1600214819photoionization of polyvalent ions

MATERIALS SCIENCE RESEARCH HORIZONS

MATERIALS SCIENCE RESEARCH HORIZONS

HANS P. GLICKEDITOR

Nova Science Publishers, Inc.New York

Copyright 2007 by Nova Science Publishers, Inc.All rights reserved. No part of this book may be reproduced, stored in a retrieval system or transmitted in any form or by any means: electronic, electrostatic, magnetic, tape, mechanical photocopying, recording or otherwise without the written permission of the Publisher. For permission to use material from this book please contact us: Telephone 631-231-7269; Fax 631-231-8175 Web Site: http://www.novapublishers.com NOTICE TO THE READER The Publisher has taken reasonable care in the preparation of this book, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained in this book. The Publisher shall not be liable for any special, consequential, or exemplary damages resulting, in whole or in part, from the readers use of, or reliance upon, this material. Independent verification should be sought for any data, advice or recommendations contained in this book. In addition, no responsibility is assumed by the publisher for any injury and/or damage to persons or property arising from any methods, products, instructions, ideas or otherwise contained in this publication. This publication is designed to provide accurate and authoritative information with regard to the subject matter covered herein. It is sold with the clear understanding that the Publisher is not engaged in rendering legal or any other professional services. If legal or any other expert assistance is required, the services of a competent person should be sought. FROM A DECLARATION OF PARTICIPANTS JOINTLY ADOPTED BY A COMMITTEE OF THE AMERICAN BAR ASSOCIATION AND A COMMITTEE OF PUBLISHERS. LIBRARY OF CONGRESS CATALOGING-IN-PUBLICATION DATA Materials science research horizons / Hans P. Glick (editor). p. cm. Includes index. ISBN-13: 978-1-60692-751-9 1. Materials science. I. Glick, Hans P. TA403.M347155 2006 620.1'1--dc22 2006032477

Published by Nova Science Publishers, Inc.

New York

CONTENTSPreface Chapter 1 Chapter 2 Photoionization of Polyvalent Ions Doris Mncke and Doris Ehrt Growth and Characterization of -Bi2O3 Thin Films by Chemical Vapour Deposition under Atmospheric Pressure T. Takeyama, N. Takahashi, T. Nakamura and S. Itoh Porous Materials: The Mathematical-Physical Expressions for Some Properties of Three-Dimensional Reticulated Porous Metallic Materials in the Same Analytical Model System P.S. Liu Influences of Process Parameters, Inclusion and Void in Copper Wire Drawing Somchai Norasethasopon Development of Hardfacing for Fast Breeder Reactors A. K. Bhaduri and S. K. Albert Tissue Engineering of Cartilage in Bioreactors Nastaran Mahmoudifar and Pauline M. Doran Heterogeneous Combustion Synthesis Hung-Pin Li Recycling of Ecocompatible Treated Red Mud and Compost from SS-MSW: Examples of Use on Sediment and Mine Soil Samples P. Massanisso, E. Nardi, R. Pacifico, L. DAnnibale, C. Cremisini and C. Alisi Formation and Adjustment of Bubbles in a Polyurethane Shape Memory Polymer W.M. Huang, B. Yang, L.H. Wooi, S. Mukherjee, J. Su and Z.M. Tai vii 1

57

Chapter 3

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Chapter 4

109 149 171 193

Chapter 5 Chapter 6 Chapter 7 Chapter 8

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Chapter 9

235 251

Index

PREFACEMaterials science includes those parts of chemistry and physics that deal with the properties of materials. It encompasses four classes of materials, the study of each of which may be considered a separate field: metals; ceramics; polymers and composites. Materials science is often referred to as materials science and engineering because it has many applications. Industrial applications of materials science include processing techniques (casting, rolling, welding, ion implantation, crystal growth, thin-film deposition, sintering, glassblowing, etc.), analytical techniques (electron microscopy, x-ray diffraction, calorimetry, nuclear microscopy (HEFIB) etc.), materials design, and cost/benefit tradeoffs in industrial production of materials. This book presents new research directions in this rapid-growing field. Chapter 1 - The effect of polyvalent dopants on photoinduced defect formation was studied in different glasses. Ionization of the glass matrix results in intrinsic defects, positively charged hole and negatively charged electron centers. Polyvalent dopants can be photooxidized or photoreduced. These extrinsic defects might replace selectively one or several intrinsic defects and / or cause an increase in the number of opposite charged defects. Photoionization can also result in unusual dopant valences otherwise not observed in glasses. The systematic comparison of different dopants and glass systems irradiated by excimer lasers helps to understand defect generation processes and might eventually help in the design of UV-resistant or UV-sensitive glasses. Defect formation occurs in the ppm range and was analyzed by optical and EPR spectroscopy. A series of polyvalent dopants such as typical trace impurities, glass or melt additives and typical dopants used for optical components like filter glasses, optical sensors, fluorophores or photochromes, were studied. Distinct melting conditions give rise to different valences of various dopants and as a consequence different photoinduced redox-reactions might be observed after irradiation. Qualitative and quantitative changes in the defect formation rates depend on the: kind and concentration of the dopant, c was varied from 50 to 5000 cation ppm. radiation parameters such as wavelength, or power density of the excimer lasers used. glass matrix; (fluoride-)phosphate and borosilicate glasses give rise to different intrinsic defects of varying stability. The matrix determines also the initial incorporation like valence or coordination of the dopants and stabilizes or destabilizes photoionized dopant species.

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initial transmission of the glass sample, which also depends on the dopant (kind, valence, coordination), its concentration, and the thickness of the sample plate, d was varied from 0.5 to 2mm. Some dopants are photooxidized while others are photoreduced Some defects recombine easily or transform into more stable defects while others are stable for months or years. Chapter 2 - Bismuth oxide (Bi2O3) thin films are interesting materials within the class of oxide semiconductors, owing to a variety of physical properties determined by its many polymorphs. This semiconductor is characterized by significant values of band gap, dielectric permittivity and refractive index as well as marked photosensitivity and photoluminescence. These properties make Bi2O3 films well suited for many applications in various domains such as microelectronics, sensor technology and optical coatings. However, the characteristics of this film strongly depend on its crystal phases: its electrical conductivity may vary by over 5 orders of magnitude, while its energy gap may change from around 2 to 3.96 eV. Therefore, it is required to manufacture high-quality Bi2O3 films with a single phase. Thin films of -Bi2O3 were prepared on the sapphire (0001) and the yttria-stabilized zirconia (YSZ) (111) substrate by means of chemical vapour deposition under atmospheric pressure. X-ray diffraction measurement revealed the deposited -Bi2O3 films on the YSZ (111) substrates have good crystal quality and a flat surface. The full width at half maximum value of out-of-plane rocking curve is 0.0260 (93.6 arcsec.). An optical band gap of 3.28 eV was estimated by the optical transmittance measurement. Spectroscopic ellipsometry shows that the refractive index n of the single crystalline -Bi2O3 film at 800 C is 2.4940 with 632.80nm. We believe this is the first time to investigate the optical properties of -Bi2O3 thin film. Chapter 3 - New developments are ceaselessly gained for the preparation, the application and the property study of porous materials. As to the theories about the structure and properties of porous materials, the famous classical model - Gibson-Ashby model has been being commonly endorsed in the field of porous materials all over the world, and is the theoretical foundation widespreadly applied by numerous investigators to their relative researches up to now. But there are some shortcomings in this classical model in fact, e.g., the impossible close-packed of pore units and the unequivalent struts. In this chapter, another model for porous materials are introduced which is put forward by the present author, and this new model can make up those shortcomings existed in Gibson-Ashby model. More importantly, the mathematical-physical expressions, which are well in agreement with the relevant experimental results, can be smoothly acquired for some properties of threedimensional reticulated foamed materials using this new model. These expressions include the relationship between electrical resistivity and porosity, the relationship between tensile strength and porosity, the relationship between relative elongation and porosity, and the relationship between biaxial loading strength and porosity. The experimental results showed that, the obtained mathematical-physical relations from this new model are obviously more excellent than that from Gibson-Ashby model when applying into the porous materials. Chapter 4 - In the copper fine wire drawing, the breakage and defect of the wire were fatal to the success of quantitative drawing operations. The first part of this paper shows how three of the main process parameters, the die half-angle, reduction of cross-sectional area and numbers of the drawing pass influenced drawing stress and internal defect by experiment. The influences of a non-central inclusion and void in the single-pass copper shaped-wire drawing were investigated by 2D FEM. The effects of the lateral and longitudinal sizes of a

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central inclusion in the multi-pass copper shaped-wire drawing were also investigated. Based on the experimental data of the optimal die half angle, wire deformation, plastic strain, hydrostatic stress and drawing stress of the copper shaped-wire containing a non-central inclusion and void were calculated. The copper shaped-wire that contained a central inclusion and void was also calculated. During drawing a wire containing a non-central inclusion, necking, bending and misalignment occurred. However, only necking occurred in the case of the central inclusion wire. In the case of the non-central inclusion wire, inclusion rotation occurred. For the same inclusion size, the inclusion size strongly influenced drawing stress but the eccentric distance slightly influenced drawing stress. The drawing stress of the copper shaped-wire that contained a central inclusion was greater than the case of the wire that contained a non-central inclusion. The drawing stress decrement due to a void and the opposite deformation behaviour between the wire that contained a central void and inclusion were found. The effects of the lateral and longitudinal sizes of a central inclusion and void on the drawing and the maximum hydrostatic tensile stress during the multi-pass copper shapedwire drawing were also carried out. The present paper also shows how two of the inclusion parameters, the size and aspect ratio of the elliptical inclusion, influenced drawing stress and maximum hydrostatic stress of the copper shaped-wire during drawing. It was found that the maximum drawing stress increased as the longitudinal inclusion size and aspect ratio increased. Both longitudinal inclusion size and aspect ratio influenced the inclusion leading edge location where the maximum hydrostatic tensile stress was induced. The necking due to a central inclusion in copper shaped-wire drawing occurred on some parts of the wire surface in front of and nearby the inclusion and the lateral neck size decreased when the longitudinal and lateral inclusion sizes increased as the inclusion passed through the die. The maximum hydrostatic tensile stress directly increased as the inclusion aspect ratio increased for the small and medium inclusions but it inversely increased for the large inclusion. It was mostly found where the inclusion leading edge was located in the drawn zone. The influences of a central inclusion on the plastic deformation, hydrostatic stress and drawing stress in the round-to-round copper wire drawing were also investigated by 3D FEM. Chapter 5 - Various components of the Fast Breeder Reactors encounter wear of adhesive or abrasive nature and sometimes erosion. Hardfacing by weld deposition have to be used to improve the resistance to high temperature wear, especially galling, of mating surfaces in sodium. Based on radiation dose rate and shielding considerations during maintenance, handling and decommissioning, nickel-base E NiCr-B hardfacing alloy was chosen to replace the traditionally used cobalt-base Stellite alloys. Studies, on the effect of long term ageing of NiCr hardface deposits on austenitic stainless steel substrate, demonstrated that E NiCr-B deposits after exposure at service temperatures up to 823 K would retain adequate hardness well above RC 40 at end of the components designed service-life of up to 40 years. Further, based on detailed metallurgical studies, including residual stress measurements after thermal cycling, the more versatile plasma transferred arc welding (PTAW) process was chosen for deposition of the E NiCr-B hardfacing alloy, so that the width of the dilution zone could be controlled by optimising the deposition parameters. This paper outlines the adaptation of technology for hardfacing with the E NiCr-B alloy using the selected PTAW process, through collaborative efforts with industries, for development of hardfacing technology for the various components of PFBR.

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Chapter 6 - The main goal of cartilage tissue engineering is to generate three-dimensional cartilage and osteochondral tissues for use in repair of large cartilage injuries. Cartilage constructs are generated by seeding and culturing viable cells in biodegradable polymer scaffolds under conditions suitable for tissue formation. In this chapter, current developments in cartilage tissue engineering are reviewed, focusing on the source of cells, the polymer scaffolds, seeding systems, bioreactors and application of mechanical stimulation for cell differentiation and tissue production. The generation of cartilage tissue constructs in the laboratory using a bioreactor system is also described. Chondrocytes were isolated from human foetal epiphyseal cartilage, expanded in monolayer, dynamically seeded into poly(glycolic acid) (PGA) polymer scaffolds and cultured in recirculation bioreactors. Composite scaffolds were used to improve the initial distribution of cells within the scaffolds and to develop cartilage constructs that were homogeneously cartilaginous throughout their thickness. The quality of the engineered cartilage was assessed after 5 weeks of bioreactor culture in terms of tissue wet weight, cell, glycosaminoglycan (GAG), total collagen and collagen type II contents, histological analysis of cell, GAG and collagen distributions, immunohistochemical analysis of collagen types I and II, and ultrastructural analysis using transmission electron microscopy. Chapter 7 - Many exothermic non-catalytic solid-solid or solid-gas reactions, after being ignited locally, can release enough heat to sustain the self-propagating combustion front throughout the specimen without additional energy. Since the 1970s, this kind of exothermic reaction has been used in the process of synthesizing refractory compounds in the former Soviet Union. This novel technique, so-called Combustion / Micropyretic synthesis or Selfpropagating High-temperature Synthesis(SHS), has been intensively studied for process implication. This technique employs exothermic reaction processing, which circumvents difficulties associated with conventional methods of time and energy-intensive sintering processing. The advantages of combustion synthesis also include the rapid net shape processing and clean products. In addition, the combustion-synthesized products have been reported to possess better mechanical and physical properties. Heterogeneous distributions of reactants, diluents, and pores are common during combustion synthesis when powders are mixed, and this directly leads to the variations of the thermophysical / chemical parameters of the unreacted compacts. Since combustion synthesis is sustained by the sequences of the local chemical reactions, the propagation manner is strongly dependent on the parameters of each portion of the reactants. Thus, the variation of thermophysical / chemical parameters of reactants caused by heterogeneities in composition and porosity is thought to significantly change the processing parameters, such as combustion temperature and propagation velocity; and further affect the product properties. This chapter systematically introduces the impact of heterogeneities during combustion synthesis with Ni + Al. Correlations of heterogeneities in the reactants and a diluent with the propagation velocity and combustion temperature are discussed. In addition, a map, considering concurrent heterogeneities in the composition and porosity, has been generated to provide a better understanding of the change in propagation velocity on account of the heterogeneous combustion synthesis. Chapter 8 - Ecological restoration of polluted areas is an increasing necessity for many countries around the world. Current technologies used to recover polluted soil and sediment are in general too costly. Recently, on-site approaches such as metal trapping and phytoremediation have attracted attention for their ability to meet criteria of economicity.

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Metal trapping is based on the diminution of metal mobility and availability as a result of applying soil amendments, for example particular industrial residues. Phytoremediation is an appealing environmental cleanup technology but a deeper understanding of the complex interactions in the soil-plant system is still needed. In this study, the effect of adding treated red mud (BauxsolTM - material with the potential to immobilise metal) on mine soil and on sediment (from a volcanic coastal lagoon in Southern Italy) and of adding both red mud and compost (produced from Source-Separated Municipal Solid Waste) on trace elements fractionation and mobility, have been investigated. Barley (Hordeum vulgare) was used as a plant model to follow any change in matrices phytotoxicity: seedlings were transplanted in pots containing the contaminated mine soil or sediment and a mixture of the investigated matrices with different percentages of treated red mud and compost. Plant growth was studied also by controlling the total protein content, biomass and enzyme activity. The knowledge of trace elements mobility and speciation in contaminated soils and sediments is an important requisite for any further environmental evaluation and these features can be evaluated through leaching tests or by "sequential extraction procedure". In this work, total concentration of selected trace elements, their fractionation by sequential extraction procedure (BCR standardised) and leaching batch tests using a kinetic approach, were studied. The most evident result in the soil trials was that the utilization of amendments, used both separately and in a mixture, always improved the growth of barley plants. In particular, barley seedlings were practically not able to grow on the polluted mine soil and the simple adding of red mud resulted in a significant improvement in plant development. An even more drastic improvement was obtained with the addition of compost and compost plus treated red mud. In the sediment trials, the best yield in plant growth was obtained in the pot with the addition of treated red mud alone. The necessity of a delicate compromise between the maintaining of an acceptable plant viability and the control of metal mobility seems to be achievable through a careful balancing of the percentages of compost and red mud utilized as amendments. Chapter 9 - Two approaches are proposed for realizing porous polyurethane shape memory polymers using water as a non-harm foam agent. We show that it is possible to control the bubbles by varying the moisture ratio and heating procedure. We demonstrate that one can further modify the size of bubbles by further heat treatment. As such, one can make resizable micro bubbles and even channels.

In: Materials Science Research Horizons Editor: Hans P. Glick pp. 1-56

ISBN 978-1-60021-481-3 2007 Nova Science Publishers, Inc.

Chapter 1

PHOTOIONIZATION OF POLYVALENT IONSDoris Mncke and Doris EhrtOtto-Schott-Institut fr Glasschemie, Friedrich-Schiller-Universitt, Fraunhoferstr.6, D-07743 Jena, Germany

ABSTRACTThe effect of polyvalent dopants on photoinduced defect formation was studied in different glasses. Ionization of the glass matrix results in intrinsic defects, positively charged hole and negatively charged electron centers. Polyvalent dopants can be photooxidized or photoreduced. These extrinsic defects might replace selectively one or several intrinsic defects and / or cause an increase in the number of opposite charged defects. Photoionization can also result in unusual dopant valences otherwise not observed in glasses. The systematic comparison of different dopants and glass systems irradiated by excimer lasers helps to understand defect generation processes and might eventually help in the design of UV-resistant or UV-sensitive glasses. Defect formation occurs in the ppm range and was analyzed by optical and EPR spectroscopy. A series of polyvalent dopants such as typical trace impurities, glass or melt additives and typical dopants used for optical components like filter glasses, optical sensors, fluorophores or photochromes, were studied. Distinct melting conditions give rise to different valences of various dopants and as a consequence different photoinduced redox-reactions might be observed after irradiation. Qualitative and quantitative changes in the defect formation rates depend on the: kind and concentration of the dopant, c was varied from 50 to 5000 cation ppm. radiation parameters such as wavelength, or power density of the excimer lasers used. glass matrix; (fluoride-)phosphate and borosilicate glasses give rise to different intrinsic defects of varying stability. The matrix determines also the initial incorporation like valence or coordination of the dopants and stabilizes or destabilizes photoionized dopant species.

Tel.: +49-3641-948511 / 948506; fax: +49-3641-948502, [email protected] or [email protected].

2

Doris Mncke and Doris Ehrtinitial transmission of the glass sample, which also depends on the dopant (kind, valence, coordination), its concentration, and the thickness of the sample plate, d was varied from 0.5 to 2mm.

Some dopants are photooxidized while others are photoreduced Some defects recombine easily or transform into more stable defects while others are stable for months or years.

INTRODUCTIONSolarization in glasses was first described by Faraday in 1825. The effect of irradiation induced transmission changes has been investigated since for its scientific and technological significance [1-31]. Pelouze recognized already in 1867 that a change in the oxidation state of the typical glass impurities Fe and Mn can cause strong solarization effects [2]. UV-radiation excites valence electrons in the irradiated material and complicated photoreaction processes lead subsequently to the formation of irradiation induced defects. Defects are generated in ppm concentrations and occur in pairs of negative electron centers (EC) and positive hole centers (HC). While intrinsic defects arise from the glass matrix itself are extrinsic defects connected to dopants or impurities. The formation of defects may result in transmission changes but also in changes of the refractive properties of the material. Studying the processes and mechanisms that govern defect formation requires more attention as stronger lamps and lasers, which work at increasingly shorter wavelengths, are more and more utilized. This knowledge can than be exploited in the development of photosensitive or photoresistant appliances. Because of their strong electronic transitions in the ultraviolet (UV) and visible (VIS) spectral range were, in analogy to similar absorbances in crystals, defects initially called color centers. Optical spectroscopy is the method of choice when studying defect formation; however, optical spectra of doped glasses are often dominated by transitions of the dopants that overlay the defects bands [3-18]. Complementary information on these defects, even regarding their detailed structure, can be derived from EPR-spectroscopy as many defects are paramagnetic [7-23, 3-18] The addition of polyvalent ions often initiates or enhances considerably the formation of defects in a glass sample. Extrinsic defects can form in addition to intrinsic defects and thus cause the increased generation of reversibly charged intrinsic defects. On the other hand can extrinsic defects substitute selectively one or more intrinsic defects of like charge [3, 19-23]. Irradiation induced defects can further be classified according to their stability in transient or stable defects. Some initially formed defects transform rapidly into more stable defects, sometimes even during the irradiation process. The transformation of defects in thermodynamically more stable centers can then again be hindered kinetically. Defect formation is a dynamic process and the kind and rate of defect development depends on many factors, e.g. the glass matrix, the concentration and species of any dopants, the initial transmission of the sample, or on the radiation parameters. For example excites UV-radiation only valence electrons while X-ray radiation detaches even the inner electrons in the material. Accordingly are different defects initiated by different radiation sources [24].

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This chapter intends to compare the role of a wide range of polyvalent metals in the formation of irradiation induced defects. All glass samples were irradiated with excimer lasers in the UV and the laser induced defects were characterized by optical and EPR spectroscopy. Only defects stable at room temperature are discussed. Even these relative stable defects show transformation and recombination reactions when the samples were stored at room temperature in the dark. The glass types studied were selected for their high transparency in the deep ultraviolet (0~160-185 nm) and consequently their application in high performance optics [25-31, 35].

2. EXPERIMENTAL SECTIONFluoridephosphate (FP) and metaphosphate glasses (NSP) were selected as primary matrix glasses for the irradiation experiments. The generated extrinsic defects were characterized by EPR and optical spectroscopy and when possible identifified yb type and charge. Additional experiments using borosilicate type samples were added later in order to study the dopants effect on defect formation in an entirely different glass matrix. All samples were prepared and irradiated under defined and comparable conditions.

2.1. Sample PreparationThe preparation of the different high purity glasses has been described in detail before [718, 25-31]. Only high purity reagents were used for all glasses. The iron content of the duran type borosilicate glass was < 1 ppm, of NSP ~ 5ppm and of FP10 < 10 ppm. The total iron content was analyzed by wet-chemical analysis. The Fe3+ content was also determined from the absorption of its CT-transition at 250 nm in the optical spectra [9, 27, 28-31]. The high purity dopant components were added in various amounts between 50 and 5000 ppm (cation %). The fluoroaluminate FP10 has the synthetic composition [10 P2O5 90 (AlF3, CaF2, SrF2, MgF2) mol%] and was melted at 1100C under air in platinum crucibles. In order to obtain reduced dopant species were some samples also remelted under reducing melting conditions under argon atmosphere in glass-carbon crucibles. The metaphosphate glass NSP [10 Na2O 40 SrO50 P2O5 mol%] was melted under air at 1300C in SiO2-crucibles. The dopants were reduced by the addition of 0.2 to 1 wt% carbon to the batch. Low alkaline borosilicate samples of the duran type [82 SiO212 B2O35 (K/Na)2O1 Al2O3 mol%] were prepared under air at 1650C. 250 to 1000 g batches were processed for 35 hours. For some samples were oxidizing or reducing conditions established by using the corresponding nitrate or tartrate salts of the reagents. All melts were cast in preheated graphite moulds and annealed from 500 or 550 C to room temperature with a cooling rate of 30 K / h.

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2.2. Radiation SourcesPolished samples plates of the dimensions 10 x 20 mm and a thickness of either 0.5, 1, or 2 mm were irradiated with excimer lasers. The sample thickness was chosen for each sample in accordance to the initial absorbance at the irradiation wavelength. Excimer lasers working at 193 nm (ArF-laser), 248 nm (KrF-laser), and 351 nm (XeFlaser) were used. The power density of the laser was 200 mJ/cm per pulse by a pulse duration of ~ 20 ns. The optical spectra were taken with increasing accumulated pulse numbers (at 10, 100, 1000, 5000, and 10000 pulses). The final pulse number normally suffices to reach the saturation level. EPR spectra of the samples were taken once after the final irradiation. Further optical spectra were obtained at increasing time intervals in which the irradiated samples were stored in the dark at room temperature. Some samples were irradiated by a high pressure mercury lamp or HOK lamp with a spectral power density of 1 kW with a wide continuos spectrum from 190 nm in the UV to the NIR.

2.3. UV-VIS-SpectroscopyUV-VIS-NIR spectra were taken in the range from 190 to 3000 nm. A double beam spectrophotometer (UV-3102, Shimadzu) recorded the absorbance A=lg(T0/T) with an error 1.0. Wire breakage occurred when Di/Do = 0.4 in the fifth pass drawing and Li/Do = 0.3 and 0.4 in the fourth and fifth pass drawing. The drawing pass numbers strongly influenced maximum hydrostatic stresses in the fourth and fifth pass drawing. They inversely and directly influenced drawing stress in the first-to-fourth and fifth pass drawing, respectively. Central Void Effects: Multi-Pass Copper Shaped-Wire Drawing. Necking occurred on the copper shaped-wire surface at the wire portion that contained a void and the lateral neck size decreased as the void size increased. The maximum hydrostatic tensile stress was not found at the wire centreline but was found on the wire surface at the neck. The maximum hydrostatic tensile stress directly increased as the void size increased. The void was transformed to be an almond-shaped void and the sharp-edge of a transformed void point in the opposite direction of the metal flow and also was transformed to be a linear crack or "pipe". The wire breakage due to the high hydrostatic tensile stress and large plastic deformation of the copper matrix in the neck was found when the void size was very large. The maximum drawing stress was normally equal to the drawing stress of the copper shaped-wire drawing without void. The minimum drawing stress of copper shaped-wire that contained the largest void, the lateral void size was equal to 1.0 (lateral wire size equal to lateral void size), was obtained and was equal to 0.0. Inclusion Size and Aspect Ratio Effects: Single-Pass Copper Shaped-Wire Drawing. The necking behavior is the same as described above. The lateral neck size decreased in accordance with the increase in longitudinal and lateral inclusion sizes while the inclusion passed through the die. The medium inclusion (a/h ratio was approximately 0.2 to 0.5) strongly influenced the maximum hydrostatic tensile stress and it rapidly increased as longitudinal inclusion size increased in this range. It directly increased as the inclusion aspect ratio increased (elliptical inclusion approach to be a circular inclusion) for the small (a/h ratio was approximately 0.0 to 0.2) and medium inclusion. It inversely increased as the inclusion aspect ratio increased for the large (a/h was approximately 0.5 to 0.8) inclusion. The maximum hydrostatic tensile stress was found where the inclusion leading edge was located in the drawn zone and was far away from the die exit. It was not found in the case of a/h = 0.4 where b/a = 0.6, a/h = 0.2 where b/a = 0.8 and a/h = 0.2 where b/a = 1.0. The maximum drawing stress occurred when the inclusion passed through the die and increased as the longitudinal inclusion size and aspect ratio increased. It was found when the inclusion leading edge was located at the inclusion displacement ratio equal to 0.55 in the reduction zone. Inclusion Size Effects: Single-Pass Round-to-Round Copper Wire Drawing. The inclusion size directly strongly influenced necking and maximum hydrostatic tensile stress of the copper wire. The maximum hydrostatic tensile stress was found when the inclusion leading edge was located around the die exit. But the maximum drawing stress was found when the inclusion leading edge was located around the reduction zone center.

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ACKNOWLEDGMENTThe author wishes to express his appreciation to the Director of the National Metal and Materials Technology Center (MTEC), National Science and Technology Development Agency, Thailand, for his support and assistance in many details of the finite element program "MSC.MARC" for this problem simulation. The author would like to thank Prof. Dr. Yoshida, K., Department of Precision Mechanics, School of Engineering, Tokai University, Japan, and Nissapakul, P., Tangsri, T., and Pramaphant, P., Department of Mechanical Engineering, Faculty of Engineering, King Mongkuts Institute of Technology Ladkrabang, Thailand, for giving him valuable discussion and comment.

REFERENCES[1] [2] Morton, J., (1999). Thomas Bolton and Sons and the rise of the electrical industry. Engineering Science and Education Journal, 5-12. Miyashita, K., Sugiyama, K., Moriai, H., Kamata, K., Tachikawa, K., and Fukuda, K., (1999). Electromagnetic Properties of Bronze Processed Nb3Sn Superconducting Wires and Multi-strand Cables for A.C. Use with a Cu-Sn-X (X,Ge,Ni,Mn,Si) Matrix and a Nb-Ta Core. IEEE Transactions on Applied Superconductivity, Vol. 9, No. 2, 709-712. Mielnik, E. M., Metalworking Science and Engineering; McGraw-Hill, Inc.; New York, 1991, pp 397-462. Raskin, C., (1997). Proceedings of the WAI International Technical Conference. Italy. Amstead, B. H., Ostwald, P. F., and Begeman, M. L., Manufacturing Processes; John Wiley and Sons, Inc.; Singapore, 1987, pp 1-687. Johnson, H. V., Manufacturing Processes; Bennett and McKnight; USA, 1984, pp 14581. Kutz, M., Mechanical Engineers Handbook; John Wiley and Sons, Inc.; New York, 1998, pp 3-1205 Colangelo, V. J., and Heiser, F. A., Analysis of Metallurgical Failures; John Wiley and Sons, Inc.; Singapore, 1989, pp 240-322. Avitzur, B., Metal Forming: Processes and Analysis; McGraw-Hill; New York, 1968, pp. 153-258. Revised edition reprinted by Robert Krieger Publishing Co., Inc.; Huntington, N.Y., 1979. Avitzur, B., Study of Flow Through Conocal Converging Dies; Metal Forming; A. L. Hoffmanner (ed.), Plenum Press, New York, 1971, pp 1-46. Campos, H. B., and Cetlin, P. R., (1998). The influence of die semi-angle and to the coefficient of friction on the uniform tensile elongation of drawn copper bars. Journal of Materials Processing Technology, Vol. 80-81, 388-391. Campos, H. B., Castro, A. L. R., and Cetlin, P. R., (1996). Influence of die semi-angle on mechanical properties of single and multiple pass drawn copper. Journal of Materials Processing Technology, Vol. 60, 179-182. Norasethasopon, S., and Tangsri, T., (2001). Experimental Study of the Effect of a Half-Die Angle on Drawing Stress during Wire Drawing. Ladkrabang Engineering Journal, Vol. 18, 134-139.

[3] [4] [5] [6] [7] [8] [9]

[10] [11]

[12]

[13]

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[14] Norasethasopon, S., and Yoshida K., (2003). Influence of an Inclusion on Multi-Pass Copper Shaped-Wire Drawing by 2D Finite Element Analysis. International Journal of Engineering, I. R. Iran, Vol. 16, No. 3, 279-292. [15] Norasethasopon, S., and Yoshida, K., (2006). Influences of inclusion shape and size in drawing of copper shaped-wire. Journal of Materials Processing Technology, Vol. 172, No. 3, 400-406. [16] Norasethasopon, S., and Yoshida, K., (2006). Finite-element simulation of inclusion size effects on copper shaped-wire drawing. Materials Science and Engineering: A, Vol. 422, No. 1-2, 252-258.



Chapter 5

DEVELOPMENT OF HARDFACING FOR FAST BREEDER REACTORSA. K. Bhaduri and S. K. AlbertMaterials Joining Section, Materials Technology Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603102, India

ABSTRACTVarious components of the Fast Breeder Reactors encounter wear of adhesive or abrasive nature and sometimes erosion. Hardfacing by weld deposition have to be used to improve the resistance to high temperature wear, especially galling, of mating surfaces in sodium. Based on radiation dose rate and shielding considerations during maintenance, handling and decommissioning, nickel-base E NiCr-B hardfacing alloy was chosen to replace the traditionally used cobalt-base Stellite alloys. Studies, on the effect of long term ageing of NiCr hardface deposits on austenitic stainless steel substrate, demonstrated that E NiCr-B deposits after exposure at service temperatures up to 823 K would retain adequate hardness well above RC 40 at end of the components designed service-life of up to 40 years. Further, based on detailed metallurgical studies, including residual stress measurements after thermal cycling, the more versatile plasma transferred arc welding (PTAW) process was chosen for deposition of the E NiCr-B hardfacing alloy, so that the width of the dilution zone could be controlled by optimising the deposition parameters. This paper outlines the adaptation of technology for hardfacing with the E NiCr-B alloy using the selected PTAW process, through collaborative efforts with industries, for development of hardfacing technology for the various components of PFBR.

1. INTRODUCTIONThe Indian 500 MWe Prototype Fast Breeder Reactor (PFBR) is a pool-type liquidsodium-cooled reactor having two separate sodium circuits with the intermediate heat exchanger (IHX) providing thermal contact between the primary pool and the secondary circuit. The secondary sodium circuits transfer heat from the IHX to the steam generator, the

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A. K. Bhaduri and S. K. Albert

steam from which drives the conventional steam turbines. The minimum sodium temperature in the primary pool during normal operation is 673 K while the mean above-core temperature is 823 K. The minimum and maximum sodium temperatures in the secondary circuit are 628 and 798 K, respectively. The steam temperature is 763 K at a pressure of 16.6 MPa. In the PFBR, 316L(N) austenitic stainless steel (SS) has been chosen as the structural material for components operating above 673 K. The liquid sodium coolant acts as a reducing agent and removes the protective oxide film present on the SS surface of the in-sodium components. Many of these components would be in contact with each other or would have relative motion during operation, and their exposure at high operating temperatures (typically 823 K) coupled with high contact stresses could result in self-welding of the clean metallic mating surfaces. In addition, the relative movement of mating surfaces could lead to galling, a form of high-temperature wear, in which material transfer occurs from one mating surface to another due to repeated self-welding and breaking at contact points of mating surfaces. Further, susceptibility to self-welding increases with temperature for 316 SS [1]. Hardfacing of the mating surfaces has been widely used in components of water-cooled and liquidsodium cooled FBRs to avoid self-welding and galling [2, 3]. Cobalt-base hardfacing alloys (e.g. Stellite) have been traditionally used very extensively for high temperature application in many critical hardfacing applications due to their excellent wear-resistance properties [4]. However, when cobalt-base alloys were used in a nuclear reactor environment, the cobalt-60 isotope formed due to irradiation enhances the radiation dose rate to operating personnel during handling, maintenance or decommissioning of the hardfaced components. Hence, there is an emerging trend of avoiding the use of cobaltbase alloys for hardfacing of nuclear power plant components. Nickel-base hardfacing alloys (e.g. Colmonoy) were developed mainly to replace the cobalt-base alloys for avoiding induced radioactivity problems in thermal and FBR applications. Accordingly, for the PFBR, selection of suitable hardfacing materials for various components was preceded by detailed induced radioactivity, dose rate and shielding computations to ensure that induced radioactivity from hardfaced components is kept to the minimum for maintenance and decommissioning purposes, and also to reduce the shielding thickness required for the component-handling flask, which in turn would reduce the flask weight, size of handling crane and loads on civil structures [5].

2. SELECTION OF HARDFACING MATERIALSelection of hardfacing material was based on, for the first time, detailed calculations of induced radioactivity and radiation shielding during maintenance, handling and decommissioning for each of the PFBR components that are to be hardfaced [5]. For these computations, replacement of Stellite 6 and Stellite 12 by same amount of Ni-base E NiCr-B (Colmonoy 5) hardfacing alloy (nominal compositions given in Table 1) was considered. Based on these calculations, for the components of PFBR, E NiCr-B (Colmonoy) was chosen as the hardfacing material to replace the traditionally used Stellites. Colmonoys have already been used in FBRs with satisfactory results. Tests on six liquid sodium pumps, with 304 SS bearings hardfaced with Colmonoy 6 and the shafts/journals

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hardfaced with Colmonoy 5, operating at 748798 K, had accumulated 20000 h for each pump without failure of the bearing area [6]. Table 1. Nominal compositions (in wt. %) of the hardfacing alloys consideredAlloy Stellite 6 Stellite 12 Colmonoy 5 B 2.5 C 1.0 1.8 0.65 Cr 27.0 30.0 11.5 Co 60.0 52.2 < 0.25 Fe < 2.5 < 2.5 4.25 Mn 1.0 1.0 Ni < 2.5 < 2.5 77.10 Si 1.0 1.0 3.75 W 5.0 9.0

The Rapsodie, Hallam, Fermi and EBR-II reactors used Colmonoy-faced sleeves and shafts in the hydrostatic sodium-lubricated pumps [7]. Bearing operation of the Hallam, Fermi and EBR-II pumps had been satisfactory, but all operations were below 813 K. However, a seizure occurred on the Rapsodie intermediate (secondary) pump before attaining an operating temperature of 823 K. The cause of the failure is not known. Another Rapsodie primary pump seizure occurred sometime later and its probable cause was lack of wear resistance in the bearing material. The temperature of the pumps was then limited to 723 K. All previous prototype bearings were made of Stellite but for the Rapsodie pumps, a change to Colmonoy was made. The use of the proven material, Stellite, might have eliminated the seizures [7]. To alleviate the main anxiety with NiCr hardface deposits, namely reduction in its hothardness, for the first-time, the hardness of long-term aged NiCr hardface deposits was studied using the Larsen-Miller parametric approach. For this purpose, a 316 SS plate was hardfaced with E NiCr-B alloy rods of 4 mm diameter by the gas tungsten arc welding (GTAW) process, with the hardface deposit thickness being about 2 mm. Samples with a hardface deposit thickness of 1.5 mm were cut from this hardfaced plate, and were subjected to ageing at three different temperatures (823, 873 and 923 K) for five different durations (200, 500, 1000, 2000 and 5000 h) at each temperature. The Vickers hardness (HV) of the asdeposited and all the aged hardface deposits were measured at room temperature (RT = 300K) using a load of 10 kg. These hardness values were then analysed to predict the hardness of the E NiCr-B deposit after long-term ageing at the service temperatures of 673 and 773 K. The hardness of as-deposited and all aged hardface deposits, measured at RT, are presented in Fig. 1. The timetemperature correlation for these hardness values were obtained using the Larson-Miller parametric approach, given by LMP = T(C + log t), where LMP is the Larson-Miller parameter, T is the temperature in Kelvin, t is the time in hours, and C is a constant. The constant C was determined as 14.4 for E NiCr-B deposit by least-square fitting with R2 of fit being about 0.97. Using C as 14.4, the RT hardness of E NiCr-B after ageing at 823 K for the service-life of the various PFBR components was estimated. Fig. 2 shows the estimated hardness after simulated service exposure of the E NiCr-B deposit for 2, 3, 5, 10, 15, 20, 25, 30, 35 and 40 years.

152800


Vickers hardness (HV, 10 kgf)

750 700 823 K 650 600 550 923 K 500 450 400 0 1000 2000 3000 4000 5000 873 K

Ageing time (h)

Figure 1. Variation of hardness at RT of E NiCr-B deposit with duration of ageing at 823-923 K.800

Vickers hardness (HV, 10 kgf) at RT

750 700 650 600 550

HV (at RT) = -0.0872 * LMP + 1947.2697 R2 = 0.9745

2 yr

3 yr

5 yr

10 yr

20 yr 25 yr 15 yr 30 yr

500 450 400 13.0

823 K 873 K 923 K Estimated at 823 K Linear (LMP, C= 14.4)13.5 14.0 14.5 15.0 15.5

16.0

40 yr

35 yr

16.5

17.0

Larson-Miller parameter = T(C + log t), T in K, t in h

Figure 2. Variation of hardness at RT of E NiCr-B deposit with Larson-Miller Parameter.

To estimate the hot-hardness of E NiCr-B deposit on prolonged exposure at the different operating temperature of the various PFBR components, namely 673 and 823 K, the average hot-hardness values of unaged E NiCr-B (Colmonoy 5) and Stellite 6, as shown in Fig. 3 [8], were used. The temperature dependence of the hardness of these hardface deposits was determined by an Arrhenius-type plot of ln(hardness at RT/hardness at temperature) vs. 1/T (K1), as given in Fig. 4. Using the relationships for both hardfacing alloys over specific temperature ranges as in Fig. 4, the hardness of E NiCr-B deposit at 673 and 823 K was estimated for prolonged exposure at 823 K, as presented in Fig. 5. The hardness values of asdeposited Stellite 6 at 300, 673 and 823 K are also presented in Fig. 5 for comparison.

Development of Hardfacing for Fast Breeder ReactorsVickers hardness (HV, 10 kgf)600 500 400 300 200 100 273 Colmonoy 5 Stellite 6 373 473 573 673 773 873 973 1073

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Temperature (K) of hardness measurement

Figure 3. Variation in average hot-hardness of unaged (as-deposited) of Stellite 6 and E NiCr-B (Colmonoy 5) deposits with temperature [8].0.45 0.401000 K 667 K 500 K 400 K 333 K

Colmonoy 5 (300-589K) Colmonoy 5 (589-922K) Stellite 6 (300-700K) Stellite 6 (700-922K) 0.30 0.25 0.20 0.15 0.10 0.05 0.00 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5

ln [HV at RT / HV at T(K)]

0.35

1000/T(K)Figure 4. Arrhenius-type plot showing temperature dependence of hot-hardness of Stellite 6 and E NiCr-B (Colmonoy 5) deposits.

Figs. 2 and 5 show that although there is expected to be about 43% reduction in the hardness of E NiCr-B deposit after 40 years of exposure at 823 K, the hardness of E NiCr-B deposit is expected to remain sufficiently higher than the hardness of as-deposited Stellite 6. Hence, E NiCr-B deposits are expected to retain adequate hardness of about 516 HV at RT and about 430 HV at 823 K after 40 years of exposure (ageing) at 823 K, i.e. up to the end of the components designed service-life.

154Estimated Vickers hardness (HV, 10 kgf)800


700

600

Colmonoy 5 at 300 K Colmonoy 5 at 673 K Colmonoy 5 at 823 K Stellite 6 at 300 K Stellite 6 at 673 K Stellite 6 at 823 K

500

400

300

200 0 5 10 15 20 25 30 35 40

Time of exposure (ageing) at 823 K (years)

Figure 5. Estimated hot-hardness of E NiCr-B (Colmonoy 5) deposit after ageing at 823 K.

3. SELECTION OF HARDFACING PROCESS AND HARDFACING ALLOY TYPENiCr hardfacing alloys, which contain high chromium and boron, form very hard chromium borides and carbides that contribute to their high hardness in addition to the solid solution strengthening by the alloying elements [9]. The abrasive resistance of the NiCr alloys is a function of amount of hard borides present in the matrix. During deposition, dilution from the substrate material occurs and this could significantly alter the microstructure and mechanical properties of the hardface deposits near the deposit/substrate interface [10]. Further, the coating thickness is optimised from the consideration that, due to differential thermal expansion of the deposit and substrate, an increase thickness would cause an increase in the residual stress and the tendency of the deposit to crack and spall under thermal cycling conditions. Also, radiation-induced damage can aggravate the integrity of the hardface coatings. Finally, when designing coatings for wear resistance, corrosion resistance and other high temperature properties, the finished coating thickness is so chosen that it is greater than the permitted wear tolerance, especially for nuclear components in which refurbishing or repair is not envisaged. While the undiluted hardface deposit provides the required wear resistance, the dilution zone at the deposit/substrate interface partially accommodates the stresses that arise during deposition or due to differential thermal expansion of the deposit and substrate during high temperature service. It is for these considerations that the best deposition process has to be adopted so that the width of the dilution zone is optimum and sufficient undiluted zone is available within the desired deposit thickness. Hardfacing with NiCr alloys by weld deposition is usually carried out using GTAW process, for which no major technological development is involved. A major problem with weld deposition by GTAW is the high dilution and tendency for cracking of the weld deposit, necessitating stress relieving at high temperature. One possible way to alleviate these problems, at least partially, is to deposit


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thinner coatings using the plasma transferred-arc welding (PTAW) process. However, the problems associated with the weld deposition of the nickel-base NiCr hardfacing alloys include low fluidity, generation of residual stress in the weld deposits that can lead to cracking, hard microstructure and significant dilution of deposit by the substrate material due to the large difference in their respective melting points. Since, the cracking resistance of hardfacing alloys is very poor, preheating and controlled slow cooling often needs to be adopted to avoid cracking. Selection of hardfacing process also depends on the form of filler material available. However, non-conventional weld deposition techniques like laser welding and PTAW are found advantageous over the other processes that generally used for hardfacing. Hence, the effect of GTAW and PTAW processes on the dilution, and the effect of stress relieving (SR) heat treatment on the properties of NiCr hardfacing alloys deposited on 316L SS were studied. For this purpose, E NiCr-A (Colmonoy 6, C-6) and E NiCr-B (Colmonoy 5, C-5) rods were deposited by the GTAW process and E NiCr-A (WT-60), E NiCr-B (WT-50) and E NiCr-C (WT-40) powders were deposited by the PTAW process. Specimens for metallography, hardness measurements and SR heat treatment (at 1123 K for 4 h) were extracted from the deposits. The effect of dilution on microstructure of hardface deposits was characterised by scanning electron microscopy (SEM), energy dispersive analysis of X-rays (EDAX) and electron probe micro-analysis (EPMA). The hardness profiles across the interface of GTA deposits (Fig. 6a) show that asdeposited hardness on the top surface of the C-5 deposit is 673 HV, while that of the C-6 deposit is 803 HV. However, the hardness of the C-5 deposit over a distance of about 1.5 mm from the substrate/deposit interface is only 350400 HV, which increases to 550650 HV over the next 1.5 mm of the deposit. For the as-deposited C-6 deposit, the hardness is about 575 HV over a distance of about 2.5 mm from the substrate/deposit interface, about 650 HV over the next 2.5 mm and about 800 HV over the remaining thickness of the deposit. In both the C-5 and C-6 GTA deposits, SR treatment does not seem to affect their hardness. The hardness profiles, across the interface of PTA deposits (Fig. 6b) show that in the asdeposited condition, the hardness of WT-40 deposit is 251 HV98N at the substrate/deposit interface and 350360 HV over the rest of the deposit. Similarly, the hardness of WT-50 deposit is 317 HV at the substrate/deposit interface and 445-454 HV over the rest of the deposit. The corresponding values for WT-60 deposit are 437 and 612-663 HV, respectively. It is obvious that variation in hardness with increasing distance from the interface is much less in the PTA deposits than in the GTA deposits. A marginal decrease in hardness is observed after SR heat treatment of the WT-50 and WT-60 PTA deposits. SEM images for C-6 GTA deposit with increasing distance from the interface are shown in Fig. 7. The microstructure of the deposit at 1 mm from the interface is significantly different from that near the top (8 mm from the interface). The volume fraction of blocky (dark) precipitates is very low near the interface, while both the volume fraction and the size of these precipitates increase with increasing distance from the interface. Further, near the interface, a eutectic mixture with a lamellar-like structure is present that disappears as the distance from the interface increases.

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A. K. Bhaduri and S. K. AlbertAs welded (C-5) As welded (C-6)900 800 700

SR (C-5) SR (C-6)

(a)316L SS

Colmonoy-6

Hardness, HV

600 500 400 300 200 100 -4 -2 0 2 4 6 8 10 12

Colmonoy-5

Distance Across the interface, mmAs welded (WT-60) As welded (WT-50) As welded (WT-40)700 600

SR (WT-60) SR (WT-50) SR (WT-40)

(b)316L SS

WT-60

Hardness, VH10

500 400 300 200 100 -2 -1 0 1

WT-50

WT-40

2

3

4

Distance Across the interface, mm

Figure 6. Variation in hardness across (a) GTA and (b) PTA deposits of NiCr hardface alloys.

Figure 7. SEM micrographs of E NiCr-A (Colmonoy 6) GTA deposit at different distances from deposit/substrate interface of: (a) 0 mm (at interface); (b) 1 mm; (c) 3.5 mm; (d) 8 mm.

X-ray intensity profiles for Fe and Ni of the as-deposited C-5 and C-6 GTA deposits across the 316L SS/hardface deposit interface were obtained by EPMA. In C-5 GTA deposit (Fig. 8), the average Fe count of 119 was higher over a distance of about 1.5 mm from the substrate/deposit interface than in the rest of the deposit (46 counts), while the average Ni count of 257 was lower over a distance of about 1.5 mm from the interface than in the rest of the deposit (308 counts). In the C-6 deposit (Fig. 9), the average Fe count of 75 was higher


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over a distance of about 2.5 mm from the interface than in the rest of the deposit (< 50 counts), while the average Ni count of 400 was lower over a distance of about 2.5 mm from the interface than in the rest of the deposit (500 counts). Thus, the X-ray intensity profiles for Fe and Ni across the substrate/GTA deposit interface confirmed dilution from the 316L SS substrate significantly affects the chemistry of these NiCr hardface deposits to the extent of about 1.5 mm into the E NiCr-B deposit and about 2.5 mm into the E NiCr-A deposit.

X-ray Intensity Fe100mSec (for 100 msec) Intensityfor

500 400

(a)

Base metal

300 200 100 0

Diluted layer of deposit

Colmonoy-5 deposit

0.5 -500

0 0

0.5 500

1.0 1500 1.5 2000 2.0 2500 2.5 3000 3.0 3500 3.5 4000 4.0 1000

the fusion interface (mm) Distance from the interface (microns) Distance across the line(mm)X-ray Intensity Intensity for Ni (for 100 msec) 100mSec400 350 300 250 200 150 100 50 00.5 -500 0 0 0.5 500 1.0 1500 1.5 2000 2.0 2500 2.5 3000 3.0 3500 3.5 4000 4.0 1000

Base metal


Colmonoy-5 deposit (b)

Distance across the interface (mm) Distance from the the interface (microns) Distance across fusion line(mm)Figure 8. X-ray intensity profiles for (a) iron and (b) nickel across E NiCr-B (Colmonoy 5) GTA deposit/316L SS substrate interface.

The microstructure of WT-60 PTA deposit at the interface is different from those at different distances away from the deposit/substrate interface (Fig. 10). However, there is no significant difference in the microstructure at about 2 mm from interface and at the top of the deposit (about 3.5 mm from interface), with the microstructure consisting of dendrites, carbides, borides and eutectic carbides. With increasing distance from the interface, the volume fraction of eutectic carbides decreases. The microstructures of WT-40 (Fig. 11a) and WT-50 (Fig. 11b) PTA deposits are considerably different from that of the WT-60 deposit.

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The microstructure of WT-40 deposit consists of a pro-eutectic dendritic matrix with interdendritic precipitates with rod-like precipitates being practically absent. In the case of the WT-50 deposit, the volume fraction of the eutectic phase is significantly larger than in the WT-40 deposit, with precipitates having fish-bone morphology being observed. As in the WT-60 PTA deposit, in these PTA deposits also no significant variation in the microstructure is observed with increasing distance from interface. The microstructure primarily consists of hypereutectic carbides, borides and a matrix with dendritic morphology. A comparison of the microstructure of the WT-50 PTA deposit after 1123 K/4 h SR heat treatment (Fig. 11c) with that of the as-deposited WT-50 (Fig. 9b), reveals that the SR heat treatment causes significant microstructural changes in the deposit, with the dendritic structure breaking down and the fish-bone type precipitates remaining unaltered.X-ray Intensity Fe Counts for(for Fe100 msec)300 250 200 150 100 50 0 -7 -6 -5 -4 -3 -2 -1 0 1

Colmonoy-6 deposit


(a)

Distance across the interface (mm)X-ray Intensity Ni (for 100 msec) Counts for Ni600 500 400 300 200 100 0 -7 -6 -5 -4 -3 -2 -1 0 1

Colmonoy-6 deposit


(b)Base metal

Distance across the interface interafce (mm)Figure 9. X-ray intensity profiles for (a) iron and (b) nickel across the E NiCr-A (Colmonoy 6) GTA deposit/316L SS substrate interface.

Base metal


159

Figure 10. Microstructure of E NiCr-A (WT-60) PTA deposit at different distances from deposit/ substrate interface of: (a) 0 mm (interface); (b) 2 mm; (c) 3.5 mm.

Figure 11. Microstructure of PTA deposits at 3.5 mm from deposit/substrate interface for: (a) asdeposited E NiCr-C (WT-40); (b) as-deposited E NiCr-B (WT-50); (c) 850C/4 h SR heat treated E NiCr-B (WT-50).

Results from the GTA deposits clearly indicate that both the microstructure and hardness variation observed with increasing distance from the interface can be attributed to dilution of the deposit by the substrate. The step-wise increase in hardness can be attributed to the deposition of multiple layers. Dilution from the substrate is the maximum in the first layer and hence its hardness is the lowest. During the deposition of the second layer, the molten metal mixes with the re-melted diluted first layer of the deposit and hence the effect of dilution is reduced. The results of EPMA studies (Figs. 8 and 9) are in agreement with the results from microstructural examination and hardness measurements. There is almost one to one correspondence between the hardness profile and the EPMA profiles for elements Fe and Ni for both C-5 and C-6 GTA deposits. In C-6 GTA deposits; the high-Fe and low-Ni region extends over a distance of about 2.5 mm from deposit/substrate interface, indicating the extent of dilution of the C-6 GTA deposit by the substrate material. This is approximately the same distance over which the hardness was low in this deposit. Results for C-5 GTA deposits are also similar except that the distances over which these changes are observed are lower at about 1.5 mm. The reason for the differences in the width of the diluted zones for the two deposits was not clear. Being multipass welds, it can be that these distances correspond to thickness of the first layer of the deposit at the location where both hardness measurements and EPMA analysis were carried out. As already stated, it is the first layer of the deposit that is most affected by dilution from the substrate. In contrast to the results obtained for GTA deposits, the hardness and microstructural changes in the PTA deposits are confined predominantly over a short distance of about 0.5 mm near the interface. Fairly uniform microstructure and hardness beyond this distance suggests that dilution from the substrate material is significantly low in these PTA deposits.

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Considering the various design requirements such as (i) minimising residual stresses (due to differential thermal expansion) both during deposition and service, (ii) avoiding cracking, (iii) ease of deposition and (iv) post-deposition machining etc., the thickness of the deposit recommended for finished components of the PFBR is 1.5 mm. From the results discussed above, it is clear that the hardness of GTA deposits of 1.5 mm thickness would be much lower than the minimum hardness achievable in the undiluted hardface alloy deposits. Hence, the PTAW process has been selected for hardfacing of the components. Among the hardfacing alloys considered for deposition by the PTAW process, the hardness of E NiCr-C (WT-40) alloy is quite low while that of E NiCr-A (WT-60) alloy is too high. Also, the poor weldability of Ni base alloys makes it very difficult to achieve crack-free deposits using the E NiCr-A (WT-60) alloy. Hence, hardfacing alloys conforming to AWS specification E NiCr-B has been chosen for hardfacing of the PFBR components. The hardness of the E NiCr-B hardfacing alloy also meets the minimum hardness requirement (40 RC 392 HV) specified for the hardface deposits of the PFBR components. The SR heat treatment at 1123 K for 4 h is specified for many of the hardfaced components of the PFBR to ensure dimensional stability of these components during final machining and high temperature exposure during service. Since it was reported that hightemperature hardness of NiCr hardfacing alloys reduces significantly with increase in temperature above 723 K, it was required to ensure that SR heat treatment at 1123 K does not adversely affect the hardness of the hardface deposit. The hardness of GTA deposits subjected to SR heat treatment indicates that this heat treatment does not have any adverse effect on the properties of the GTA deposits (Fig. 6a). The small differences in hardness observed between the as-deposited and SR heat treated GTA deposits, is attributed to non-uniform distribution of precipitates. However, SR heat treatment of the PTA deposits seems to have some effect on its hardness and microstructure. As seen in Fig. 11(c), the dendritic microstructure of the matrix breaks down resulting in a slight reduction in hardness (Fig. 6b). However, as the hardness reduction after SR heat treatment is only marginal, it is unlikely that the performance of hardfaced components would be adversely affected.

4. HARDFACING OF TAPER ROLLER BEARINGS OF THE TRANSFER ARMThe Transfer Arm of the PFBR, which is removable for maintenance, is designed for normal operation at 523 K and for exposure at 823 K during reactor operation. As part of the development of high-temperature liquid-sodium bearing for the Transfer Arm, the surfaces of the roller bearings were to be hardfaced for imparting adequate wear resistance to the contacting surfaces. Using a suitably optimised deposition procedure, 4 sets of the taper roller bearings (4 cups and 4 cones) were the first to be hardfaced with E NiCr-B alloy by the selected PTAW process by an indigenous manufacturer. The cups and cones of the taper roller bearing, made of 316LN SS, were received in the pre-machined condition, and Fig. 12 shows the drawings of these components along with the locations on the outer diameter (OD) of the cones and the inner diameter (ID) of the cups that were to be hardfaced. The E NiCr-B (WT-50) hardfacing alloy, with hardness of 52 RC, was used in the form of powders of size 150/+53. As the dimensions of cups and cones were


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small, no preheating of the components was carried out prior to hardface deposition by automatic PTAW process. However, the interpass temperature was meticulously maintained during the deposition and also after the completion of deposition. The deposition in each of the components was completed in about 340 seconds using two passes with 50% of the second pass deposit overlapping the first pass deposit. As the final thickness of the hardface coating is specified as 23 mm, a hardface deposit of 3.03.5 mm was made to provide allowances for rough and final machining in each component.111 2+0

65 0

+1

85 0

+2

15

+0 2

100 2 130 78 0+2

+0

35 0

+1

(a) (b)

Figure 12. Drawings of the pre-machined (a) cups and (b) cones of the taper roller bearing, with the locations to be hardfaced indicated as JJJ

After completion of deposition the components were cooled very slowly in vermiculitepowder. Then the hardface deposits on all the cups and cones were inspected by liquid penetrant test (LPT) and were found to be free of cracks. Subsequently, all the hardface deposits were rough machined and subjected to the SR heat treatment at 1123 K for 4 h. This was followed by a final round of LPT. All the hardface deposits were also inspected by ultrasonic testing, and found to be free of defects. Dimensional measurements carried out on all the hardfaced cups and cones were found to be acceptable. Fig. 13 shows the hardfaced cups and cones in the as-deposited condition.

Figure 13. Outer surface of cones (left) and inner surface of cups (right) of a taper roller bearing hardfaced with E NiCr-B alloy by the PTAW process.

5. HARDFACING OF INNER SURFACE OF GRID PLATE SLEEVESOne of the critical components of the PFBR, the Grid Plate (GP) sleeves made of 316L(N) SS that holds the core subassemblies, are to be hardfaced to prevent galling, minimize wear caused by subassembly insertion and removal and erosion due to high velocity

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of liquid sodium at 673 K. The hardface deposit on the sleeves must have good thermal shock resistance for reliable operation during the 40-years design life of the reactor, during which they would be subjected to a large number of thermal cycles due to shut downs and reactor scrams. The sleeve, an internally bored tube of about 1000 mm length, are to be hardfaced at two locations where it comes in contact with the core subassembly one on the top chamfered portion, and the other on the inner diameter at a depth of about 500 mm from either ends with the internal bore diameter at the location of hardfacing being less than 80 mm (Fig. 14).

Figure 14. Drawing of Grid Plate sleeve, showing the two hardfacing locations

When technology development of hardfacing of grid plate was taken up there was no process or equipment commercially available to carry out the job. Even attempts to fabricate the sleeve with internally hardfaced ring electron beam welded on either side could not achieve the required dimensional tolerance during welding. It was at this stage that an indigenous manufacturer designed and developed a suitable miniature PTAW torch for hardfacing of the internal surface of the sleeve. Hardfacing on the ID of the sleeve was simulated by hardfacing a 316 SS mock-up sleeve on its inside surface with E NiCr-B hardfacing alloy powder using the PTAW process (Fig. 15a). After hardfacing, the sleeve was cooled slowly in vermiculite powder, machined to the required thickness of 1.5 mm, and examined by LPT. This hardfaced mock-up hardfaced sleeve was used to study the effect of thermal cycling during service on the residual stress distribution. For this purpose, the hardfaced sleeve was cut into two halves along AB (Fig. 15a). One half (with location C at the middle) was given a SR heat treatment at 1123 K for 35 min, using a heating rate of 150 K/h and holding time of 2.5 min/mm of thickness. The other half of the sleeve (with location D at the middle) was retained in the asdeposited condition, for comparison. Both the sleeve-halves were then subjected to thermal cycling between 473 and 823 K for 20 cycles, with holding duration of 1 h at both the temperatures and immediate transfer of samples between furnaces maintained at the two temperatures. The thermal cycling temperatures of 473 and 823 K that were used correspond to the minimum and maximum temperature of liquid sodium that would be encountered. Residual stress measurements were carried out by X-ray diffraction technique on the substrate


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and across the coating. For both the half-sleeves, in-plane residual stress measurements were carried out in axial direction across the coating at three locations (Fig. 15b) in the asdeposited and SR conditions, and after 5 and 20 thermal cycles.

Figure 15. (a) E NiCr-B (WT-50) hardfaced mock-up sleeve of 316 SS and (b) one half of the hardfaced sleeve, with arrow showing direction of residual stress measurements

Axial residual stresses at all three locations across the hardface deposit on both halfsleeves (Fig. 16) showed very high compressive residual stress in as-deposited condition, due to difference in the coefficient of thermal expansion (CTE) between the E NiCr-B deposit (14-15 m/m/K) and the 316 SS substrate (17-18 m/m/K). During post-deposition cooling, the austenitic SS substrate shrinks more due to its higher CTE resulting in tensile residual stresses in the substrate and balancing compressive residual stresses in the deposit. The residual stress at the centre of the deposit (location 2 in Fig. 15b) is higher than those at the periphery of the deposit (locations 1 and 3 in Fig. 15b), because the total restraint of the substrate and deposit is higher at the centre than at the periphery. The SR heat treatment at 1123 K significantly reduces compressive residual stresses across the hardface deposit at all locations (Fig. 16a) as tensile thermal stresses generated during SR heat treatment offsets compressive stresses present in the as-deposited condition. Thermal cycling reduces peak compressive residual stress and residual stress gradient across the deposit (Fig. 16). Local yielding due to repeated expansion and contraction during thermal cycling relaxes prior residual stresses resulting in smoothening of residual stress distribution. After thermal cycling, compressive residual stresses increase at peripheral locations in the SR deposit. Differential shrinkage between coating and substrate, which depends on the cooling rate and difference in CTE, increases the compressive residual stress. On the other hand, local yielding decreases compressive residual stresses. The combined effect of these two factors results in the observed changes in residual stress in the peripheral locations. However, these changes in residual stress distribution on thermal cycling did not have any adverse effect on the integrity of the deposit, as LPT, UT and radiography of the hardfaced sleeves showed no evidence of cracking either in the deposit or at the deposit/substrate interface. The microhardness profile across the E NiCr-B deposit/316 SS substrate interface (Fig. 17) shows an appreciable rise in hardness over a distance of 0.4 mm across the interface from about 175 VHN in the substrate to about 475 VHN in the deposit. The short distance over which the hardness rises across the interface is indicative of the narrow dilution zone obtained in this PTA deposit. It is also observed that the hardness of the E NiCr-B PTA deposit varies from about 475 VHN near the interface to 500530 VHN in undiluted deposit.

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Figure 16. Residual stresses across E NiCr-B deposit in (a) half-sleeve C and (b) half-sleeve D in asdeposited, SR (only in half-sleeve C) and thermal cycled conditions.

The indigenously designed and developed miniature PTAW torch was successfully demonstrated for hardfacing deep inside the inner surface of the GP sleeves (Fig. 18). To eliminate the risk of micro-cracking and delamination of the deposit, and to minimise the magnitude of residual stress, an optimised PTAW deposition procedure was qualified. By controlling deposition parameters, groove design, preheat temperature etc. it was possible to avoid any for cracking, debonding and other form of defects on the hardface deposit. Subsequently, as a part of technology development in collaboration with fabricators, a large number of these sleeves were successfully hardfaced with E NiCr-B alloy by this procedure.

Development of Hardfacing for Fast Breeder Reactors550 500 450

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Hardness (VHN)200 gm

400

316 SS350 300

Delero-50

200 150 100 -4 -3 -2 -1 0

Fusion line1

250

Distance across the fusion line (mm)Figure 17. icrohardness profile across the 316 SS substrate/E NiCr-B (Deloro 50) deposit interface in the as-deposited condition in the grid plate sleeve.

Figure 18. Hardfacing of the grid plate sleeves using specially designed miniature PTAW torch.

6. HARDFACING OF THE BOTTOM PLATE OF THE TECHNOLOGY-DEVELOPMENT GRID PLATEThe Grid Plate was one of the components selected for technology development prior to construction of the PFBR. One of the important manufacturing activities was the hardfacing with E NiCr-B alloy by the PTAW process. Hardfacing on the inner surface of sleeves (discussed above) and the bottom plate were among the most difficult challenges that had to overcome during technology development of this component. For the bottom plate, a welded circular plate of diameter 6830 mm and thickness 65 mm, an annular outer ring of about 21 m circumferential length and about 40 mm width had to be hardfaced. The sheer area and quantum of deposition were challenging. The PTAW process was used for deposition of E NiCr-B (Deloro 50) powders. Although the hardfacing procedure for the bottom plate was qualified, when in collaboration with fabricators, hardfacing of the bottom plate of actual dimensions was taken up (Fig. 19) difficulties had to be overcome at various stages during the hardfacing of the bottom plate due to large volume of the hardface deposit.

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6.1. Initial Hardfacing on the Bottom PlateFor the initial hardfacing of the bottom plate, a 45 mm wide annular ring of OD 6420 mm and ID 6330 mm was hardfaced. For deposition of the hardfacing alloy, a groove of depth 6 mm with groove angle of 30 from the normal was machined on the bottom plate. The entire bottom plate was preheated and maintained at 723 K prior to hardfacing in a special electrical furnace. During hardfacing, three PTA machines were used simultaneously to deposit in three equally divided sectors, and the deposition was completed as a single layer using four passes. After completion of hardfacing the entire bottom plate was cooled slowly. However, LPT of the hardface deposit revealed transverse cracks at many locations during deposition, repair and SR heat treatment.

Figure 19. Bottom plate of technology-development grid plate after hardfacing of annular ring.

6.2. Modifications in Hardfacing Procedure and Groove DesignTo reduce cracking susceptibility of the hardface deposit during deposition, repair and SR heat treatment, modifications were carried out in the groove design and hardfacing procedure. To confirm the adequacy of these modifications, a mock-up circular plate of diameter 980 mm and thickness 50 mm was hardfaced. As this mock-up piece was considerably smaller than the actual bottom plate, the hardfacing was carried out in 360-mm long sectors on diametrically opposite sides leaving a gap of 100 mm, which were filled after all the 360mm sectors were deposited. LPT of the hardface deposit before SR heat treatment did not reveal any cracks; however, some porosity clusters were found at locations where the deposit sectors overlapped. The mock-up piece was then subjected to SR heat treatment at 1123 K for 2 h, and subsequent LPT revealed only one crack close to a deposit-overlap location. This crack was repaired using a pre-qualified repair welding procedure using the GTAW process. After the repair, the mock-up piece was subjected to another SR heat treatment at 1123 K for 2 h. Inspection using both LPT and UT, after rough machining, did not reveal any unacceptable indications. Based on the feedback from successful hardfacing of the mock-up


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piece, including demonstration of GTAW-based repair procedure, and discussions with the fabricators, additional modifications were made to the PTA machine for hardfacing of deposit-overlap regions, and SR heat treatment and preheat temperatures.

6.3. Second Hardfacing on Bottom PlateA second hardfacing on the bottom plate was carried out in another area an annular ring just inside the initial hardfaced ring using all the modifications to the hardfacing procedure as also the experience of the successful hardfacing of the mock-up piece. After completion of hardfacing, SR heat treatment at 1023 K was immediately carried out without allowing the job to cool down to room temperature. LPT of the deposit after SR heat treatment revealed only a few cracks. These cracks were repaired using the GTAW-based repair procedure already qualified during hardfacing of the mock-up piece. After all the cracks were repaired, the bottom plate was directly heated to 1123 K for carrying out the SR treatment. LPT after the SR heat treatment showed that no cracks were present.

Figure 20. Fabrication sequence for E NiCr-B (Colmonoy 5) alloy bushes

7. FABRICATION OF HARDFACING ALLOY BUSHESWear-resistant bushes for high-temperature application, made of hardfacing alloys are required in various components for in-sodium service in the PFBR. To substitute for very expensive import of precision castings of E NiCr-B hardfacing alloy bushes, they were fabricated using a novel procedure involving weld deposition of the hardfacing alloy on austenitic SS rods by GTAW process followed by precision machining of the hardface

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deposits (Fig. 20) [11]. Ultrasonic examination, hardness measurements, dimensional stability on high-temperature ageing, as also achieving the dimensional tolerance and surface finish on the bushes as per specification, confirmed the success of this fabrication procedure. This procedure has been successfully implemented for fabricating wear-resistant bushes for the Transfer Arm gripper assembly, and is now being transferred to industry.

8. CONCLUDING REMARKSThe developments in hardfacing technology have gained from developments worldwide, and in turn have contributed significantly to these technologies. Many challenges were faced while evolving a robust hardfacing strategy for the components of PFBR. At first, based on radiation dose rate and shielding considerations during maintenance, handling and decommissioning, nickel-base E NiCr-B hardfacing alloy was chosen to replace the traditionally used cobalt-base Stellite alloys. Also, it was demonstrated that the hot-hardness of E NiCr-B deposits after exposure at service temperatures would retain adequate hardness at end of the components designed service-life of up to 40 years. Further, based on detailed metallurgical studies, including residual stress measurements after thermal cycling, the more versatile PTAW process was chosen for hardfacing, so that the width of the dilution zone could be minimised. Hardfacing with E NiCr-B alloy by the selected PTAW process was first successfully implemented on the taper roller bearings of the Transfer Arm. Hardfacing deep inside the inner surface of the sleeves and on the bottom plate of the Grid Plate were among the most difficult challenges that were overcome during technology development, involving hardfacing inside the sleeves using an indigenous miniature PTAW torch and hardfacing of an annular ring of about 21 m circumferential length on the bottom plate. A novel procedure, involving hardfacing alloy deposition followed by precision machining, was also developed for fabrication of high-temperature wear-resistant hardfacing alloy bushes. Thus, adaptation of the hardfacing technology for PFBR, through collaborative effort with industries, has to use of semi-automatic PTAW process that has now been qualified and demonstrated for hardfacing of various technology-development components of the PFBR.

REFERENCES[1] E.Yoshida, Y.Hirakawa, S.Kano and I.Nihei, Proceedings of International Conference on Liquid Metal Technology, Societ Francaise d Energie Atomique, Paris (1988) 5021. R.A.Douty and H.Schwartzbart, Welding Journal 51 (1972) 406s. E.Lemaire and M.Le Calvar, Wear 249 (2001) 338. S.K.Albert, I.Gowrisankar, V.Seetharaman and S.Venkadesan, Proceeding of National Welding Seminar, Indian Institute of Welding, Bangalore (1987) A1. A.K.Bhaduri, R.Indira, S.K.Albert, B.P.S.Rao, S.C.Jain and S.Asokkumar, Journal of Nuclear Materials 334 (2004) 109. N.J.Allnatt and G.R.Bell, Proceedings of International Colloquium on Hardfacing Materials in Nuclear Power Plants, Avignon (1980).

[2] [3] [4] [5] [6]

Development of Hardfacing for Fast Breeder Reactors [7]

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Large Sodium Pump Design Study, Report no. WARD-3762-1, Westinghouse Advanced Research Division, USA. [8] Deloro Stellite Limited, Tribaloy Product Catalogue, Deloro Stellite Limited, Swindon, UK. [9] ASM Metals Handbook, Volume 6, 9th edition, ASM International, Materials Park, Ohio, USA (1993) 794. [10] C.R.Das, S.K.Albert, A.K.Bhaduri, C.Sudha and A.L.E.Terrance, Surface Engineering 21 (2005) 290. [11] C.R.Das, S.K.Albert, A.K.Bhaduri and G.Kempulraj, Journal of Materials Processing Technology 141 (2003) 60.



Chapter 6

TISSUE ENGINEERING OF CARTILAGE IN BIOREACTORSNastaran Mahmoudifar and Pauline M. DoranSchool of Biotechnology and Biomolecular Sciences University of New South Wales Sydney NSW 2052, Australia

ABSTRACTThe main goal of cartilage tissue engineering is to generate three-dimensional cartilage and osteochondral tissues for use in repair of large cartilage injuries. Cartilage constructs are generated by seeding and culturing viable cells in biodegradable polymer scaffolds under conditions suitable for tissue formation. In this chapter, current developments in cartilage tissue engineering are reviewed, focusing on the source of cells, the polymer scaffolds, seeding systems, bioreactors and application of mechanical stimulation for cell differentiation and tissue production. The generation of cartilage tissue constructs in the laboratory using a bioreactor system is also described. Chondrocytes were isolated from human foetal epiphyseal cartilage, expanded in monolayer, dynamically seeded into poly(glycolic acid) (PGA) polymer scaffolds and cultured in recirculation bioreactors. Composite scaffolds were used to improve the initial distribution of cells within the scaffolds and to develop cartilage constructs that were homogeneously cartilaginous throughout their thickness. The quality of the engineered cartilage was assessed after 5 weeks of bioreactor culture in terms of tissue wet weight, cell, glycosaminoglycan (GAG), total collagen and collagen type II contents, histological analysis of cell, GAG and collagen distributions, immunohistochemical analysis of collagen types I and II, and ultrastructural analysis using transmission electron microscopy.

Correspondence to: Nastaran Mahmoudifar, telephone: +61-2-9385-2086; fax: +61-2-9313-6710, e-mail: [email protected]

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Nastaran Mahmoudifar and Pauline M. Doran

INTRODUCTIONFunctional cartilage and osteochondral tissues are needed for implantation to repair large or full-thickness cartilage injuries. Cartilage in adults has a very limited capacity for self repair once it is damaged due to injury or disease; conventionally, autografts or allografts are implanted to repair the damage. However, the limited availability of autografts and the problems of immunorejection and transmission of infectious disease in the case of allografts have made tissue engineering of cartilage a promising alternative. The success of repairing small cartilage injuries by injecting autologous cartilage cells (chondrocytes) into the damaged site has been encouraging for tissue engineering, which aims to generate threedimensional cartilage and osteochondral biomaterials by seeding and culturing viable cells in biodegradable polymer scaffolds. The role of the polymer is to provide an initial scaffold for cell attachment and production of cartilage extracellular matrix (ECM). The polymer gradually dissolves and disappears as the tissue is formed. The main goal for engineered cartilage tissues is the repair of articular cartilage; however, other applications include plastic and reconstructive surgery of ears and noses.

CELL SOURCEDifferentiated chondrocytes or undifferentiated stem cells may be used to generate tissueengineered cartilage. Consistent with the clinical practice of injecting autologous chondrocytes into damaged joints to treat small articular cartilage injuries (Brittberg et al., 1994), chondrocytes isolated from native cartilage have been applied in most tissue engineering studies. As indicated in Table 1, chondrocytes from a variety of animal sources have been tested experimentally; human cells have also been used for cartilage generation. Most researchers isolate chondrocytes from foetal or juvenile individuals for cartilage production in vitro. Although better results in terms of cartilage ECM development have been reported using immature rather than adult chondrocytes (Carver and Heath, 1999b), the presence of undesirably high levels of collagen type I in engineered cartilage has been attributed to the use of foetal cells and the developmental plasticity of foetal chondrocytes in the production of both bone and cartilage tissues (Mahmoudifar and Doran, 2005a). The multipotency of adult stem cells is being exploited increasingly to produce tissueengineered cartilage. An important advantage of using stem cells rather than autologous chondrocytes for cartilage engineering is that removal of healthy cartilage from the patient is not required, thus eliminating the risk of morbidity at the donor site. Mesenchymal stem cells are present in many tissues including synovium, muscle, adipose, bone marrow and bone (Jorgensen et al., 2004; Tuli et al., 2003) and have the capacity to differentiate along multiple lineages to form chondrocytes, osteoblasts or adipocytes under the direction of appropriate differentiation factors (Awad et al., 2003; Johnstone et al., 1998; Pittenger et al., 1999; Winter et al., 2003; Zuk et al., 2001). Typically, chondrogenesis is induced using high-density cell culture in a three-dimensional environment and supplementation of the medium with growth factors from the transforming growth factor beta (TGF-) family; insulin and de

1600214819photoionization of polyvalent ions

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