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2.

Experimental The nanosilver colloidal solution of concentration of 5000 ppm was prepared by

chemical reduction of silver ion with aid of physical method, reducing agent and stabilizers. First, 31.5 g silver nitrate (AgNO3, Junsei, 99.8 %) was dissolved in distilled water of 3.7 l and 40g poly-(N-vinyl-2-pyrrolidone) (PVP, Aldrich) as stabilizer was added. Second, 1 g sodium borohydride (NaBH4, Jensei, 98.0 %) as reducing agent was dissolved in distilled water of 0.2 l and this solution was slowly dropped in silver ion/PVP solution under sonication. After dropping, 28.5 g quaternary ammonium chloride (Cluster Instruments Co., 80 %) as another stabilizer was dissolved and vigorously stirred for 1 h. The particle size of nanosilver and UV-visible spectrum of nanosilver colloidal solutions was characterized by Transmission electron microscopy (TEM) and UV spectrometer, respectively. The antifungal effects of nanosilver solution carried out at a commercial greenhouse (Sung-Ju Farm), located at SungJu (Gyeongsangbuk-Do, Korea), an important area for cut rose production. Rose plants, belonging to the 'Suncity' cultivar, were grown according to the cultural practices normally adopted by local growers. The nanosilver solution of 500 kg with concentration of 10 ppm was sprayed at large area of 3306 m2 polluted by the rose powdery mildew. The antifungal effects were observed by an optical microscope and photographs. 3.

Results and discussion TEM images of nanosilvers revealed that the average size was 1.5 nm with size distribution

of 1-5 nm (Fig. 1). In addition, TEM images showed that the nanosilvers were densely and well dispersed in the colloidal solution. It can be expected the small size and high dispersity nanosilver lead to have high effective antifungal properties at low concentration. The absorption spectra of a very dilute colloidal solution (1 and 5 ppm) were obtained at 410nm which was a typical absorption band of nanosilver (Fig. 2). The absorbances of 1 and 5 ppm aqueous solution at 410 nm were 0.09 and 0.45, respectively. The increment of absorbance based on the concentration and non-shift of band position showed that the nanosilver was very well dispersive and highly stabilized. The antifungal effects of nanosilver against rose powdery mildew were examined in large area of 3306 m2. The white rose powdery mildew faded out above 95 % after 2 days and was not recurred for a week. Fig. 3 showed that the nanosilver solution was very adhesive and harmless on the leaves. The strong interaction of nanosilver solution with leaf surface offers the possibility of coating of nanosilver. Therefore, the nanosilver could pass into the leaf surface and take off the stabilizers by the interaction between the micelle and plant cells or between the micelle and rose powdery mildew. The stripped nanosilver was adhered on the leaf surface and the fungus cell. And then, the nanosilver could produce the reactive oxygen causing death of fungus cell and catalyze the complete destructive oxidation of microorganisms.

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Fig. 1. TEM images of the nanosilver in colloidal solution of 1000 ppm

(a)

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Fig. 2. Absorption spectra of a nanosilver colloidal solution after dilution; (a) only stabilizers, (b) 1 ppm, (c) 5 ppm.

(b)

(c)

Fig. 3 mildew; (a) before treatment, Fig. 1. Photographs of rose with powdery (b) immediately after treatment and (c) 2 days after.

Also, it was generally believed that heavy metals reacted with proteins by combining the thiol (-SH) groups and desulfide bonds (-S-S-) of bacteria, which lead to the inactivation of the proteins [5]. The nanosilver as microorganism catalyst could have the two possibilities that were a production of reactive oxygen and a inactivation of the proteins of fungus by adhesion. However, the mechanism for the antifungal effects of nanosilver was not fully investigated. Fig. 5 showed the growth inhibition effect of nanosilver colloidal solution against Sphaerotheca pannosa var. rosae. Sphaerotheca pannosa var. rosae was completely maintained for a week at 10 ppm of nanosilver concentrations. This result supported that one possible reason for the antifungal activity of our nanosilver colloid solution might be their adsorption on fungus surface. Because the Sphaerotheca pannosa var. rosae of rose was a microparasite, this strong adhesion could provide a high antifungal effectiveness and an extermination of fungus. Results shown at Fig. 4 and 5 supported that the nanosilver had the high effectiveness and durability. In addition, the nanosilver did not have phytoxicity about the plants cell of leaves, stem and bud.

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(a)

(b)

Fig. 4. of leaves with Fig. 4. Photographs powdery mildew; (a) before treatment and (b) after a week.

4.

(a)

(b)

Fig. 5. Fig. 5. Optical microscope images of powdery mildew on leaf of rose; (a) before treatment and (b) after a week (80 magnification).

Conclusion To investigate antifungal activity of nanosilver against Sphaerotheca pannosa var. rosae,

we prepared their colloidal solution doubly stabilized with PVP and quaternary ammonium chloride. The double-capsulized nanosilver showed high dispersity and stability. The photographic results showed that the effects of nanosilver colloidal solution against rose powdery mildew were very high and durable for a week. In addition, the nanosilver did not have phytoxicity about the plants cell of leaves, stem and bud of roses. As a result, we knew that a well dispersive and stabilized nanosilver could be recommended as new fungicide for powdery mildew. Acknowledgements This work was supported by grant No. RTI04-03-02 from the Regional Technology Innovation Program of the Ministry of Commerce, Industry and Energy(MOCIE) References [1] Horst, R.K., 1983. Compendium of Rose Diseases. Americn Phytopathological Society, St Paul, USA. [2] De Vries, D.P., Dubois, L.A.M., 2001. Developments in breeding for horizontal and vertical fungusresistance in roses. In; van Huylenbroeck, J., VanBockstaele, E., Debergh, P. (Dds.), Proceedings of the Eucarpia symposium on new ornamentals. Acta Hort., 552. 103. [3] Agrios, G.N., 1978. Plant pathology, second ed. Academic Press, Inc., New York, USA. [4] M. Uchida, Chem. Ind. 46 (1995) 48, N. Grier, Silver and Its Compounds, Disinfection, Sterilization and Preservation, Lea and Febiger, Philadelphia, 1983, 375. [5] A. Lehninger, D. Nelson, M. Cox, Principles of Biochemistry, seconded., Worth, New York.

Solid State Phenomena Vol. 135 (2008) pp 19-22 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.19

Fabrication of oriented TiO2-based nanotube array thin films Yupeng Guo1, Nam-Hee Lee1, Hyo-Jin Oh1, Cho-Rong Yoon1, Chang-Kyu Rhee2, Kyung-Sub Lee3, and Sun-Jae Kim1* 1

Institute/faculty of Nanotechnologies and Advanced materials Engineering, Sejong University, #98 Gunja-dong, Gwangjin-gu, , Seoul 143-747, Korea 2 Nuclear Nanomaterials Development Laboratories, Korea Atomic Energy Research Institute, P.O. Box 105, Yuseong-gu, Daejeon 305-353, Korea 3 Devision of Materials Science & Engineering, Hanyang University, #17 Haengdang-dong, Seongdong-gu, Seoul 133-791, Korea

*[email protected] Keywords: TiO2, nanotube, arrays, thin films

Abstract In this study, we have successfully developed the technology to grow the nanotube array thin films from dip-coated titania using hydrothermal method. The nanotube array thin film strongly adhered onto the substrate, was formed in short time reaction at 140°C. Even a self-supporting films, consisted of vertically aligned nanotube in large part, were formed after long time reactions at 140°C with 10 µm in thickness. The most probable formation mechanism of TiO2-based nanotube array thin films is discussed. Introduction Titanium dioxide is a particularly versatile material with technological application as a photocatalyst, photovoltaic material, gas sensor, structural ceramic, electrical circuit varistor, biocompatible material for bone implants, and more. Therefore, the ability to control the architecture of titania down to nanoscale dimensions can be expected to positively impact a variety of economically important technologies. As such, there have been significant efforts to develop nanotubular titania, including anodic oxidation [1], template-based approaches [2] and the wet chemical method [3]. The hydrothermal technique is widely employed to prepare titanate nanotubes by treating the TiO2 powder precursors in concentrated NaOH. While the crystalline structure, compositions (TiO2, Na2Ti2O5, H2Ti2O5, A2Ti3O7, etc.), thermal stability and the function of post-treatment is still in controversy. Most of the TiO2 based nanotube prepared by this method is powder type. However, thin films and coatings of oriented nanostructures are often more desirable for applications involving catalysis, filtration, sensing and photovoltaic cells. Tian et al. reported that hydrothermal and seeded growth method to prepare H2Ti3O7 nanotubes and continuous films [4]. However, data is still scarce for further discussing the effect of temperature, reaction time, concentration of NaOH and orientations on the substrate.

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The aim of the present study is to find the effect of hydrothermal conditions. The formation mechanism of titanate nanotubes is also to be discussed in terms of the detailed observation of the products by electron microscopy. Experimental Procedure The process describing the preparation of TiO2 based nanotubes is thoroughly discussed in our previous work [5]. To prepare nanotube arrays and thin films, we first prepared a dilute TiO2 suspension by dispersing 1.0 g of Degussa P25 TiO2 powder in 50 g of deionized (DI) H2O. Titanium (99.7% pure) flake as a substrate was degreased by sonicating in acetone and ethanol, followed by rinsing with DI water. TiO2 nanoparticles were deposited on a titanium flake through dip coating from this TiO2 suspension at room temperature. Upon withdrawing from the sol, the substrates were dried at 60°C for 20 min. TiO2 layers on substrates could be thickened by means of consecutive dip-coating processes. The Ti flake containing the predeposited TiO2 thin films was then reacted with an alkaline solution in an autoclave reactor (internal capacity of 1000 ml), containing 10 M NaOH solution. The reaction temperature ranged was 140 °C. After the reaction, the Ti flake, covered with the newly formed film, were soaked and washed with DI H2O, and then dried in air. Characterization of these thin films was carried out using scanning electron microscopy (SEM, Hitachi S-4700), energy dispersive x-ray spectroscopy (EDX, attached to the SEM) and powder x-ray diffraction (Simens-D50050D). Results and Discussion The XRD patterns of the starting material, Ti flake and some of samples are shown in Figure 1. It can be seen from the figure that the XRD patterns of the TiO2 deposited Ti flake are similar to pure Ti flake. Dip coating methods is a thin-film coating method, so only the prominent appearance of the anatase TiO2 (101) peak around 25o are significantly. The broad peak at 20o and peaks around 21o and 23o (Figs, 1b, 1c) represent some kinds of titanium oxide formed on the flake surface. After the hydrothermal process at 140°C for 12 h, these peaks still exist, and the diffractogram of the nanotube arrays film shows characteristic peaks around 10o, 24o and 28o, which are similar to our previous work using anatase TiO2 sol as a precursor. These peaks have been assigned to the diffraction of titanates such as A2Ti2O5⋅H2O, A2Ti3O7, and lepidocrocite titanates. The anatase TiO2 (101) peak around 25o is disappear, that indicates all the precursors have transformed to nanotubes. In the present work, the preparing condition is at 140 o for 12 h. After 6 h at 140°C, all the TiO2 particles with 20-30 nm (Fig. 2b) became naotubes, continuous thin film, less than 1 µm in thickness (Fig. 2d), formed. Sample reacted for 6 h shows some sheetlike structures, involving organized nanotubes. Reference 6 results revealed the pathways of how this folding can take place. After 12 h of the reaction, the nanotubes grew very long in length, and a continuous film, 10 µm in thickness, formed (Figs 3a, 3b). The film can be removed form the substrate as a free-standing film. The high magnification FESEM image of the samle reveals that most nanotubes near the film base are vertically aligned (Fig. 3c), the top surface of the films is similar to that discussed above (Fig. 2d).

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2θ (deg.)

Figure 1. Comparative XRD patterns of the P25,Ti flake, TiO2 coating Ti flake and titanate nanotube arrays and thin films;(a) P25, (b)Ti flake, (c) Ti with TiO2 deposited, and (d) sample after 12 h at 140°C; (A) anatase, (R) rutile, and (Ti) titanium metal.

Figure 2. FESEM micrographs of (a) Ti flake, (b) Ti with P25 deposited, (c) Top view of the film A after 6 h at 140°C, and (d) A cross-section image of the sample.

Figure 3. FESEM images of continuous nanostrctured film B after 12 h of reaction at 140°C. (a) Low magnification face-on image of the free-standing film. (b) A low magnification cross-section image of the film. (c) A cross-section image of oriented nanotbes near the substrate (d) Top view of the film

Sample C reacted for 3 h at 140°C shows similar oriented nanostructures, but these films contain half way curled nanofoils (arrows) as well as nanotube formation involves folding of sheetlike structures (Fig. 4). The probable formation mechanism of the titanate nanotubes is elucidated. At a very early stage, TiO2 reacts with NaOH solution to form a layered intermediate phase containing Ti, O and Na, which cannot roll up completely for 3h at 140°C, so the edges of layers are bent (Fig.4 arrows). With a longer time, the intermediate layers can be peeled off and curved naturally like wood shavings, forming nanotubes. Energy dipersive X-ray spectroscopy (EDX) studies show that Ti, Na and O

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Nanocomposites and Nanoporous Materials VIII

(except possibly H, which cannot be detected by EDX) are present in the nanotubes films (Fig. 5a). In the sample after treatment with 0.1 M HCl for 24 h, only Ti and O are present (Fig. 5b). By controlling the conditions of acid treatment, the amount of residual Na+ ions changes. It is expected that a new type of titanate nanotube films having new properties will be formed by controlling the amount of residual Na+ ions and by replacing residual Na+ ions with other ions.

Figure 4. FESEM images of continuous nanostrctured film C after 3 h of reaction at 140°C.

Figure 5. EDX spectra obtained from film B washed with (a) only DI water, and (b) HCl for 24 h at 25°C.

Conclusion The nanotube arrays films were successfully obtained directly from titania using dip-coating and hydrothermal method. The X-ray diffraction patterns and the high-resolution SEM results were proposed for the Na-tianate naotubes. Acknowledgments This work was supported by grant No. 2006-N-PV03-P-02-000-2004 from the Basic Research Program of the Korea Science and Engineering Foundation. References 1. C. Ruan, M. Paulose, O. K. Varghese, G. K. Mor, and C. A. Grimes, J. Phys. Chem. B 109 (2005) 15754. 2. P. Hoyer, Langmuir 12 (1996) 1411. 3. T. Kasuga, M. Hiramatsu, A. Hoson, T. Sekino, and K. Niihara, Langmuir 14 (1998) 3160. 4. Z. R. Tian, J. A. Voigt, J. Liu, B. Mckenzie, and H. Xu, J. Am. Chem. Soc.125 (2003) 12384. 5. M. Qamar, C. R. Yoon, H. J. Oh, D.H. Kim, J. H. Jho, K. S. Lee, W. J. Lee, H. G. Lee and S. J. Kim, Nanotechnology 17 (2006) 5922. 6. S. Zhang, Q. Chen, and L. M. Peng, Physical Review B 71 (2005) 014104-1.

Solid State Phenomena Vol. 135 (2008) pp 23-26 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.23

Formation of lanthanum hydroxide and oxide via precipitation So-Jin Kim1, a, Won-Kyu Han1, b, Sung-Goon Kang 1, c, Min-Su Han2, d, Young-Hun Cheong 3, e 1

Division of Materials Science and Engineering, Hanyang University, Seoul 133-791, South Korea

2

Conservation Science Division, National Research Institute of Cultural Heritage, Daejeon 305-380, South Korea 3

Machinery and Materials Center, Korea Testing Laboratory, Seoul 152-718, South Korea a

[email protected], [email protected], [email protected], d

[email protected], [email protected]

Keywords: La(OH)3; La2O3; Precipitation; XRD; TG/DTA.

Abstract. Lanthanum hydroxide and oxide were prepared by the precipitation method in an aqueous medium at room temperature. The precipitate was examined using thermal analysis, X-ray diffraction and Scanning Electron Microscopy to investigate the phase evaluation and the thermal transformation by decomposition. The as-precipitated powder from the precipitation method was hexagonal La(OH)3. The lanthanum hydroxide was decomposed to oxide in two-steps as La(OH)3 → LaOOH + H2O and 2LaOOH → La2O3 + H2O. Introduction Lanthanum oxide and hydroxide are of great research interest because of their prospects as catalytic materials and for optical, electrical and magnetic applications [1-5]. La(OH)3 has used as the catalyzer and sorbent. La2O3 is used to produce automobile exhaust-gas convectors, ceramic superconductors, positive temperature coefficient in resistance (PTCR) BaTiO3 and colossal magneto resistance (CMR) manganites [6]. Moreover, it is also used as a strengthening agent in structural materials [7]. The sensitivity of gas sensitive ceramics could be improved when La2O3 nanophase was added into the ceramics [8]. Recently, it was found that La2O3 had the greatest effect as an additive in improving sintering resistance as well as in lowering the thermal conductivity of yttria stabilized zirconia (YSZ) coatings in thermal barrier coatings (TBCs) [9]. The properties of ceramics in general are greatly affected by the characteristics of the powder, such as particle size, morphology, purity and chemical composition. Wet chemical synthesis of fine ceramic powders has several advantages. Using these methods, such as precipitation, sol-gel and the hydrothermal process have been shown to control the morphology, chemical composition of preparation powder and the increased homogeneity and high surface area of powders [6]. This paper describes the results of precipitates from precipitation method. The thermogravimetric analysis (TGA), differential thermal analysis (DTA) and X-ray diffraction (XRD) were employed to evaluate the transformation from La(OH)3 to La2O3. And BET surface area measurements and Scanning Electron Microscopy (SEM) also were used to observe the particle size and morphology of lanthanum hydroxide and oxide. Experimental Lanthanum nitrate (La(NO3)3 · 6H2O, 99.99%, GFS Chemicals) was used as starting materials and ammonia as precipitant. Lanthanum nitrate was dissolved in distilled water to form 0.04M and then that was stirred for 3h. To this solution, 2M ammonia is added dropwise to precipitate lanthanum until pH ~10 under stirring was reached to ensure the hydroxide. The precipitates were continuously stirred for 3h. Then the gel was filtered and washed with distilled water three times and with ethanol

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Nanocomposites and Nanoporous Materials VIII

two times. The washed precipitates were dried at 80°C, for 24 h. The dried precipitates were calcined at 600°C, 700°C and 800°C for the formation of the La2O3 phase. The thermal decomposition of the precipitates was analyzed with TGA and DTA measurements with the Simultaneous Thermal Analyzer (TA instrument-SDT 2960, England). The powder was heated at each temperature in air at a heating rate of 10°C/min. The XRD data of calcined powders at different temperatures were obtained using Mac Science-MXP18VA (Japan) with Cu Kα radiation. The samples were scanned in the 2θrange of 10-90° for a period of 5s in each step. The morphology of La(OH)3 and La2O3 nanoparticles was examined by SEM (Hitachi S-7400, Japan). Results and discussion The precipitates of lanthanum was identified as hexagonal trihydroxide, La(OH)3. La hydroxide was precipitated according to the following reaction. La(NO3)3·6H2O + 3NH4OH → La(OH)3 ↓+ 3NH4NO3 + 6H2O

(1)

La(OH)3 was decomposed with two steps of weight loss. The TG/DTA results of prepared La(OH)3 by precipitation from room temperature to 1500°C are described in Fig. 1. The mechanism of the decomposition process for lanthanum hydroxides can generally be described by two equations [10-12]. La(OH)3 → LaOOH + H2O

(2)

2LaOOH → La2O3 + H2O

(3)

The TGA curve was decreased at 100°C due to a removal of surface absorptive water of the powder, and the next weight loss started at 350°C, then decreasing step at 600°C. The total weight loss was 22.38%. The first step of weight loss was 4.37%, the second step was 10.08%, and the third step was 7.91%. The corresponding endothermic peaks on the DTA curve appeared at 348 °C (2) and 701 °C (3). The weight loss and endothermic reactions described the decomposition reaction of lanthanum compound. After heating at 350°C, LaOOH formed. At 650°C, La2O3 appeared.

Fig. 1. TG/DTA curves of as-precipitated powder, La(OH)3.

The evolution of the X-ray diffraction pattern is shown in Fig. 2. The XRD patterns of the as-precipitated and annealed powders at 600°C, 700°C and 800°C for 2 h. La(OH)3 was identified as

Solid State Phenomena Vol. 135

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a hexagonal structure. The as-heated La(OH)3 powder developed crystallinity. After the heat treatment of the precipitates at 600°C, the XRD peaks are very similar to those of La2O3. Since lanthanum is a basic element, its oxide and hydroxide can react with CO2 in air to form a carbonate that is not so crystalline [13]. The powder of heat treatment at 800°C showed La2O3 as a hexagonal structure.

Fig. 2. XRD patterns of the as-precipitated La(OH)3 powder and heated powder: (a) as-precipitated, (b) heated powder at 600°C, (c) 700°C, (d) 800°C for 2 h. Fig.3 shows SEM images of the as-precipitated and calcined powders at 800°C for 2 h. The La(OH)3 displayed the uniform morphology of nanorods with 15-20 nm in width and 500-700 nm in length. In case of catalyzer, the enlargement of surface area can improve the performance. The surface area of as-precipitated powder, La(OH)3 was measured to be 102.5 m2/g. Hence, the nanocrystal of La(OH)3 could be beneficial to the catalytic capability. The particle size of calcined powder, La2O3 was 30-40 nm. This shows that the particle size increases with rising calcinations temperature.

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Nanocomposites and Nanoporous Materials VIII

Fig. 3. SEM micrographs of the La(OH)3 (a, b) and La2O3 (c,d) Summary Lanthanum hydroxides and oxides were obtained by homogeneous precipitation. The thermal decomposition of La(OH)3 was investigated by TG/DTA studies and these confirmed the two-step mechanism of La(OH)3 → LaOOH + H2O and LaOOH → La2O3 + H2O. The XRD patterns of La(OH)3 annealed temperatures were shown as phase transformations of La(OH)3 powder. The average diameter size of La(OH)3 was 15-20 nm, and the average particle size of La2O3 at 800°C was 30-40 nm. References [1] S. Sampath, N.k. Kulkarni, M.S. Subramanian, N.C. Jayadevan: Carbon 26 (1988), p.129 [2] K.R. Barnard, K. Foger, T.W. Turney, R.D. Williams: J. Catal. 125 (1990), p. 265 [3] L. Chen, Z. Xu, L. Lin, X. Li, Ranliao Huaxue Xuebao 22 (1994), p.337 [4] S.L. Li, S.X. Zhang: J. Catal. 25 (2004), p. 762 [5] P.N. Babin, A.Kh. Akishev, Z.K. Kairbaeva, N.V. Kirchanova: Kompleksnoe Ispol’zovanie Mineral’nogo Syr’ya (USSR) 2 (1979), p. 55 [6] A. Vadivel Murugan, S.C. Navale, V. Ravi: Mater. Lett. 60 (2006), p.848 [7] J.C. Yang, Z.R. Nie, Y.M. Wang: Appl. Surf. Sci. 215 (2003), p.87 [8] Gao Yuzen, Li Yonghong, Zhang Taisong: J. Rare Earths 14 (1996), p. 206 [9] M. Matsumoto, H. Takayama, D. Yokoe, K. Mukai, H. Matsubara, Y. Kagiya, Y. Sugita: Scripta Mater. 50 (2006), p.2035 [10] E. Ino, K. shimizu, T. yamate: Zairyo 25 (1976), p.1165 [11] M.P. Rosynek, D.T. Magnuson: J. Catal. 46 (1977), p.402 [12] A. Neumann, D. Walter: Tagungsband zur Wohler-Tagung (2004), p.4 [13] M. Ozawa, R. Onoe, H. Kato: J. Alloy. Compd. 408-412 (2006), p. 556

Solid State Phenomena Vol. 135 (2008) pp 27-30 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.27

Colloidal Crystal Templating of Two-Dimensional Ordered Macroporous SiCN Ceramics In-Hyuck Song1,a, Yong-Jin Kim1,b, Hai-Doo Kim1,c, Dong-Pyo Kim2,d 1

2

Center of Powder Processing, Korea Institute of Machinery and Materials, Chang-Won, Kyung-Nam 641-010, South Korea

Department of Fine Chemical Engineering and Chemistry, Chungnam National University, Taejon 305-764, South Korea a

[email protected], [email protected], c

[email protected], [email protected]

Keywords: Porous material, SiCN, Silica template, Preceramic polymer

Abstract. In this study, two-dimensional (2D) ordered macroporous SiCN ceramics were prepared by infiltrating sacrificial colloidal silica (SiO2) templates with the low molecular weight preceramic polymer, polysilazane. This was followed by a thermal curing step, pyrolysis at 1250°C in a N2 atmosphere, and finally the removal of the templates by etching with diluted HF. In particular, 100large-scale monolayer silica crystals were prepared on the cleaned Si substrates by spin coating. Two-dimensional SiCN ordered pore arrays were fabricated by a solution-dipping template strategy. Introduction Colloidal crystals, as templates, have attracted much attention in fabrication of versatile specially structured functional materials [1-7]. In this fabricating process, the target product materials are ‘fixed’ into the voids of a template. After suitable treatment, the target composition with the designed morphologies and structures is finally replicated after removing the template. Based on this procedure, the periodic sphere-inversed structure is attracting much attention, with much having been accomplished, e.g. three-dimensional ordered macroporous (3DOM) materials and twodimensional ordered macroporous (2DOM) materials. Three-dimensional (3D) colloidal crystals are typically used for the preparation of highly periodic porous materials [1-4]. They have many potential applications, including utilization as photonic crystals, bio-sensors, waveguide materials, and optoelectronics. Two-dimensional (2D) ordered macroporous arrays have attracted much attention as a result of their applications in biosensors, catalysts, etching masks, magnetic materials, and phonon diffraction gratings [5-7]. In particular, it is anticipated that non-oxide eventually will be used as catalytic supports that operate under harsh conditions due to their high thermal, mechanical, and chemical stabilities. However, only a few reports have been issued regarding periodic nanostructured non-oxide ceramics, although the polymeric route to non-oxide ceramics via wet fabrication and subsequent pyrolysis offers a unique opportunity to manufacture conventional ceramic components [8-9]. Here, a simple method to prepare long-range ordered macroporous SiCN ceramics using a solution dipping method based on the 2D SiO2 monolayers is presented. To the best of the authors’ knowledge, this is the first report regarding two-dimensional ordered SiCN ceramics directly using the 2D SiO2 monolayers as template. Experimental Procedure In order to make the two-dimensional (2D) colloidal crystals, monodisperse silica spheres (99.9%) were bought from Lancaster. The sphere size was 500 nm in a diameter. A silica sphere suspension (10wt% in alcohol) was used. Ordinary Si substrates were ultrasonically cleaned in

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Nanocomposites and Nanoporous Materials VIII

ethanol for 1 hour. Large-scale monolayer SiO2 colloidal crystals were prepared on the cleaned substrates by spin coating (3000rpm) on a custom-built spin-coater. The monolayer SiO2 colloidal crystals on the Si substrates were put into the furnace and sintered at 900°C for a certain time, then taken out and cooled room temperature. Subsequently, a droplet of SiCN precursor was applied onto the sintered colloidal monolayer with a quantitative pipette; it was able to infiltrate into the interstices among the SiO2 colloidal crystals. Polysilazane (PSZ, Kion Co.) as a preceramic polymer was used to prepare the SiCN ceramics, using a method reported elsewhere [8]. A low viscous precursor, with no dilution, was infiltrated into the silica template. The infiltrated polymer of the sample was cured at 300°C for 12 hours. This was followed by heating to 1250°C at a heating rate of 2oC/min in N2 maintained for 0.5h. The ceramic composite was then dipped into a 5% aqueous hydrofluoric acid (HF) solution for 6h to etch out the silica sphere. Finally, the archived porous product was obtained and dried in an oven at 110°C overnight. Fig. 1 shows a schematic drawing of the fabrication process of the 2D SiCN ordered pore array. The synthesized sample was characterized by field emission scanning electronic microscopy (FE-SEM). Results and Discussion Fig. 2(a) shows a SEM image of a monolayer colloidal crystal with a diameter of 500nm of silica fabricated via a spin-coating method. It is observable that the spherical particles are in contact with each other quasi point style in the monolayer. However, a long-range ordered particle array was not obtained. Fig. 2(b) shows the 2D SiCN pore film by monolayer colloidal template of Fig. 2(a). Each pore exhibits a truncated-hollows sphere shape. The height of the inversed porous SiCN film is approximately 300nm from the cross-sectional FE-SEM image with 45o tilting. As can be seen in the notation of Fig. 2(b), hexagonal and cubic array are mixed in the pore structure. The array of the sphere is important when colloidal crystal templates are used to make photonic crystals in threedimensional ordered material, and less critical in certain other applications such as catalysis support in two-dimensional ordered materials. The morphology of the cubic array is identical to (110) planes in a BCC structure and (100) planes in a FCC structure. Based on theoretical considerations, the equilibrium structure of an arrayed hard sphere is a closed packed hexagonal array [10]. Previous researchers reported that a cubic array was observed under special circumstances from a diluted suspension [10, 11]. However, it is unclear why the cubic array forms. Fig. 3 shows representative FE-SEM images of a 2D sphere-inversed porous SiCN film in a high magnification of Fig. 2(b). Pore sizes were approximately 510 nm in the hexagonal array (Fig. 3(a)) and about 515nm in the cubic array (Fig. 3(b)). The pore size indicates the average size of the truncated-hollows sphere. The pore size increased slightly compared to the 500 nm size of the original sacrificial silica sphere templates, which provides a means of tailoring the pore size of porous materials. Previous researchers [8] reported that the pore size decreased by 10-15% compared to the sizes of the original sacrificial silica sphere in three-dimensional ordered macroporous (3DOM) materials, as the silica spheres and polymer shrink during the pyrolysis stage.

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However, in the monolayer colloidal crystal, shrinkage did not occur because the bonding force between the silica sphere and silicon substrate was more effective than the shrinkage force between the silica spheres during the pyrolysis stage. In contrast, it appeared that when the SiCN precursor was dropped into the silica spheres, the liquid precursor made loose contacts between the silica spheres and increased the average pore size. The surface of the strut wall was not smooth, as shown in Fig. 3. Previous research has reported that meso-sized pores (3-5 nm) occurred [8]. It is believed that these mesopores were formed by etching out the diffused silica in the SiCN framework of the sphere template under hightemperature pyrolysis. In other words, smaller particles with higher surface activity may lead to more significant diffusion which develops pores in the walls of the SiCN frameworks, resulting in a higher surface area.

Fig. 2. SEM images of (a) 2D ordered colloidal monolayer on the Si substrate and (b) 2D SiCN pore film fabricated using a monolayer colloidal template with 500nm silica in the diameter pyrolized at 1250°C.

Fig. 3. FE-SEM images of the 2D SiCN pore film: (a) hexagonal array (b) cubic array Fig. 4 shows a schematic illustration of the 2D SiCN pore film. When colloidal crystals were filled with polymer solutions using capillary forces in this system, small vacancies could be observed at the triangular interstices of the pore walls. These vacancies are believed to be caused by incomplete filling of the template by the precursor solution. In the case of a closed packed array, a triangular interstice is too small to make a secondary pocket; however, a cubic array has enough space to make a secondary pocket in the rectangular interstice. The relationships between the radius (r) of the interstice and the radius (R) of the SiO2 sphere were calculated. A closed packed array

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Nanocomposites and Nanoporous Materials VIII

results in r=0.154R and a cubic array results in r=0.414R. The interstice of a cubic array is larger than that of a closed packed array. As can be seen in Fig. 3(b), we could observe the pocket as a vacancy in the interstice site of the cubic array.

Fig. 4. Schematic illustration of the 2D SiCN pore film: (a) hexagonal packed array; r= 0.154R (b) cubic array; r= 0.414R (R: radius of the SiO2 sphere, r: radius of the interstice) Conclusions In summary, a simple method to prepare ordered macroporous SiCN ceramics using a solution dipping method based on the 2D SiO2 monolayers is presented. Polysilazane as a preceramic polymer was used to make the SiCN ceramics. Each pore exhibits a truncated-hollows sphere shape, and hexagonal and cubic array are mixed in the pore structure. The relationships between the radius (r) of the interstice and the radius (R) of the SiO2 sphere were calculated. We could observe the pocket as a vacancy in the interstice site of the cubic array. References [1] A. Stein and R. C. Schroden, Current Opioon in Solid State and Material Science Vol. 5 (2001), p. 553 [2] O. D. Velev and E. W. Kaler, Adv. Mater. Vol. 12 (7) (2000), p. 531 [3] J. H. Moon, G. R. Yi and S. M. Yang, J. Col. Inter. Sci. Vol. 287 (2005), p. 173 [4] J. Li, J. Fu, Y. Cong, Y. Wu, L. Xue, and Y. Han, Applied Surface Science Vol. 252 (2006), p. 2229 [5] Y. Li, W. Cai, B. Cao, G. Duan, and F. Sun, Polymer, Vol. 46 (2005), p.12033 [6] Y. Li, W. Cai, G. Duan, F. Sun, B. Cao, and F. Lu, Mater. Lett. Vol. 59 (2005), p.276 [7] Y. Li, W. Cai, G. Duan, B. Cao, F. Sun, and F. Lu, J. Col. Inter. Sci. Vol. 287 (2005), p.634 [8] H. Wang, S. Y. Zheng, X. D. Li, and D. P. Kim, Micro. Meso. Meter. Vol. 80 (2005), p.357 [9] G. Gregori, H. J. Kleebe, H. Brequel, S. Enzo, G. Ziegler, J. Non-Crystal. Sol. Vol. 351 (2005), p.1393 [10] A. Stein, Micro. Meso. Meter. Vol. 44-45 (2001), p.227 [11] P. B. Landon, C. L. Gilleland, B. Jarvis, B. D. Waters, K. Inoue and R. Glosser, Colloids and Surfaces A: Physicochemical and Engineering Aspects, Vol. 259 (1-3) 2005, p. 31

Solid State Phenomena Vol. 135 (2008) pp 31-34 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.31

Reflectometry Studies of Mesoporous Silica Thin Films Young Kyu Hwang,1*A. S. Mamman,1 K. R. Patil,1 Lee-Kyung Kim,2 Jin-Soo Hwang,1 Jong-San Chang1* 1

Research Center for Nanocatalysis, Korea Research Institute of Chemical Technology (KRICT), Yusung, Daejeon 305-600, Korea. Email: [email protected] and [email protected]

2

System Group, Korea Materials & Analysis Corporation,Yusung, Daejeon, 308-380, Korea

Keywords: Mesoporous silica film; thickness; Spectroscopic ellipsometry; Reflectometry, porosity.

Abstract. Reflectometry technique has been successfully applied to investigate the correlation between the porosity and optical property (refractive index) of the ordered mesoporous thin film deposited on silicon wafer substrates. The measured optical spectra were simulated by the Effective Medium approximation model. The reflectometry technique has been found to be appropriate for the measurement of thickness of thin films as well as thick layer films. The mesoporous silica films prepared from tri-block copolymer (F-127) as a surfactant and polypropylene oxide as a swelling agent were subsequently exposed to the ammonia vapors to enhance thermal stability and shrinkage minimization of the film that results in increased film thickness. Introduction The porous silica films are gaining importance in membrane separations, chemical and optical sensors and especially optical devices such as low-k dielectric materials [1]. Currently, thin films used in advanced devices are often required to have non-destructive specifications approaching such as thickness measurement and porosity analysis. For these porous thin films owing to the limited quantity of samples available from mesoporous thin films, it is very difficult to obtain film characteristics such as surface area, porosity, and pore volume by the BET analysis. Several reports have appeared in the literature for numerous experimental techniques such as ellipsometry [2], reflectometry [3], alpha-step [4], X-ray photoelectron spectroscopy (XPS) [5], fourier transform infrared spectroscopy [6], atomic force microscopy (AFM) [7], secondary ion mass spectroscopy (SIMS) [8]. The most adopted measurement techniques such as spectroscopy ellipsometry and laser reflectometry need no reference samples for calibration, measure two independent parameters, viz. ψ and ∆ at each wavelength, and are noninvasive and nondestructive ones [9]. All other methods either need to be calibrated against other techniques or require assumptions about substrate system, such as electron mean free path for XPS, surface damage in alpha-step and AFM. Among all these methods, laser reflectometry is able to focus on a small spot (micro-area) for optical measurement equipped with a measuring speed of less than 0.5 second. The present investigation has been undertaken with an endeavor to assess the suitability of the optical reflectometry in measuring the thickness of mesoporous silica film at room temperature by using analytical representations for their optical responses to correlate film porosity with optical property, refractive index

Experimental Mesoporous silica thin films which were fast condensed by using ammonia vapor were prepared by reported procedure [10]. The samples thickness and porosity were analyzed by laser reflectometry and by spectroscopic ellipsometry. Reflectometry measurement was carried out with

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Nanocomposites and Nanoporous Materials VIII

fiber optic reflectometer (K-MAC ST-2000 instrument). In order to compute thickness and optical property (refractive index and extinction coefficient), the models were used for fitting consisted of a porous silica film deposited on silicon substrate. The thickness, refractive index, and extinction coefficient of the top layer are fitted with parameters. The optical parameters were modeled by Bruggeman Effective Medium Approximation (EMA) [11]. The root mean squared error (RMSE) was used as a figure of merit function. The thickness of the films is determined (after cross-sectioning) by SEM (Philips XL-30) for their comparison with the fitted values. The surface roughness is measured using atomic force microscopy (AFM Nanoscope III Digital instrument) in a tapping mode. Optical ellipsometric measurements are carried out with a phase-modulated ellipsometer (UVISEL-Jobin-Yvon) using the spectral range of 400-750 nm and an incidence angle of 70 o. The details of methodology are provided elsewhere [10]. Results and Discussion A synthetic strategy adopted for preparing thin silica film results in highly condensed mesosporous structures. The development of film organization resulting in the formation of mesostructures in the thin silica film has been discussed elsewhere [10]. As the condensation progresses of silica precursor, the silica network becomes rigid as a result of enhancement of H+ concentration and water content. The competitive processe of reorganization and condensation is arrested due to rigidification of an inorganic network. Table 1 shows a comparison of refractive index and film thickness of NH3-treated and calcined mesoporous materials. The optical properties of silica materials are changed due to ammonia treatment. The ammonia vapors penetrate the thin film at the initial stages of its formation and react with inorganic skeleton to initiate soft condensation while the surfactant template remains intact within the pores. This results in the minimization of contraction acting normal to the surface of the film. Table 1. Film refractive index, thickness and porosity from reflectometry and ellipsometry. Sample No.

Reflectometry Lorentz oscillation model

EMA model

Ellipsometry

SEM

EMA model

Cross section

tavg (nm)

navg

Porosity

tavg (nm)

navg

Porosity

tavg

navg

(nm)

Z-AB1

520.3

1.415

12

521.3

1.435

7

-

-

-

Z-AB2

669.2

1.242

52

655.0

1.240

52

-

-

-

Z-AB3

409.3

1.193

63

392.0

1.210

59

399.8

1.233

407.0

Z-AC3

253.5

1.336

31

251.4

1.340

30

236.5

1.369

277.0

* Z-AB1 (as-prepared) , Z-AB2 (NH3-treated) , Z-AB3 (NH3-treated and calcined film with polyproyleneoxide, PPO) and Z-AC3 (NH3-treated and calcined film without PPO), tavg=thickness, navg=refractive index

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Thickness of the mesoporous silica films, as determined by reflectometry, was found to be in a good agreement. The thin silica film treated with ammonia vapors shows an increase in its thickness by about 30%. The exposure of the films to ammonia vapors prior to its rigidification increases the rate of condensation which results in the formation of uniform highly ordered silica network of mesoporous structure. Under the influence of ammonia vapors, the rate of condensation is faster than the rate at which organization of spherical mecelles. The penetration of NH3 vapors has densified the silica wall to enhance thermal stability of the silica film and minimized the shrinkage of the film. The effect of NH3 vapors and thermal treatment on porosity of the film was calculated as summarized in Table 1. A decrease in its thickness by 20% was observed after calcinations (Table 1) due to elimination of surfactant phases from the film, matrix densification, and a partial collapse of pores thus leaves a porous structure behind. Film thickness obtained from Lorentz oscillation and EMA is in good agreement with each other. Table 1 also shows that the addition of poly (propylene oxide) increases the thickness of the film, thus suggesting that such addition of swelling agent leads to film expansion due to an increase in the interaction of silicate species with nonionic block copolymer (F127). This, in turn, increases porosity and pore diameter of the film. The refractive index is directly correlated with the porosity because the volume porosity of thin film can be estimated from its refractive index by using the following Lorentz-Lorenz equation.13 Porosity (volume) = 1 – [{(n2 – 1)/(n2 + 2)}•{(no2 + 2)/ (no2 – 1)}] where no is the refractive index of the dense silica (1.46), and n is the measured values of the films [13]. An additional evidence for the variation in film structure from as prepared, NH3-treated film and additional swelling agent was provided by cross-sectional SEM. Fig. 1 shows the cross-sectional SEM micrographs of the same films used in thickness measurement. All these films examined by SEM had uniform thickness consistent with those values determined by reflectometry and ellipsometry. Table 1 and Fig. 1 show the SEM thickness within 10% of the film thickness measured by reflectometric and ellipsometric techniques. The close correspondence of the SEM, ellipsometric and reflectometric thickness is critical because it confirms the reflectometric analysis that yielded the low refractive index for these films. These results demonstrate that the film thickness obtained from reflectometry measurements is consistent with one from cross section SEM image.

(a)

(b)

Figure 1. Cross-sectional SEM micrographs of the same film used in thickness measurement (a) with and (b) without poly (propylene oxide) as a swelling agent at 450o C, Z-AB3 and Z-AC3, respectively. Summary The refractive index of these films is affected by NH3 vapour treatment and poly (propylene oxide) resulting in an increase in thickness. Calcinations reduced the amount of organic material within these films, decreased the film thickness, and produced silica with porosities as high as 63%

34

Nanocomposites and Nanoporous Materials VIII

and the refractive index as low as 1.19. A parametric dispersion models like Lorentz oscillation and EMA models are employed considering the multi layer models and obtained a good agreement for film the thickness. The film thickness obtained from reflectometry, ellipsometry and SEM cross-section also shows consistencies. Acknowledgment This work was supported by the Korea Ministry of Commerce, Industry and Energy (MOCIE) through the Research Center for Nanocatalysts. The authors thank Dr. Y. J. Ko and Miss H. –K. Kim, for the Reflectometry analysis. A. S. Mamman gratefully acknowledges the financial support by KOFST. References [1] (a) Chen, F. L., Liu, M. L., Chem. Commun., 1829 (1999), (b) Miller, R. D., Science, 286, (1999) 421 . [2] (a) Mate, C. M., Lorenz, M. R., Novotny, V. J., J. Chem. Phys., 90 (1989) 7550. (b) M. Losurdo, D. Barreca, P. Capezzuto, G. Bruno, E. Tondello, Surf. Coat..Tech., 151-152 (2002) 2. (c) T. F. Stoica, V. S. Teodorescu, M. G. Blanchin, T. A. Stoica, M. Gartner, M. Losurdo, M. Zaharescu, Mater. Sci. Eng. B101 (2003) 222. (d) P. Petrik, M. Fried, T. Lohner, R. Berger, L. P. Biro, C. Schneider, J. Gyulai, H. Ryssel, Thin Solid Films, 313-314 (1998) 259. (e) B. Gruska, Thin Solid Films 364 (2000) 138. [3] (a) R. S. Balmer, C. Pickering, A. J. Pidduck, T. Martin, J. Crystal Growth, 245 (2002) 198. (b) R. S. Balmer, C. Pickering, A. M. Kier, J. C. H. Birbeck, M. Saker, T. Martin, J. Crystal Growth 230 (2001) 361. (c) S. Nair, M. Tsapatsis, Microporous Mesoporous Mater., 58 (2003) 81. [4] J. Rodriguez, M. Gomez, J. Ederth, G. A. Niklasson, C. G. Granqvist, Thin Solid Films, 365 (2000) 119. (b) B. Muller, M. Jager, Y, Tao, A. Kundig, C. Cai, C. Bosshard, P. Gunter, Optical Materials 12 (1999) 345. [5] (a) M. Hoshino, Y. Kimachi, J. Electron Spectrosc. Related Phenomena, 81 (1996) 79 (b) E. A. Irene, Solid State Electronics 45 (2001) 1207. [6] D. D. Saperstein, Appl. Spectrosc, 43 (1989) 481. [7] P. Petrik, M. Fried, T. Lohner, R. Berger, L. P. Biro, C. Schneider, J. Gyulai, H. Russel, Thin Solid Films, 313-314 (1998) 259. [8] M. R. Lorenz, V. J. Novotny, V. R. Deline, Surface Science 250 (1991) 112. [9] (a) D. E. Aspnes, A. A. Studna, E. Kinsborn, Phys. Rev. B29 (1984) 768. (b) K. Vedam, P. J. Mc Marr, J. Narayan, J. Appl. Phys. Lett. 47 (1985) 339. [10] (a) D. Grosso, A. Balkenende, P. A. Albouy, A. Ayral, H. Amenitsch, F. Babonneau, Chem. Mater. 13 (2001) 1848. (b) Y. K. Hwang, K. R. Patil, S. H. Jhung, S.-E. Park, J.-S. Chang, Microporous Mesoporous Mater. 78 (2005) 245. [11] B. Stjerna, E. Olsson, C. G. Granqvist, J. Appl. Phys., 76 (1994) 3797. [12](a) R. M. A. Azzam, N. M. Bashara, ‘Ellipsometery and Polarized Light.’ North-Holland Elsevier, Amsterdam, 1989. (b) P. A. Cuypers, J. W. Corsel, M. P. Janssen, J. M. M. Kop, W. T. Hermens, H. C. Hemker, J. Biol. Chem. 258 (1983) 2426. [13] (a) L. Hu, T. Yoko, H. Kozuka, S. Sakka, Thin Solid Films 219 (1992) 18; (b) M. Born, E. Wolf, principles of Optics, 4th ed., Pergamon Press, Oxford, 1970, Chapter 2.

Solid State Phenomena Vol. 135 (2008) pp 35-38 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.35

Conductive property of carbon-nanotube dispersed nanocomposite coatings for steel Yunshik Yang 1,2,a, Myeong-Jun Kim 1,b, Youngchul Lee 1,c and Si-Tae Noh 2,d 1 Green Engineering Team, Korea Institute of Industrial Technology (KITECH), 2

a

Chonan 330-825, KOREA Department of Chemical Engineering, Hanyang University, Ansan, Korea.

[email protected], [email protected], [email protected], [email protected]

Keywords : carbon-nanotubes, alkyd, nanocomposites, conductivity. Abstract : Nanostructured modification of polymers has opened up new perspective for multifunctional materials. Carbon-nanotubes have the potential to increase the conductivity of their composite, with improved or retaining mechanical performance. This study focuses on the evaluation of the thermal and electrical conductivities of carbonnanotube filled alkyd resins for steel coatings. Polymer/Carbon-nanotube nanocomposites have been prepared by mixing commercial multiwall carbon-nanotubes with alkyd resins and by curing. The thermal and electrical conductivities of carbon-nanotubes filled nanocomposite was found to be increased comparing with the original resin without any fillers or with the resin with carbon-black or carbon-nanofiber. 1. Introduction The outstanding thermal, electrical and mechanical properties of carbon-nanotubes(CNTs) offer the possibility of a new generation of thermal interface materials(TIMs) for thermal management of high density electronics. The conductivity and mechanical properties of individual CNTs is significantly higher than that of traditionally utilized fillers. CNTs provide a more efficient network for heat flow inside the polymer matrix because of their high aspect ratio.[1-2] This study aimed to investigate the effect of the CNT on the thermal and electrical conductivities of alkyd-based composites and to compare the effects of other fillers such as carbon-black, carbon-nanofiber and acid-treated CNT. In this study, efforts were focused on the characterization the conductivities of coated composite films on steel. 2. Materials and Experimental Materials A commercialized mutiwalled-carbonnanotube(MWNT) was obtained from Iljin-nanotec Inc. The MWNT was produced by the chemical vapour deposition (CVD) process. In order to compare the effects of fillers on conductivities, MWNT, carbon-black(CB, Degussa), carbon-nanofiber(CNF, Nanostructured & Amorphous Materials Inc.) and MWNT treated by acid( H2SO4 : HNO3 = 1: 3)

36

Nanocomposites and Nanoporous Materials VIII

were used. The acid treatment was intended to shorten the carbon-nanotubes and the length and diameter of acid treated CNT appears to be shortened and unchanged, respectively according to the transmission electron micrographs in Fig 1. The basic data of the nanofillers investigated in this study are given in Table 1 and the micrographs of their structure are taken and shown in Fig 1. The polymer matrix used for the composites was a short oil type alkyd resin (Aekyung chemical Inc.) and cured with a methoxymethyl melamine, supplied by Cytec Inc. Table 1. Specification of used nanoparticles.(from the suppliers catalogues) Purity Diameter Length Density Shape (um) (gcm-3) (%) (nm) MWNT 10~15 10~20 0.1 95 tube Acid-MWNT 10~15 tube Carbon-Black 19 0.17 95 sphere Carbon100~200 5~40 0.19 95 fiber nanofiber

20 nm

20 nm

2020nm nm

50 nm

Fig 1. TEM photographs of nanofillers: (a) MWNT, (b) Acid-treated MWNT, (C) Carbon-black, (d) Carbon-nanofiber Experimental The composites were produced using an identical processing condition. However the filler for composites and its content were varied. Homogeneous and translucent alkyd/fillers mixture was prepared using both sonication mixing and high speed shear mixing. Briefly, the fillers were dispersed by using ultra sonication for 20 min at the room temperature. The hardener was added to a filler/resin mixture and mixed using a high speed mixer at 3000rpm for 2hr to ensure uniform dispersion within the sample. The nanocomposite was coated on steel by using a bar coater and was cured at 180°C for 20 min. As substrates, steel plates having dimension of 150mm × 75mm × 0.5mm, were used after cleaning with acetone. The thermal conductivity was measured using a thermal analyzer TC-30 prove (Hot wired, Mathis), which is based upon a modified transient hot wired technic. The surface resistivity was measured using Loresta-GP-MCP-T600(4pin prove, Mitsubishi chemical) at a room temperature. More than five individual measurements were performed on random spots of coated steel for both thermal conductivity and resistivity.

Solid State Phenomena Vol. 135

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3. Results & Discussion The fractured surface of MWNT/Alkyd nanocomposites was investigated by SEM images and the dispersion of MWNT in alkyd resin was shown in Fig 2. MWNT can be identified as white spots in SEM micrographs, MWNT loadings were increased up to 1.5 vol % and MWNTs appear to be well-dispersed. (a)

(b)

1 µm

(d)

1001 µm

(e)

1㎛

(c)

1 µm

100 nm

(g)

(f)

1 µm

1 µm

1 µm

Fig 2. SEM of fractured surface of Alkyd/MWNT nanocomposites : (a)No CNT, (b)CNT contents 0.1vol%, (C)CNT contents 0.3vol%, (d)CNT contents 0.5vol%, (e)CNT contents 0.7vol%, (f)CNT contents 1vol%, (g) CNT contents 1.5vol%, The enhancement of the thermal conductivity with increasing volume content of the nanofillers was shown in Fig 3. The incorporation of MWNT into polymers resulted in a slight enhancement in the thermal conductivity comparing with other fillers. The transport of thermal energy in composites can be related to a phonon conduction mechanism.[3] In the MWNT the phonons can be carried in the inner walls without hindrance comparing with other fillers. It is considered to be that the phonon vibrations in the spherical-shaped carbon-black and stiff carbon-fiber are dampened by the matrix interaction. The higher aspect ratio and formation of a percolation network of CNTs seems to increase the heat conduction.[4] The surface resistivity of coated steel as a function of the volume content of the nanofillers is shown in Fig 4. Electrical percolation of CNT composite seems to occur at a low filler content comparing with other fillers. The addition of carbon nanotubes to the alkyd induces an electrical conduction even at a low CNT concentration, which can be explained by their high aspect ratio. As expected, the percolation threshold occurred at a lower content for fiber-shaped fillers (high aspect ratio) than for spherical particles. Electrical percolation of MWNT composite seems to occur at a lower filler content than acid-treated MWNT composite due to a higher aspect ratio. Any kind of treatments leading to a reduction of the aspect ratio may increase the percoration threshold since

38

Nanocomposites and Nanoporous Materials VIII

the percoration threshold concentration is mainly determined by aspect ratio and shape of fillers. 7e+5

No filler MWNT Acid-MWNT CB CNFs

3.45 3.40 3.35

MWNT Acid-MWNT CB CNFs

6e+5

surface resisitivity (ohm)

Thermal conductivity(W/mk)

3.50

3.30 3.25 3.20

5e+5 4e+5 3e+5 2e+5 1e+5 0

3.15 0

2

4

6

8

10

12

14

16

18

20

Filler contents(vol%)

Fig 3. Thermal conductivity of coated steels as a function of filler contents (vol%)

2

4

6

8

10

12

14

16

18

Filler contents(vol%)

Fig 4. Surface resistivity of coated steels as a function of filler contents (vol%)

4. Conclusions In this study, the thermal and electrical conductivities of composites containing CNTs were investigated. According to the SEM characterization, the carbon-nanotubes are nanoscopically welldispersed in the alkyd matrix. The incorporation of CNTs into alkyd resin resulted in an enhancement of the thermal conductivity comparing with other fillers. The higher aspect ratio and formation of a percolation network of CNTs seem to have a certain influence on the enhanced heat conduction. Electrical percolation of CNT composite seems to occur at a lower filler content comparing with other fillers. MWNTs exhibited the highest potential for an efficient enhancement of the electrical conductivity, mainly due to a high aspect ratio and the tubular shape. The aspect ratio of the filler and its dispersibility were found to be crucial parameters for the conductive composites at low filler contents. Improvements in dispersibility of CNTs and a higher aspect ratio of CNT seem to be required in order to develop thermally and electronically conducting CNT-filled coating materials. Acknowledgement : The funding of this research by the Ministry of Commerce, Industry and Energy through High Performance Nano Composites Program is gratefully acknowledged. References [1] A. Yu, M. E. Itkis, E. Bekyarova, and R. C. Haddon, Phys. Rev. Lett. 89, 133102(2006). [2] L. E. Nielsen, Ind. Eng. Chem. Fundam. 13, 17 (1974) [3] F. H. Gojny, M. H. G. Wichmann, B. Fiedler, I. A. Kinloch, and K. Schulte, Polymer. 47, 2041(2006). [4] J. K. W. Sandler, J. E. Kirk, L.A. Kinloch, A. H. Windle, Polymer. 44, 5893(2003).

Solid State Phenomena Vol. 135 (2008) pp 39-42 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.39

Preparation of Platinum-Ruthenium Nanoparticles on Graphite Nanofibers Seok Kim1 and Soo-Jin Park2,a 1

Advanced Materials Division, Korea Research Institute of Chemical Technology, P.O. Box 107, Yuseong, Daejeon 305-600, South Korea 2

Department of Chemistry, Inha University,

Yonghyun-dong, Nam-gu, Incheon 305-600, South Korea a

e-mail : [email protected]

Keywords: nanoparticles, graphite, nanofibers, catalysts, electroactivity

Abstract: Electroactivity of graphite nanofibers (GNFs)-supported PtRu particles was examined for their application as DMFCs anode. In this work, composites of PtRu nanoparticles of 2-8 nm size and graphite nanofibers were prepared by the electrodeposition methods. As a result, the methanol oxidation current for graphite nanofibers-supported PtRu catalysts was investigated by changing a deposition time. The electroactivity could be attributed to the particle size, particle dispersion ability, and deposition level.

Introduction Direct methanol fuel cells (DMFCs) are attractive portable power source because of their high energy density, easy fuel handling and lower operating temperature [1-4]. However, DMFC has two most serious drawbacks. One is the slow kinetics of methanol oxidation on platinum. Various Ptbased binary, ternary, and even quaternary compounds have been investigated intensively to increase the catalytic activity of methanol oxidation [5]. In addition, the development of a supporting material is essential to minimize noble metal loadings and achieve optimum catalytic performance. Carbon is a widely used supporting material, and different kinds of carbon materials such as mesoporous carbons, carbon nanotubes, and graphitic carbon nanofibers were reported, [6-8]. The other issue in DMFCs is the methanol crossover from anode to cathode side across the membrane. It is well known that methanol crossover lowers fuel utilization and causes cathode over-potential. In order to block or reduce methanol crossover, a commercial Nafion membrane has been modified by various technique and many new composite membranes have been proposed as alternatives. Even methanol-tolerant cathode catalysts, which are inactive to methanol cross-over were suggested, but their activity is still not satisfactory. In this work, graphite nanofibers were selected as catalyst supports to overcome the low electronic conductivity problem, and their composites with PtRu particles were examined to evaluate their electrocatalytic activity. Composites of PtRu nannoparticles and graphite nanofibers were prepared by the electrodeposition methods. Experimental Preparation of nanoparticles: Deposition and electrocatalytical properties of PtRu nanoparticles

40

Nanocomposites and Nanoporous Materials VIII

on GNFs have been investigated by an Autolab with PGSTAT 30 (Eco Chemie B.V.; Netherlands). A standard three-electrode cell was employed. GNFs electrodes with a definite area of 1.70 cm2 were used as the working electrode. A platinum wire as the counter electrode and a saturated calomel electrode (SCE) was used as the reference electrode. PtRu nanoparticles were electrodeposited on the GNFs electrode by step (interval time: 0.06 sec) potential plating method from distilled water with ruthenium chloride and chloroplatinic acid. In the deposition solution, the concentration of Pt and Ru was kept to be constant (20mM). GNFs electrodes were cycled in the range of -0.3 V to -0.7 V in deposition solution for a different plating time. Characterization: Electrochemical behaviour of the PtRu/GNFs catalysts was studied in 1.0 M CH3OH + 0.5 M H2SO4 aqueous solutions by cyclic voltammetric technique. All experiments were carried out at room temperature (25°C). The crystallinity of PtRu/GNFs catalysts was evaluated using X-ray diffraction (XRD) performed on a Rigaku D/MAX-ΙΙΙB X-Ray diffractometer using a CuKα source. The X-ray diffractograms were obtained for 2θ values varying between 30°and 85°. The mean crystalline sizes of the particles were determined from the X-ray dffractograms, using a Scherrer equation. L =

0.9λ B 2θ cos θ max

where λ is the X-ray wavelength (1.54056 Å for the CuKα radiation), B2θ is the width of the diffraction peak at half-height and θmax is the angle at the peak maximum position. The deposition content of catalysts was obtained by Inductively Coupled Plasma-Atomic Emission Spectroscopy (ICPAES) methods. Results and discussion Nanocomposite structures: The crystalline structures of the PtRu/GNFs catalysts have been investigated by XRD. Figure 1 shows the XRD patterns for the PtRu/GNFs catalysts prepared by changing the plating time. The peaks at 2θ = 40°, 47°, 68°, 82° are associated with the (111), (200), (220), and (311), respectively.

Pt (111) Pt (200)

Intensity (a.u.)

Pt (220)

Pt (311) (36)

(d)

(24)

(c)

(12) (b) (6)(a) 40

60

80

2θ

Fig. 1. XRD patterns of PtRu/GNFs catalysts as a function of plating time of (a) 6, (b) 12, (c) 24, and (d) 36 min

Solid State Phenomena Vol. 135

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The diffraction patterns of PtRu/GNFs catalysts were similar to those of the Pt. The XRD result confirms obviously that Pt is co-electrodeposited on the surface of GNFs. The average size of the Pt particles was calculated from broadening of the (220) diffraction peak. The average crystalline sizes calculated from the XRD using a Scherrer equation were given in Table 1. The average crystalline sizes were between 2 and 8 nm. In the viewpoint of catalyst sizes, 24 min plating time was the best condition for obtaining the nano-sized catalysts.

Table 1. Average crystalline sizes and contents of PtRu/GNFs catalysts prepared by changing plating time Plating Time (min)

Crystalline Size (nm)

Deposition Content (%)

6

7.31

2.4

12

5.17

3.4

24

2.33

6.4

36

2.63

12.4

15

d

I(mA/cm2) I (mA/mg)

10

c

5

b a

0 -5 -10 -15 200

400

600

800

mV (Ag/AgCl)

1000

1200

Specific Current Density (mA/mg-catalyst)

Electroactivity: Figure 2 gives the cyclic voltammograms and current densities of graphite nanofibers-supported PtRu catalysts in 1 M CH3OH and 0.5 M H2SO4 solution at room temperature. Until 12 min plating, an oxidation peak is not clearly shown. In the case of 24 min and 36 min plating, the oxidation peak was found at ~900mV and the peak current was increased. The specific current densities were calculated by obtaining the oxidation peak current at 900 mV by dividing the weight of particle catalysts. By comparing these values, we could determine the electroactivity of catalyst electrodes. The specific current density after 24 min plating had been largely increased, when compared to that of 12 min plating. However, in the case of 36 min plating, this value had been slightly decreased, probably due to the increased particle size.

100

50

0 6

12

18

24

30

36

Plating time

Fig. 2. Cyclic voltammograms and specific current density of PtRu/GNFs as a function of plating time of (a) 6, (b) 12, (c) 24, and (d) 36 min

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Nanocomposites and Nanoporous Materials VIII

The methanol oxidation current for graphite nanofibers-supported PtRu catalysts was highest in the case of 24 min deposition time. The enhanced current density of PtRu catalysts was attributed to the smaller particle size and the better particles dispersion ability. Summary In the present study, the preparation and characterization of PtRu/GNFs catalysts had been investigated. PtRu particles were electrodeposited on GNFs by step potential method from aqueous solution containing ruthenium chloride and chloroplatinic acid. The crystalline size of PtRu catalyst was about 2-8 nm. The specific current density for graphite nanofibers-supported PtRu catalysts was increased until 24 min deposition time. After that time, the current density was decreased. The electrochemical activity was controlled with the change of the plating time. The enhanced current density of PtRu catalysts was attributed to the smaller particle size and the better particle dispersion ability. Acknowledgement This paper was performed for the Hydrogen Energy R&D Center, one of the 21st Century Frontier R&D Program, funded by the Ministry of Science and Technology of Korea. References [1] C. A. Bessel, K. Laubernds, N. M. Rodriguez, and R. T. Baker: J. Phys. Chem., Vol. 105 (2001) p. 1115. [2] M. Watanabe, M. Uchida, and S. Motoo: J. Electroanal. Chem., Vol. 229 (1987) p. 395. [3] S. Kim, M.H. Cho, J.R. Lee, S.J. Park: J. Power Sources, Vol. 159 (2006), p. 46. [4] J. Lee, S. Han, T. Hyeon: J. Korean Ind. Eng. Chem., Vol. 15 (2004) p. 483. [5] A. Chu and R. Jiang: Solid State Ion., Vol. 148, (2002) p. 591. [6] S.H. Joo, S.J. Choi, I. Oh, J. Kwak, Z. Liu, O. Terasaki, R. Ryoo: Nature, Vol. 412, (2001) p. 169. [7] C.Y. Chen, P. Tang, J. Power Sources, Vol. 123, (2002) p. 37. [8] S. Kim, S.J. Park, Electrochim. Acta, Vol. 52, (2007) p. 3013.

Solid State Phenomena Vol. 135 (2008) pp 43-46 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.43

Thermal Behaviors and Fracture Toughness of Polyurethane-dispersed Difunctional Epoxy Resins Soo-Jin Park1* and Jae-Rock Lee2 1

Department of Chemistry, Inha Univ., 253, Nam-gu, Incheon 402-751, Korea (South) 2

Advanced Materials Division, Korea Research Institute of Chemical Technology P.O. Box 107, Yuseong, Daejeon 305-600, Korea (South) *[email protected]

Keywords: Difunctional epoxy, Polyurethane modified epoxy, Mechanical interfacial properties

Abstract. In this work, the thermal and mechanical interfacial properties of diglycidylether of bisphenol A (DGEBA)/polyurethane modified epoxy (UME-305) blends were investigated. 4,4’-Diaminodiphenyl methane (DDM) was used as a curing agent, and the content of UME-305 in the mixture was 0, 20, 40, 60, 80, and 100 wt%. The cure behaviors of DGEBA/UME-305 blends were studied by DSC. The mechanical interfacial properties were confirmed by critical stress intensity factor (KIC) at 77K and 298K. As a result, the exothermic peaks in DSC results were shifted to higher temperature region as increasing the UME-305 content in the blends. The KIC was also enhanced with increasing the UME-305 content and showing a maximum value at 60 wt.% UME-305.

Introduction Epoxy-based materials have been widely used in coatings, as encapsulant for electronic components, as adhesives, as foams used to produce low weight castings for electronic applications and for coating textiles because of their outstand mechanical, thermal, and electrical properties[1,2]. However, highly cross-linked epoxy resins are rigid and brittle in nature and have poor crack resistance, which limit their many end-use applications, such as structural materials. Thus, a wide variety of fillers have been added to the epoxy resins in order to achieve an improvement of some properties, such as fracture toughness. Recently, interpenetrating polymer networks (IPNs) that use a rubbery polyurethane (PU) phase have been considered for such applications [3,4]. Polyurethane elastomers, which are linear alternating block copolymers, have excellent elasticity, abrasion resistance, impact strength, and low-temperature performance and the properties of both rubber and plastics. Typically, these materials consist of a continuous amorphous soft domain and a discrete glassy or crystalline hard domain. The soft segment domains impart elastomeric properties, whereas the hard segment domains provide rigidity and the resulting mechanical properties. The objective of this work is to study the effects of polyurethane in the epoxy on the cure behaviors, mechanical and properties of the diglycidyl ether of bisphenol A (DGEBA)/UME-305 blends. Experimental Epoxy resins used in this study were diglycidylether of bisphenol A (DGEBA) supplied by Kukdo Chem. of Korea (YD-128), which had an epoxide equivalent weight of 185-190g eq-1and a density of about 1.16 g cm-3 at 25°C. Diaminodiphenyl methane (DDM) purchased from Aldrich Chem. was selected as a curing agent. The DGEBA was heated to melt at 80°C for 1 h in the beaker. After the epoxy resin was melted, the UME-305 was added into the beaker. The content of UME-305 was varied within 0, 20, 40, 60, 80, and 100 wt% to DGEBA. The mixtures were degassed in a vacuum oven at 60°C for 1 h before being poured into the mold, and cured at 120°C for 1 h, 150°C for 2 h, and at 180°C for 1 h in a convection oven. The cure behaviors of the epoxy blends were investigated by a

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Nanocomposites and Nanoporous Materials VIII

DSC (Perkin-Elmer DSC-6) in the temperature range from 30 to 300°C under a nitrogen flow of 30 ml min-1. The KIC value based on the single-edge-notched (SEN) beam fracture toughness test in a 3-point bending flexural was conducted on an Instron 1125 mechanical tester according to ASTM E-399, which was carried out at room and cryogenic (77K) temperatures to determinate the performance of UME-305 at room and low temperatures, respectively. The fractured surfaces were examined using a scanning electron microscope (JEOL Model 540A, SEM). Results and Discussion Cure behaviors. The cure behaviors of the DGEBA/UME-305 blends cured with DDM were evaluated using DSC. Fig. 1 shows the dynamic DSC curves at a heating rate of 10°Cmin-1 for the cure of DGEBA/UME-305 blends. These thermograms provide information for determining the condition of cure reactions. The exothermic peak temperature of the blends was shifted to higher temperature as the UME-305 contents increased, due to the higher intermolecular interactions in the DGEBA/ UME-305 blends. 25

Heat flow(mW)

20

15 UM E-305 Content 0% 20% 40% 60% 80% 100%

10

5

0 50

100

150

200

250

300

o

Temperature( C)

Fig. 1. Dynamic DSC thermo grams of DGEBA/UME-305 blend. Mechanical interfacial properties. The fracture behaviors of materials depend on the stress level, flaw concentration, material properties, and failure mechanism. The critical stress intensity factor (KIC) was used to investigate the fracture toughness of brittle materials [5-7]. Fig. 2 shows the evolution of KIC in the flexure of the DGEBA/UME-305 blends as a function of UME-305 content. Fracture toughness at room and cryogenic temperatures increased with increasing UME-305 content and reached a maximum value at 60 wt% of UME-305. It can be explained by the increase in hydrogen bonding between the hydroxyl group in DGEBA and the isocyanate group in UME-305. It is interesting to note that higher KIC values are found at a 77K test, indicating that excellent low-temperature performance can be introduced by adding UME-305, attributed to the increased flexibility of the polyurethane [8]. Fig. 3 is the results of impact strength of the DGEBA/UME-305 blends as a function of UME-305 content. Ductile toughness of the blends was increased with increasing UME-305 content up to 60 wt.%. This is probably due to the enhancement of interpenetration between DGEBA and UME-305 chains, resulting in the high compatibility and crosslinking density. Fig. 4 shows SEM results of the fractured surfaces of neat epoxy and the blends after the fracture tests.

Solid State Phenomena Vol. 135

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room temp. cryogenic temp.

6.0

1/2

KIC(MPa m )

5.5 5.0 4.5 4.0 3.5 3.0 2.5 0

20

40

60

80

100

UME-305 Contents (wt%)

Fig. 2. Critical stress intensity factor of DGEBA/UME-305 blends.

Impact strength(kgf)

6

5

4

3

2 0

20

40

60

80

100

UME-305 Contents (wt%)

Fig. 3. Impact strength of DGEBA/UME-305 blends as a function of UME-305 content.

(a)

(b)

(c)

(d)

(e)

(f)

Fig. 4. SEM micrograph of DGEBA/UME-305 blends: (a) 0%, (b) 20%, (C) 40%, (d) 60%, (e) 80%, (f) 100%.

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Nanocomposites and Nanoporous Materials VIII

Conclusion In this work the cure behaviors, and mechanical interfacial properties of DGEBA/UME-305 blends cured with DDM were studied. As a result, the KIC and impact strength of the DGEBA/UME-305 blends showed enhanced in the presence of UME-305, and showed maximum value at 60 wt% UME-305 content. These results probably occurred because of intermolecular hydrogen bonding between the hydroxyl group in DGEBA and the isocyanate group in UME-305, resulting in improved compatibility of the components within the interpenetrating polymer networks. Moreover, a significant improvement in the KIC at cryogenic temperature of epoxy resins could be achieved as UME-305 was mixed with DGEBA. References [1] J. J. Meister: Polymer Modification (Marcel Dekker Inc., New York 2000). [2]

R. S. Bauer, Epoxy Resin Chemistry (Advanced in Chemistry Series, No. 114, American Chemical Society, Washington, DC, 1979).

[3] B. C. Kim, D. S. Lee, and S. W. Hyun: J. Ind. Eng. Chem. Vol. 7 (2001), p. 449. [4] S. J. Park, K. Li, and F. L. Jin: J. Ind. Eng. Chem. Vol. 11 (2005), p. 720. [5] D. W. Chung, J. P. Kim, and D. Kim: J. Ind. Eng. Chem. Vol. 12 (2006), p. 783. [6] S. J. Park, M. K. Seo, T. J. Ma, D. R. Lee, J. Colloid Interface Sci. Vol. 252 (2002), p. 249. [7] B.C. Kim, D.S. Lee, S.W. Hyun, J. Int. Eng. Chem. Vol.7 (6) (2001), p. 449. [8] S. R. Sandler, F. R. Berg: J. Appl. Polym. Sci. Vol. 9 (1965), p. 3909.

Solid State Phenomena Vol. 135 (2008) pp 47-52 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.47

Influence of Multiwalled Carbon Nanotube on Rheological Behavior of Mesophase Pitches Young-Seak Lee1 and Soo-Jin Park2,* 1

Dept. of Fine Chemicals Engineering and Chemistry, Chungnam National University, Yuseong, Daejeon 305-764, Korea (South) 2

Dept. of Chemistry, Inha Univ., 253, Nam-gu, Incheon 402-751, Korea (South) *[email protected]

Keywords: rheological behavior, mesophase pitch, multi-wall carbon nanotubes

Abstract. The rheological behaviors of mesophase pitch containing different contents (0, 1.0, 2.0 wt%) of multi-wall carbon nanotubes (MWNTs) were studied by using ARES cone-plate rheometer. The dynamic response of mesophase pitch containing MWNTs was different from that of pure mesophase pitch due to the MWNTs as a suspension in viscous pitch. The dynamic viscosity increased with increasing the amount of MWNTs, which is a clear evidence of the interruption of MWNTs in mesophase pitch. Also, the phase angle result indicates that mesophase pitch containing MWNT had less elastic nature than pure mesophase pitch. Introduction Carbon fibers are commonly produced by using melt-spinning method and their properties depend on both the fiber precursor and the processing conditions used during fiber formation and heat treatment [1]. In previous research, it is reported that good spinnability of precursor leads to enhance mechanical properties of the fibers, particularly tensile strength [2]. Spinnability is, in turn, related to the rheological behavior of the pitch precursor at the spinning temperature. The die-swell phenomenon is of particular interest since it is related to the viscoelastic nature of the pitches. Since the discovery of carbon nanotubes (CNTs) in 1991 by Iijima [3], their unique mechanical properties, such as the high strength and stiffness and their large aspect ratio make them a potentially attractive structural element for improving (fracture-) mechanical properties [4]. It might be expected that when CNTs are added to mesophase pitches, the rheological properties of melt will be significantly changed. In this work, in order to study the rheological characteristics of CNTs-mesophase pitches, storage and loss modulus are determined at different temperatures to examine elastic and viscous nature of pure and MWNT-mesophase pitches. Also, the effect of CNTs in molten pitches is studied by the analysis of phase angle and dynamic viscosity. Experimental The mesophase pitches were produced by the Mitsubishi Gas Chemical Company (MGCC) and had a designation of ARA24R lot 6U21. The Mettler softening point was found to be 296.2°C. Carbon nanotubes (multiwalled carbon nanotubes, MWNTs) used as filler were manufactured by CVD process (MWNTs were supplied from Iljin Nanotech Co. of Korea, degree of purity: >90-94%, length: 10-50 µm, diameter: 10-20 nm). In order to get a homogeneous dispersion of the MWNTs in the mesophase pitch matrix, the following procedure was used. The mesophase pitches were first refluxed in a reactor at 150°C under a nitrogen flow. Then the MWNTs were progressively added to the reactor. The MWNTs-mesophase pitch mixture was stirred for 1 h and the mixtures containing 1 and 2 wt% MWNTs were finally obtained. The ARES rheometer was used to measure the dynamic response of pure and mesophase pitches containing different contents of MWNT (referred to as MWNTs-mesophase pitches). A 25 mm cone-plate fixture was used with 0.1 radian of cone angle. The motor was kept on dynamic mode, and the transducer was set to a low torque range of 0.2-200 g·cm.

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The samples were presheared for 60 s at a shear rate of 0.1 sec-1. Approximately, 5 mm diameter pellets of pure mesophase pitches and MWNTs-mesophase pitches were prepared and then placed between cone and plate fixture. The sample was loaded in the enclosure oven, and then the rheological experiment was carried out. Testing temperatures were set ~ 30ºC above the softening point of the pure pitches. Each test was run over temperature range of 300~320 °C for pure mesophase pitches, 310~330°C for 1 wt% MWNTs-mesophase pitches, 320~340°C for 2 wt% MWNTs-mesophase pitches. Results and Discussion Fig. 1 shows the results of frequency sweep experiment of pure, 1 wt% MWNTs, and 2 wt% MWNTs-mesophase pitches. The results show similar trends for all experimental conditions. Both storage (G’) and loss modulus (G”) increase with increasing the frequency and then decrease with increasing the temperature. At low temperature, both G’ and G” go through a maximum value in all the samples and the plot lines of G’ and G” are roughly parallel to each data at intermediate and high temperature. Master curves are often used to compare the elasticity of pure mesophase pitches and MWNTs-mesophase pitches. The master curves of pure and MWNTs-mesophase pitches were generated from the individual storage and loss modulus curves. The procedure involves using a reduced temperature to relate all data sets. The reduced temperatures were chosen by considering that the dynamic viscosity is about 20 Pa⋅s that is typical shear viscosity range for melt spinning of mesophase pitches. The dynamic response of pure, 1 wt% MWNTs, and 2 wt% MWNTs-mesophase pitches is shown in Fig. 2. The master curves of pure and 1 wt% MWNTs-mesophase pitches are well behaved, and there is no significant discontinuity in the data. However, the master curve of 2 wt% MWNTs-mesophase pitches shows several discontinuities in the plot of G”. Therefore, it is assumed that the MWNTs influence on dynamic response of mesophase pitches. The phase angle, delta (δ), between the imposed strain and the resulting stress provides very useful information about the elastic nature of a given material. Cheung et al [5] suggested that the typical response of viscoelastic materials is observed at both low and high frequencies. The low frequency contribution is due to the diffusion mechanisms, whereas the high frequency contribution is due to the molecular motions of side chains. When the phase angle equals to zero, the materials are purely elastic. In contrast, an increase in phase angle indicates that the materials are becoming less elastic and more viscous. The phase angle of 90° indicates purely viscous materials, i.e., Newtonian fluid. The phase angle and dynamic viscosity of pure and MWNTs-mesophase pitches are shown in Fig. 3. It is common that the phase angle does not show significant difference by 20 rad/s of frequency. At the high frequency region, the pure mesophase pitches and MWNTs-mesophase pitches exhibit nearly the Newtonian behavior [6]. The value of delta approaches approximately 85° for the pure mesophase, whereas the phase angle for the MWNTs-mesophase pitches is slightly less (about 75°). Therefore, it is assumed that MWNTs-mesophase pitches are more elastic than the pure mesophase pitches. It can be assumed that the dynamic viscosity probably displays similar trend to the apparent viscosity based on steady sweep experiments, as commonly known in the literature [7-9]. As the frequency increases, the dynamic viscosity decreases and it attains a constant value after the frequency of 100 rad/s. As mentioned previously, the main difference of flow behavior between the pure and the MWNTs-mesophase pitches appears to be caused by the MWNTs in mesophase pitches. The MWNTs appear to behave as a suspension and influence the flow of mesophase pitches during measuring dynamic test. Therefore, the MWNTs-mesophase pitches are more elastic than pure mesophase pitches. However, it is difficult to explain the exact differences observed in dynamic test. Also, it is postulated that the spinning temperature of MWNTs-mesophase pitches should be higher as the amount of MWNTs increases in the pitches. Conversely, it is expected that the spinnability of MWNTs-mesophase pitches will be decreased as the amount of MWNTs is increased.

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Nanocomposites and Nanoporous Materials VIII

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Conclusion In summary, MWNTs-mesophase pitches exhibited discontinuous master curves, indicating that the MWNTs act as a suspension in mesophase pitches and interrupt the rheological behavior of the mesophase pitches. The trend of dynamic viscosity of mesophase pitches is similar for both the pure and the MWNTs-mesophase pitches. However, the phase angle of pure mesophase pitches indicates that the pure mesophase pitches are less elastic than the MWNTs-mesophase pitches. Acknowledgement This work was performed for the Hydrogen Energy R&D Center, one of the 21st Century Frontier R&D Program, funded by the Ministry of Science and Technology of Korea. References [1] L. H. Peebles: Carbon fibers: formation structure, and properties (CRC Press, Boca Raton, 1995). [2] J. K. W. Sandler, S. Pegel, M. Cadek, F. Gojny, M. van Es, J. Lohmar, W. J. Blau, K. Schulte, A. H. Windle, and M. S. P. Shaffer: Polymer Vol. 45 (2004), p. 2001. [3] S. Iijima: Nature Vol. 354 (1991), p. 56. [4] M. M. J. Treacy, T. W. Ebbesen and J. M. Gibson: Nature Vol. 381 (1996), p. 678. [5] T. Cheung, M. Turpin and B. Rand: Carbon Vol. 34 (1996), p. 265. [6] M. K. Seo and S. J. Park: Chem. Phys. Lett. Vol. 395 (2004), p. 44. [7] S. I. Bakhtiyarov and R. A. Overfelt: Acta Materialia Vol. 47 (1999), p. 4311. [8] F. F. Fang and H. J. Choi: J. Ind. Eng. Chem. Vol. 12 (2006), p. 843. [9] S. S. Ray: J. Ind. Eng. Chem. Vol. 12 (2006), p. 811.

Solid State Phenomena Vol. 135 (2008) pp 53-56 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.53

Preparation and Characterization of AuNP/Al2O3 with Bimodal Nanoporous Structure Younghun Kim1,a*, Ji Bong Joo2,b, Wooyoung Kim2,c, Jeongjin Lee2,d, and Jongheop Yi2,e 1

Department of Chemical Engineering, Kwangwoon University, Wolgye-dong, Nowon-gu, Seoul 139-701, Korea 2

School of Chemical and Biological Engineering, Seoul National University, Shinlim-dong, Kwanak-ku, Seoul 151-742, Korea a

[email protected], [email protected], [email protected], [email protected], e

[email protected]

Keywords: bimodal pore, PROX, gold nanoparticle, nanoporous alumina, water gas shift

Abstract. AuNP (gold nanoparticle) supported by oxide shows a high reactivity for a PROX (preferential oxiation) reaction at low temperature. Au catalysts were usually prepared by conventional methods such as precipitation, impregnation and vapor phase grafting. In this study, we developed a novel method for the preparation of AuNPs supported on a bimodal nanoporous alumina. The AuNPs were prepared in a toluene phase by the modified Brust method. The metal particle size was able to be controlled from 2 to 50 nm via the control of the surfactant concentrations. The resulting materials were characterized by BET, FE-SEM, TEM, and XRD analyses. After calcinations at 700oC, AuNP/Al2O3 catalyst revealed a bimodal nanoporous structure, with the pore sizes of 3.5 and 7 nm, and demonstrated both a high surface area (350 m2/g) and pore volume (0.9 cm3/g). Introduction Development of new energy resources, such as fuel cells and hydrogen production, is of great interested as an alternative petroleum resource. The WGS (water gas shift) and PROX reaction are known to be a crucial process for the production of hydrogen, and has been accomplished mainly by conventional copper-based catalysts, such as Cu/Al2O3 and Cu/CeO2 [1]. Interestingly, a recent review showed that AuNP supported by oxides demonstrates high reactivity for CO oxidation, however, much debate occurred in the literature in regarding on the activity and stability of supported gold catalysts [2-4]. A common feature among gold catalysts that show high activity is the presence of gold primarily in a zerovalent, metallic state at a nanosize. The catalytic properties of oxide-supported gold catalysts are strongly dependent upon the preparation methods of the catalysts, which determine both the cluster size and the nature of the active sites. Oxide-supported gold catalysts were prepared usually by conventional precipitation, electro deposition, and grafting of AuNPs. Martin et al. reports that the AuNP (ca. 8 nm) supported AAO (anodic aluminum oxide) membrane with uniform pores were prepared by electrodepositing gold within the pores of AAO [3]. Datye et al. prepared AuNPs in a mesoporous silica (MCM-41) by grafting gold onto the amine groups in the pores [4]. In this study, we prepared AuNPs/Al2O3 catalysts with a bimodal nanopore structure, which has a potential application to the PROX reaction. Here, nanosized gold particles were prepared in a toluene phase, and then used as a gold source and to form a secondary pore for support. Hierarchical channels, or well-connected small- and large-pore networks, show multiple advantages for the application to catalysis or adsorption process. Primary nanopores provide high surface area with size or shape selectivity for a guest molecule, while additional nanopores reduce mass transport limitations [5].

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Experimental Bimodal nanoporous AuNP/Al2O3 was prepared by mixing AuNPs in toluene with a suspension of aluminum oxide. Self-stabilized gold sol in toluene was prepared by modified Brust procedure [6]. [CH3(CH2)7]4N(Br), HAuCl4, and NaBH4 were used as a surfactant, a gold precursor, and a reducing agent, respectively. After a color change to dark ruby, the resulting solution was subsequently washed in H2SO4, Na2CO3, and H2O, followed by drying in MgSO4. The resulting sol contained approximately 12 mM of gold with a particle mean diameter of 6-8 nm [6,7]. To prepare nanoporous alumina, stearic acid and aluminum sec-butoxide were used as porogen and aluminum precursors, respectively. Alumina sol was prepared by a templating method via post-hydrolysis [7]. AuNPs in 20 ml of toluene was mixed with alumina sol, followed by further agitation. The product was then dried at room temperature for 48 hr, then it was sequentially calcined for 3 hr at 350oC in air and 5 hr at 700oC in vacuum. Both stearic acid and AuNPs in toluene acted as pore-generating agents. Stearic acid formed the primary nanopore, while the AuNPs/toluene formed the secondary nanopore. Calcined materials in air and in vacuum are referred to as AuMA-a and AuMA-v, respectively. The pore properties of the final products were analyzed with ASAP-2010 (Micromeritics) apparatus, and surface morphologies were characterized with FE-SEM (JSM-6700F, Jeol) and TEM (JEM-2000EXII). X-ray diffraction (XRD, M18XHF-SRA, MAC/Science) measurements were carried out in order to investigate the phase transformation of supports. Results and discussion As shown in Fig. 1, pore properties of AuNP/Al2O3 were analyzed with an N2 adsorption/desorption test. As a first step, the materials prepared were calcined at 350oC in air in order to remove surfactants and organics, followed by the second calcinations at 700oC in vacuum for the phase change of the aluminas and improve thermal stability. Calcined materials are colored violet in the AuNP solution. When the organic material was not fully removed in the first calcinations step, the residual organics caused carbonization in the second calcinations step, which was characterized as a black powder. Calcined materials in both the air and the vacuum procedures represented the two distinctive nanopores. When the pore properties of AuMA-a were analyzed using the BJH model, two pore sizes were appeared at 3.5 and 5.8 nm, while AuMA-v had 3.5 and 6.9 nm. The primary pore (3.5 nm) has the typical size of a unimodal nanoporous alumina prepared using stearic acid as a template [8]. The secondary pore (6-7 nm) is formed due to the presence of AuNPs in the solution. Toluene in the solution penetrates into the inside (hydrophobic part) of the micelle and then acts as a swelling agent to expand the primary pore. It should be noted that the framework structure of the aluminum hydroxide (i.e., alumina sol) was firmly formed by hydrolysis and condensation reactions during agitation for 24 hr. Thus, the swelling of the framework pores would not be induced by the addition of toluene. Consequently, the secondary pores were formed by the permeation of AuNP/toluene into the primary particluated sol, as seen in Fig. 2. AuMA-v has a textural porosity, such as a void fraction from primary particles, induced by AuNP/toluene penetration into the alumina sol. In the vacuum calcinations step, the primary pores were maintained at a typical size (3.5 nm), while the second ones were expanded as much as about 1 nm, due to the sintering of both the alumina framework and of the AuNPs. The scheme is proposed in Fig. 2. Table 1. Pore properties of AuNP/Al2O3 calcined in air and vacuum. Calcination condition

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345

0.82

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0.91

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353

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The BdB-FHH model is known to be more accurate than the BJH model for spherical pores [5]. As summarized in Table 1, the second pore measured by the BdB-FHH model showed approximately 8 nm, which was similar to the size of the AuNP. The surface area and pore volume of AuMA-v was 350 m2/g and 0.9 cm3/g, respectively, which was similar to that of pure mesoporous alumina [8].

Fig. 1. Pore size distribution of AuNP/Al2O3 calcined in air and vacuum, and N2 isotherm of the calcined material in vacuum.

Fig. 2. Scheme of the formation of secondary pore in AuNP/Al2O3; (a) pore swelling and (b) penetration of AuNP solution into alumina sol. The surface morphology was analyzed with TEM and SEM. In Fig. 3a, AuMA-v showed a distribution of pores with a uniform pore size, which was confirmed by an N2 isotherm analysis. The pore sizes measured are uniform at 3.5 and 7 nm. Interestingly, nanoporous alumina showed typical wormhole-like pore structures [8], which could facilitate the penetration of target materials onto active sites in the inner surface for the applications to the catalytic reaction and the adsorption process. As shown in Fig. 3b, AuMA-v displayed a sponge-like morphology. Note that when only stearic acid is used for the preparation of nanoporous alumina, SEM images appear as hard particles with a smooth surface [9].

Fig. 3. (a) TEM and (b) SEM images of AuNP/Al2O3 calcined in vacuum. Distribution of AuNP either surface or inside the support was analyzed. The loading amount of Au supported on alumina was 5 wt%. The Au contents analyzed with SEM-EDS were approximately 0.8wt%, which is exposed to the surface. Therefore, approximately 4wt% of AuNP might be located at the interior of the support. It should be noted that initial Au concentration prepared by the Frens method is 12 nM, and 10-6 times smaller than that prepared by the Brust method. The amount of AuNP loading on the support was adjustable by controlling the initial concentration of AuNPs by Brust method. The crystalline structure of AuMA-v was characterized by XRD analysis

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(Fig. 4b). The four diffraction peaks can be indexed to the (111), (200), (220) and (311) planes of FCC gold, respectively (JCPDS 04-0784). As compared with the XRD pattern of pure nanoporous alumina (Fig. 4a), it was confirmed that the framework structure of AuNP/Al2O3 has an γ-alumina phase. Since γ-alumina is mainly used as a catalyst support, the resulting material obtained in this study has a potential application to catalytic reactions. The prepared material was tested in PROX reaction as a model reaction. The feed of H2/CO/O2 with equimolar ratio was injected in quartz reactor with 100 ml/min rate at 150oC. The resulting conversion of CO is 64.1% with 0.28 of selectivity.

Fig. 4. XRD patterns of (a) γ-alumina and (b) AuNP/Al2O3. Conclusions To prepare an alumina-supported nanosized gold catalyst, an AuNP solution prepared by the modified Brust method was used as a nano-gold source. The secondary pore was formed using an AuNP solution during the formation process of alumina sol. Bimodal AuNP/Al2O3 showed nanopores of 3.5 and 7 nm, due to the stearic acid and AuNP solutions, respectively, in the toluene phase. In model reaction of PROX, the resulting activity was low, but we found that AuNP/Al2O3 was possible to use in PROX reaction. This reaction system should be more improved with adjustment of Au contents and reaction condition to obtain the high activity of PROX reaction of AuNP/Al2O3. We gratefully acknowledge the financial support of the Basic Research Program (R01-2006-000-10239-0) of the Korea Science & Engineering Foundation. This research was conducted through the Realistic 3D-IT Research Program of Kwangwoon University under the National Fund from the Ministry of Education and Human Resources Development (2006). References [1] C. Zerva and C. J. Philippopoulos, Appl. Catal. B, 67 (2006) 105. [2] R. Burch, Phys. Chem. Chem. Phys., 8 (2006) 5483. [3] G. L. Hornyak, C. J. Patrissi and C. R. Martin, J. Phys. Chem. B, 101 (1997) 1548. [4] M. T. Bore, H. N. Pham, T. L. Ward and A. K. Datye, Chem. Commun., (2004) 2620. [5] Y. Kim, C. Kim and J. Yi, Mater. Res. Bull., 39 (2004) 2103. [6] N. Fishelson, I. Shkrob, O. Lev, J. Gun and A. D. Modestov, Langmuir, 17 (2001) 403. [7] M. Brust, D. Bethell, C. J. Kiely and D. J. Schiffrin, Langmuir, 14 (1998) 5425. [8] Y. Kim, C. Kim. I. Choi, S. Rengaraj and J. Yi, Environ. Sci. Technol., 38 (2004) 924. [9] Y. Kim, B. Lee and J. Yi, Korean J. Chem. Eng., in press (2007).

Solid State Phenomena Vol. 135 (2008) pp 57-60 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.57

The Effect of Physicochemical Treatment on Pd Dispersion of Carbon-supported Pd Catalysts Yong Tae Kim1, Eun Duck Park1, a, Min Kang2 and Jae Eui Yie2 1

Division of Energy Systems Research and Division of Chemical Engineering and Materials Engineering, Ajou University, Suwon, 443-749, Republic of Korea 2 Department of Applied Chemistry, Division of Biotechnology & Nanotechnology, Ajou University, Suwon, 443-749, Republic of Korea a

[email protected]

Keywords: Pd, carbon, surface oxygen, CMK-5, mesoporous carbon, XANES, EXAFS.

Abstract. Carbon-supported palladium catalysts were prepared by an impregnation method using palladium chloride and different carbon supports such as activated carbons with different surface oxygen concentrations and a mesoporous carbon, CMK-5.. The different degree of surface oxidation was achieved by the nitric acid and high-temperature heat treatment. The molecular PdCl2 species was stabilized by adding an HCl in an impregnation step.. As the reduction temperature increased, the Pd dispersion decreased for all Pd catalysts. There was no noticeable difference in Pd dispersion among Pd catalysts supported on carbon supports with different physicochemical properties when the reduction temperature was 423 K. Pd catalysts supported on the carbon support with a high concentration of surface oxygen groups showed a better dispersion than did other Pd catalysts when they were reduced at 573 K. The maximum Pd dispersion was observed over Pd catalyst supported on carbon supports with the highest surface area when the reduction temperature was higher than 573 K. Introduction Carbon materials such as activated carbon, carbon black, and graphite have been widely utilized as a support for precious metal catalysts because of their some advantages over other conventional metal oxide supports such as alumina and silica. Activated carbon, the most frequently used support, is stable in both acidic and basic media and can be burnt off to lead an efficient recovery of precious metal from spent catalysts. About 30% of powder precious metal catalysts are known to be supported palladium catalysts [1]. Most of them are being used for hydrogenation reactions, e.g. the synthesis of amines from nitro compounds or the saturation of carbon-carbon and carbon-heteroatom multiple bonds [1-3]. The surface properties of carbon can be modified by several methods including chemical and thermal treatments [4].. Oxidation in the gas or liquid phase can be used to increase the concentration of surface oxygen groups, while heating under inert atmosphere may be used to selectively remove some of these functions. Carboxyl, carbonyl, phenol, quinine, and lactone groups have been identified on carbon surfaces. A variety of experimental techniques have been used to characterize these functional groups, such as chemical titration methods, temperature-programmed desorption (TPD), X-ray photoelectron spectroscopy (XPS) and infra-red spectroscopy methods (FTIR, DRIFTS) [4].. Generally, the interaction between functional groups of the support surface and metal complexes in the impregnation step has been considered to be important in determining the catalytic properties of the final catalyst such as an average metal size and a particle size distribution. Carbon-supported palladium catalysts have been studied in this approach. Although there are some previous reports on carbon-supported Pd catalysts prepared from different kinds of palladium precursors such as H2PdCl4, [Pd(NH3)4](NO3)2, Pd(CH3COO)2 [5], the study on the interaction between the Pd precursor and the support has been limited on Pd catalysts supported on carbon black [6,7]. The partial reduction of PdCl2 during an impregnation step has been reported to be the cause of poor dispersion in the final Pd catalyst. The mesoporous carbon has recently attracted much attention for its application to the catalytic support because it has a high surface area and a relatively large pore

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size [8]. In this work, palladium catalysts supported on activated carbons with different concentrations of surface oxygen groups and a mesoporous carbon, CMK-5, were prepared from palladium chloride as a palladium precursor. The molecular PdCl2 species was stabilized by adding an HCl in an impregnation step. The interaction between palladium precursor and the support was characterized with an X-ray absorption fine structure (XAFS). The effect of reduction temperatures on the Pd dispersion was also studied. Experimental Activated carbon (AC) was obtained from Aldrich and its particular size was about 100 mesh. To prepare nitric acid-treated activated carbon (NT-AC), activated carbon was treated with 7 N HNO3 at 353 K for 3h, washed with distilled water, and dried at 423 K for 12 h to remove the remaining nitric acid. To remove surface functional groups, activated carbon was pretreated with helium at 1273 K (HT-AC). CMK-5 was prepared following the reported procedure [8]. The catalyst was prepared by a wet impregnation method to impregnate supports with an aqueous solution of palladium. The support was dried at 423 K for 12 h in an oven before impregnation. 88mg of Palladium chloride (Sigma, 99.9%) was dissolved in the mixture of 1ml of 37% HCl solution and 10 ml of acetone. 1 g of support was added to this solution, mixed for 10 h at room temperature, and then evaporated to remove solvent. The temperature programmed desorption (TPD) was conducted in a He stream from 303 K to 1303 K at a ramping rate of 10 K/min monitoring gas products with a mass-selective detector (HP5971 MSD). The X-ray absorption fine structure (XAFS) spectra were taken in a transmission mode for the K-edge of Pd at beamline 7C of the Pohang Light Source (PLS). The measured spectra were analyzed by using ATHENA [9]. CO chemisorption on Pd was carried out at 300 K with a pulse injection method [10]. Result and Discussion

80x103

AC NT-AC HT-AC CMK-5

60x103

40x103

20x103

CO Abundance / a.u.

CO2 Abundance / a.u.

The BET surface area of AC, NT-AC, HT-AC, and CMK-5 was determined to be 850±16m2/g, 848±17m2/g, 901±18m2/g, and 1751±29m2/g, respectively. The fresh activated carbon and modified ones have almost the same BET surface area. Hence, modification has been successfully made without disrupting the original texture of AC. TPD was conducted to find out the qualitative amount of surface oxygen groups on each carbon supports as shown in Fig.1. NT-AC produced the large amount of CO2 and CO during TPD. This can be interpreted that lots of carboxylic and carbonyl groups were formed during the liquid phase oxidation. The concentration of surface oxygen groups appeared to decrease in the order: NT-AC >> AC > CMK-5 > HT-AC. 80x103

AC NT-AC HT-AC CMK-5

60x103

40x103

20x103

0

0

200 400 600 800 1000 1200 1400

200 400 600 800 1000 1200 1400

Temperature / K

Temperature / K

Fig.1 TPD patterns of various carbon supports

Normalized Absorbance / a.u.

Solid State Phenomena Vol. 135

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f

F.T. Magnitude / a.u.

e d c b a

60

f 40

e 20

d c b a

0

243202434024360243802440024420

0

1

Energy / eV

2

3

4

Radial Distance / A

Fig. 2.. Pd K-edge XANES spectra and Fourier transforms of k3-weighted EXAFS oscillations of Pd K-edge for Pd/AC (a), Pd/NT-AC (b), Pd/HT-AC (c), Pd/CMK-5 (d), PdCl2 (e), and Pd foil (f). All catalysts were dried at 373 K before the measurement.

TCD Signal/ a.u.

After PdCl2 was impregnated on each support, Pd K-edge XAFS spectra were obtained as shown in Fig. 2. In X-ray absorption near edge structure (XANES) spectra, Pd was present as PdCl2-like structure for Pd/AC, Pd/NT-AC, and Pd/HT-AC. This was also supported by the extended X-ray absorption fine structure (EXAFS) analysis. The only interaction due to Pd-Cl was observed and its coordination number was almost same to that of PdCl2. The presence of Pd metal has been reported for Pd/C catalyst after an impregnation step [6]. This difference must be due to the presence of hydrogen chloride in the Pd precursor solution in this work. The excess chloride can inhibit the formation of Pd metal in the impregnation step. However, different XANES spectra were observed for Pd/CMK-5, which showed the presence of Pd metal. The linear combination XANES fitting was conducted for Pd/CMK-5 and the weight fraction of Pd metal and PdCl2 was determined to be 0.13 and 0.87, respectively. The presence of Pd metal was also confirmed by the EXAFS analysis. The coordination number of Pd-Cl and Pd-Pd was calculated to be 2.7±0.4 and 0.8±1.4, respectively. This result revealed that CMK-5 had stronger reduction power than did activated carbons. The temperature-programmed reduction (TPR) experiment was conducted to find out the reducibility of PdCl2 on carbon supports. As shown in Fig. 3, the maximum reduction peak was observed around 420 K with a shoulder around 460 K. No noticeable difference in TPR patterns among Pd catalysts supported on different carbon supports was observed.

300

350

400

450

500

550

600

Temperature / K Fig. 3. The TPR pattern of Pd/AC..

650

700

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Nanocomposites and Nanoporous Materials VIII

Table 1 Pd dispersion over supported Pd catalysts reduced at different temperatures. Catalysts

Reduction

Pd

Catalysts

temperature dispersion

Reduction

Pd

temperature dispersion

5%Pd/AC

423 K

60%

5%Pd/HT-AC

423 K

53%

5%Pd/AC

573 K

15%

5%Pd/HT-AC

573 K

12%

5%Pd/AC

773 K

2%

5%Pd/HT-AC

773 K

2%

5%Pd/NT-AC

423 K

57%

5%Pd/CMK-5

423 K

59%

5%Pd/NT-AC

573 K

22%

5%Pd/CMK-5

573 K

29%

5%Pd/NT-AC

773 K

3%

5%Pd/CMK-5

773 K

9%

Table 1 shows the Pd dispersion in supported Pd catalysts reduced at different temperatures. The almost same dispersion was obtained for Pd catalysts supported on different carbons after reduction at 423 K. Among Pd catalysts, Pd/CMK-5 showed the highest Pd dispersion after reduction at high temperatures such as 573 and 773 K. The palladium catalyst supported on the carbon support with a high concentration of surface oxygen groups showed a better dispersion than did other Pd/C catalysts when they were reduced at 573 K. The carboxylic groups on carbon supports appeared to inhibit the sintering of palladium metal particles. Conclusion The Pd dispersion of the carbon-supported Pd catalyst is dependent on the reduction temperature as well as the physical property of carbon support such as its surface area and the concentration of surface oxygen groups. There was no noticeable difference in the Pd dispersion when the Pd catalyst was reduced at a low temperature such as 423 K. The surface oxygen groups inhibit the agglomeration of Pd particles. Compared with other Pd/C catalysts, Pd/CMK-5 was more resistant to sintering of Pd particles even after a high-temperature reduction because of its high surface area. Acknowledgement Experiments at PLS were supported in part by MOST and POSTECH. References [1] E. Auer, A. Freund, J. Pietsch, T. Tacke, Appl. Catal., A, 1998, 173(2), 259. [2] M.L. Toebes, J.A. van Dillen, K.P. de Jong, J. Mol. Catal., A, 2001, 173, 75. [3] H.-U. Blaser, A. Indolese, A. Schnyder, H. Steiner, M. Studer, J. Mol. Catal., A, 2001, 173, 3. [4] J.L. Figueiredo, M.F.R. Pereira, M.M.A. Freitas, J.J.M. Orfao, Carbon, 1999, 37, 1379. [5] M. Gurrath, T. Kuretzky, H.P. Boehm, L.B. Okhlopkova, A.S. Lisitsyn, V.A. Likholobov, Carbon, 2000, 38, 1241. [6] P.A. Simonov, A. V. Romanenko, I.P. Prosvirin, E.M. Moroz, A.I. Boronin, A. L. Chuvilin, V.A. Likholobov, Carbon, 1997, 35(1) 73. [7] S.D. Lin, Y.-H. Hsu, P.-H. Jen, J.-F. Lee, J. Mol. Catal. A: Chem., 2005, 238, 88. [8] S.H. Joo, S.J. Choi, I. Oh, J. Kwak, Z. Liu, O. Terasaki, and R. Ryoo, Nature, 2001, 412, 169. [9] B. Ravel, M. Newville, J. Synchrotron Rad., 2005, 12, 537. [10] H.Y. Song, E.D. Park, and J.S. Lee, J. Mol. Catal. A: Chem., 2000, 154, 243.

Solid State Phenomena Vol. 135 (2008) pp 61-64 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.61

The Effect of Si/Al ratio on Selective Catalytic Reduction of NOx with NH3 over Pt/Al-SBA-15 Min Kang1, Jae Hyun Park2, Eun Duck Park2, Ji Man Kim3, Dae Jung Kim4 and Jae Eui Yie1,a 1

Department of Applied Chemistry, Ajou University, Suwon 443-749 Republic of Korea

2

Division of Energy Systems Research, Ajou University, Suwon 443-749 Republic of Korea 3

Department of Chemistry, Sungkyunkwan University, Suwon 440-746 Republic of Korea

4

Department of Chemistry & Biotechnology, Texas Tech University, Lubbok, TX 79409-1061, USA a

[email protected],

Keywords: NH3-SCR, Pt, Al-SBA-15, aluminum content, low temperature,

Abstract. The selective catalytic reduction of NOx with NH3 (NH3-SCR) was investigated over Pt catalysts supported on various supports such as alumina, ZSM-5, SBA-15, and Al-SBA-15 with different amounts of alumina. Among them, Pt catalysts supported on Al-SBA-15 showed the higher NOx conversion at low temperatures than those of others. These also showed high NOx conversions over a wide reaction temperature. As the Si/Al ratio in Al-SBA-15 decreased, the NH3-SCR activity increased. This was closely related to the amount of strongly-adsorbed NH3. Introduction Selective catalytic reduction (SCR) of nitrogen oxides (NO, NO2, N2O) with NH3 is a well-known reaction of post-combustion technology for the removal of nitrogen oxides from stationary sources and chemical plants [1]. Although V2O5/TiO2-based catalysts have been commercialized as medium temperature (573 ~ 723 K) de-NOx catalysts, there still exist challenges and opportunities to further extend the operation temperature window toward lower (423 ~ 573 K) operating temperature to meet the needs of industrial applications [2]. Catalysts based on transition metal-containing zeolites have been intensively investigated in the first phase, but the strong decrease in the activity in the presence of H2O and/or SO2 in the reactant [3] led to the development of alternative catalysts such as Pt/ZSM-5 [4] and Pt group metals supported on SiO2 and Al2O3 [5]. However, these catalysts have some limitations for the practical application because they have a narrow temperature window showing high NOx conversions. SBA-15 [6] is a mesoporous material, which represents an interesting catalyst support for metal oxides, metals and organometallic compounds due to the combination of good accessibility, uniform pore size and high surface area. However, only a few works have been reported regarding catalysts based on SBA-15 for the application to SCR [7,8]. In order to combine the advantages of zeolites and oxide based systems, we studied the catalytic properties of platinum supported on mesoporous molecular sieves, SBA-15. This study elucidates the effect of alumina contents on the activity of Pt/SBA-15 catalysts for NH3-SCR to widen the operation temperature showing a high NOx conversion. Experimental A mesoporous SBA-15 silica was prepared according to the typical synthetic process reported by literature [9]. Triblock copolymer P123 and 2M HCl were dissolved in water to form a clear solution. Then TEOS was added to that solution under stirring condition then further stirred for 24 h at 313K and then for 24 h at 373K under static condition. The pure-silica materials were aluminated via

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Nanocomposites and Nanoporous Materials VIII

post-grafting technique to prepare Al-SBA-15 (Si/Al = 20, 50 and 100). 1.0 wt% of platinum was loaded onto the Al-SBA-15 by an incipient wetness impregnation method using tetraamineplatium(II) nitrate as the Pt precursor. Subsequently, the Pt-supported catalysts were dried at 373 K, calcined at 773 K for 2 h in air. Catalytic activities were measured over a fixed bed of catalysts in a tubular flow reactor (8 mm I.D.). Reactant gases were fed to the reactor by means of electronic mass flow controller (MKS). The total flow rate of gases is maintained at 200 ml/min at room temperature with a resulting GHSV of 50,000 h-1.. The reactant gas typically consisted of 500 ppm NO, 500 ppm NH3 and 5 vol% O2 in N2 flow. The NOx concentration in the inlet and outlet gas was analyzed by means of a NO/NO2 combustion gas analyzer (Euroton). The steady-state NOx conversion was measured at each reaction temperature. Results and Discussion The nitrogen adsorption-desorption isotherms and the corresponding pore size distributions of the catalysts are shown in Fig. 1. All the catalysts give type IV with hysteresis loops at relative pressures from 0.3 to 0.7. Table 1 summarizes the surface area, total pore volume and pore size of Pt/SBA-15 catalysts. The BET surface area and pore volume of the SBA-15 are slightly decreased with increasing the content of aluminum incorporated. Table 1 and Fig. 1 indicate that the pore structure is nearly unchanged even after the alumination and platinum impregnation.. Table 1. Physicochemical properties of catalysts SBET (m2/g) Pt/SBA-15 598.3 a Pt/Al-SBA-15 (100) 519.8 Pt/Al-SBA-15 (50) 519.3 Pt/Al-SBA-15 (20) 518.5 Pt/ZSM-5 320.7 Pt/Al2O3 109.2 a Numbers in parentheses are Si/Al ratios

VP (cc/g) 1.05 0.89 0.89 0.86 0.21 0.31

DP (nm) 6.8 6.5 6.5 6.3 0.4 6.9

0.6 Pt/SBA-15 Pt/Al-SBA-15(100) Pt/Al-SBA-15 (50) Pt/Al-SBA-15 (20)

500

-1

0.5 0.4

400 0.3 300 0.2

200

0.1

100 0

3 -1

600

dV/dD (cm g nm )

3

Volume adsorbed (cm /g)

700

0.0 0.0

0.2

0.4

0.6

0.8

Relative pressure (P/P 0 )

1.0 0

5

10

15

20

Pore diam eter (nm )

Fig. 1. N2 adsorption-desorption isotherms and the corresponding pore size distributions of the catalysts Pt/SBA-15 and Pt/Al-SBA-15 catalysts

Solid State Phenomena Vol. 135

Pt/Al-SBA-15 (20)

Amount of NH3 adsorbed

Pt/Al-SBA-15 (20)

Intensity (a.u.)

63

Pt/Al-SBA-15 (50)

Pt/Al-SBA-15 (100)

Pt/Al-SBA-15 (50)

Pt/Al-SBA-15 (100)

Pt/SBA-15

Pt/SBA-15

1

2

3

4

5

300

400

500

600

700

800

2θ θ

Temperature (K)

Fig. 2. XRD Patterns of Pt/SBA-15 and Pt/Al-SBA-15 catalysts

Fig. 3. NH3-TPD profiles of Pt/SBA-15 and Pt/Al-SBA-15 catalysts

Fig. 2 shows powder X-ray diffraction (XRD) patterns for the Pt/SBA-15 and Pt/Al-SBA-15 catalysts. XRD patterns for the mesoporous materials indicate that all the materials have well defined 2-d hexagonal mesostructures. It is noteworthy that the mesopore structure is retained after the incorporation of aluminum and platinum impregnation. This is in good agreement with the N2 adsorption-desorption isotherms and the corresponding pore size distributions for the Pt/SBA-15 and Pt/Al-SBA-15 catalysts. The surface acidity is one of the most important properties required for a SCR catalyst. From the NH3-TPD profiles in Fig. 3, it can be seen that the incorporation of aluminum into SBA-15 materials significantly enhanced the surface acidity of the catalyst compared to the original SBA-15. The weak acidity at low temperature (373 ~ 393 K) did not change systematically, while the strong acidity at high temperature (583 ~ 603 K) was increased with increasing the aluminum content of Al-SBA-15. The more aluminum contained on the catalyst surface, the more NH3 adsorbed on the surface. Selective catalytic reduction of NOx with NH3 over Pt/SBA-15, Pt/Al-SBA-15 (Si/Al = 20, 50 and 100) catalysts and Pt/γ-Al2O3 and Pt/ZSM-5 as references were studied in the temperature between 373 and 723 K. The space velocity was 50,000 h-1. The activities of NO reduction as a function of temperature are shown in Fig. 4.. The order of activity of the catalysts is Pt/Al-SBA-15 (20) > Pt/Al-SBA-15 (50) > Pt/Al-SBA-15 (100) > Pt/SBA-15 based on the NOx conversion at all the temperature. It is also the order of the acidity of the catalysts at high temperature. However, Pt/γ-Al2O3 and Pt/ZSM-5 catalysts present a maximum of NOx conversion at 473 K followed by a drastic decrease in NOx conversion at higher reaction temperatures, which is typically observed for Pt based de-NOx catalysts. Pt/Al-SBA-15(20) catalyst showed the highest NOx conversion over a wide reaction temperature from 423 to 723 K.. This improved catalytic performance of these catalysts can be related to the participation of aluminum and the high dispersion of platinum particles in the SBA-15 supports. The acidic sites formed by the incorporation of alumina onto SBA-15 can help the

100

100

80

80

60

60

40

40 Pt/SBA-15 Pt/Al-SBA-15(100) Pt/Al-SBA-15 (50) Pt/Al-SBA-15 (20)

20

20

Pt/Al-SBA-15(20) Pt/γγ-Al2O3

NOx Conversion (%)

Nanocomposites and Nanoporous Materials VIII

NOx Conversion (%)

64

Pt/ZSM-5

0

0 373

423

473

523

573

623

Temperature (K)

673

723 373

423

473

523

573

623

673

723

Temperature (K)

Fig. 4. Catalytic activities of NOx reduction over supported platinum catalysts basic reductant, NH3, adsorb on the surface. This may also explain why the NOx conversions increase at low temperatures as the aluminum content of the catalyst increased. This catalyst showing high NOx conversion over a wide reaction temperature can give us the flexibility in operating De-NOx system. Conclusions Among Pt catalysts supported on various supports such as alumina, ZSM-5, SBA-15, and Al-SBA-15 with different amounts of alumina, Pt catalysts supported on Al-SBA-15 showed the higher NOx conversion at low temperatures than those of others. These catalysts also exhibited high NOx conversions over a wide reaction temperature. The amount of strongly-adsorbed NH3 increases as the Si/Al ratio in Al-SBA-15 decreases, which resulted in the increasing NH3-SCR activity over Pt/Al-SBA-15. References [1] R. M. Heck, Catal. Today, Vol. 53 (1999) p. 519. [2] R. M. Heck, R. J. Rarrauto, Catalytic Air Pollution Control – Commercial Technology (Van Nostrand Princeton, NJ, 1995) [3] H. Y. Chen, W. M. H. Sachtler, Catal. Lett., Vol. 50 (1998), p. 125. [4] M. Iwamoto, H. Yahiro, H. K. Shin, M. Watanabe, J. Guo, M. Konmo, T. Chikahisa, T. Murayama, Appl. Catal. B, Vol. 5 (1994), p. L1. [5] R. Burch, P. J. Millington, Catal. Today, Vol. 26 (1995), p. 185. [6] G. Q. Lu, X. S. Zhao, Nanoporous Material: Science and Engineering (Imperial College Press, London, 2004).. [7] L. Chmielarz, P. Kuśtrowski, R. Dziembaj, P. Cool, E. F. Vansant, Appl. Catal. B, Vol. 62 (2006), p. 369. [8] Y. Segura, L. Chmielarz, P. Kustrowski, P. Cool, R. Dziembaj, E. F. Vansant, Appl. Catal. B, Vol. 61 (2005), p. 69. [9] J. M. Kim and G. D. Stucky, Chem. Comm., (2000), p. 1159.

Solid State Phenomena Vol. 135 (2008) pp 65-68 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.65

Effect of ball-milling method on the formation of ODS Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3 powders J.-H. Ahn1*, B. H. Park1 and J. Jang2 1

Dept. of Materials Engineering, Andong National University, Andong, Gyungbuk 760-749, Korea

2

Korea Atomic Energy Research Institute, Nuclear Materials Technology Development Div. P.P. Box 105, Yusung, 305-600, Daejeon, Korea (* [email protected])

Keywords: Mechanical alloying, High-energy ball-milling, Ferritic stainless steels, Nano-dispersion

Abstract. Mechanical alloying (MA) process has been examined to synthesize ferritic stainless steel powder dispersed with nano-sized Y2O3 particles. A pilot-scale horizontal mill was fabricated and compared with laboratory-scale ball mills and an attrition mill. Horizontal milling resulted in a much better distribution of particle size and dispersoids than other milling methods. Although horizontal milling is considered as a low-energy process requiring long times, processing time could be markedly reduced with increasing the diameter of horizontal mill with a proper control of milling parameters.

1. Introduction Ferritic stainless steels have recently attracted much attention for high temperature applications such as interconnects for oxide fuel cells [1-2] and in-core and out-core materials for the Generation IV nuclear reactors [3-6]. Conventional ferritic stainless steels which are generally strengthen by precipitates such as carbides, have been considered as inadequate materials for high temperature uses due to rapid degradation of mechanical strength. This is mainly due to the coarsening of carbides at elevated temperatures. Therefore, there has been a need for advanced ferritic stainless steels without carbon addition. These new ferritic stainless steels are strengthened by fine oxides instead of carbides and considered as strong candidate materials for the Generation IV nuclear reactors such as the Liquid Sodium Fast Reactor or the Supercritical Water Cooling Reactors. The advanced ferritic stainless steels with oxide dispersion are expected to have much better creep strength than austenitic stainless steels under high neutron irradiation up to 200 dpa [7]. For the fabrication of the advanced ferritic stainless steels with nano-oxide dispersion, conventional casting metallurgy followed by metal working is inadequate due to non-wettability between metals and oxides. One of the most effective methods for the uniform dispersion of fine oxide particles in the matrix is MA process invented by International Nickel Company. However, the efficiency of the process as well as the quality of synthesized powders is strongly dependant on the employed milling method. For instance, attrition milling which is one of the most powerful milling methods and an original technique of MA, is sometimes unsuitable for mass production due to the difficulties in scale-up. Furthermore, this method frequently resulted in the formation of inhomogeneous powders even with a small variation of milling parameters. In the present work, we have fabricated a pilot-scale horizontal mill to employ horizontal ballmilling for mass producyion of MA powders instead of attrition milling. We have synthesized a new ODS (oxide-dispersion strengthened) ferritic steel powder, Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3, named as ANU-1. A comparison of this pilot-scale horizontal mill was made with other mills such

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Nanocomposites and Nanoporous Materials VIII

as laboratory-scale ball mills and an attrition mill. Various milling parameters and milling efficiency were examined. 2. Experimental procedures We have examined four different milling methods for mechanical alloying, using different mills: a Szegvari Attrition Mill (Model 1S) (designated as the mill ‘A’ in this paper) and three home-made horizontal mills, a pilot-scale mill (mill ‘B’) and medium and small size laboratory mills (mills ‘C’ and ‘D’). The details of milling parameters were presented in Table 1, and the photographs of the mill are shown in Fig.1. Both milling vial and balls are made from hardened stainless steels. Using these mills, elemental powder mixture with a composition of Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3. was high-energy ball-milled. The resulting powders were examined by SEM(JF-6300), TEM (EM2010) and particle sized analyzer (Malvern, model: Mastersizer 2000). Table 1 : Milling condition for the preparation of Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3 powders

Fig. 1 : Photograph of employed ball-mills described in Table 1. 3. Results and Discussion The parameters of horizontal ball-milling (B, C and D in Table 1) were chosen to give a friction mode of balls inside a milling vial rather than an impact mode, as indicated in Fig.2. In this mode, the milling of powders occurred mainly at the bottom part of the inner wall of vial. The effectiveness of friction mode during ball-milling was discussed in detail in the previous work[8]. Figure 3 shows the changes in particles size with increasing ball-milling time for various milling methods for the preparation of Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3 powders. To reach a final stage of milling with a mean particle size between 10-14µm, the attrition mill (A) required the shortest time (20 h), while the longest for the small sized horizontal mill (D) (600 h). This time reduced markedly as the diameter of mill increased, and is ~200 h for the pilot scale mill. This is due to the increased weight acting on the balls at the bottom part of vial, facilitating more effectively the friction mode.

Solid State Phenomena Vol. 135

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It is worth noting that milling with austenitic stainless balls such as SUS 304 required a more milling time than that with hardened matrensitic stainless steels which we used in the present work.

Fig.2 : Two modes of milling. (a) Friction and (b) impact mode of balls inside milling vial.

Fig. 3 : Changes in particles size with increasing ball-milling time of Fe-14Cr-2Al-1Si-0.3Ta1Y2O3 powders, prepared by milling methods A~D.

The mean particle size after the final stage of steady state was similar in all cases : 12.2, 11.6, 12.8 and 12.6 µm for ball-milling with the mills A, B, C and D, respectively (Fig. 4). The pilot scale mill (B) produced a slightly finer mean particle size than the other mills. The mill B also resulted in the narrowest distribution of particles size. This can be clearly seen from the SEM microstructures. In all cases, the powder undergoes a typical three stages of mechanical alloying: the initial stage where cold welding of powders is a dominant event, the fracturing-dominant second stage, followed by the final stage of steady state where the both events are balanced and the mean particle size remains unchanged upon further milling. Figure 5 shows the SEM morphology of Fe-14Cr-2Al-1Si0.3Ta-1Y2O3 obtained after the final stage of different ball-milling methods. As shown in the micrographs, the powders prepared by the mill B are more homogeneous than those prepared by the other mills. The attrition milling (A) resulted in the poorest distribution of powder, although it was the most powerful milling method with the shortest processing time. The distribution of nano-sized Y2O3 dispersoids was checked by a TEM for the samples prepared by the hot isostatic pressing (HIPing) of the MA powders at 1100oC for an hour. A typical example of TEM microstructure prepared from the powder processed with the mill B is shown in Fig.6. The micrograph shows that nano-scale Y2O3 particles (5~ 20 nm) are quite homogeneously distributed in the ferritic matrix, showing the effectiveness of the milling.

Fig. 4: Distribution of particle size after the final steady-state of ball-milling of Fe-14Cr-2Al-1Si-0.3Ta-1Y2O3, prepared by different milling methods A~D.

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Nanocomposites and Nanoporous Materials VIII

Mill A

Mill C

Mill B

Mill D

Fig. 5 : SEM image of Fe-14Cr-2Al-1SiFig. 5 TEM micrograph of Fe-14Cr-2Al-1Si 0.3Ta-1Y2O3 after the final stage of different ball -0.3Ta-1Y2O3 after HIPing at1100oC with milling methods. powder processed with the mill B.

4. Conclusions In this short paper, we examined various ball-milling methods to synthesize nao-sizsed Y2O3 dispersed ferritic stainless steels using different mills. A pilot-scale horizontal ball-mill which was fabricated in the present work, was compared with laboratory-scale horizontal mills and an attrition mill (Szegvari Attrition Mill). The horizontal mill with large diameter produced a more homogeneous distribution of particle size and dispersoids than the other mills. Although horizontal milling is considered as a low-energy process requiring long milling times, increasing mill diameter could markedly reduce processing times and enhance the homogeneity of powder. Acknowledgements This work was supported by a grant for the Atomic Energy R&D Program from the Korea Science & Engineering Foundation. References [1] H. Yokokawa, N. Sakai, T. Horita and K. Yamaji, Fuel Celles, 1, No. 2 (2001) 117 [2] S. Elangovan, S. Balagopal, and I. Bay, J.of Materials Engineering and Performance, Vol. 13, No. 3, (2004) 6. [3] M.L. Hamilton, D.S. Gelles, PNLL-13168 (2000) [4] D. J. Edwards, E. P. Simonen and S. M. Bruemmer, J. Nucl. Materials, 317 (2003) 13 [5] D. J. Edwards, E. P. Simonen, F. A. Garner, B. A. Oliver and S. M. Bruemmer, J. Nucl. Materials, 317 (2003) 32 [6] S. Ukai, M. Harada, H. Okada, M. Inoue, S. Nomura, S. Suikakura, K. Asabe, M. Fujiwara, J. Nucl. Mater., 204 (1993) 65 [7] J. J. Fischer, U.S. Patent 4,075,010 issued Feb 21, 1978 [8] J-H Ahn and Y-K Paek, J. Mater. Sci. Lett 18 (1999) p.17-19

Solid State Phenomena Vol. 135 (2008) pp 69-72 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.69

Nano-sized yttria dispersed ferritic stainless steels for SOFC interconnect applications J.-H. Ahn1*, H.-J. Kim2, I.-H. Oh2, B.H. Park1, S.-H. Jang2 1

Andong National University, Dept. of Materials Engineering, Andong, Gyungbuk 760-749, Korea 2

Korea Institute of Industrial Technology

992-32 Techno-Park Song-do, Dongchun-dong, Incheon 406-130, Korea (*[email protected])

Keywords: Solid oxide fuel cells, Interconnects, Stainless steels, Ferritic steels, Nano-dispersion, High-temperature alloys, Electrical conductivity

Abstract. In the present work, we have developed two ferritic stainless steels, Fe + 20%Cr + 0.8%Mn + 0.2%Ti + 0.5%Y2O3 and Fe + 17%Cr + 0.1%C + 1%Si + 1%Mn + 0.5%Y2O3 for oxide fuel cell (SOFC) interconnects. Nano-sized Y2O3 particles were dispersed in the metallic matrix by high-energy ball-milling. The alloys exhibited excellent electrical conductivity and thermal expansion coefficient suitable for SOFC applications. The finely dispersed nano-sized yttria is thought to be attributed to high conductivity.

1. Introduction The interconnect is a critical component in SOFCs. Its roles are the electrical connection between cells and the physical separation between air and fuel in the cell stack. The interconnect material must have compatibility with all of the cell components as well as stability with respect to oxidizing and reducing gases at high temperatures. So far, ceramic based interconnect materials such as La1-x(Sr, Ca)xCrO3 have been used in conventional SOFCs which operate at high temperatures near 1000oC. Recently, progress has been made to develop advanced SOFCs operating at intermediate temperatures (~800oC) with the cell capacity and durability equivalent to hightemperature SOFCs [1-2]. This enables SOFCs to use metallic interconnects which have many advantages over conventional ceramic interconnects such as lower cost of raw materials and fabrication, better electrical and thermal conductivity and easier shape-forming ability, etc [3-6]. Among various candidate materials, ferritic stainless steels have been considered as one of the most promising interconnect materials for SOFCs operating at temperatures near 800oC. However, several problems such as low conductivity due to the formation of oxide layer at high temperatures, poor durability with cell components and thermal mismatch between ceramic electrodes and metallic interconnects should be improved for practical application of the metallic interconnects. For this, Honegger at al. [7] suggested that a small addition of oxides of rare earth elements may improve the performance of metallic interconnects. It is expected that a small addition of Y2O3 enhances the ionic conductivity, chemical durability at high temperatures, thermal stability and mechanical properties of interconnects. In the present work, we have developed new ferritic stainless steels dispersed with nano-sized yttria. One of the the objectives of the present work is to verify the improvement of properties by nano-sized yttria dispersion in ferritic stainless steels as SOFC interconnects.

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Nanocomposites and Nanoporous Materials VIII

2. Experimental precedures We have prepared two alloy compositions : Fe-20Cr-0.8Mn-0.2Ti-0.5Y2O3 (Alloy A) and Fe17Cr-0.1C-1Si-1Mn-0.5Y2O3 (Alloy B). To disperse nano-sized Y2O3 particles (~20 nm) in the metallic matrix, ball-milling was employed. Elemental powders with high purity were mixed and ball-milled in an Ar atmosphere using a horizontal ball mill. The size of the milling jar was Φ 13 x 16mm. Both the jar and milling ball (Φ 6.3) were SUS 304 stainless steel. The weight ratio of ball to powder was 50:1 and the rotating speed of jar was 86 r.p.m. The ball-milled powders were hotconsolidated by conventional sintering or hot pressing. Hot pressing was carried out at 1000°C for an hour at a pressure of 14 MPa in an Ar atmosphere. For comparison, some powders were sintered at 1250°C for 90 minutes. The coefficient of thermal expansion (CTE) of the alloys was measured from room temperature to 1000oC (5°C/min) in air by a Shimadzu TMA-60H analyzer. For electrical resistance measurement by a four-probe method, the alloy plates (10x10x2 mm) were preoxidized in air at 800°C(5°C/min) for 100h. Both surfaces of pre-oxidized alloy plate were then covered by Pt paste, followed by pressing Pt meshes on the top of the plates as current collector to which Pt wires were spot-welded. 3. Results and discussion The powder ball-milled to the final stage of steady-state had a mean particle size of about 13µm. Only broadened crystalline peaks of Fe appeared by X-ray diffraction after ball milling, indicating that additive alloying elements were completely dissolved in ferritic solid solution and/or in amorphous phases. After, hot consolidation such as sintering or hot pressing, a phase separation took place as shown in Fig.1. Both dark and bright phases in the microstructure were α-ferrite with b.c.c. structure, but the darker phase was a chromium-richer phase. Much finer microstructure and less phase segregation were observed in Alloy A than Alloy B. Sintered density in both cases was less than 80 % theoretical density due to low green density caused by a poor compressibilty of work-hardened MA powder. Hot-pressing at 1000oC for an hour resulted in a slight improvement in density. Although consolidated density was not high, such moderately porous alloys may have a beneficial effect to buffer and minimize thermal cracks which may produce at the interface between the alloy and oxide layer (scale) during heating and cooling. A

Hot pressed

Density : 86% Hardness : 250Hv

B Density: :>80% Density below80% Hardness Hardness : :370Hv 370Hv

(1000C, 1 h, 14 MPa)

Sintered (1250 ℃, 90m, Ar) 1㎛

x 4000

Fig. 1: SEM microstructures of hot-pressed or sintered Fe-20Cr-0.8Mn-0.2Ti-0.5Y2O3 (Alloy A) and Fe-17Cr-0.1C-1Si-1Mn-0.5Y2O3 (Alloy B).

Solid State Phenomena Vol. 135

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In fact, Alloy A which is denser than Alloy B developed more pronounced cracks at the interface. (Fig.2) In the micrographs, the dark area to the right side is polymeric resin used for sample preparation. Although the cause of the observed cracks is under investigation whether they were produced during the drying of resin or during the heating-cooling cycle, the appearance of less abundant cracks in Alloy A indicates a possible beneficial effect of crack inhibition at the oxide layer/alloy interface for slightly porous alloys. Fig. 3 shows thermal expansion behaviour of two alloys measured from room temperature to 800oC at a heating rate of 5oC/min. Alloy A exhibited unstable (non-linear) plot line compared to Alloy B. This is thought to be related to abundant pores of Alloy A. In both alloys, the values of coefficient of thermal expansion (CTE) were compatible with those of YSZ (yttria stabilized zirconia)-based electrolytes or oxide electrodes, showing a suitability of these alloys for SOFC interconnects. The values at 800oC were 12.99x10-6 and 12.99x10-6/K for Alloy A and B, respectively. Conventional La-based ceramic interconnects have CTEs between 9.5-13.1 x10-6/K . [8] Fig.4 shows the change in electrical conductivity measured from room temperature to 800oC. Because the electrical conductivity of these alloys is determined by that of their oxide layers (scale) at the surface, the alloys were pre-oxidized at 800oC for 100 hours to form oxide layer. As shown in

Fig. 2: SEM microstructures of surface oxide scales in Fe-20Cr0.8Mn-0.2Ti-0.5Y2O3 (alloy A, left) and Fe-17Cr-0.1C-1Si-1Mn0.5Y2O3 (alloy B, right) after oxidation in air at 800oC for 100 h.

25℃~1000℃ 12.99x10-6/K

25℃~1000℃ 11.33x10-6/K

Fig. 3: Dependence of coefficient of thermal expansion on temperature in Fe-20Cr-0.8Mn-0.2Ti0.5Y2O3 (alloy A) and Fe-17Cr-0.1C-1Si-1Mn-0.5Y2O3 (alloy B).

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Nanocomposites and Nanoporous Materials VIII

4

A (Fe-20Cr-0.8Mn-0.2Ti-0.5Y2O3) B (Fe-17Cr-0.1C-1Si-1Mn-0.5Y2O3)

Fig. 4: Dependence of electrical conductivity of Fe-20Cr-0.8Mn-0.2Ti0.5Y2O3 (alloy A) and Fe-17Cr-0.1C-1Si1Mn-0.5Y2O3 (alloy B) on temperature.

Conductivity (S/cm)

3

2.54 S/cm 2

1

0.6 S/cm

0 0.8

1.0

1.2

1.4

1.6

1.8

-1

1000/T (K )

Fig.4, the electrical conductivity of Alloy A is much higher than that of Alloy B at high temperatures (left side values in the figure). This might be due to silicon-rich oxide layer at the surface in the Si-containing Alloy B. It is known that the silicon oxide has lower conductivity than Cr-based oxides. The morphology of surface oxide can be seen in Fig.2. The oxide layers consists of two parts, one Cr-rich inner layer and Si-rich outer layer. We have initially added Si to enhance oxidation resistance of alloy, but it degraded electrical conductivity. In spite of that both Alloys A and B exhibit relatively high electrical conductivity at 800oC: 2.54 and 0.60 S/cm for Alloy A and B, respectively. These values are higher than those of conventional ferritic stainless steel [8] without Y2O3 nano-dispersion. 4. Conclusions In the present work, we have developed two ferritic stainless steels, Fe + 17%Cr + 0.1%C + 1%Si + 1%Mn + 0.5%Y2O3 and Fe + 20%Cr + 0.8%Mn + 0.2%Ti + 0.5%Y2O3 for SOFC interconnects. Ball-milling was employed to disperse Y2O3 nano-sized. The hot-consolidated alloys exhibited excellent electrical conductivity and thermal expansion coefficient suitable for SOFC applications. The finely dispersed nano-sized yttria might be attributed to the observed high conductivity.

Acknowledgements: This work was supported by a grant from the Korea Institute of Industrial Technology (Core Technology 06-006, Gwangju). References [1] S. de Souza, S.J. Visco, L.C. De Jonghe, Solid State Ionics, 98 (1997) 57. [2] S. de Souza, S.J. Visco, L.C. De Jonghe, J. Electrochem. Soc. 144 (1997) L35. [3] H. Ishihara, H. Matsuda, Y. Takita, J. Am. Chem. Soc. 116 (1994) 3801. [4] M. Feng, J.B. Goodenough, Eur. J. Solid State Inorg. Chem. (1994) 663. [5] P. Huang, A. Petric, J. Electrochem. Soc. 143 (5) (1996) [6] K.Q. Huang, R. Tichy, J.B. Goodenough, J. Am. Ceram. Soc. 81 (1998) 2565. [7] K. Honegger, A. Plas, R. Diethelm, W. Glatz, Solid Oxide Fuel Cells VII, eds. H. Yokohama et al. The Electrochem. Soc. Proc. Pennington, NJ, PV2001-16 (2001) p.803 [8] H. U. Anderson et al., High Temperature Solid Oxide Fuel Cells, eds. S.C. Singhal et al., Oxford UK (2004) p.177

Solid State Phenomena Vol. 135 (2008) pp 73-76 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.73

Methane storage on surface modified activated carbons

Seok-Min Yun1,a, Ju-Wan Kim2,b , Hangkyo Jin3,c, Young-Ho Kim1,d , and Young-Seak Lee1,e , * 1

Department of Fine Chemical Engineering and Chemistry, Chungnam National University, Daejeon 305-764, Korea

2

Department of Nanotechnology, Chungnam National University, Daejeon 305-764, Korea

3

Korea Research Institute of Chemical Technology, P.O. Box 107, Daejeon 305-600, Korea a

[email protected], [email protected], [email protected], [email protected], e,

*[email protected]

Abstract. In this study, methane gas absorption ability of activated carbon (AC) with surface functional group effect, was evaluated after nitric acid and urea treatment of AC surface. Specific surface area and pore distribution of AC were studied through nitrogen absorption isotherm at 77 K. Micro pore volume was calculated through H-K method. Absorbed methane amount was evaluated through volumetric method at room temperature by using auto absorption apparatus. Absorbed methane amount of AC was found to increase with specific surface area increase. Correlation proposed between the methane absorption amount and surface nature indicates that the surface nature plays an important role on the absorption amount at a given temperature. Keywords: Activated carbons, methane absorption, surface modification, porosity

Introduction The interest in natural gas (NG) as a possible alternative for transportation fuel has grown considerably since 1980s [1]. Worldwide effects have been made to make natural gas vehicles (NGVs) competitive with current ones using conventional fuels. The reason is that NG is much cheaper than conventional petroleum-based gasoline natural abundance and especially known for its clean burning [2, 3]. Many researchers reported absorption and storage of NG, methane in different porous materials, including activated carbon, carbon nanotubes and activated carbon fibers with different specific surface areas and pore sizes [4]. Activated carbon develops micropore, during activation process due to which oxygen functional groups are introduced on carbon surface. In general, gas absorption capacity depends not only on porosity of absorbent, but also on their surface nature. Surface nature of activated carbon can be changed by various techniques, such as chemical treatment, heat treatment, fluorination and plasma treatment etc. We assume that interactions between surface of absorbent and gas molecules are changed by surface modifications. In this study, we used AC modified by chemical treatment using nitric acid and urea for methane storage. Experimental study was conducted in terms of absorption isotherm of N2 at 77K. The absorption isotherms of methane from 263 K to 298 K were investigated. The relationships between the absorption capacity with respect to temperature and the changed surface nature and specific surface area (SSA) of the ACs are discussed.

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Nanocomposites and Nanoporous Materials VIII

Experimental Sample preparation. Three carbon samples were investigated in this study for methane storage. The raw materials for surface modification were coconut based granular type activated carbons, purchased from Yu-rim carbon chemical Co, Korea. Before surface modification, the as such received activated carbon (raw) was washed with distillated water to remove impurities, and dried in oven at 383 K for 24 h. and named as R-AC. The R-AC was then chemically treated by nitric acid and urea. In order to introduce oxygen functional groups, activated carbons were modified in 1 M nitric acid solution at 373 K for 1 h. After acid treatment, residual nitric acid was removed by distillated water washing and dried at 383 K for 24 h. Acid treated activated carbon was marked as A-AC. Urea can endow with basic functional groups to surface of activated carbon. 1 M urea solution was prepared, and R-AC was impregnated in urea solution for 6 h by mixing in shaking machine. Heat treatment of urea-treated activated carbon was carried out in tube furnace at 723 K for 1 h. This sample was named as U-AC. Analysis. The nitrogen absorption isotherms of the three activated carbons were measured on ASAP 2020 (Micromeritics Ins. Corp.) at liquid nitrogen temperature. All samples were degassed at 423 K for 3 h before analysis. The specific surface area and pore structure were obtained by Langmuir equation and Horvath-Kawazoe method. The changes of surface functional groups and total acidity of activated carbons after modification were investigated by Boehm titration. One gram activated carbon was mixed with 100 ml 0.1 N NaHCO3, Na2CO3 and NaOH solution, respectively. The flasks were covered with parafilm and shaken for 48 h at 200 rpm. The solutions were filered by membrane filter (Ø = 0.45 µm) and titrated with 0.1 N HCl solution. Total basicity of samples was obtained by NaOH titration. The method was same with Bohem titration except that 0.1 N HCl solution mixed with samples and titrated with 0.1 N NaOH solution. Further, pH change of samples was measured. To determine the contents of C, H, O and N in activated carbons used in this study, elemental analysis was carried out with EA1110 (CE Instrument, Italia). Methane absorption. All samples were out-gased at 473 K for 6 h and then methane absorption isotherms were obtained from volumetric method using auto absorption apparatus (Autosorb-1, Quantachrome, USA) at 263, 273 and 298 K. Results and discussion Characterization of ACs by absorption isotherm of N2 at 77K. Nitrogen absorption isotherms at 77 K on the three activated carbon samples are shown in Fig. 1. Detailed information on the textural properties of the samples is presented in Table 1. All of isotherms of samples represent type I (Langmuir type) in IUPAC classification. In case of physical absorption, Langmuir isotherm pattern indicates that the nature of material are microporous ( Fig.1 and Table 1)

Amount adsorbed (mmol/g)

16

R-AC A-AC U-AC

12

8

0.0

0.2

0.4

0.6

0.8

1.0

Relative pressure ( P/Po)

Fig. 1. Nitrogen absorption isotherm of activated carbons.

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Table 1. Pore textural characteristics of activated carbons a

Sample

S (m2/g)

b

Vt (cm3/g)

c

Vm (cm3/g)

d

Wa (nm)

R-AC

1249

0.46

0.44

0.74

A-AC

1245

0.47

0.44

0.74

U-AC

1310

0.48

0.46

0.73

a: Langmuir specific surface area,

b: Total pore volume at P/Po = 0.98

c: Micropore volume by H-K method,

d: Average pore width = 2Vt/S

Boehm Titration. Table 2 furnishes the Boehm titration results for the surface properties of the ACs used in this study. It can clearly be seen that the chemical treatments influence the pH, acid– base surface values of the ACs. This seems to be a consequence of the changes of functional group of the chemically treated ACs. As presented in Table 2, the pH, base surface value of the A-AC sample show prominent changes compared with R-AC and U-AC. This result indicates that a strong acid–base reaction is occurred on the original base-like carbon surface. By contrast, U-AC shows an increased base value. Table 2. pH and acid-base values of activated carbons Samples

pH

Carboxyl (meq/g)

Lactonic (meq/g)

Phenolic (meq/g)

Total acidity (meq/g)

Total basicity (meq/g)

R-AC

7.66

0.01

0.12

0.73

0.86

0.31

A-AC

4.34

0.47

0.26

1..31

2.04

0.11

U-AC

8.58

0.09

0.02

0.47

0.57

0.54

Elemental analysis. Table 3 lists the different elemental analysis results of activated carbons treated by various methods. Decrease in carbon content with corresponding increase in oxygen value of AC is noticed may be due to introduction of oxygen by HNO3. Table 3. Elemental analysis of activated carbons C(wt%)

H(wt%)

N(wt%)

Oa(wt%)

O/C

H/C

N/C

R-AC

93.6

0.5

0.01

5.89

0.06

0.005

0.0001

A-AC

88.2

0.6

0.43

10.77

0.12

0.007

0.0049

U-AC

94.3

0.2

0.27

5.23

0.06

0.002

0.0029

Sample

a: The oxygen is assessed by difference Methane absorption isotherms on ACs. Fig. 2 illustrates methane absorption isotherms of samples at various temperatures. From the experimental results, it is found that the methane absorption amount increases monotonically with pressure at the three temps studied, which is quite similar to the Langmuir isotherms given by activated carbon [3]. Methane absorption amounts of UACs were the largest at all temperatures whereas A-ACs was the least of samples on all of the absorption temperature. This result suggests that the methane absorption amount depend strongly on surface nature. U-ACs has small acid and large basicity. To investigate the relations of surface nature and methane absorption amount, we calculated the absorbed amount of methane versus total acidity and basicity, and depicted in Fig. 3, respectively. As the correlation coefficient in Fig. 3, the amount of methane absorption is more dependent on basicity than acidity onto ACs. The increase in oxygen of AC leads to less storage for methane.

Nanocomposites and Nanoporous Materials VIII

48

32

(b)

R-AC A-AC U-AC

60

3 Adsorbed amount ( cm /g)

3 Adsorbed amount ( cm /g)

(a)

R-AC A-AC U-AC

40

48

36

24

36

24

16

24

12

8

12

0

0 0

20

40

60

80

0 0

100

(c)

R-AC A-AC U-AC

60

3 Adsorbed amount ( cm /g)

76

20

40

P (KPa)

60

80

100

0

P (KPa)

12

24

36

48

60

P (KPa)

Fig. 2. Methane absorption isotherm of activated carbons: (a) 298 K, (b) 273 K, and (c) 263 K.

64

3 Adsorbed amount ( cm /g)

acidity basicity

56

y = 46.5x + 35.7 R2 = 0.9997

48

y = -8.3x + 46.2 R2 = 0.8348

40

32

0.0

0.5

1.0

1.5

2.0

acid-base values (meq/g)

Fig. 3. Amount of methane absorbed according to acid-base values. Conclusions Absorbed methane amount was evaluated through volume method at room temperature (298 K) by using auto absorption apparatus. Absorbed methane amount of AC was increased according to specific surface area. Correlation proposed between the methane absorption amount and surface nature indicates that the surface nature plays an important role on the absorption amount at a given temperature. The amount of methane absorption is dependent on total basicity rather than acidity. Acknowledgment This work was supported by the Korea Research Foundation Grant funded by the Korean government (MOEHRD, Basic Research Promotion Fund) (KRF-2005-005-J00403). Reference [1] D. Lozano-Castello, J. Alcaniz-Monge, M.A. de la Casa-Lillo, D. Cazorla-Amoros and A. Linares-Solano: Fuel Vol. 81 (2002), p. 1777. [2] J. Alcaniz-Monge, M.A. Dela Casa-Lillo, D. Cazorla-Amoros and A. Linares-Solano: Carbon Vol. 35 (1997), p. 291. [3] X. Shao, W. Wang, X. Zhang: Carbon Vol. 45 (2007), p.188 [4] X. R. Zhang and W.C. Wang: Fluid Phase Equil. Vol.194 (2002), p. 289.

Solid State Phenomena Vol. 135 (2008) pp 77-80 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.77

Nickel decoration on multi-walled carbon nanotubes using multi-step impregnation method Shin Dong Kim1,a, Soo Jin Park2,b, Young Seak Lee1,c, * 1

Dept. Fine Chemical Engineering and Chemistry, Chungnam National University, 220, Gung-dong, Yuseong-Gu, Daejeon 305-764, Korea 2

Dept. Chemistry, Inha University, 253, Nam-Gu, Incheon 402-751, Korea a

[email protected], [email protected], c,* [email protected]

Keywords: multi-walled carbon nanotubes, nickel decoration, multi-step impregnation, dispersion

Abstract In this work, nano-sized nickel particles were dispersed on multi-walled carbon nanotubes using multi-step impregnation method, to use them as hydrogen storage media. The dispersion degree of nickel particles on multi-walled carbon nanotubes is inversely proportional to the nickel concentration of solution. It was observed that the low nickel concentration is efficient to decorate nickel particles into the inner space. Multi-step impregnation method of MWNTs through several times with low nickel concentration is more efficient to manufacture Ni-MWNTs having well dispersed metallic nickel particles.

Introduction Hydrogen is one of the cleanest and idealized energy sources. Hydrogen storage is widely recognized as a critical enabling technology for the successful commercialization and market acceptance of hydrogen powered vehicles. However, there are three primary barriers that must be overcome to enable industry commercialization of hydrogen fuel cell vehicles: (1) on-board hydrogen storage system are needed that allow a vehicle driving range of greater than 300 miles (500 km) while meeting vehicle packaging, cost and performance requirements; (2) fuel cell system cost must be lowered to $30 per kilowatt by year 2015 while meeting performance and durability requirements and (3) the cost of safe and efficient hydrogen production and delivery must be lowered to be competitive with gasoline (a target of $2.00 to $3.00 per gasoline equivalent, delivered, untaxed, by year 2015) independent of production pathway and without adverse environmental impact [1]. In order to increase the hydrogen storage capacity of carbon materials, several t studies were carried out all over the world. Especially, carbon nanotubes (CNTs), having high surface-to-volume ratios, are ideal for fast kinetics due to their reversible characteristics during hydrogen charging/discharging. Unfortunately, there was essentially no hydrogen stored in pure single walled nanotubes (SWNTs) at room temperature, while metal alloy-doped SWNTs were observed to store 2-3 wt% hydrogen [2]. Kim and co-workers reported 6 wt% nickel nanoparticle-dispersed multiwalled carbon nanotubes (MWNTs) released 2.8 wt% hydrogen amount [3]. In this paper, nickel-loaded MWNTs were prepared and characterized by TEM, XRD and EDS, in order to investigate the effect of impregnation time on particle size and impregnation amount of nickel on MWNTs. Also, effective impregnation method for nickel loading on MWNTs is discussed.

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Nanocomposites and Nanoporous Materials VIII

Experimental In order to decorate nickel particles on MWNTs [4], two methods were used. First method is that 0.5 g of MWNTs treated with acids [5] was immersed in 110 mM nickel nitrate solution of 20 ml. This method was named as one-time impregnation method (Ni-MWNTs: 50-1). This solution was sonicated for 1 h at room temperature for good dispersion, and then, was evaporated in oil evaporator until solution was perfectly vaporized. After vaporization, vaporized powder was dried at 383 K overnight. On the other hand, Second method is that 0.5 g of MWNTs immersed into 22 mM nickel nitrate solution of 20 ml. This solution was sonicated for 1 h at room temperature for good dispersion, and then, was evaporated in oil evaporator until solution was perfectly vaporized. After vaporization, vaporized powder was dried at 383 K overnight. This procedure was repeated with four times. This method was named as five-time impregnation method (Ni-MWNTs: 10-5). After drying, the samples were reduced in flowing 5 % H2/Ar as the reducing gas, during which reduction conditions were first heated by 5 K/min up to 773 K and kept for 3 h, respectively. To estimate the crystallinity of the samples, X-ray diffraction (XRD, Rigaku, D/Max III) analysis was used. To investigate the quantity of nickel, energy dispersive spectrometry (EDS) were used. Field emission transmission electron microscopy (FE-TEM, JEOL model JEM-2100F), installed at Korea Basic Science Institute (KBSI), was used to measure the wall thickness and the particle size of nickel. A small quantity of MWNTs was dispersed in ethanol, a drop of which was then deposited on a 400 mesh carbon-coated copper grid which was then allowed to dry at room temperature prior to imaging. Results and discussion

(a)

(b)

Fig. 1. Low resolution TEM images of Ni-MWNTs: (a) 50-1, (b) 10-5. Fig. 1 shows TEM images of Ni-MWNTs obtained from 1 time impregnation of 110 mM denoted as 50-1 and 5 times impregnation of 22 mM denoted as 10-5. As can be seen in Fig. 1, nickel particles Ni-MWNTs (50-1) are more aggregated than that of 10-5. It may be due to the high initial nickel concentration of 50-1. In order to investigate more specifically, TEM images were magnified. As shown in Fig. 2 (b), in case of Ni-MWNTs (10-5), small nickel particles were observed in the inner wall of MWNTs, interestingly. Also, Kim and co-workers suggested that nano-sized nickel particle were more efficient to store hydrogen on MWNTs [3]. This result suggests that impregnation method using low nickel concentration through several times can penetrate nickel partilces into the inner wall.

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In order to investigate the rough efficiency of nickel content on MWNTs, EDS spectra were recorded and displayed in Fig. 3. It is shown that the intensity of Ni-MWNT (10-5) is higher than that of 50-1. This result supports the observations make by high resolution TEM images (cf. Fig. 2).

(b)

(a)

CuKa

600

CKa

600

400

0

NiKb CuKb

NiKa

200

CuKa

NiKb CuKb

400

NiKa

200

CuKa

800

(b)

OKa

1000

CuKa

Counts

800

(a)

OKa

1000

CKa

Fig. 2. High resolution TEM images of Ni-MWNTs: (a) 50-1, (b) 10-5.

0 1

2

3

4

5

6

Energy (keV)

7

8

9

10

1

2

3

4

5

6

7

8

9

10

Energy (keV)

Fig. 3. EDS spectra of Ni-MWNTs: (a) 50-1, (b) 10-5. Fig. 4 shows the XRD patterns obtained from Ni-MWNTs. Prior to measuring XRD, NiMWNTs were reduced at 773 K for 5 h with 5 % hydrogen stream balanced with argon. As shown in Fig. 4, it is possible to observe two diffraction peaks corresponding to the (1 1 1) to (2 0 0) planes of metallic nickel at 44.5 and 52.5 o (JCPDS, 65-2865), respectively. In addition, from the inset of these XRD patterns, it is clear that impregnated nickel particles on MWNTs were thoroughly reduced to metallic nickel due to the intensity of 44.5 and 52.5 o. The intensity of Ni-MWNTs (501) is much higher than that of Ni-MWNTs (10-5), which means agglomerated nickel or large particles reflected x-ray much more than that of dispersed nickel particles on Ni-MWNTs (10-5 and 10-3). It is reasonable that the peak of Ni-MWNTs (10-3) is lower than that of Ni-MWNTs (10-5), since total nickel amount decorated on MWNTs was low.

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Nanocomposites and Nanoporous Materials VIII

(a)

Intensity (a.u.)

(b) (c) 43

44

45

46

2 Theta (deg.)

V

(a) O

I

(b)

(c) 10

20

30

40

50

60

2 Theta (deg.) Fig. 4. XRD patterns of Ni-MWNTs: (a) 50-1, (b) 10-3, (c) 10-5 (↓: Ni, V, │: Ni2O3 , □:graphite Carbon , O: NiOOH ) and inset : XRD patterns of Ni-MWNTs from 43 to 46. Summary In this study, nano-sized nickel particles were successfully impregnated on the surface and inner wall of MWNTs using multi-step impregnation method. Multi-step impregnation method using low concentration of nickel solution in this study was proved to be a good method to penetrate the solution containing nickel cation into the inner space of MWNTs. This indicates that it is difficult to penetrate the solution containing nickel cation into the inner space of MWNTs by one time. Metallic nickel can adsorb hydrogen molecules. Therefore, Ni-MWNTs of 10-5 may be effective as a good hydrogen storage media. Acknowledgement This study was performed for the Hydrogen Energy R&D Center, one of the 21st Century Frontier R&D Program, funded by the Ministry of Science and Technology of Korea. References [1] S. Satyapal, J. Petrovic, C. Read, G. Thomas and G. Ordaz: Catal. Today Vol.120 (2007), p.246 [2] Information on http://www.hydrogen.energy.gov/annual_progress04_storage.html#carbon [3] H.S. Kim, H. Lee, K. S. Han, J.H. Kim, M.S. Song, M.S. Park, J.Y. Lee and J.K. Kang: J. Phys. Chem. B Vol.109 (2005), p.8983 [4] S.H. Cho, D.Y. Kim, J.K. Heo, Y.H. Lee, K.H. An, S.D. Kim and Y.S. Lee: Carbon Science Vol.7 (2006) p.267 [5] S.D. Kim, J.W. Kim, J.S. Im, Y.H. Kim and Y.S. Lee: J. Fluorine Chemistry Vol.128 (2007) p.60 [6] JCPDS (2002) No. 65-2865, 06-0140, 75-1621, 14-0481

Solid State Phenomena Vol. 135 (2008) pp 81-84 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.81

Preparation and characterization of electrospun carbon nanofibers with Na2CO3/H3PO4 activation Ji-Sun Im1,a, Soo Jin Park2,b, and Young-Seak Lee1,3,c, * 1

Department of Nanotechnology, Chungnam National University, Daejeon 305-764, Korea 2

Department Chemistry, Inha University, 253, Nam-Gu, Incheon 402-751, Korea

3

Department of Fine Chemical Engineering and Chemistry, Chungnam National University, Daejeon 305-764, Korea a

[email protected], [email protected], c,*[email protected]

Keywords: Carbon nanofiber, Electrospinning, Chemical activation, Na2CO3, H3PO4

Abstract. In this work, carbon nanofibers(CNFs) were prepared by using electrospinning method. Phosphoric acid and sodium carbonate activation of CNFs were conducted to increase surface area and pore volume. Pore structures of activated CNFs were developed with increasing surface area and pore volume through activation. Specific surface area increased about 60 times and total pore volume developed around 120 times. Activated CNFs have different pore distribution with different chemical agent. Introduction Activated carbon materials have been used extensively in industrial purification and chemical recovery operations due to their large specific surface area and pore volume [1]. Chemical activation is an effective method to prepare activated carbons with high specific surface area among a wide variety of activation methods. Chemical activation is characterized by an advantage with so simple procedure of activation. Chemical activation can lead to unique pore structures according to chemical agents. In this paper, pore structure was studied by comparing Na2CO3 and H3PO4 activated CNFs. Experimental Polymer solution was prepared by dissolving polyacrylonitrile(PAN, d = 1.184, 181315, Aldrich) in N,N-Dimethyl formamide(DMF, d = 133, 766137, Fisher). Electrospun fibers were obtained from polymer solution by using electrospinning method. Polymer solution was ejected from tip to collector by high electric voltage. Electrospinning was carried out with the following experimental conditions [applied voltage : 18 kV, tip to collector distance(TCD) : 13 cm, needle : 18 gage (inner diameter 1.27 mm), collector rpm : 200, and syringe pump rate : 1 ml/h] [1]. Before carbonization, oxidation is necessary to change the thermoplastic character into thermosetting character, because electospun materials can not keep their fiber form at high temperature. They would be soften or melt [2]. Oxidization step was carried out at 250 °C for 8 h. Carbonization was carried out with the following conditions [heating rate : 10 °C/m, temperature : 1050 °C, and holding time : 1 h]. This sample was named as R.

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CNFs(R) were chemically activated by using Na2CO3/H3PO4 solutions. CNFs(R) were immersed in six solutions [Na2CO3(1M and 2M) and H3PO4 (20, 40, 60,and 80 wt%)] for 8 h with shaking. These six samples were named N 1, N 2, H 20, H 40, H 60, and H 80. These wet CNFs were transferred to alumina boat for activation. Chemical activation was conducted at 750 °C for 3 h in nitrogen atmosphere. Activated CNFs were washed with distilled water several times and dried over night. Results and Discussion Fig. 1 shows the morphology of all CNFs. In case of non-activated CNFs shown in Fig. 1(a), fibers are straight having the diameter of fiber about 250 nm. After activation, fibers are curved with slightly decreased diameter of fibers due to following chemical activation mechanism [3, 4]. Carbon monoxide and carbon dioxide generated from carbon of fibers. (M = Na) 6MOH + C ↔ 2M + 3H2 + 2M2CO3 M2CO3 + C ↔ M2O + 2CO M2CO3 ↔ M2O + CO2 2M + CO2 ↔ M2O + CO M2O + C ↔ 2M + CO

(a)

(d)

(b)

(e)

(1) (2) (3) (4) (5)

(c)

(f)

(g)

Fig. 1. SEM images of samples; (a) : R, (b) : N 1, (c) N 2 ,(d) : H 20, (e) : H 40, (f) : H 60,(g) : H 80. The summary of textural properties is presented in Table 1. Specific surface area and pore volume were increased after activation. In case of sodium carbonate activation, BET specific surface area of N 2 was increased about 9 times from 11.8 up to 99.5 m2/g, comparing with non-

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activated sample R. After phosphoric acid activation, BET specific surface area of samples(H 20, H 40, H 60,and H 80) was increased from 282.1 up to 602.1 m2/g. Pore volume was increased also through activation. Sample H 80 has the total pore volume showing 0.483 cc/g. All of activated CNFs are mesoporous as shown the micro pore volume percentage of total pore volume less than 13%. Table 1. The summary of textural characteristics of CNFs R

N1

N2

H 20

H 40

H 60

H 80

VT (cc/g)

0.004

0.133

0.245

0.213

0.271

0.329

0.483

VM (cc/g)

0.001

0.004

0.007

0.011

0.017

0.024

0.061

VM/VT X 100 (%)

25

3.01

2.86

5.16

6.27

7.29

12.63

2

11.8

46.8

99.5

282.1

357.4

560.6

602.1

2

1.7

3.4

7.9

11.7

19.8

27.8

49.7

14.41

7.26

7.94

4.15

5.54

4.96

8.25

ST (m /g) SM (m /g) SM/ ST X 100 (%)

VT: Total pore volume, VM: Micropore volume, ST: BET total surface area, SM: HK micropore surface area

Fig. 2 shows the density functional theory (DFT) pore size distribution of sodium carbonate activated CNFs. The DFT can determine the successive pore size distribution curve from micropore, mesopore to macropore. DFT assumes each pore acts independently. Each pore size presents then contributes to the total adsorption isotherm in proportion to the fraction of the total area of the sample that it represents [5]. There are pores in the pore diameter range (from 2 to 5.5 nm). Regarding the mesopore diameter range (from 2 to 50 nm), they are small mesopores. N 2 has more pore volume and surface area except for the pore diameter range (from 3.7 to 4.3 nm). There are peaks in the pore diameter range (from 2.7 to 3.3 nm) in Fig. 2 (a) and (b). 12

Pore surface area (m2/g)

Pore volume (cc/g)

0.016

R N1 N2

0.014 0.012 0.010 0.008 0.006 0.004 0.002 0.000

R N1 N2

10 8 6 4 2 0

1

2

3

4

Pore diameter (nm)

(a)

5

6

1

2

3

4

5

6

Pore diameter (nm)

(b)

Fig. 2. Pore size distribution of sodium carbonate activated CNFs; (a) : pore volume distribution, (b) pore surface area distribution. The pore size distribution of phosphoric acid activated CNFs is illustrated in Fig. 3. According to the increment of phosphoric acid solution concentration, the pore volume and pore surface area increased. All pores of phosphoric acid activated CNFs have diameter range from 2.3 to 4.9 nm. Phosphoric acid activated CNFs have also small mesopores. H 80 shows the highest pore volume and surface area comparing with other samples.

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Nanocomposites and Nanoporous Materials VIII

0.04

0.04

Pore volume (cc/g)

0.03

0.02

0.01

Pore surface area (m2/g)

0.00 2.0

2.5

3.0

3.5

4.0

4.5

0.02

0.01

2.5

3.0

3.5

4.0

Pore diameter (nm)

(a)

(b) H 80 H 40

15 10 5

2.5

0.03

Pore diameter (nm)

20

0 2.0

H 60 H 20

0.00 2.0

5.0

3.0

3.5

4.0

Pore diameter (nm)

(c)

4.5

5.0

Pore surface area (m2/g)

Pore volume (cc/g)

H 80 H 40

4.5

5.0

H 60 H 20

20 15 10 5 0 2.0

2.5

3.0

3.5

4.0

4.5

5.0

Pore diameter (nm)

(d)

Fig. 3. Pore size distribution of phosphoric acid activated CNFs; (a) and (b) : pore volume distribution, (c) and (d) pore surface area distribution. Summary As a carbon source, CNFs were obtained by using electrospinning. To develop pore structure of CNFs, Na2CO3 and H3PO4 chemical activation was conducted. Even though pore structure of CNFs was developed with increasing specific surface area (max 60 times) and pore volume (max 120 times), they have quite different pore size distribution according to chemical agent. Acknowledgement This research was performed for the Hydrogen Energy R&D Center, one of the 21st Century Frontier R&D Program, funded by the Ministry of Science and Technology of Korea. References [1] J.S. Im, S.J. Park, Y.S. Lee, Journal of Colloid and Interface Science Vol. 314 (2007), p.32 [2] K. Eiji, H. Takashi, H. Kaoru, K. Hiroyuki, JP 01-156513, (1989). [3] J.S. Jang, Y.S. Lee, I.K. Kim and G. Yim: Carbon Science Vol. 1 (2000), p. 69 [4] M.A. Lillo–Rodenas, J. Juan–Juan, D. Cazorla–Amoros, A. Linares–Solano: Carbon Vol. 42 (2004), p. 1371 [5] M.M. Dubinin, G.M. Plavnik, Carbon, Vol. 6 (1968), p. 183

Solid State Phenomena Vol. 135 (2008) pp 85-88 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.85

Removal of Hexavalent Chromium on Chitosan-deposited Activated Carbon Jeong-Min Lee 1, a, K.Palanivelu2, 3, b and Young-Seak Lee1, 3, c, * 1

Department of Nano Technology, Chungnam National University, Daejeon 305-764, 220, Gungdong, Yuseong-Gu, Daejeon 305-764, Korea Korea 2 Centre for Environmental Studies, Anna University, Chennai 600 025, India 3 Department Fine Chemical Engineering and Chemistry, Chungnam National University, 220, Gung-dong, Yuseong-Gu, Daejeon 305-764, Korea a

[email protected], [email protected], c, * [email protected]

Keywords: Activated carbon, Chitosan, Adsorption, Cr (VI) removal.

Abstract. This research involved the deposition of chitosan on the surface of activated carbon(AC) to get an adsorbent(CS-AC) to removal Cr(VI) from aqueous solution. Each adsorbent used in this study was characterized using BET specific surface analyzer, FT-IR and elemental analyzer. Batch experiments were conducted to determine the removal efficiency of Cr(VI). Chitosan was physically adsorbed on the surface of the AC and the removal efficiency of CS-AC was greater than that of raw activated carbon(RAC), suggesting that the surface modification of AC by chitosan provides more adsorption sites on their solid surface for Cr(VI) adsorption.

Introduction Since the recent industry advances, the occurrence of wastewater with the heavy metal is also increasing. Especially, in plating industry, wastewater containing the various heavy metals has to be discharged. Among the heavy metal ions, Cr (VI) is one of priority list of toxic pollutants. It is necessary to remove Cr (VI) from wastewater effectively. There are a variety of methods using adsorbent on the removal of Cr (VI) from wastewater. A series of adsorbents such as activated carbon, zeolite and ion exchange resin have been investigated. Among these adsorbent, the activated carbon has been attractive to various areas. Owning to its economic and environmentally friendly properties, it is used to control air pollution and water pollution widely [1, 2]. Specially, since AC has very high specific surface area and advanced pore structure, not only adsorption capacity is large, but also adsorption rate is fast. To use AC as an adsorbent, it was researched about manufacturing methods and purification by surface properties [3, 4]. In this work, the surface of AC was chemically modified with chitosan and characterized. The adsorption of Cr(VI) on the adsorbents was evaluated. Experimental Materials Activated carbon (AC) used in this study was the coconut-based granular activated carbon (60×100 mesh) which was purchased from Eulim Carbon Co. Granular AC was washed with the distilled water and then dried in an oven at 110 °C for 24h. The chitosan has low molecular weight (150,000)that was obtained from Sigma Chemical Co. Ltd. Chitosan (1−10 wt%) was added to 2% acetic acid solution and stirred in a rotary shaker at 200 rpm for overnight. Activated carbon was dipped into chitosan solution and stirred for 1 h at 200 rpm. This mixture was evaporated at 70 °C in a vacuum and dried at 80 °C in oven for overnight.

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Nanocomposites and Nanoporous Materials VIII

The elemental analysis of adsorbents was carried out with elemental analyzer(EA1110, CE Instrument, Italy). Specific surface area, pore size distribution and pore structure were analyzed by using N2 adsorption (Autosorb-1, Quantachrome, USA) at 77 K. IR spectra were recorded in FT-IR Prestige-21(Simadzu, Japan). The pH measurement of the prepared adsorbent samples was done as per ASTM D 3838 method. 0.5 g of sample was immersed in 20 ml of distilled water at pH 7, mixed for 12 h, and filtered. The pH of the filtrate was measured with a pH meter. Cr(VI) adsorption studies Potassium dichromate (AR grade) was used as the source of Cr (VI). The Cr stock solution of 100 ppm with pH 3 was prepared. pH of Cr(VI) stock solution was adjusted to known value with 0.1 N NaOH or 0.1 N HCl. To evaluate the adsorption ability of CS-AC for Cr (VI), sample of 1 g adsorbent was immersed in Cr (VI) solution (100 ml) and stirred at room temperature for 12 h. Cr (VI) removal efficiency of CS-AC was analyzed by ICP-AES(Inductively Coupled Plasma - Atomic Emission Spectrometer, Atomscan-25, Thermo Jarrell Ash, USA Shimadzu, Japan ) Results and Discussion Chitosan is a high molecular substance and it is deposited on the surface of AC after dissolving in 2% acetic acid solution. The deposited chitosan content on samples was varied from 0 to 10 wt % chitosan. The results of elemental analysis are summarized in Table 1. As shown in Table 1, the nitrogen and oxygen content increased with increase of chitosan concentration. Also, the amount of hydrogen is proportional to the content of chitosan. In case of high chitosan content of 10 wt%, therefore, AC has the highest contents of nitrogen, oxygen and hydrogen. From these results, it is clear that chitosan was well deposited on AC. Table 1. Elemental analyses of CS-AC Type of Nitrogen Carbon Hydrogen adsorbent (%) (%) (%) RAC 0.01 93.63 0.47

a

Oxygen (%) 5.90

1 %CS-AC

0.20

86.85

0.66

12.29

2 %CS-AC

0.27

87.00

0.77

11.96

5 %CS-AC

0.52

84.07

1.24

14.17

10 %CS-AC

0.85

79.26

1.70

18.19

a : The oxygen is assessed by difference

On the other hand, in order to compare the textural properties of RAC and CS-ACs, the parameters obtained from BET specific surface area and pore volume are listed in Table 2. As can be seen in Table 2, there is little change of specific surface area on various adsorbent samples. It indicates that chitosan deposition apparently did not significantly affect the specific surface area of CS-AC due to its a little difference compared to specific surface area of RAC. In general, when some materials are deposited on porous materials such as AC, specific surface area is decreased. In this study also the surface area deceased with increase in the concentration of chitosan used for deposition. However, the decrease was only slight. In order to characterize the functional groups of AC surfaces, playing an important role for the adsorption of heavy metal ions, FT-IR spectroscopy was used. IR spectra obtained from as-received and 5 wt% CS-AC are shown in Fig. 1. According to Mahajan's research [5], in case of the adsorption of heavy metal ions, the chemical surface properties including functional groups are more important than the structural properties like specific surface area or micropore.

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Table 2. Textural properties of RAC and CS-AC Sample Sa (m2/g) Vb (cm3/g) Vc (cm3/g) R AC 1147 0.494 0.455 1 wt% CS-AC 1111 0.479 0.442 2 wt% CS-AC 1097 0.472 0.436 5 wt% CS-AC 1090 0.467 0.434 10 wt% CS-AC 1096 0.473 0.435

Wd (nm) 0.77 0.76 0.76 0.77 0.76

a: BET specific surface area, b: Total pore volume c: Micropore volume by H-K method, d: Average pore width by H-K method

As shown in Fig. 1, the peak around 3450 cm-1 corresponds to hydroxyl (-OH) and amino (-NH) stretching bands. In addition, the peak around 1100 cm-1 represents carboxyl (C-O) stretching bands. As a result, there are two peaks in IR spectra of RAC and 5 wt% CS-AC. This means that chitosan deposition provided the functional groups such as hydroxyl, amino and carboxyl band. This result is in good agreement with elemental analysis summarized in Table 1. Adsorption isotherms of as-received and chitosan deposited activated carbons are represented in Fig. 2. All of adsorption isotherms show Langmuir type. In the physisorption, Langmuir type isotherm indicates that adsorbent is microporous material. As presented in Table 2, microporosities of activated carbons used in this study are about 0.92. According to these results, porosity and isotherm type are not significantly changed after chitosan deposition on activated carbons.

5 wt%CS-AC

C=N

N-H str. O-H str. 4000

3500

3000

2500

C-O str.

2000

C-O str. C-N str.

1500

1000

Wave number ( cm-1)

Fig. 1. FT-IR spectra of RAC and CS-AC.

Volume adsorbed ( cm3/g)

320

RAC

300 280 260

RAC 1wt% CS-AC 10wt% CS-AC

240 220 200 0.0

0.2

0.4

0.6

0.8

1.0

Relative pressure ( P/Po)

Fig. 2. N2/77K adsorption isotherm of the activated carbons studied as a function of chitosan content

In addition to IR spectra, pH of RAC and CS-AC surfaces are furnished in Table 3. Otake and Jenkins [6] have shown the acidity of the carbon increased with the oxygen content. As seen in Table 3, the more the content of chitosan on the surface of AC, the lower pH is. It is natural that acid value of surface lowers pH.

Sample pH

Table 3. pH measurement of CS-AC 1 wt% 2 wt% 5 wt% RAC CS-AC CS-AC CS-AC 6.65 3.88 3.80 3.64

10 wt% CS-AC 3.55

In order to evaluate Cr (VI) adsorption capacity of CS-AC, adsorption experiments of Cr (VI) solution with pH 3 at the room temperature were carried out. The removal efficiency of Cr (VI) is

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Nanocomposites and Nanoporous Materials VIII

depicted in Fig 3 (a). In case of 2 wt% CS-AC, the removal efficiency of Cr (VI) was maximum as approximately 99.5% after 12 h. On the other hand, in cases of 5 and 10 wt% CS-AC, the removal efficiency slightly decreased to 98.3 and 95.0 %, respectively. These results mean that the deposited chitosan on AC contributed to adsorb Cr (VI) in low impregnation condition, whereas much impregnated chitosan covered the available adsorption sites being on AC surface. Therefore, in case of high impregnated chitosan on AC, the adsorption of Cr (VI) was prevented due to the sites adsorbed with chitosan. In addition, adsorption kinetics of Cr (VI) with 2 wt% CS-AC is shown in Fig. 3 (b). As can be seen in Fig. 3 (b), this adsorption isotherm is similar to Langmuir type. This type shows rapid adsorption at initial time. It indicates that optimum time of 2 wt% CS-AC for Cr (VI) removal is 3 h, with 100 ppm, at room temperature (25°C). Removal efficiency (%)

Removal efficiency (%)

100 99 98 97 96 95 94 93 92 91

0

(a)

2

4

6

8

Chitosan content (wt%)

10

100

80

0 0

2

(b)

4

6

8

10

12

Time (h)

Fig. 3. Removal Cr(VI) by ICP-AES : (a) by chitosan content (wt%) and (b) by time (h).

Summary The functional groups and adsorption properties of CS-AC were investigated with respect to chitosan concentration variation. The CS-AC has the various functional groups on the surface of AC such as hydroxyl, amino and carboxyl group. CS-AC as an adsorbent was used to remove Cr (VI) from waste water. The following conclusions have been drawn. 1. CS-ACs can remove Cr (VI) ions in wastewater. 2. The more the chitosan content was impregnated, the more the amount of nitrogen, oxygen and hydrogen of the surface increased. 3. Specific surface area of ACs was not changed so much by chitosan deposition. 4. The chitosan content of AC is inversely proportional to pH of activated carbon due to acidic functional groups of chitosan. 5. Adsorption kinetics of 2 wt% CS-ACs is almost similar to Langmuir adsorption type. Acknowledgement This work was supported by the Korea Research Foundation Grant funded by the Korean government (MOEHRD, Basic Research Promotion Fund) (KRF-2005-005-J00403) and BK21E2M program. References [1] R.C. Bansal, J.B. Donnet and F. Stoeckil: Active carbon, New York, Marcel Dekker (1988) [2] A. A. Attia, M. A. Shouman, S.A. Khedr, T.E. Nabarawy: Carbon Sci. Vol. 7 (2006), p.249 [3] Y. Onganer and C. Temur: J. Colloid Interface Sci. Vol. 205 (1998), p. 241 [4] M. Suzuki: Carbon Vol. 32 (1994), p. 577 [5] P. Mahajan, A. Youssef, P.L. Walker: Sep. Sci. Technol. Vol. 13 (1978), p. 487 [6] Y. Otake, R.G. Jenkins: Carbon Vol. 31 (1993), p. 109

Solid State Phenomena Vol. 135 (2008) pp 89-92 © (2008) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.135.89

Oxidation of Sulfur Components in Diesel Fuel with tert-Butyl Hydroperoxide Using Chromium Containing Catalysts Kwang-Eun Jeong1, a, Ho-Jeong Chae1, b, Ung-Chul Kim1, c, Soon-Yong Jeong1, d, Wha-Seung Ahn2, e 1

New Chemistry Division, Korea Research Institute of Chemical Technology, 100, Jang-dong, Yuseong-gu, Daejeon 305-600, Korea 2

Department of Chemical Engineering, Inha University, Incheon 402-751, Korea

a

[email protected], [email protected], [email protected], [email protected], e

*

[email protected]

Corresponding author: [email protected]

Keywords: Hexagonal Mesoporous Aluminophosphate; AlPO-5; Oxidative desulfurization; Chromium

Abstract. Transition metal (iron, cobalt, and chromium) substituted Hexagonal Mesoporous Aluminophosphate (HMA), and AlPO4-5 were studied for the oxidative desulfurization (ODS) reaction of refractory sulfur compounds such as DBT (dibenzothiophene) and 4,6-DMDBT (4,6dimethyldibenzothiophene). Selective oxidation activity varies with the kind of transition metal employed. The ODS activity of HMA substituted catalyst follows the order: CrHMA >> CoHMA ≥ FeHMA. Comparison between HMA and AlPO-5, CrHMA catalyst shows better performance than Cr-AlPO4-5 catalyst when tert-butyl hydroperoxide is used as oxidizing agent. Introduction Sulfur in transportation fuels is a major source of air pollution. Currently, ultra-deep desulfurization of fuel oil is attracting wide interest because of more stringent U.S. environmental regulations limiting the sulfur level in diesel to