Continuous Casting

III

Continuous Casting Proceedings of the International Conference on Continuous Casting of Non-Ferrous Metals

Edited by H. R. Müller

Deutsche Gesellschaft für Materialkunde e.V.

VI

I

Continuous Casting Proceedings of the International Conference on Continuous Casting of Non-Ferrous Metals Edited by H. R. Müller

Deutsche Gesellschaft für Materialkunde e.V.

II

Further titles of interest: A. Hazotte (Ed.) Solid State Transformation and Heat Treatment ISBN 3-527-31007-X D. M. Herlach (Ed.) Solidification and Crystallization ISBN 3-527-31011-8

III

Continuous Casting Proceedings of the International Conference on Continuous Casting of Non-Ferrous Metals

Edited by H. R. Müller

Deutsche Gesellschaft für Materialkunde e.V.

IV Editor: Dr. H. R. Müller Wieland-Werke AG Graf-Arco-Str. 36 89079 Ulm Germany

All books published by Wiley-VCH are carefully produced. Nevertheless, editor, authors, and publisher do not warrant the information contained in these books, including this book, to be free of errors. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.

Library of Congress Card No.: applied for British Library Cataloguing-in-Publication Data: A catalogue record for this book is aailable from the British Library Bibliografic information published by Die Deutsche Bibliothek Die Deutsche Bibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliografic data is available in the Internet at . © 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim Printed on acid-free paper Printed in the Federal Republic of Germany All rights reserved (including those of translation in other languages). No part of this book may be reproduced in any form – by photoprinting, microfilm, or any other means – nor transmitted or translated into machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, even when not specifically marked as such, are not to be considered unprotected by law. Composition: W.G.V. Verlagsdienstleistungen GmbH, Weinheim Printing: Strauss GmbH, Mörlenbach Bookbinding: J. Schäffer GmbH, Grünstadt ISBN-13: 978-3-527 31341-9 ISBN-10: 3-527-31341-9

V

Preface In 1976 the Committee for Continuous Casting of the Deutsche Gesellschaft für Materialkunde (DGM) organised for the first time a continuous casting conference for non-ferrous metals. A periodical update every 5 years has in the meantime been proven to be useful. Following the very successful previous meeting in 2000 in Frankfurt, Germany, this year's conference reviews the complete range of the processing chain covering both melt treatment and casting technology as well as specific measures for micro-structural control. A focal point of the programme deals with modelling and simulation that has become an integral part of modern manufacturing. Current progress also includes spray forming as an increasingly important processing option. For the first time poster presentations are included. Experts from the manufacturing industry, researchers and scientists from university and industry as well as suppliers of equipment and ancillary products present information on most recent technical and economical developments. The organising committee thanks all authors for their contribution and last but not least the audience for their discussion and comments.

Dr. Hilmar R. Müller

Chairman of the Organising Committee

VI

VII

Contents Melt Treatment

1

Melt Treatment of Copper and Aluminium – The Complex Step Before Casting Friedrich, B., Kräutlein, C., Krone, K., IME Process Metallurgy and Metal Recycling RWTH Aachen, Germany

3

A Study on Surface Defects Caused by Grain Refiners Keles, O., Dundar, M., Assan Aluminum, Istanbul, Turkey

23

Effect of Grain Refiner on Surface Crack of 3004 Alloy during DC Casting Morishita, M., Tokuda, K., Kobe Steel, Ltd, Moka-city, Tochigi, Japan

29

Investigation of Factors Affecting the Extent of Microporosity in an Aluminum Casting Savas, O., Kayikci, R., Sakarya University, Sakarya, Turkey Model Studies of Gas Bubbles Physical Characteristics at Inert Gas Purging into Molten Metals and Alloys Stefanoiu, R., Geanta, V., Voiculescu, I., Politehnica University of Bucharest, Bucharest, Romania

36

42

Casting Technology

49

Remarks about Process and Technology of Continuous Casting Schliefer, H., Khoury, A., Porten, M., Norddeutsche Affinerie AG, Hamburg, Germany; Wolber, P., Boller, K.H., SGL, Bonn, Germany; Dürrschnabel, W., Müller, H.R., Wieland-Werke AG Ulm, Germany; Schneider, St., Deutsche Gießdraht GmbH, Emmerich, Germany; Müller, W.H., Schwarze , M., SMS-Meer, Mönchengladbach, Germany; Oelmann, H., Rode, D., Frankenberg , R., KME, Osnabrück, Germany

51

Continuous Strip Casting of Magnesium Alloy by a Horizontal Twin Roll Caster Watari, H., Oyama National College of Technology, Oyama, Japan; Haga, T., Osaka Institute of Technology, Osaka, Japan; Koga, N., Nippon Institute of Technology, Saitama, Japan; Davey, K., The University of Manchester, Manchester, UK

70

Strip Casting of Mg-Al based alloy with Ca by Twin Roller Caster Matsuzkai, K., Hatsushikano, K., Torisaka, Y., Hanada, K., Shimizu, T., National Institute of Advanced Industrial Science and Technology(AIST), Tsukuba, Japan

77

VIII New Strip Casting Process for Magnesium Alloys Bach, Fr.-W., Hepke, M., Rossberg, A., Institute of Materials Science (IW), University of Hanover, Germany

81

Production of Twin Roll Cast AA6016 for Automotive Applications Dündar, M., Keles, Ö., Assan Aluminum, Ýstanbul, Turkey; Anger, G., AMAG Automotive GmbH, Ranshofen, Austria

87

Magnesium Upward Direct Chill Casting Bach, Fr.-W., Schacht, S., Rossberg, A., Institute of Materials Science (IW), University of Hanover, Germany

95

Spray Forming of Advanced High Strength Aluminum Alloys Krug, P., Commandeur, B., PEAK Werkstoff GmbH, Velbert, Germany

101

A Method of VDC Hot Top Mould Design and Setting of Process Conditions Bainbridge, I.F., Cooperative Research Centre for Cast Metals Manufacturing (CAST), Division of Materials, The University of Queensland, Brisbane, Australia; Grandfield, J.F., Cooperative Research Centre for Cast Metals Manufacturing (CAST), CSIRO, Division of Manufacturing and Infrastructure Technology, Preston, Australia.

106

Continuous Casting of Non Ferrous Metal Micro Wrought Shapes Bast, J., TU Bergakademie Freiberg, Germany; Bombach, E., Deutsche Solar AG Freiberg, Germany

112

Influence of Quality of Water and Surface Roughness on Quenching Rate Król, J., Specht, E., Otto-von-Guericke-University, Magdeburg, Germany

118

Electromagnetic Casting of Aluminum and Steel Billet Using Slit Mold Park, J., Kim, M., Research Institute of Industrial Science and Technology; Jeong, H., Kim, G., POSCO

124

Aluminium Alloy Strip Casting Using an Unequal Diameter Twin Roll Caster Haga, T., Osaka Institute of Technology, Osaka, Japan; Watari, H., Oyama National College of Technology, Oyama city, Japan; Kumai, S., Tokyo Institute of Technology, Kanagawa, Japan

131

Fabrication of High Purity Copper Rod with Unidirectional Solidification Structure by Continuous Casting Using Cooled Mold Hoon Cho, Duck-young Hwang, Han-shin Choi, Shae K. Kim, Hyung-ho Jo, Korea Institute of Industrial Technology,Yeonsu-gu Incheon, Korea

137

IX High Speed Roll Casting of Al Alloy and Mg Alloy Strips Haga, T., Osaka Institute of technology, Osaka city, Japan; Watari, H., Oyama national collage of Technology, Oyama city, Japan; Kumai, S., Tokyo Institute of Technology, Kanagawa, Japan

143

Simulation / Modeling

149

State-of-the-Art in the Modelling of Aluminium and Copper Continuous Casting Processes Drezet, J.-M., Computational Materials Laboratory, Ecole Polytechnique Fédérale de Lausanne and Calcom-ESI SA, Lausanne, Switzerland; Gremaud, M., Calcom-ESI SA, Parc Scientifique, Lausanne, Switzerland; Rappaz, M., Computational Materials Laboratory, Ecole Polytechnique Fédérale de Lausanne, Lausanne, Switzerland Numerical Simulation of the Growth of Interacting, Equiaxed Dendrites with a Two-Scale Model Jurgk, M., Max-Planck-Institut fur Physik komplexer Systeme, Dresden; Emmerich, H., Siquieri, R., RWTH Aachen, Germany Monte Carlo Simulation of Grain Growth in Three Dimensions Zöllner, D., Streitenberger, P., Otto-von-Guericke-Universität Magdeburg, Fakultät für Naturwissenschaften, Magdeburg, Germany Thermal Conductivity of Ternary and Multi-Component Aluminum Alloys up to and above the Melting Temperature Brandt, R., Neuer, G., Institute for Nuclear Technology and Energy Systems (IKE), University of Stuttgart, Stuttgart, Germany; Bender, W., Grün, G.-U., Hydro Aluminium GmbH, Bonn, Germany

151

162

168

174

FEM Simulation of Near Net Shape DC Billet of High Strength Al-Mg-Si Alloy Nagaumi, H., Takeda, Y., Nippon Light Metal Company Ltd., Kambara, Japan; Suvanchai , P., Umeda, T., Chulalongkorn University, Bangkok, Thailand

182

Numerical Simulation of DC casting Boender, W., Burghardt, A., van Klaveren, E.P., Corus RD&T, IJmuiden, The Netherlands

189

Continuous Casting of Hypermonotectic AlBiZn Alloys: Experimental Investigations and Numerical Simulation Gruber-Pretzler, M., Mayer, F., Wu, M., Ludwig, A., University of Leoben, Leoben, Austria; Moiseev, J., Tonn, B., Clausthal University of Technology, Clausthal-Zellerfeld, Germany

194

X Numerical Simulation of the Upward Continuous Casting of Magnesium Alloys Landaberea, A., Pedrós, P., Anglada, E., Garmendia, I., INASMET Foundation, San Sebastian, Spain New Possibilities in the Simulation of Continuous Casting Processes with WinCast-Conti Ricken, H., Technische Universität München, Institute of Metal Forming and Casting (utg), Garching, Germany; Honsel, C., RWP GmbH, Roetgen, Germany

202

209

Modeling Continuous Casting of Metal Wire Rods Chang Hung-Ju, Hwang Weng-seng, Department of Material Science and Engineering, National Cheng Kung University, Tainan, Taiwan; Chao Long-sun, Pan Wensen, Department of Engineering Science, National Cheng Kung University, Tainan, Taiwan; Lai Yi Lin, Metal Industries Research & Development Center

213

Modeling of Macrosegregations in Continuous Casting of Sn-Bronze Gruber-Pretzler, M., Mayer, F., Wu, M., Ludwig, A., Christian-Doppler Laboratory for Multiphase Modeling of Metallurgical Processes, Department of Metallurgy, University of Leoben, Leoben, Austria

219

Continuous Casting Simulation: From Solidification and Fluid Flow to the Calculation of Grain Structures Eberle, R., Wieland-Werke-AG, Ulm, Germany Mould Temperature Fields during Continuous Casting of DHP-Copper Mäkinen, M., Helsinki University of Technology, Espoo, Finland; Uoti, M., Outokumpu Copper R&D, Pori Finland Simulation of Heat Transfer and Solidification in Continuous Casting of Copper Alloys and the Effect of Fluid Flow Vapalahti, S., Louhenkilpi, S., Mäkinen, M., Väyrynen, P., Helsinki University of Technology, Laboratory of Metallurgy, Espoo, Finland

226

234

240

Micro / Macro Structure

247

Spray Forming and Post Processing of Superalloy Rings Uhlenwinkel, V., Ellendt, N., University Bremen, Bremen, Germany; Walter, M., Böhler Edelstahl, Kapfenberg, Austria; Tockner, J., Böhler Schmiedetechnik, Kapfenberg, Austria

249

Macro- and Microstructure of Spray-Formed Tin-Bronze V. Kudashov, D., Müller, H.R., Zauter, R., Wieland-Werke AG, Ulm, Germany

256

XI Influence of the Crystallization Conditions on the Microstructure and Mechanical Properties of TiAl- and Ti3Al-Based Alloys Greenberg, B.A., Kazantseva, N.V., Volkov, A.E., Akshentsev, Yu.N., Institute of Metal Physics, Ural Division, Russian Academy of Sciences, Ekaterinburg, Russia Effects of Process Parameters on the Characteristics of the Billet Sump and Related Defect Formation during DC Casting of Aluminum Alloys Eskin, D.G., Netherlands Institute for Metals Research, Delft, The Netherlands; Katgerman, L., Delft University of Technology, Dept. Materials Science and Engineering, Delft, The Netherlands Effect of Casting Speed and Grain Refining on Macrosegregation of a DC Cast 6061 Aluminum Alloy Kumar Nadella, R., Eskin, D., Netherlands Institute for Metals Research, Delft, The Netherlands; Katgerman, L., Delft University of Technology, Department of Materials Science and Technology, Delft, The Netherlands Effect of Melt Flow on Macrostructure and Macrosegregation of an Al–4.5% Cu Alloy Turchin, A.N., Eskin, D.G., Netherlands Institute for Metals Research, Delft, The Netherlands; Katgerman, L., Delft University of Technology, Department of Material Science and Engineering, Delft, The Netherlands Quenching Study on the Solidification of Aluminum Alloys Ruvalcaba, D., Eskin, D., Netherlands Institute of Metals Research, Delft, The Netherlands; Katgerman, L., Kiersch, J., Delft University of Technology, Delft, The Netherlands Numerical Study of the Influence of an Applied Electrical Potential on the Solidification of a Binary Metal Alloy Nikrityuk, P.A., Eckert, K., Grundmann , R., Institute for Aerospace Engineering Dresden University of Technology, Dresden, Germany Microstructure and Strain Distribution Influence on Failure Properties in Eutectic AlNi, AlFe Alloys Olaru, P., INAV-S.A., Bucharest, Romania; Gottstein, G., IMM, RWTH-Aachen, Germany; Pineau, A., ENSM, Paris, France

265

271

277

283

290

296

309

Dendrite Coarsening and Embrittlement in Continuously Cast Tin Bronzes Virtanen, T., Tampere University of Technology, Tampere, Finland

314

Continuous Casting of Tin Containing Alloys and their Transformation Lebreton, V., Sadi, F., Bienvenu, Y., Centre des matériaux P.M.Fourt, Ecole Nationale Supérieure des Mines de Paris, Evry, France

320

XII Suplier Session

327

Horizontal Casting Technology for Copper Products Müller, W., Schneider, P., SMS Meer GmbH, Mönchengladbach, Germany

329

Horizontal Direct Chill (HDC) Casting of Aluminium – The HE Universal Caster Niedermair, F., Zeillinger, H., Hertwich Engineering, Braunau, Austria

336

Aluminium-Semi-Continuous Casting Technic, State of the Art Brockmann , G.J., Maerz-Gautschi Industrieofenanlagen GmbH., Tägerwilen, Switzerland

344

OCP Crucible Monitoring System in Long-Term Tests Schmitz, W., Donsbach, F., Otto Junker GmbH, Simmerath, Germany; Hoff, H., Lios Technology GmbH, Cologne, Germany

353

A New Continuous Casting Process Sommerhofer, H., Sommerhofer, P., Sommerhofer Technologies, Graz, Austria

368

Author Index

377

Subject Index

379

39

Melt Treatment

40

3

Melt Treatment of Copper and Aluminium – The Complex Step Before Casting B. Friedrich, C. Kräutlein, K. Krone IME Process Metallurgy and Metal Recycling RWTH Aachen

1

Introduction

Aluminium and Copper are second and third in metal production world wide after iron and steel. But in melt treatment technologies they are probably the number one. Although Copper is the metal that is longest used technological by mankind and Aluminium one of the shortest, in both industries a variety of melt treatments and sometimes similar technologies were investigated over time. But as for most industries a look into the technology of a non competing neighbour often is missing, because contact is lost due to a lack of connection points. This plenary paper tries to bridge this gap, here the most important melt treatment techniques are presented and compared between these two metals. It is accentuated what the industries can learn from each other. To see the differences for Aluminium and Copper the most important properties of Copper are summarised in Table 1. Table 1: Properties of Copper and Aluminium [2] Property Density Melting Point Elastic modulus Ultimate yield strength (UYS) Elect.conductivity Heat Conductivity Electr. Cond./density Heat cond./density

Aluminium (high purity) 2698 kgm–3 660°C 70.3 GPa 90–100 MPa 40 MS 237 Wm–1K–1 14.8*10–3 MSm3kg–1 87,8*10–3 Wm–1K–1m3 kg–1

Copper (high purity) 8960 kgm–3 1084°C 128 GPa 210–230 MPa 64.5 MS 401 Wm–1K–1 7.2*10–3 MSm3kg–1 44.7*10–3 Wm–1K–1m3kg–1

The major difference of Aluminium and Copper is the affinity to oxygen. While Aluminium is a very un-noble element and its melt forms unsoluble oxides rapidly, Copper is considered a half noble metal but with a high solubility for oxygen in the liquid state. The major similarity is the outstanding heat and electrical conductivity of both metals. Although Copper has a 50 % better conductivity than Aluminium the conductivity to density ratio is in favour of Aluminium. This is especially in mobile applications of interest as for example heat exchangers in automobiles. In comparison Copper heat exchangers are preferably used in stationary and elevated temperature applications. For both metals partly similar melt treatment techniques were developed. Impurities like dissolved gases and solid inclusions are battled with the same principles, whereas dissolved metallic impurities have to treated differently. In this paper an overview on the melt treatment techniques of Aluminium and Copper melts are given and gas purging, slag treatment and filtration as examples are explained in more detail.

4

2

Applications of Aluminium and Copper

For increasing applications the Aluminium quality has to meet high performance specifications like foils, sheets for can bodies and offset plates as well as parts for the production of CDs. The requirements on the material in respect to cleanliness are very high, e.g. as the thickness of a can body sheet is nowadays only <300 μm. Therefore the number of defects caused by inclusions and gas pores has to be decreased drastically. More than 60 % of the Aluminium production is used for packaging from which a large part is Aluminium foils. Aluminium foils are used for the protection of food, e.g. in combination with plastic or paper for the production of juice containers. Inclusions of a size >10 μm lead to holes in the foil and cause spoiled products. Lithographic sheets for offset plates have to have a perfect surface.

Figure 1: High-tech application for Aluminium (Cans, CD, Lithography sheet, beverage containers)

The highest demands in respect to cleanliness exist for material used for the production of computer discs. At CDs Aluminium is used as material which has to reflect the laser beam. The reflective material has to have constant reflective properties so that the information stored at the disc can be read out correctly. The spot where the laser reads the information is only ~ 0.5 μm wide, so the size of surface defects should be as small and their amount as low as possible. Copper is mainly used for electrical conductivity applications. Besides the day-to-day household uses there are some fields where extreme product cleanliness is necessary. One example is the use as a cladding material for superconductors as shown in Figure 2. In case of cooling failure (most super conductors still need very low temperatures) the Copper matrix takes over the

5 conduction because of good deep temperature conductivity until the system is shut off. Contrarily to the super conductors Copper does not reveal a sharp resistance step at a certain temperature.

Figure 2: Copper in printed circuit boards (left) and super conductors applications (right)

Another example for Copper in the high-tech industry is the application in the micro electronics industry. The electronic industry boomed with the development of the so called printed circuit boards (PCBs) as shown in Figure 2: a Copper foil is applied onto a non conducting substrate; alternatively the Copper is deposited electrolytically on the substrate. Then the circuit lines are printed with a special colour, the excess Copper is etched away and the conductor lines remain. Nowadays layers of only 50 μm are applied. The boards with these thin Copper layers can also be made flexible. This is a great challenge for the material: it has to combine a high conductivity with good mechanical properties and elasticity. One last example for a Copper high-tech application is the use in connector pins, an every day’s but often not recognised application. In modern automobiles numberless electronic helpers are connected to one central site. An example of connector pins is shown in Figure 3 Therefore tens of connectors are combined in one connector system. The pins need a high conductivity but also

Figure 3: Copper strip and connectors (Bilder DKI, KME)

6 flexibility that allows to assemble and disassemble them often. At the same time they should be stiff enough to guarantee the electrical connection even at temperatures in an engine compartment and should have the same lifetime expectance as the automobile. These three fields of application are, of course, only a very small extract of the Copper world. Copper and its alloys are also used in vacuum switches, vacuum capacitors, electron beam tubes, welding electrodes, heat exchangers, as moulds for continuous casting of steel, Aluminium and Copper, in generators of power plants and wire. Aluminium and Copper finished products are sometimes of comparable shape although their application fields are different as for example Aluminium foil for food containers and Copper foil for conducting purposes. Nevertheless both products require an extremely high amount of knowledge for their production and because of the effort that was put in their development they are called high-tech products

3

Impurities in Aluminium and Copper Melts

Impurities in Aluminium melts can be divided into ”solid inclusions” and ”dissolved impurities”. Solid impurities in Aluminium have different sources. The exogenous inclusions may come from the melt environment as the refractory linings of furnaces, ladles, reactors or launders etc. Mainly these are simple oxides as Al2O3 and MgO, K-, Ca- and Al- silicates, Na-, Ca- Mg- aluminates, spinels like Al2O3.MgO or TiB2 cluster originating from grain refining. The endogenous inclusions for e.g. Al3C4, AlN or AlB2 are formed in the melt during production, e.g. in the electrolysis cell, at the melt treatment operations esp. during gas purging, or during storage and cooling down steps of the melts. Depending on the material produced the most important inclusions are Al2O3, MgO and Al4C3. Dissolved impurities may be foreign metals and dissolved gas. Foreign metals in potroom metal are Na, Li, and Ca coming from the electrolyte. Remelted metal may contain Fe, Si, and Cu as impurities. These metals can not be removed industrially and must be diluted by the addition of pure Aluminium or corresponding alloys in the casting furnace. The only dissolved gas in Aluminium melts is hydrogen, because it does not form compounds with Aluminium as other gases (e.g. nitrogen forms AlN, oxygen forms Al2O3). Compared with iron and Copper Aluminium has a rather low solubility for hydrogen (at 660 °C liquid Aluminium dissolves 0.69 ppm H and solid Aluminium only 0.039 ppm H). Hydrogen has to be removed, because bubbles originating during solidification lead to unacceptable gas pores in the produced material. Due to the rather small solubility of hydrogen in Aluminium melts its removing is a demanding task. Elements and compounds typically occurring in Aluminium melts are summarised in Table 2. The issue of impurities in Copper can be separated in two parts: impurities in primary Copper remaining or collected after refining electrolysis and impurities in secondary not electrorefined Copper scrap. The refining electrolysis produces cathodes with min. 99.995 wt. % Cu, the major remaining impurities are silver, sulphur, nickel and iron. But the contents are usually so small that they are not detrimental to the properties of Copper. The more critical elements in this sense namely hydrogen and oxygen as well as inclusions enter the primary Copper usually during the remelting and casting process.

7 In secondary materials the impurity matter is more complicated. Remelting of Copper scrap makes ecological and economical sense because the material does not have to be lead back into the energy intensive primary electrolysis. There are two types of scrap, the sorted and mostly clean production scrap which is easily reusable and end of lifecycle scrap (=”old scrap”) consisting often of a mixture of different alloys or even compounds with other metals and materials. In the process of producing a clean and specified alloy the undesired elements either have to be removed or diluted. They can e.g. form intermetallic phases in the Copper matrix and as a result decrease the mechanical properties like the ultimate yield strength and the ductility. Due to the noble character of Copper, most elements like Silicon, Aluminium and iron can easily be removed from a Copper melt by selective oxidation up to a very low concentration (activity). Physically and chemically more similar elements like nickel, cobalt, tin and lead have to be treated with more attention. Elements and compounds typically occurring in Copper melts are summarised in Table 2. Dissolved metallic impurities in small amounts mostly are not detrimental to the properties of Copper, but some elements as for example Lead and Arsenic precipitate at the grain boundaries of the Copper materials and lead to embrittlement of the material. Generally oxygen and hydrogen pick-up can lead to very negative effects. The two gases have a high solubility in liquid Copper that decreases sharply during solidification. This can lead to a bubble formation, i.e. porosity in the solid material. Oxygen can also form cuprous oxide (Cu2O) above its solubility level that immediately reacts with the moisture of the air forming water vapour during annealing or welding, this phenomenon is called hydrogen illness. Dissolved hydrogen and oxygen (or Cu2O) will react to water under extreme pressure in the lattice and will form cracks and lead to embrittlement. Solid inclusions like intermetallics or oxides from alloying elements or the refractory material usually do not have a negative impact on Copper and Copper alloys. Because the density difference between the Copper melt and the particles is very high the particles tend to float to the surface (e.g. the density of Copper at 1100°C is 7.96 g/cm³ while iron oxide has a density of 5.25 g/cm³ [2]). However Stokes law predicts that even at high density differences very small particles tend to stay suspended (In case of FeO particles the diameter smaller than 10 μm rise with 0.59 m/h). This leads to problems where wires with a very small diameter are drawn, for thin Copper foils or where small connector pins are made from thin Copper strip, being etched or punched. Generally impurities in metals can be distinguished in dissolved gases, dissolved metals and non metals as well as solid inclusions, Table 2 summarizes the typical impurity elements and compounds for Aluminium and Copper. Table 2: Comparison of impurities in Aluminium and Copper Impurity type Gas Dissolved metals

Inclusions

Less noble

Aluminium H2 Li, Na, Mg, K

Copper H2, O2 Pb, Sn, Ni, Zn, Fe, Si, Cr, etc.

More noble

Fe, Mn, Si, Cu, Ni

Au, Ag, PGM

Exogenic

Al3C4, Al2O3, MgO, Silica- SiO2, Al2O3, SiC etc. tes, Aluminides Al2O3, MgO, AlB3 Cu2O, MexOy

Endogenic

8 For both metals Hydrogen is a major problem as a dissolved gas, whereas oxygen is unsoluble in Aluminium and forms immediately solid compounds. In Copper melts the oxygen concentration can exceed 1 wt.% and is the second major problem due to the reaction with hydrogen to water vapour or with carbon to CO/CO2. Dissolved metallic impurities are generally not a problem as long as they are less noble than the target metal. But as Aluminium is one of the least noble elements the variety and amount of more noble metals is much greater than in Copper and Aluminium is difficult to clean. In Copper mainly noble metals like Silver, Gold and PGMs can not be removed from the melt, but also metals with low activities at low concentrations like lead or nickel, if very high purities have to be obtained. The high oxygen affinity of Aluminium leads to a vast formation of oxides that can harm the products. In Copper ceramic impurities are mainly from the refractory and from less noble alloying elements not being transferred to the slag.

4

Purity Requirements for High Performance Materials

Potroom Aluminium contains up to 0.3 ppm H, 150 ppm Na, 20 ppm Li, up to 5 ppm Li and more than 1000 ppm of inclusions, mainly as Al4C3. The ppm-/ppb-/ppt-concentrations of inclusions are defined as the total volume of inclusions taken as Al2O3 related to the volume of 1 kg liquid Aluminium. The impurity content of remelted materials sums up to max. 0.6 ppm H2, to 40 ppm Ca, 10 ppm Na, and some 1000 ppms of inclusions mainly Al2O3, MgO, MgO.Al2O3,Al4C3 and TiB2. Alloyed Aluminium melts supplied to the casthouse shall contain not more than 1 ppm inclusions, pure Aluminium melts up to 100 ppb inclusions. After filtration this content can be reduced to less than 10 ppb. Alloyed material for extrusions can contain between a 10–100 ppb inclusions. The highest requirements exist in respect to melt cleanliness for materials used for computer disc production. They may contain only between approx. 100–1000 ppt of inclusions. The alkali-metal content has to be lowered to a few ppms, while the hydrogen content has to be decreased to smaller than 50 ppb (0.05 ppm). The requirements of different Aluminium applications is shown in Figure 4. Examples for the increased quality requirements of Aluminium melts are can body sheets and foils. In the early 1960´s at the beginning of the production of Aluminium can bodies the thickness was approx. 0.5 mm, today it is thinner than 0.3 mm. This means a decrease of 40 % in thickness with a corresponding reduction in weight. Today Aluminium foil is rolled out to a thickness of only 6 μm, which is in the magnitude of order of typical inclusions. Without an effective melt cleaning by gas purging and filtration the requirements for the production of these materials can not be met. The most frequent impurity in refined Copper is oxygen. If the alloy is not molten under oxygen free atmosphere, Copper melts will always pick up oxygen. Copper melts hyper saturated with oxygen lead to the so called hydrogen illness during casting and welding in case hydrogen is present in the material or the atmosphere. Dissolved oxygen can also oxidise less noble elements which leads to the formation of brittle oxides in the material. In technically alloys, an oxygen content of a few hundred ppm is accepted whereas in special grade materials the content has to be lowered below one ppm, some grades and their respective tolerated oxygen content are shown in Figure 5.

9

Figure 4: Requirements on Aluminium cleanliness

Figure 5: Requirements for oxygen contents in Copper

In technical alloys the DIN limit for specified metallic impurities mainly varies between 0.1 wt. % and 0.5 wt. %. The sum of all not specified remaining elements should not exceed 0.5 wt. %. However some elements are more detrimental as for example Aluminium, arsenic and lead, their content must be kept below 0.01–0.05 wt % [3]. For high quality products limits of < 90 ppm Pb, < 50 ppm Sn and < 150 ppm Ni are required. Solid inclusions are especially a problem in thin wire drawing and thin sheets rolling. Equation 1 describes the critical inclusion size, any particle bigger will initiate wire breaks:

d critical t

D˜K Tn

(1)

10 with dcritical: inclusion size that leads to a break, D: wire diameter, K is a material constant, T: applied pulling tension, n: reduction ratio. That means that the possibility for a wire break increases sharply with the size reduction ratio and the pulling tension. In foils production inclusions lead in the worst case to a foil break. Smaller inclusions lead to surface defects and elongated holes in the material.

5

Processing of Molten Aluminium and Copper

The cleaning of Aluminium melts starts with a simple ladle treatment esp. of potroom metal for the removal of alkaline metals, before the melt is transferred into the casting furnace. There the alloying is carried out and a further settling operation may take place. From the casting furnace the molten metal is fed via a launder to the degassing unit for the removal of hydrogen. Grain refining is carried out by wire injection between the gas purging unit and the filtration station. Sometimes gas purging is combined with a filter in one unit. After the melt treatment the liquid metal is cast in a DC casting unit to billets, cakes or slugs [5, 6]. The processing of primary Copper usually starts with the melting of Copper cathodes in a shaft furnace. For melting of Copper scrap in general induction furnaces or drum furnaces are applied. After the preliminary melting furnace the melt is either casted directly or subjected to a casting furnace where the melt is stored, (alloyed) and heated to casting temperature. For Copper the continuous wire casting by casting wheels or Hazelett casters is especially important. Besides this also vertical and horizontal slab casting is applied as well as mould casting. For treatment of Aluminium melts a variety of methods are in industrial use. A cheap and simple settling procedure in the casting furnace is an easy but ineffective method to clean an Aluminium melt. Solid inclusions settle down depending on size, form and density. Because

Figure 6: Processing of molten Aluminium

11

Figure 7: Processing of molten Copper

there is only a small density difference between inclusions (e.g. Al2O3) and liquid Aluminium their settling is very slow. Small inclusions do not settle at all. According to the difference in partial pressure between hydrogen dissolved in the melt and hydrogen resp. water vapor within the atmosphere hydrogen can be removed. The return reaction:

H 2O  2Al œ 6H Al  Al 2O3 Kp

aH pH3 2 O

(2)

leads to a new hydrogen pick up and oxide formation. Hydrogen and solid inclusions can be removed only partially using this method. So settling is only rarely used as a preliminary step to treat Aluminium melts. By a ladle treatment alkaline and earth alkaline metals can be removed by mechanically stirring in salts into the Aluminium melt. Different technical solutions are in industrial application (e.g. the TAC-process). Today ladle treatment is replaced by the development of the RFI processes (see below). Gas purging removes hydrogen as well as solid inclusions, latter only partially by flotation. Also alkaline and alkaline earth metals are removed if chlorine is added to the purging gas. Melt filtration is used extensively for the separation of solid particles. Elements like Fe, Si, Mn and Cu, which may be contained in remelted metal in forbidden concentrations, cannot be removed at all and have to be diluted by the addition of pure Aluminium or corresponding alloys in the casting furnace. The ”classical” melt treatment of Copper is the oxidation by air through oxygen injection or top blowing. With this technique elements that are less noble than Copper can be removed from

12 the melt. Today this technique is often combined with a specific slag that can take up the impurity oxides and supplies a certain oxygen potential to improve the impurity separation. To remove dissolved gases especially hydrogen and oxygen the oldest method is first of all the right melt handling. Hydrogen can be removed by an excess of oxygen in the melt, oxygen is usually added by blowing air on the melt surface. After the removal of hydrogen by oxidising conditions, the melt has to be treated under reducing conditions to remove the oxygen. For the reduction since former times the so called poling by tree trunks especially birch is used [4]. This technique is still used today in some places. A new pick-up of hydrogen has to be avoided by shielding of the melt. An advantage of the tree trunks is that they have a CO2 emission of zero, as plants are considered to be regenerative. Alternatively to this classical procedure, modern techniques reduce the partial pressure in the surrounding atmosphere; this led to vacuum and gas purging technologies like for Aluminium. The removal of dissolved metals in Copper is an upcoming problem because of the increasing recycling material volume not being treated by refining electrolysis. Metals like Zinc, Arsenic and Antimony can be evaporated by a vacuum treatment, for others like Nickel, Cobalt, Tin and lead a special slag treatment is more economic. Nevertheless very often the high specifications of high tech applications cannot be met using only one refining technique. To remove solid particles from Copper melts a simple settling is usually enough for standard qualities. But the increasing requirements, e.g. for wire production, led to the development of filtration and flotation techniques. All the melt treatment techniques are actually batch processes. They have to be implemented in the existing process routes in a way that their effect is not lost before casting and solidification of the metal. That means that after deoxidation, a pick up of oxygen has to be avoided by proper casting gutters protected by a coal or coke cover or a shielding gas. After filtration a laminar flow of the melt through the launder has to be assured to avoid turbulences that promote abrasion of the refractory and formation of oxides. Although the impurities in Aluminium and Copper are different, the applicable melt treatment principles are almost the same in general gas purging, vacuum treatment, filtration and settling can be applied. The reactants of the different melt treatment technologies are shown in Table 3. Table 3: Comparison of melt treatment principles Melt Treatment

Aluminium Reactant Cl2, Ar, N2

Copper Removal of Reactant Gas purging H2, alkalines and floatation Ar, CH4, CO, O2 of inclusions Chlorination Cl2 Alkalines , Mg – – Vacuum treatment – Mg, Zn, Pb, H2 Filtration Ceramic Foam Inclusions Ceramic foam (Al2O3) (SiC) Settling ”Time” ”Time” Inclusions (Al2O3, MgO etc.) Slag treatment NaCl, KCl Inclusions (Al2O3, MgO etc.)SiO2 etc.

Removal of H2, O2,less noble metals (Fe, Pb, etc.) – Zn, As, Sb, H2, O2 Inclusions Inclusions (Cu2O, MexOy) Pb, Ni, Mn, Fe

From both metals dissolved gases can be removed by inert gas purging, where the principles are actually the same. Most dissolved impurity metals are removed from Aluminium by a treat-

13 ment with Chlorine while in Copper melts an oxidising procedure is favoured. With a vacuum treatment dissolved gases as well as highly volatile elements like Zinc, Arsenic etc. can be removed. The filtration removes solid inclusions in both metals only the type of filter material is different reflecting the different properties of the metals. Settling is a simple but standard technique to avoid solid inclusions in the melt.

5.1

Gas Purging

First mentioning of gas purging of metallic melts goes back to 1856 for steel [7]. The gas purging of Aluminium melts was mentioned first by D.R. Tullis in 1928, who used pure chlorine. Mixtures of chlorine with inert gases esp. nitrogen were developed very soon. In 1931 Koch proposed the use of a mixture of chlorine and nitrogen for the removal of Fe and Si from commercial Aluminium alloys [8]. Based on the basic work of Röntgen and Haas the chlorine/nitrogen converter was developed and set in operation in 1948 [9]. A couple of those units have been used up to the early 1960s in Europe. In 1964 the trigas mixture was developed. Carbon-monoxide was added to the chlorine-inert gas mixture to lower the Aluminium-oxide formation at the inner surface of the bubble. This enhanced the transport of hydrogen through the gas-melt interface. After the development of the gas mixtures the research was focused on the technology of gas purging. In the early times of Aluminium melt treatment simple tube lances were used to introduce the purging gas into the melt. Jet injection technology was developed already in the 1970ies where using a high-speed jet of gas is injected into the melt via nozzles. The gas is dispersed into fine bubbles and distributed in the reactor. Porous plugs were introduced in the Aluminium metallurgy in 1973. Porous plugs are mounted into the furnace technology, so their application is limited. But, they are widely spread in the Aluminium industry. At the beginning of the 1990s porous plugs were placed in launders. In the middle of the 1970ies the rotary gas injection (RGI) technology was developed by different companies nearly at the same time. The principle of this technology is the fact, that a gas stream introduced into a melt via a high speed rotor is disintegrated into very small bubbles by shearing forces. A couple of different in–line systems were developed which differ mainly in the design of the rotor [10]. These units are built in form of boxes, which can be fitted into the melt treatment line easily. They are in use in casthouses worldwide. The latest development in the middle of the 1990ies was the rotary flux injection (RFI) technology, in which salts, replacing chlorine in metallurgy, are added to the inert purging gas and are injected via a rotor into the Aluminium melt. The main target of this development is a decrease of chlorine consumption and emission. In the course of the development of gas treatment systems the chlorine consumption decreased from up to 0.7 kg Cl/t Al using lances, over 0.1– 0.2 kg/t with the RGI-system down to 0.05 kg/t in the RFI-systems. Gas purging is based on the difference in the partial pressures of hydrogen dissolved in the melt and within the bubbles of the purging gas. The purging gas, usually nitrogen or argon, is introduced into the melt by lances, nozzles, porous plugs or high–speed rotors. A bubble formed e.g. at a pore of a porous plug has a hydrogen partial pressure of nearly zero. Hydrogen atoms dissolved in molten Aluminium are transported to the bubble by convection and via diffusion through the melt-gas boundary layer. There the dissolved hydrogen atoms combine to gaseous hydrogen by chemical reaction. The ascending bubble becomes larger because the metallostatic

14

Figure 8: Principle of gas purging

pressure decreases and hydrogen is taken up, theoretically until the thermochemical equilibrium is reached. Normally the retention time of the gas bubbles in the melt is not long enough to reach equilibrium conditions. The effectiveness of gas purging operations depend on the kinetics of the reactions during the degassing process. The speed of the hydrogen removal can be described by a first order reaction and roughly by the equation: d cH dt

E ˜ x ˜ A V

(3)

Therefore the decrease of the hydrogen concentration in the melt c depends on • • • •

the retention time t of the bubble in the melt, the mass-transfer coefficient E, the melt volume V, and the mass-transfer area A (most important).

The mass-transfer area A is the total surface of the bubbles in the melt during gas purging. Consequently the formation of as many and small bubbles as possible in units is essential. Furthermore the depth of the melt is important, because the retention time of the bubbles in the melt is determining too. While using porous plugs a careful adjustment of the gas-throughput is necessary. Only at slow gas velocities small bubbles are formed; at high velocities rather large bubbles are produced because the whole plug surface acts as a bubble source (so called ”flooding”). The smallest bubbles can be produced by the application of high-speed rotor systems Lances are almost ineffective for gas purging operations.

15

Figure 9: Gas bubble surface versus bubble diameter

Figure 10: Efficiency of gas purging methods

State of the art in Aluminium melt treatment is the application of degassing boxes which are used worldwide. They are installed in-line between casing furnace and grain refining unit. Further developments of the RGI technology are the launder resp. through degassing units using also rotor systems. Compared to the earlier used degassing boxes the launder (compact) degassing technology has following advantages: • Reduced production costs by – diminished metal losses, – decreased process gas consumption and – decreased depreciation due to less expensive equipment.

16

Figure 11: In line Aluminium gas purging: SNIF Box and Alcan compact degasser

• Increased cleanliness of the melts by – decreased hydrogen content to below 0.05 ppm and – improved inclusion content of lower than 20 ppb. • Reduced chlorine emissions. • Reduced space consumption. All further developments must have the following targets: • Increasing the effectiveness of hydrogen (and inclusion) removal by – lowering the bubble size, – increasing the bubble residence time and – improving the bubble distribution in the reactor • Further decreasing of operation costs • Decreasing the chlorine emissions to zero, which is the main challenge for the future. The purging technologies of Copper alloys can be divided in those with inert gases and those with reactive gases. Purging with inert gases is based on a low partial pressure of the gas that needs to be removed. This process is diffusion controlled, i.e. the speed depends on the diffusion constant and the specific surface area of the melt-bubble interface. A diffusion controlled mass transfer can be influenced by rising temperatures (technically not feasible) and a decrease in the thickness of the Nernst layer. Such the process can significantly be intensified by an increase of the surface of the gas bubbles by appropriate gas supplying technique. The principles for the gas purging was described above. Argon and nitrogen are appropriate inert gases for purging Copper, especially for hydrogen removal. Their solubility in liquid Copper is negligible. Also gas purging methods for Copper deoxidation were developed to replace the poling with tree trunks that goes back at least to the year 1200 AD when it was mentioned in ”De Re Metallica”. In the 1960s extensive research was conducted on reactive gases with different gaseous and solid reducing agents. Reducing gases were tested like different carbon-hydrates [11], carbon monoxide as well as hydrogen. Ammonia [12] for gas purging was investigated in the

17 1970s. Also oil or coal dust were used as reducing agents, the problem of these substances was their impurity content, especially sulphur that enriches in the Copper melt. Natural gas as a reducing agent was tested by Klein [13] but the process was found to be very ineffective and slow. In comparison a reformed, partially oxidised natural gas where the carbon hydrates are reacted to carbon monoxide and hydrogen leads to a fast deoxidation of the Copper melt. This process was implemented at the Phelps Dodge refining works in the 1960s according to US-patent 2,989,397. An intensive investigation on the kinetics of the Copper deoxidation by carbon monoxide was carried out by Andreini et al. in 1977 [14]. The chemical gross equation of the deoxidation with carbon monoxide is: 2OCu  CO œ CO2

(4)

It was found that the oxygen diffusion in liquid Copper to the gas bubbles is the rate controlling step, valid in a concentration interval of 50 to 1000 ppm. Below 50 ppm the kinetics of the oxygen removal decreases sharply. This can not be explained by equilibrium reasons, it is possibly due to interactions with sulphur in the melt. This study agrees fairly good with an older study from Nanda et. al. [15] which also found a sharp decrease of the deoxidation speed below a concentration of 50 ppm oxygen. The lowest oxygen concentration in liquid Copper that could be achieved by this process was 10 ppm. As an alternative to carbon hydrates ammonia was investigated as a possible reducing agent by Henych et al. [12]. The gross equation of this reaction is: 3O2  4NH 3

2N 2  6H 2O

(5)

The reducing element in this case is the hydrogen. The developing nitrogen is an inert gas and does not dissolve in the Copper melt. The lowest oxygen concentration in liquid Copper that could be achieved by this process was 200 ppm. This technique never reached industrial scale. Up till now the favoured gas injection technology in the Copper industry are tuyeres/injectors and top blowing lances even though in other industries different gas supplying techniques are used as rotary vaned dispersers and porous plugs

5.2

Slag Treatment

A slag treatment of Aluminium and its alloys is usual only necessary when very fine particles like chips are molten. The natural Aluminium oxide skin prevents the burning of Aluminium melts up to 700 °C. In Aluminium recycling where fine fractions are molten usually salts on the basis sodium chloride, potassium chloride are used. The mixture depends on the local deposits and suppliers. This system is selected, because the melting temperature is close to the melting temperature of Aluminium but at the same time the system has a rather high evaporation point. This mixture has a better wetting behaviour for oxides that for Aluminium metal and therefore takes up oxides during a melt treatment and last but not least it is readily available and cheap. Usually to the salts on chlorine base fluoridic compounds like AlF3 or CaF2 etc. are added. The fluorides accelerate the cracking of the Aluminium oxide layer and therefore improve the coalescence [15]. Oxidic slags are not used in the Aluminium industry because of their high melting

18 point and also because Aluminium is very un-noble and tends to reduce most of the slag oxides as for example Silicon oxide. Slags in Copper refining have the task to take up the impurity oxides, formed during the oxidation procedure. Their properties should be: • • • • •

High solubility for impurity oxides Low solubility for Copper and Copper oxide Melting temperature close to the melting temperature of Copper High thermal stability Low interaction with the refractory material

Before or during a slag treatment the Copper melt has to be oxidised for example by top blowing of air. A slag that is industrially used is the so called ”Fayalite” slag, which is also applied in the primary Copper metallurgy. It is based on the system FeO–Fe2O3–SiO2 and effective for the gross removal of Cd, Fe, Pb, Sn and Zn especially for elements with the valence of two that can be trapped as silicates. This slag can be used in furnaces with silica refractories [17]. This is the slag used mainly today. On laboratory scale other slag types were investigated as well, for example Calcium-ferrite slags which are based on the ternary system FeO-Fe2O3-CaO. They remove the elements Al, As, Fe, Sb and Sn, especially elements which exhibit an acidic character in a slag [17] at their highest oxidation levels. This slag can be used with base refractories. Another system is based on CaF2-CaO-MgO-SiO2 this slag shows the same behaviour as the ”Calcium-ferrite” slag but with a lower solubility for Copper oxide. Also investigated by researchers were salt slags that are mainly based on Sodium carbonate (Na2CO3) but also on other alkaline carbonates like Lithium carbonate (Li2CO3) andPotassium carbonate (K2CO3). They are very effective refining slags, but they all attack ”usual” refractories of the Copper industry. It was found that by fluxing with a Sodium carbonate slag the amount of arsenic and antimony could be lowered below 0.1 ppm. The problem that occurred was, that the binary solution of Sodium carbonate and Antimony oxide let to an increase of solubility of Copper oxide in the slag and therefore to high Copper losses [18]. For fluxes on other salt base mainly fluorides can be suitable because their melting and evaporation point is very high. Here especially systems based on Calcium fluoride (CaF2) mixed with Aluminium fluoride (AlF3) or Sodium fluoride (NaF) are possible candidates. All fluxes operate in the best way when the reaction surface between metal and slag is increased for example by solid flux injection in the metal melt [19] instead of an addition to the surface.

5.3

Filtration

In 1935 a procedure was proposed for the filtration of light metal melts by DEGUSSA, which was transferred to Aluminium melts very soon. The bed filtration (BF) was developed using bulk petrol coke and/or ceramic particles by ALCAN in the 1940´s. The development of ceramic foam filters (CFF) started in the beginning of the 1970´s by SELEE. First rigid media filters (RMF), which are called also bonded particle tube filters (BPF), appeared on the market in the 1980ies, but were initially not accepted by the aluminum industry. In the 1990ies two stage fil-

19

Figure 12: Principles of melt filtration

ter systems were developed having a much better particle removal efficiency. The latest advances in filtration technology is the development of surface active filter systems starting in the mid 1990s. By the formation of active surfaces inside the filter itself the effectiveness for the separation of small inclusions was significantly improved. For the filtration of molten metals the same laws apply as for aqueous dispersions. Two different kinds of filtration have to be distinguished: cake and bed filtration. Usually both filtration types occur combined and happen successively. In the case of cake filtration the filtration process itself happens at least at the beginning by sieve effects. First inclusions larger in size as the pore diameter of the filter settle on the filter surface forming a thin layer. The thickness of the cake increases as more melt flows through and more inclusions are separated. For Aluminium melt treatment cake filtration is rather unusual and limited to melts with high inclusions contents (> 200 ppm) and larger inclusions. Bed filtration is the common mechanism used for Aluminium melts. In this case the separation of inclusions from the melt is rather complex. It happens mainly by direct collision or adhesion of particles to/at the filter surface, sedimentation by gravity as well as by inertia forces, collision of particles by Brown´s movement or/and fluid dynamic effects. Up to now no closed theory exist of the filtration of Aluminium melts [20]. So a mathematical modeling, which would allow calculating filtration efficiencies, filtration times, filter sizes etc. is not yet possible. The filter materials are generally refractory material, preferentially Al2O3. They can be distinguished between bed filters (BF), ceramic foam filters (CFF) and rigid media resp. bonded particle tube filters (RMF resp. BPF). Bed filters (BF) are bulks built of Al2O3-balls or chips with a size of 2–8 mm. Bulks of carbon or coke are not more used. BF´s are separate in-line units, which need rather much space. They are built in externally heated boxes and are suited for the throughput of large amounts of melts up to 1000 t. Particle form and size, layer thickness, and the sequence of different layers are varied to improve the filtration effectiveness. BP filters are suited for the separation of small inclusions < 20 μm from melts with low inclusion concentrations.

20

Figure 13: Filter devices for Aluminium melt filtration (from left to right: CFF, RMTF, BF)

Ceramic foam filters (CFF) consist of a labyrintic structured ceramic material in which a very effective cleaning of the Aluminium melt happens by deep bed filtration effects. They are produced by the infiltration of a ceramic sludge into porous polyurethan foam. During firing the plastics decomposes and the porous ceramic remains. CFF are also built in separate boxes, which must not be heated externally. Filter plates are commercial available in different sizes, thicknesses (normally 50 mm) and pore sizes between 10 to 80 ppi (pores per inch). They are one-way products and rather cheep so that the operational costs are low. This filter type is used most often in Aluminium metallurgy [21]. Rigid media filters (RMF) consist of porous ceramic tubes, which are built in, as BF´s in external heated boxes in form of pipe bundles. The melt flows from the outside to the inside of the pipes. The filtration processes are very similar to the CFF`s. Because they have a smaller pore size there is a higher pressure drop and they have a shorter lifetime. RMF´s are rather expen-

Figure 14: Technology of Aluminium melt filtration

21 sive in respect to investment and operation cost, therefore their application is limited to special applications. State of the art is the application of deep bed filtration in casthouses, where large amounts of the same alloy have to be cleaned. For general purposes CFF´s are used. Only for special applications RMF´s are in operation because they are most efficient for the removal of very small particles. In normal filter systems single CMF plates or combination of CMF´s with different pore sizes are built in in separate boxes which must not be heated externally. They are positioned in-line directly before the DC unit. The melt is allowed to run through the filter plate, mostly downwards. CMF´s are used in combination with BF´s and with degassing units, too. Additional ceramic filter clothes may be used at the DC casting unit to retain coarse impurities which can come into the melt after filtration. Targets for the filter development in future are filters that can remove even finer particles with high efficiency at a reasonable pressure drop and with minimized metal losses. Filtration of Copper is not very common, because the density difference between the Copper melt and the characteristic solid inclusions is big enough that they settle in the casting furnace. But because of the increasing demands on the semi finished products as for example Copper strip with 50 μm the amount of very small particles < 10 μm that do not settle easily has to be reduced. In 1981 the filtration of Copper melts was first mentioned by Chia et. al. Today filtration of Copper melts is mainly applied for material that is drawn into wire [22]. The filtration of Copper melts will increase with a further demand on wire and strip thinness.

6

Summary

In this paper the today melt treatment techniques for Aluminium and Copper are presented. Emphasis was laid on gas purging, slag treatment and filtration because these techniques are widely spread in both industries. Until now the development of melt treatment techniques always was in step with the demands of the semi finished product producers and also with the demands for

Figure 15: In line Aluminium melt treatment: Hydro (VAW)-Filter

22 environmental safety. But the future holds more challenges. Because of size reduction foils and wire diameters will further decrease. For the Aluminium industry the avoidance of Chlorine is a major point that needs to be solved. For Copper the removal from dissolved elements from scrap charges has priority to avoid the energy intensive refining electrolysis. The principles of melt treatment give both industries guidelines in which direction to proceed and to produce the optimum metal quality.

7 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]

References G. Armstrong Smith, Journal of the Institute of Metals, 100, 1972, 125–130 R. Blachnik, D´Ans Lax: Tachenbuch für Chemiker und Physiker: Elemente, anorganische Verbindungen und Materialien, Minerale; Springer, Heidelberg 1998 CEN/TS 13388:2004 E. Brunhuber, Schmelz- und Legierungstechnik von Kupferwerkstoffen”; 1959, Schiele und Schön P. Waite, in Light metals 2002 (Ed.: W. Schneider), TMS, Warrendale Pa. 2002; 841 ff W. Schneider, H.-P. Krug, N. Keegan, Aluminium, 43, 1998, 40 ff Steward, Christopher, Forward, US Patent 23123 (1859) W. Koch, Zeitschrift für Metallkunde, 1931, 34, 95 ff W.Dautzenberg, Metall, 1950, 4, 125 ff K. Krone, Aluminium Recycling, VDS, Düsseldorf 2000 F.E. Brantley, C.H. Schack: Bureau of mines R.I. 6113, 1962 R. Henych, F. Kadlec, V. Sedlacek, Journal of Metals, 1965, 386–388 L. Klein, Journal of Metals, August, 1961, 545–547 R. J. Andreini, J.S. Foster, R. B. Phillips, Metallurgical Transactions B, 8B, 1977, 633–638 C. R. Nanda, G. H. Geiger, Metallurgical Transactions, 2, 1971, 1101–1106 M. Gerke, Dissertation RWTH Aachen, Shaker, 2002 J. Gortais, F. Hodaj, M. Allibert, J.-M. Welter, Metallurgical and Materials Transactions B, 25B, 1994, 645–651 H. Fukuyama, T. Fujisawa, C. Yamauchi, Metallurgical Processes for the Early First Century, The Minerals, Metals & Materials Society 1994, 443–452 B. Zhao, N. J. Themelis: “Kinetics of As and Sb removal from molten Copper”; EPD Congress, The Minerals, Metals & Materials Society, 1995, pp. 515–524 B. Hübschen Dissertation, Shaker, 2004 W. Schneider; Gießen von Aluminium-Werkstoffen; VAW Aluminium AG Chia et.al., Us Patent 4277281, 1981 B.P. Kamath, R. Adiga, L.K. Sharma, S. Gupta, Copper2003 – Cobre 2003, Fifth international Conference, 719–727, 2003

23

A Study on Surface Defects Caused by Grain Refiners O. Keles, M. Dundar Assan Aluminum, Istanbul, Turkey

1

Abstract

The role of grain refiners has been gaining importance in improving characteristics of as-cast and end wrought products. Ti-Al-B master alloys have been used for grain refining in aluminum castings. Agglomeration of the grain refiners causes streaks, which especially affects the surface quality of sheets. Nowadays, customer requirements on surface quality of aluminum sheets have been getting more and more demanding. In order continuous aluminum casting technology to compete with DC cast counterparts, products have to be free of surface defects and cost effective. Although, there are many sources for aluminum sheets surface defects, in this study the defect caused by grain refiners has been examined. Cast house practices showed that these agglomerates are likely to settle down in launders, nozzles and plugs the filters used. In this study, samples, gathered from ceramic foam filters, casting nozzle, launders as well as end product have been investigated by utilizing an optical and electron microscopy in order to show agglomerations and their effect on the surface of products.

2

Introduction

In aluminum alloys TiAl, TiAlB, and TiCAl base products have been used for grain refinement for years. There have been several theories proposed to understand the grain refinement mechanisms in aluminum alloys. Some studies have been focused on the production and application methods of these alloys as well as compositions in order to make the grain refinement process efficient and to avoid defects in cast products [1–14]. The chemistry, sizes, shapes, and distributions of grain refiner constituents, feeding rate as well as feeding position play important roles in reaction efficiency, agglomeration and settling behavior of grain refiners. It is known that, as TiBAl based grain refiners consists of TiB2 and TiAl3 phases. Even though TiAl3 dissolve in time, TiB2 phase is known to be stable and abundant in aluminum melts [1,4,9,10,15–17]. The sizes of these phases and their concentration in grain refiners vary depending on composition as well as production methods [5,7,13,17,18]. TiAlB base products have a tendency to agglomerate and to settle down after a certain period of time. In aluminum casting, continuous casting technology has become preferable due to its low capital investment costs, reasonable productivity measures and low product costs. Technological developments in continuous casting have been focused on reducing strip thicknesses. Twin roll casting technology offers to cast thin gauge strips [19–21]. In order to utilize the advantages of this technology, process and product assurances have to be provided.

24 Continuous casting being younger than direct chill casting technology, grain refinement applications are less understood [22,23]. One of the main problems seen in grain refining mechanism is the agglomeration of grain refining particles in the components of molten metal delivery system [24,25]. These agglomerates can block filters, settle down in launders and caster tips by leading to feeding problems during casting. In addition to this, even flush into cast strips, result in surface defects in aluminum sheets and foils. In this paper, to illustrate the complications arising from grain refiner agglomeration filters, caster tips and launders as well as end products are examined.

3

Experimental Procedure

In the present study, 150 tons of AA1050 and AA3003 aluminum alloy is cast by utilizing Fata Hunter“ Speed caster. Before a casting campaign is begun, launders are cleaned and a new ceramic foam filter and a caster tip are introduced. Launders, ceramic foam filters and caster tips are made of silica, alumina, alumina-silica based products respectively. During casting, TiAl5B1 grain refiner rod is fed into a molten metal at the filter box before ceramic foam filter with the feeding rate of 40 cm/min for a strip having 2120 mm width. After 150 tons casting, samples are gathered from the metal left in the launder, the ceramic foam filter and the caster tip. These samples are mounted, grinded and polished for OM examinations and then coated with gold, investigated by utilizing SEM &EDS. Furthermore, a sheet and a foil sample are investigated in order to see the faults produced by grain refiners on an aluminum sheet and also a foil.

4

Results and Discussion

In Figure 1, the pictures of the launder, the ceramic foam filter and the caster tip utilized for the present study is shown. First visual inspection of the components has shown that bottom of the launders and caster tip have experienced heavy lump and cluster formation. In all parts, their sizes differ depending on grain refiner feeding rate, casting speed, the amount of metal cast in one campaign. The sizes of clusters for this study are found to be in a range of 2–20mm in height. Since TiB2 has a greater specific density than that of aluminum and also has a tendency to agglomerate, they settle down at the bottom of the materials used for transfer. Although feeding is conducted before the ceramic foam filter, these clusters are observed in the launder as well as in

Figure 1: a) Launders b) ceramic foam filter c) caster tip

25

Figure 2: a) filter, b) tip section, c) tip corner, d) launder, e) and f) tip section polarized

the caster tip which are located after the filter. It is noteworthy to mention that during casting these clusters has no or light attachment to the transfer media. During casting, these clusters may be dislodged and introduced into the cast strips. Even small amount of particles escape into the cast strips can cause defects as the strips are rolled. High concentration of Ti compounds and other insoluble particles tend to form a reticulated lacey pattern. These pattern can be seen in each materials consumed for this study. In Figure 2 the cross section of the filter, the clusters taken from the metal left inside the tip and the launder as well as the metal left at the edge of the tip can be seen. As noticed in Figure 2, the characteristic feature in all microstructure is having a lacey network. These networks are also seen in Mabry et al.’s study [26]. In all SEM studies, these boundaries are found to have oxygen, titanium, iron, manganese and vanadium. In Figure 3, the sizes of TiB2 particles vary from 0.5–4 Pm. There are also needlelike structures identified to be Al-Fe-Mn phase (see Figure 4). Similar, phases are seen in several studies made for AA1050 alloy. Since in the campaign without changing the transfer media both AA1050 and AA3003 alloy cast, it is likely to see Al-Mn-Fe intermetallics instead of AlFe ones [15,27,28]. It is well known that fading occurs on prolonged holding times and grain refinement loss can be reversed by agitation of the melt [29,30]. With the help of agitation TiB2 particles can be suspended into the melt and start nucleation process for a second time. This could be the way to increase the efficiency of grain refiners but it is important to note that agitation can also cause to wash out the agglomerates in products cast. Decreasing the feeding rate of grain refiners could be another solution to prolong the settlement period and the quantity of agglomerates in the transfer media. The tradeoff, in this case, should be made between the quality of coils and the quantity of grain refiners. Lastly, with changing the sizes of constituents in grain refiners, the shapes of grain refiners and finding a way to suspend TiB2 little longer in the melt could make possible to increase the reaction time as well as grain refiners’ efficiency [5,30]. Due to vibration or erosion in caster tips, the titanium compounds, believed to be TiB2, flush into cast strips and hot rolled with aluminum. It is hard to distinguish these insoluble particles in

26

Figure 3: EDS analyses a) a cross section of the filter b) the metal left in the tip section

as cast strips. The processes applied in lithographic sheet production easily reveal those defects existing on the surface. They become extremely noticeable as black streaks on the surface after brushing, etching, and anodizing operations. In Figure 5 a lithographic sheet having a black streak on its surface and related optical microscope and EDS analyses are given. The width of the streak is in a range of 1–3mm and the length is approximately 150–200 mm. As the strip gauge decreases, the clusters bring about a much negative influence. Contrary to the ductile aluminum matrix that can be easily elongated during rolling, TiB2 particles have limited ability to be deformed. This mechanical incompatibility leads to premature strain accumulation within the TiB2 clusters and results in early failure initiated from those areas. While they can be observed as very small crack on the surface of sheet products, extreme cases results in

Figure 4: An EDS analysis of needlelike structures

27

Figure 5: Micrographs of a lithographic sheet. a. image scanned, b.OM picture, c. EDS analysis

Figure 6: SEM and EDS analysis of foil having a pinhole

high pinhole counts and at the end rupture in foils. In general, the pinholes display consecutive “v” shape appearances (see Figure 6 ).

5

Conclusions

The influence of AlTiB grain refiners in the filter, the launder and the tip used for the casting of AA1050 and AA3003 alloy campaign have been examined. It is concluded that: Grain refiners tend to form clusters on the filter, the launder and the tip used. The sizes of these grain refiner accumulations may change from 2 to 20 mm in height. Their size may change depending on the amount of aluminum cast and grain refiner feeding rate. These clusters have TiB2 and Al-Fe-Mn phases in the metal left after casting. TiB2 particles have formed a lacey layer structure and the intermetallic phases seen in the grains formed by TiB2 particles.

28 Especially in continuous casting, as these cluster are introduced into the products due to changes in metal flow, they cause a black streak in lithographic sheets while showing pinholes and ruptures in foils. Further studies are being made in order to eliminate or reduce the amount of clusters by understanding the role of contact time, feeding rate and grain refiner type in twin roll casting.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

References P. Schumacher, A.L. Greer, Light Metals. 1995, 869–877 J.A. Meggy, D. A. Granger, G. K. Sigworth, C. R. Durst, Light Metals. 2001, 943–949 M. A. Kerans, S. R. Thistlethwaite, P.S. Cooper, 125th Annaheim Annual Meeting. 1996 J.J. Del Campo, M. Martin, L. Galan, Light Metals. 1999, 797–802 K. Venkateswarlu, B.S. Murty, M. Chakraborty, Materials Science and Engineering. 2001, A301, 180–186 A.M. Detomi, A.J. Messias, S.Majer, P.S. Cooper, TMS. 2001 L. Peije, E.G. Kandlova, A.G. Makrenko, V.I. Nikitin, Y. Zhang, A.R. Luts, Materials Letter. 2004, 58, 1861–1864 A. Hardman, F.H. Hayes, Materials Science Forum. 1996, 217–22, 247–252 H. Jin, T. Sagstad, TaftI, P.T. Zagierski, Materials Sience Forum. 1996, 217–222, 241–246 M. Johnsson, Light Metals. 1993, 769–777 L. Backerud, S. Yidong, Aluminum 67, Jahrgang 1991,780–785 P. Cizek, B.J. Mckay, P. Schumacher, Continuous Casting. 2000, 251–256 G. T. Campbell, S. A. Danilak, Light Metals.1992 K.T. Kashyap, T. Chandrashekar, Bull Mater. Sci. 2001, 24, 345–353 X. G. Chen, Light Metals. 1999, 803–809 T.E. Quested, A.L. Greer, Acta Materiala, 2004, 52, 3859–3868 C. Limmaneevichitr, W. Eidhed, Materials Science and Engineering. 2003, A355, 174–179 Venkateswarlu, S.K. Das, M. Chakraborty, B.S. Murty, Materials and Engineering. 2003, A351, 237–243 O. Daaland, A.B. Espedal, M.L. Nedberg, I. Alvestad, Light Metals 97. 745–752 P. Vangala, R. Duvvuri, D. Smith, C. Romanowski, TMS 1993, 333–349 M. Cortes, LightMetals. 1995, 1161–1164 P. S. Cooper, P. Fisher, Light Metals 1994 M. Yun, S. A. Lockyer, J.D. Hunt, Light Metals 2000. 857–861 D. J. Pearson, P. Cooper, 6th Australian Asian Pasific Conference. 1999 S. R. Thistlethwaite, P. Fisher, TMS, 1995, 249–259 G. Mabry, J. Kaems, D. Granger, W.C. Setzer, Light Metals. 1997, 983–989 M.W. Meredith, A.L. Greer, P.V. Evans, R.G. Hamerton, Light Metals. 1995, 811–817 T. Stucczynski, M. Lech-Grega, Light Metals. 2003, 961–968 P.S. Mohanty, R.I.L. Guthrie, J. E. Gruzleski, Light Metals. 1995, 859–867 C. Limmaneevichitr, W. Eidhed, Materials Science and Engineering. 2003, A349, 197–206

29

Effect of Grain Refiner on Surface Crack of 3004 Alloy during DC Casting M. Morishita1, K. Tokuda2 1

Kobe Steel, Ltd, Kobe-city, Hyogo, Japan Kobe Steel, Ltd, Moka-city, Tochigi, Japan

2

1

Introduction

An ingot cracking of aluminum alloy during DC casting is one of the most serious problems as less productivity. Therefore to control casting conditions and micro/macro-structure becomes very important to relax stress in the ingots, because the cracking has been occurred by a solidification shrinkage, a thermal contraction and a stress concentration within a mushy state. Remarkable progress has been made over the last few years in the numerical simulation of casting and solidification process, and the methods also have had the possibilities to estimate a prediction of stress with good precision [1,2]. However, without quantitative data of the mechanical behavior of alloys at high temperatures particularly in the mushy state, it is not possible predict to whether a cracking during casting will occur for a given stress and deformation [3,4]. Generally, even if the alloys are the same, ingots of finer grain materials hardly cause cracking problems. But fracture criterion is often not described quantitatively in the literature. Therefore 3004 alloy used as can body was taken up in the present study and the objectives was settled next threefold: the first is to identify the crack initiation and clarify the cracking behavior, the second is to quantify the effect of grain size to mechanical properties in the mushy state and the third is to clarify the criterion of ingot cracking used by numerical simulation in some casting conditions.

2

Observation and Measurement of Mushy State Mechanical Properties

2.1

Observation of Surface Cracking Within a DC Ingot

A surface crack of the actual DC ingot is shown in Figure 1. A fracture surface of the ingot is shown in figure 2. The minute crack on the ingot surface becomes initiation and the surface minute crack progresses to the centerline of the ingot. Figure 2 also shows the surface cracking was occurred in the mushy state since the fracture near the ingot surface is covered by dendritic morphology. It indicates that thermal contraction at the ingot surface is higher than the elongation of a material in high temperature. It is very important for cracking prevention to control the surface strain level by casting conditions and material properties especially elongation by grain size.

30

Figure 1: Surface crack of an actual ingot 600mm x 2000mm about 7000mm length

2.2

Figure 2: Shape of fracture surface and micro structure (600mm thickness)

Micro-Structure inside Ingot Surface with/without Grain Refiner

Figure 3 shows that the difference of the cracking behavior between fine and coarse grain sample used by RING test method. Grain refiner which was Al-5%Ti-1%B rod was added 50ppm Ti into the molten metal to prepare "fine grain" sample. Even if the same test condition, „fine grain“ sample is obviously less breakable. In order to cast different grain size samples, the same grain refiner condition was applied to the actual ingot casting. Figure 4 shows the micro-structure inside ingot surface with/without grain refiner and they are also different grain size. Chemical compositions of two different grain boundaries were measured to predict melting point by Thermo-calc (Table 1). The coarse grain sample has more micro-segregation than fine one. The calculation results show that coarse and fine grain samples are about 26 °C and 15 °C

Figure 3: Difference of cracking behavior between with and without grain refiner used by RING tester

31

Figure 4: Difference of microstructure inside the DC Figure 5: Schematic of semi-solid tensile testing instruingot surface between with and without grain refiner ment (600mm x 2000mm ingot)

lower solidus temperature at grain boundary than average solidus temperature of 3004 respectively. Table 1: Chemical compositions of 3004 alloy and grain boundary, and calculated solidus temperatures (mass%) 3004 alloy (average) grain boundary

2.3

fine grain coarse grain

Fe 0.45 1.77 1.90

Si 0.25 0.57 0.88

Mg 1.20 2.16 3.37

Mn 1.10 1.87 1.77

0.2 0.1 0.0

solidus temp. (calc.) 617 oC 602 oC Cu

Experimental Procedure to Measure Mechanical Properties of Mushy State

Two different grain specimens are taken from ingot surface, and strength and elongation measurements were carried out using a developed semi-solid tensile testing instrument which is shown in Figure 5. In order to minimize microstructure changes that occur during heating, a resistance furnace is employed to attain rapid heating. Target temperature of each tensile test is ranging from 500 °C to 620 °C. Strain rates are 0.01 s–1. Dimension of a specimen is 68 mm u 8.0 mm diameter and gauge length is 45 mm. To prevent liquid dropping off from mushy area, specimens are covered by a 8.1 mm inside diameter quartz tubes with thermocouples.

32

3

Results

3.1

Mechanical Properties of Mushy State

The stress and strain values vary significantly with temperature and grain except under solidus temperature as shown in figures 6 and 7. They indicate that fine grain materials have high elongation values and coarse grain ones have no elongation between 3004 solidus temperature and segregated solidus temperature. Accordingly the finer grain material must be hardly cracked even if the ingot are distorted much. Numerical simulation which is to predict distortion and stress are carried out based on the above semi-solid tensile test data below 3004 solidus temperature.

Figure 6: Ultimate tensile strength of 3004 alloy as a function of temperatures at two different grain material

Figure 7: Elongation of 3004 alloy as a function of temperatures at two different grain material

4

Discussion

4.1

3D Numerical Simulation for Cracking

In advance of three dimensional (3D) simulation, unknown boundary conditions such as heat transfer coefficient between a mold and melt interface were fit into actual experimental data with two dimensional simulation. Then the boundary conditions are applied to 3D simulation in order to predict a thermal and stress histories. The code is ABAQUS and boundary condition is shown in figure 8. Since the distortion and stress are strongly effected by not only solidification behavior in the cross section but also vertical temperature variation, a model was designed as follows. Firstly molten metal shape into rectangular prism, secondly the rectangular prism moves through the water cooling mold region and direct chill region, and the thermal distribution are calculated in each time step and finally stress and distortion are calculated based on previous thermal results. The dimensions of ingot are 2000 mm width, 600 mm thickness and 4000 mm length. The casting rate is 50 and 70 mm/min. And amount of water flow rate is 3.9 and 1.5 l/min.cm. Figure 9 shows the example of temperature and distortion distribution. In the figure, location

33

Figure 8: Boundary conditions calcurated by 2D numerical simulation

Figure 9: Example of 3D simulation results : Rolling surface, 70 mm/min, 3.9 l/min cm(a) distortion distribution of horizontal direction (b) temperature distribution

2/4 W, 1/4 W and 0/4 W are at the center, the quarter and the edge respectively. These locations were chosen because during the casting cracks are most likely to start at these locations.

4.2

Criteria of Cracking Depend on Grain Size and Cooling Conditions

As the ingot cracking progressed from surface of the ingot, calculated surface distortion is compared with measured high temperature strain data. In this comparison, liquid at grain boundary is assumed to be movable freely, and the stress is also assumed to start to build up at 3004 solidus temperature 617 °C. Therefore calculated distortion values are reset at the temperature to

34

Figure 10: Comparison between calculated distortion value and experimental strain data (70mm/min, 3.9 /min cm)

compare with tensile test data. Rested upon the comparison, it is possible to evaluate the material stability for cracking. The comparison between ingot distortion and material strain is shown in figure 10. Whereas the strain of fine grain material is always higher than surface distortion of the ingot at all locations, the strain of coarse grain is lower than surface distortion between 3004 solidus temperature and solidus temperature of grain boundary. Consequently it is clarified that the fine grain materials are stable against cracking in high temperature and the adding grain refiner is very effective to prevent the crack. Figure 11 and 12 show results of low speed casting and low water cooling respectively. The two results indicate that both conditions will reduce the distortion. It is possible to conclude that less crack will occur when we cast by low speed and low water cooling in the mold, because of less distortion in high temperature state.

Figure 11: Comparison between calculated distortion value and experimental strain data (50 mm/min, 3.9 /min cm)

Figure 12: Comparison between calculated distortion value and experimental strain data (70 mm/min, 1.5 /min cm)

35

5

Conclusions

1. The minute crack on the ingot surface becomes initiation and the crack progresses to the center line of the ingot. 2. The fine grain materials which are with grain refiner have less micro-segregation and high elongation in high temperature. Therefore fine grain material must be hardly cracked even if the ingots are distorted much. 3. By 3D simulation, it is clarified that the fine grain materials are stable against cracking in high temperature quantitatively. And low speed and low water cooling in the mold are also very effective.

6 [1] [2] [3] [4]

References B. Hannart, F. Cialti, and R. Van Schalkwijk, Light Metals (1994) , 879 J. M. Drezet and M. Plata, Light Metals (1995), 941–950 M. G. Chu and A. Giron, Proceedings of the M.C.Flemings Symposium on Solidification Processing and Material Processing, TMS, (2000), 223 M. G. Chu and D. A. Granger, Solidificantion Processing 1997, University of Sheffield, UK, July 7-10, (1997),198

36

Investigation of Factors Affecting the Extent of Microporosity in an Aluminum Casting O. Savas, R.Kayikci Sakarya University, Sakarya, Turkey

1

Introduction

Micro porosity has detrimental effects on the tensile strength, ductility and pressure tightness of cast aluminum parts. Therefore, it needs to be predicted and controlled to minimize a level that does not hinder the required performance of the cast parts. The prediction and control of microporosity, particularly in the design stage of a cast part, is an important practical interest which in turn requires knowledge of factors influencing the microporosity formation during the solidification of the alloy being cast. Formation of microporosity in aluminum castings is generally caused by three mechanisms. (i) rejection of hydrogen from the growing solid into the adjacent liquids as a consequence of the large difference in the solubility of hydrogen between the liquid and the solid phases (gas porosity).(ii) interdendritic shrinkage during solidification without gas being present (shrinkage porosity). (iii) Combination of the first two mechanisms [1]. For an efficiently fed aluminum casting with an adequate directional solidification the general prospect of porosity formation is that for a given melt hydrogen content the pore volume fraction and the pore size of the casting decrease with increased cooling rate and similarly, for a given cooling rate the pore volume fraction and the pore size decrease with decreased hydrogen content [2–3]. A Number of work concluded that modification of A356 alloy by strontium addition has been found to increase the pore volume fraction and pore size [2,3,4]. Reports from experiments on A356 and A319 alloys indicated that gas porosity after solidification depends not only on hydrogen content in the melt, but also on metal cleanliness [5]. Porosity level increased for a similar H content when the melts are oxidized by chip addition or by stirring [5]. Liu and Samuel [6] recommended the use of ceramic filters and they claimed that oxide films have a much more deleterious effect on the mechanical properties compared to that expected from other inclusion. A number of researchers have carried out experiments to correlate various castings and solidification parameters to predict the level and distribution of porosity in aluminum castings via criteria functions [9,10,11,12,13]. In previously published literature the most frequently used parameters were thermal gradient (G), cooling rate (R), local solidification time (Ts), and solidus velocity (Vs). In this study, effects of such parameters as mold filling, liquid filtering, cooling ratio, melt hydrogen level and liquid treatments of molten alloy with Sr addition, on the final pore size and distribution in an sand cast A360 alloy were investigated.

37

2

Experimental Procedure

2.1

Experimental Design

The main aim of the present study was to apply the Taguchi method to establish the effects five casting parameters on the extent of microporosity in an aluminum alloy. The five design parameters (factors) and their levels are given in Table 1. Table 1: Design parameters and levels. Factors A- Modification B- Hydrogen level C- Solidification time D- Filtering E- Filling condition

Level 1 Unmodified Low (0.07 ml/100 gmAl) High (14.0 minutes) Unfiltered Laminar flow

Level 2 Modified (with 0.185 % Sr) High (0.22 ml/100 gmAl) Low (4.0 minutes) Filtered (with foam filter) Turbulent flow

The basic principle of the Taguchi method is to develop an understanding of the individual and combined effects of various design parameters from a minimum number of experiments. Taguchi method uses a generic signal-to-noise (S/N) ratio to quantify the present variation. There are several S/N ratios available depending on the type of characteristics, including “lower is better” (LB), “nominal is best” (NB), and “higher is better” (HB). Since the lower porosity is desirable in a casting The S/N ratio for the LB characteristics is related to the present study which is given by [14];

S N

§1 n · 10log ¨ ¦ yi 2 ¸ n © i 1 ¹

(1)

where, n is the number of repetition in a trial under the same design conditions, yi represents the measured values (percent microporosity), and subscript i indicates the number of design parameters in the orthogonal array (OA) which is shown in Table 2. Table 2: Experimental lay out and results with calculated S/N ratios for measured porosity. Experiments 1 2 3 4 5 6 7 8

A

Factors and levels B C D E

1 1 1 1 2 2 2 2

2 2 1 1 1 1 2 2

1 2 1 2 2 1 2 1

1 2 1 2 1 2 1 2

2 1 1 2 2 1 1 2

Measured porosity (%)

y1 1.44 0.27 0.98 0.07 0.09 1.01 0.26 1.47

y2

y3

1.72 0.05 0.51 0.04 0.15 0.58 0.08 1.69

0.79 0.16 0.31 0.08 0.12 0.52 0.16 0.63

S/N ratio

y 1.32 0.16 0.60 0.06 0.12 0.70 0.17 1.26

–2.75 14.67 3.58 23.66 18.20 2.63 14.82 –2.57

38 In the Taguchi method, a design parameter (factor) is considered to be significant if its influence is large compared to the experimental error as estimated by the analysis of variance (ANOVA) statistical method from the equations shown below [14]. If this is the case, the design parameter is a fundamental factor in determining the optimal solution to the design problem.

S ST

ª N S 2º T 2 « ¦ ( )i »  ¬ø 1 N ¼ N

(2)

S SA

ª K A § Ai 2 · º T 2 «¦ ¨ ¸»  «¬ ø 1 © n Ai ¹ »¼ N

(3)

vtotal

N 1

(4)

V factor

F factor

S S factor v factor

V factor V error

(5)

(6)

where, SST is the sum of squares due to total variation, N is the total number of experiments, SSA represents the sum of squares due to factor A, KA is number of levels for factor A. Ai stands for the sum of the total ith level of the factor A, nAi is the number of samples for ith level of factor A. T is the sum of total (S/N) ratio of the experiments, Qtotal is the degrees of freedom, Vfactor is the variance of the factor, SSfactor represents the sum of squares of the factor and Ffactor is the F ratio of the factor.

2.2

Materials

The casting model is shown in Figure 1 which is consisted of four different section thickness increased consecutively to obtain higher solidification times from the bottom to the top. The top end section of the casting was designed to behave an effective feeder for the lower sections via modeling a 3D computer simulation system. In some castings the running system was configured to fill the castings under turbulent flow. When necessary 10 ppi foam filters sized 55 u 55 u 20 mm, were accommodated across the horizontal runner as seen in Figure 1. Primer ingots of A360 alloy was melted in a 100 Kg electric crucible furnace. Liquid treatments such as degassing, regassing and Sr additions necessitating from the experimental lay out shown in Table 2 were carried out in a sequence from Experiment 1 to Experiment 8. Molten alloy was taken from the crucible using a 10 kg ladle which was poured into the molds at 700 °C. After cooling down the castings were sectioned and samples sized 15 u 15 u 15 mm were taken from the locations shown in Figure 1. For the microporosity measurements the samples were metallographically prepared and their micrographs were taken using an optical microscope under u50 magnification. Micro porosity measurements were carried out from the micrographs using an image analysis technique and this procedure (including grinding and polishing)

39

Figure 1: Configuration of casting model

was repeated at least three times to reduce the error arising from the photography and the measurement measurements.

3

Results and Discussion

Using the measured porosity values given in Table 2, the corresponding S/N response table was derived, as shown in Table 3. According to the principles of the Taguchi method, for a given design factor, the present study defines high influence on porosity as a maximum S/N ratio. The mean S/N values of each factor level of design parameters are shown in Table 3. Therefore, Table 3 indicates the optimal design parameters combination and the corresponding value of each factor (see Table 1), i.e. A1: unmodified alloy, B1: low hydrogen level, C2: low local solidification time (fast cooling rate), D2: filtered liquid, and E2: use turbulent filling condition. Table 3: Mean S/N values of factor (factor response). Factors A B C D E

Level 1 9.79 12.02 0.23 8.46 8.93

Level 2 8.27 6.04 17.84 9.60 9.13

ANOVA analysis was performed using Equations (2–6) and the resulting data is summarized in Table 4. The high contribution and variance of factors B and C indicate that hydrogen level of the liquid alloy and the local solidification of the casting (cooling velocity) are significant.

40 Meanwhile, the Ffactor ratio indicates that factors B and C are of 99 % confidence. On the other hand, the results reveal that design factors marked with (*) namely, factor A, modification treatment (i.e Sr additions to the liquid alloy), factor D, filtering the melt during mold filling, and factor E, laminar or turbulent filling conditions have a less significant influence upon the microporosity formation in the castings. Table 4: Analysis of variance. Sum of squares (SS) Degrees of freedom (Q) A* (Modification) 4.63 1 1 B (Hydrogen level) 71.39 C (Solidification time) 620.22 1 D* (Filtering) 2.58 1 E* (Filling condition) 0.09 1 e 8.18 2 Total 707.09 7 ep (error) 15.48 5 Factors

Variance (V) Ffactor 4.63 71.39 620.22 2.58 0.09 4.09 101.01 3.10

P ( %)

Pooled 23.06# 10 200.35# 89 Pooled Pooled 1 100

* pooling, ep # At least 99% confidence

According to Taguchi techniques to determine the optimal conditions, and to compare the result with the expected conditions, it is necessary to perform a confirmation experiment. If the generated design fails to meet the specified requirement, the process must be reiterated using a new system until the required criteria are satisfied. In the present study, the required confirmation experiment should satisfy the optimum design parameter (factor) combination as: A1 B1 C2 D2 E2. Since this combination of design parameter had already been included in the main experimental layout (see Table 2, as Experiment 4) there was no need to carry out an extra confirmation experiment. For confidence the calculated S/N value for Experiment 4 should be between 19.18 and 25.33. The S/N value of Experiment 4 (see Table 2) which is the highest value indicating the S/N value (23.66) for the lowest porosity (0.06 % mean) within the whole experiments validating the confidence of the present study.

4

Conclusions

The present study has applied the Taguchi method to investigate the factors affecting the microporosity formation in an A360 aluminum alloys. The conclusion of this study can be summarized as follows. 1. An A360 aluminum alloy was sand cast under five different design parameters with two level variations to evaluate microporosity formation within the cast part sections. Results showed that the microporosity is due to the dissolved hydrogen and ranging between 0.04 % and 1.72 % within the castings. 2. The Taguchi design method revealed that the local solidification time of a casting (i.e the cooling rate) was the most influential factor upon the porosity formation. Dissolved hydro-

41 gen level of the original liquid alloy has an also significant effect on the extent of microporosity. 3. The ANOVA results indicated that the addition of Sr into liquid alloy for modification of eutectic silicon, the pouring velocity and the filtering of the liquid alloy in the running system have a less significant effect upon the porosity level of the castings.

5

Acknowledgements

The authors thank to Finitesolutions Inc.,USA for provision of Flowcast® mold filling simulation and Erdöküm ve Makina Sanayi A.., Istanbul, Turkey, for providing materials during the casting experiments in the promises of the company.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

References J.P.Anson, J.E Gruzleski, AFS Transactions, Vol.99, pp 135–142 (1991) Q.T.Fang, D.A.Granger, AFS Transactions, Vol.87, pp 989–1000, 1989 E.N.Pan, H.S.Chiou, G.J.Liao, AFS Transactions, Vol.99, pp 605–621, 1991 R.Fuoco, H.Goldenstein,J.E.Gruzleski, AFS Transactions, Vol.102, pp 297–306, 1994 G.Laslaz, P.Laty, AFS Transactions, Vol.99, pp 83–90, 1991 L.Liu and F.H.Samuel, J. Materials Science, Vol.33, pp 2269–2281, 1998 A.M.Samuel, F.H.Samuel, AFS Transactions, Vol.100, pp 657–666, 1992 Y.W.Lee, E.Chang, C.F.Chieu, Metall. Trans.B, Vol.21B, pp 715–722, 1990 S.T.Kao, E.Chang and L.C.Chan, AFS Transactions Vol.103, pp 531–536, 1995 E.Niyama, T.Uchida, S.Saito, J.Cast metals,Vol6. pp 16–23, 1981 S.T.Kao and E.Chang, AFS Transactions, Vol.104, pp 545–549, 1996 F.Chiesa, J.Mammen and L.Smiley, AFS Transactions, Vol.106, pp 149–153, 1998 F.Chiesa, J.Mammen, AFS Transactions, Vol.107, pp 103–111, 1999 P.J. Ross, Taguchi Techniques for Quality Engineering, Loss Function, Orthogonal Experiments, Parameter and Tolerance Design, McGraw-Hill Inc., 1988

42

Model Studies of Gas Bubbles Physical Characteristics at Inert Gas Purging into Molten Metals and Alloys R. Stefanoiu, V. Geanta, I. Voiculescu Politehnica University of Bucharest, Bucharest

1

Introduction

The shape and dimensions of the contact area between gaseous and liquid phases represent the best way to stirring the metallic bath and to perform the chemical reactions by the transfer processes that take place at the gas-liquid interface. The contact performances in the mentioned system depend on the interface shape: the film-type systems are characterized a continuous interface between the two phases; a liquid phase dispersed in a fluid constitutes a drop system, while a gaseous phase dispersed in liquid is usually called bubble system. The bubble-type systems are mainly used in the metallurgical processes. These generate favorable effects due to the operational convenience, and to the dimensions of the interphase surface. However, the most important is their disturbing potential that derives from the bubbles intrinsic floating energy. The study of the gas bubbles physical characteristics (size and distribution) formed at the injection orifice is important for establishing the flow regime that takes place at the inert gas injection into metallic bath. Also, knowing the gas bubbles size and distribution is determinant in order to estimate the physico-chemical processes efficiency specific to molten metals and alloys refining procedure. Creating fine bubbles leads to intensifying the nonmetallic inclusions removing process by flotation, while creating great bubbles is efficient for thermal and compositional melt homogenization in the treatment ladle. Forming bubbles with controlled dimensions is possible through the use of a plug-type injection system having a non-oriented porosity for the first case and having an oriented porosity for the latter. The physical characteristics of the gas bubbles formed by inert gas injection in molten metals depend to a great extent on the injection orifices diameter and also on the injected gas flow. This paper presents the results of the model experimental researches regarding the inert gas bubble dimension formed inside of the metal baths at the inert gas injection through three types of injection systems. The experimental researches regarding the size and distribution of the gas bubbles formed by inert gas injection in molten metals were made in order to establish the dependence of these factors on the used injection system, and also on the technological parameters specific to metallic bath refining process by inert gas injection.

2

Experimental

Due to difficulties regarding the direct research of the metallic melt refining by inert gas injection processes in industrial conditions, physical and mathematical modeling, with low temperature models, makes most of the experimental researches.

43 2.1

Inert Gas Injection Systems used for Experimental Researches

In the experimental researches, three types of injection systems were used, defined as: • The injection system type 1 (I.S.1), which simulates a refractory porous plug having the working surface diameter (the surface in contact with the experimental fluid) of 20 mm; • The injection system type 2 (I.S.2), which simulates a refractory porous plate having the working surface diameter of 50 mm; • The injection system type 3 (I.S.3), which simulates a refractory ceramic plug having 18 different capillaries with the gas flowing surface diameter of 0.5 mm / channel. The metallic melt refining by inert gas injection is more efficient when the injection system is placed at the bottom of the ladle, than for using an injection lance introduced from above the treatment ladle. Consequently, in all three types of injection systems experimentally used, the gas injection was made from the bottom of the experimental vessel, with the help of a central injection system taking into account the vessel symmetry axis. The injection system features for the first two types used were qualitatively and quantitatively determined by microscopic analysis (using the Buehler Omnimet Enterprise, 5.0 version processing software). Thus, the average porosities and the average gas flowing diameter (orifice diameter d0) were determined for the two mentioned types. These values are shown in table 1. It should be known that the qualitative and quantitative analysis of the injection systems was made by taking several pictures of the working surface (the surface which is in contact with the experimental fluid), and in the subsequent mathematical modeling, and the obtained average values were used. Table 1: Features of the injection systems used in the experimental program Injection system Type 1 Type 2 Type 3

2.2

Average porosity [%] 17 18 –

Average diameter of the injection orifice (u10–3) [m] 0,055 0,067 0,5

Experimental Set-up

The concept shown in fig. 1 was used for the experimental researches. The device is made out of a transparent cylindrical vessel (6) filled with distilled water up to the working level. At its bottom is it placed an injection system (5), through which air with different flow rates is injected, the air being provided by the compressor (1). The determination of the gas bubbles characteristic dimensions was made by quickly photographing the area of gas bubbles forming, using a digital camera (7) placed at the bubbles forming surface level. For the researches, the three types of injection systems previous mentioned were used. The injected flow rate variation was maintained between 1.94 and 8.33 Nm3/s, these values being in the interval of working flow rate variation, used in the physical modeling of metallic

44

Figure 1: Experimental device for (air – water) model study of the gas bubbles dimensions generated in gas stirred systems: 1 – compressor; 2 – manometer; 3 – flow meter; 4 –gas feeding flexible pipe; 5 – gas injection system; 6 – experimental vessel; 7 – digital camera; 8 – lid

bath refining by inert gas injection process. Thus, five working gas flow rates were used, for each of them the gas feeding being continuous. The images were taken 5 minutes after the gas feeding had begun, in order to stabilize the working flow. For each gas flow rate and for each injection system type were taken ten different images.

3

Results and Discussion

In order to determine the gas bubbles diameter the images were processed with the help of Corel PHOTO-PAINT 12 graphical processing software, and the truthfulness confirmation of the obtained results was done using the Buehler Omnimet Enterprise, 5.0 version quantitative analysis software. The dimensional results obtained for the same flow rate and the same injection system varied in maximum ratio of ±10 %. For the gas bubbles dimension measurement, as a simplified assumption, it was considered that the bubbles formed at the injection orifice level are spherical and they are deforming under the action of the external forces after detaching from the injection orifice, this also being confirmed by the captured images. The diagram in fig. 2 shows the gas bubbles diameter variation, determined at the injection orifice level, function of the gas flow rate used for the three experimental injection systems. From the diagram analysis, it results a quasi-linear variation of the bubbles diameter with the injection gas flow rate, for all systems. By analyzing the data in figure 2, one can notice that the gas bubbles diameter increases with the increase of the injected gas flow rate, no matter the used injection system and its geometrical features. For the same gas flow, the gas bubbles diameter increases with the increase of the injection orifice diameter. The greatest values for the bubbles diameter were obtained using the I.S. 3, for which the injection orifice diameter d0 has the greatest value.

45

Figure 2: Gas bubble diameter variation function of injected gas flow rate

With the injection gas flow rate variation between 1.94 and 8.33 (*10–5) Nm3/s, the diameter of the bubbles formed at the injection orifice level is: 1.546–2.988 (*10–3) m for I.S. 1; 1.893–3.660 (*10–3) m for I.S. 2 and 4.180–5.016 (*10–3) m for I.S. 3. From the diagram in fig. 2 can be noticed that for S.I. 1 and S.I. 2 types, the slopes of the two regression curves are bigger than for using the S.I. 3 type. This fact is explained by the increasing of the injection pressure for the same gas flow rate. In the literature there are different equations that express the correlation between the diameters of the gas bubble formed at an immersed nozzle level and the technological parameters of the injection process. The empirical relation, obtained from the experimental data, that expresses the bubbles diameter variation function of the gas flow rate and the injection orifice diameter is give by the equation:

db

8.38624 ˜ Qg0.9  0.075057 ˜ d 00.4066

(1)

The equation was obtained using mathematical modeling with regression analysis by programmed active experiment, using the STATISTICA 6.0 software. In order to determine the equation, a nonlinear model was used with the independent variables Q1 (Qg) and Q2 (d0) and the dependent variable Q3 (db). Using the Levenberg-Marquardt and Gauss-Newton methods, the nonlinear estimation of the model coefficients was established. The diagram in figure 3 shows the regression lines obtained by equation (1) as well as the experimental values in the case of the three injection systems used in the researches. From fig. 3 it results that the values of the diameter of the gas bubbles obtained for the three experimental types are close to the regression curves traced for the equation 1. Figure 4 shows the 3D correlation between the gas bubbles diameter, gas flow and injection orifice diameter, variation obtained according to equation (1). From the figure it results that the diameter of the gas bubbles formed in a fluid as a consequence of gas injection varies proportionally with the injected gas flow rate and with the injection orifice diameter.

46

Figure 3: Gas bubbles diameter variation function of gas flow rate and injection orifice diameter. (Comparison between experimental values and regression curves obtained for eq. 1)

Figure 4: Gas bubbles diameter variation function of gas flow rate and injection orifice diameter

4

Conclusions

At the inert gas injection into the metal bath, the gas flow rate, the injection pressure and the gas input surface area all play an important technological role, through this parameters the particularly dynamic and transfer processes being controlled. From the experimental researches, it results that the dimensions of the gas bubbles formed at the injection point inside of the metal bath depend both on the injection orifice diameter and on the gas flow rate.

47 With the help of the presented diagrams, the gas bubbles dimension value may be determined for any gas flow rate and orifice diameter smaller than 0.5 (u10–3) m. Also, using equation (1) the diameter of the gas bubbles formed at gas injection through an immersed nozzle having the diameter between 0.05 and 0.5 mm can be obtained. At inert gas injection into the metallic melt in the treatment ladle, the interior diameter of the porous systems is between these values.

5 [1] [2] [3] [4] [5] [6] [7]

References D. Mazumdar, R.L. Guthrie, I.S.I.J. International, Japan, vol. 35, 1, 1995, 1–20 M. Sano, K. Mori, Transaction I.S.I.J., vol. 20, 1980, 675–682 R. Stefanoiu, V. Geanta, Metalurgia International, 1, 2005, 3–9 L. Wang, H.G. Lee, P. Hayes, I.S.I.J. International, Japan, vol. 36, 1, 1996, 17–24 M. B. Goldschmit, A. H. Coppola Owen, Ironmaking and Steelmaking, vol. 28, 4, 2001, 337–341 K. Mori, M. Sano, T. Sato, Transaction I.S.I.J., Japan, vol. 19, 9, 1979, 553–558 R. Stefanoiu, PhD Thesis, Chapter 5, Bucharest, 2004

48

49

Casting Technology

50

51

Remarks about Process and Technology of Continuous Casting H. Schliefer, A. Khoury, M. Porten Norddeutsche Affinerie AG, Hamburg

P. Wolber and K.H. Boller SGL, Bonn

W. Dürrschnabel and H.R. Müller Wieland-Werke AG Ulm

St. Schneider Deutsche Gießdraht GmbH, Emmerich

W.H. Müller, M. Schwarze SMS-Meer, Mönchengladbach

H. Oelmann, D. Rode, R. Frankenberg KME, Osnabrück

1

Abstracts

During the last fifty years continuous casting increase in the fields of technology and process also in the materials and “tools”. It will now given remarks about the history and the future in this production area.

2

The First Step

The first step into development from single to continuous casting is dated in the literature in 1840 and it needed nearly 100 years to become industrial practice. In the historic of the continuous casting Mr. Siegfried Junghans has build on his own a pilot plant in Schwandorf and after a lot of trials – and errors too – this new technology was born. After World War II further developments were done by a syndicate of Wieland-Werke AG, Norddeutsche Affinerie AG and Krupp. First for copper and later the steel continuous casting was created on Krupp side. So for the continuous casting the copper side was the technology entrepreneur. This was the first great step from mould to mould cast into a continuous process with all positive effects in it: from the bigger piece and length (10 times and more) to better – over all – quality, from lesser heavy work to better energy consumption, from higher casting speed to better economic terms. The important tool for the continuous casting was the heat exchanger (the mould) and included in the mould is the knowledge of • the heat transfer and the grain solidification / direction in the mould • the cooling rate and cooling zones in and under the mould • the graphite shell / layer into/on the copper jacket and the taper in the mould • the connection between graphite lining and copper jacket without air gaps • the lubrication between melt and mould • the oscillation rate and amplitude

52

Figure 1: NA cast shop from last early 1950

Figure 2: Traditional moulds for cakes and billets with graphite shell

This type of mould with graphite inlet – also only with “one way” graphite layer on the copper jacket – is also very common and in use all over the world for copper and his alloys. This figure gives a good overview of the interaction of all parameter into this kind of mould. Generally for all copper alloys – like brass, bronze or alloying of copper with multiple elements – the knowledge about casting and/or physical parameter is very small and naturally data is hardly to find in books or databases. About a lot of alloys especially the parameter – for the

53

Figure 3: Scheme transfer and direction of heat and grain growing

Figure 4: Dispersation of Fe in CuFe2P

mould construction and also for the casting technology – near the solidification temperature like i.e. high temperature strength, thermal conductivity/coefficient and also heat transfer rates, are unknown. So for the same alloy different types of moulds and cooling strategy are in use during the developments by the casting shops. As example: CuFe2P, as a connector alloy, is with nearly over 30.000 t/y produced in considerable quantities but the casting technology and the moulds are different all over the world.

54 The important parameter of this alloy is the distribution of Fe and the other physical parameters are all predetermined at the time it is cast. So the casting condition must fulfil a slab free of cracks, shrinkage cavities (pipe, blowholes) and this uniform dispersion of Fe which is absolutely necessary for the further treatment. The big improvement in this matter was done by continuously recording and analysing of all parameter during the casts, mathematical modelling of temperature, stresses – especially for the fresh solidified shell and based on this to optimise the main casting parameter like casting speed, pouring temperature, volume flow rate of metal and cooling water, different soft secondary cooling and taper in the mould. And after that, based on the quality of the products, optimising the water quality and the cooling water circuit system, installing a mould maintenance schedule and also improving the mould-construction and the mould material (i.e. ELBRODUR) for better durability, stiffness and stability. Looking on the control of casting speed and oscillation for smaller tolerances. With the real effects of this developments are the cakes cross section has increased by approx. 50 % and the reject rate has decreased by approx. 60 %. The old odds created a reasonable access by try and error in this market section and which new bets will do in the future? Today for this alloy a compromise in technical and economic terms is in use. For both we can imagine – based on all necessarily theoretical data near the solidification point like i.e. strength, conductivity, heat transfer coefficient, solidification and distribution to build up a realistic mathematical model – and for other alloys too – and based on this parameter and data – as we know this all – we can construct and build a new generation of mould. This new generation will have a pure graphite mould, indirect inner cooling circuit, inner cooling defined areas and cooling water volumes, indirect soft and direct spray cooling areas defined by localisation and cooling water pressure, a defined cooling water quality, pressure and quantities. This graphite mould with inline cooling system would create the optimal result for alloys in technical as also economic terms in any parameter, even if it is more expensive.

Figure 5: Graphite mould (cakes) and inline cooling detail

55 From the beginning this mould was and still is only used for semi-continuous casting due to the common sense, that this mould is too “fragile” to operate continuously. Nowadays this mould is established in the full continuously production for cakes. Billets are not so common because of dimension intolerances created be regular maintenance. For cakes this mould is working with very good results for surface and inner quality like shrinks and holes, for excellent casting speed, for long lifetime of the graphite blocks and inner and outer cooling devices, and more than three cooling zones with different cooling water volumes and velocity. This means different heat transfer rates can be installed in the same mould. Any cross section i.e. from 130 u 400 to 1130 u 350 mm is available. In other words, in the future the progress of this mould construction includes the best flexibility and success for the development of copper and alloy casting and additionally the lowest mould cost/t (i.e. ETP, DHP, OF, CuAg, CuFe2P). This type of mould can also be used in horizontal casting and especially for alloys (i.e. CuSn and CuNi) based on sophisticated heat transfer conditions, on complicated solidification and first shell building, on the flexibility in the inside construction for the different heat transfer areas and localisations from the sidedams to upper- and bottom-side of the mould. For all this moulds – without connection with the furnace – in the following small remarks about „tools“ are given. A lot of inventions were done and plenty of creative people have worked on the development of tools: Casting without soot means a „white casting shop and people“. The protection of the open surface in the mould is necessary for lubrication and also helping against oxidation. Protection by soot or lamp black is very common in the copper industry with all the risk of benzpyrene in the soot and also the contamination of the casting shop and workers – „going black“. One technology – invented by HITACHI and now promoted by SMS-MEER – is working with a gas mixture, which combines the surface protection with the lubrication between mould and strand. The change from gas to soot is done by temperature and in the presence of molten copper. This originally HITACHI technology works very well with the result of clean plants and workers and additionally higher casting speed, i.e. in the HALCOR plant in Greece.

Figure 6: Mould under protective gas

56

Figure 7: „Cracking“ of gas to soot on the cold borderline

Figure 8: Mould level sensor by ionisation

The simplification from the included gas generator to a simple gas mixture station will bring this technology to better economic scale. The combination with a level sensor by ionisation as shown in the picture brings together gas inlet, level control, surface protection and lubrication with only one tool. This leads to simplification and zero contamination. Graphite grades are also tools for vertical and horizontal casting. For each alloy different graphite specifications are necessary for porosity, density, metal penetration, thermal shock re-

57

Figure 9: Grade comparison continuous casting

sistance and elongation, conductivity (heat and/or electric), physical parameter like hardness and strength. This data sheet shows a rough classification. Horizontal (HC) casting needs more porosity by degassing than vertical casting (VC). A balance of fine /rough porosity approx. 10 % for CuSn and approx. 5% for brass is common. Penetration is not a big issue for VC by the borderline lubrication/protection of soot. All physical parameter have to be better in HC. Conductivity depends on mould construction and alloy, i.e. lesser than 100 (W/m K) for copper and up to 150 (W/m K) for alloys. Depending on the different mould construction and casting parameter it is not possible today to define general graphite grades for copper and this alloys. This work can and has to be done only by the casting people themselves and cannot be reduced to the answer: which graphite grade do you need ? Black graphite! This means for the future going into a real open minded teamwork to minimise the number of grades and also a big part of production cost in consideration of the competition on the global market. Booster and stirrer have a long tradition in using ultra sonic devices to create smaller grains, improve the distribution of alloys, minimise the porosity and optimise other parameters too. Coupling these ultrasonic waves to the melt is the most difficult part of this technology. A simple way is to bring US in the solid strand, but the mass of metal needs very high energy to have any effect in the melt. Another technology is to bring the US by booster directly into the melt from the top of the mould. The advantage of this technology is the direct connection with the melt and the grains. The disadvantage is the technology itself: The temperature limitation of the booster material in the melt and the cooling of the US unit above the melt as well as the connection to the booster. Trials were made now and then with common boosters i.e. delivered by Krautkrämer. Metallurgical effects were found in form of very fine grains as long as the booster was working, which means the temperature was over 1,100°C in the melt on one end and nearly room temperature on the other end. Any US is disturbed and/or deleted by temperature-differ-

58 ences in the booster material. By solving the above material/temperature problems for the direct booster input in the melt chances for the future will increase. The electromagnetic stirrer is very common in the steel industry and works very well in the long liquid centre part of the strand after the mould. In the copper industry – with comparatively short liquid sump – this technology does not create any kind of process/quality optimisation. Calculations, constructions and trials were done at the RWTH Aachen by Prof. Dr. Block with the result, that reasonable effects were found by a special constructed stirrer in the copper. Success in quality, distribution and grain size by this special designed stirrer was only achieved for a defined cross section and due to this reason this technology was “killed” by the economic terms. The use of the stirrer as a tool for copper alloys is an option for the future.

3

The Next Step

The next step – really one and half – was coming up with the Southwire/Contirod/Properzi -process to produce wire rod directly. This is a continuous casting of a cross section – like a wirebar – by a casting wheel or double belt caster. After casting in one line the direct rolling process in one heat produced wire – near net shape by rolling. The same way, but by another technology was done by the DIP FORMING technology, which is for the casting part like a “grandfather” of the UPCAST process. For the wire production copper and aluminium, and partial steel too, the before mentioned technology combines the biggest production share in the world wire market. And which advantages are to be expected in the future by this technology, to fulfil this growing market with a yearly rate near 5 % without investments in the existing plants and minimising the production cost by volume? First of all within the next 3–4 years higher capacity/plant will be coming up to approx. 70 t/h, this means approx. 500.000 t/y by bigger melting/holding furnaces, casting wheel and/or

Figure 10: Scheme of Southwire plant NA

59

Figure 11: Scheme of Contirod plant

bigger cast cross section. This includes step by step more rolling speed/power and/or one or two additional rolling units and also high speed clean up/cooling and coiling equipment. This creates new high speed inline test sensors. The coil weight and diameter (maximum truck width) will also increase to 15–22 t/coil – one truck load –, to minimise the transportation and the offtime costs in the cable plants. For market niches smaller (this is a real speed problem) and bigger diameter (this is a coiling problem) and a production of alloys (i.e. CuAg) in big lots in only o n e plant but for different companies is cost efficient and earns together the volume-effects.

Figure 12: Calibration of flat strip

60 The copper flat strip-market worldwide has, including width nearly 100 mm, a market share over 60 %, i.e. electric and/or ACR / welded tube indicate the next and second step by the production of flat material in copper wire rod plants. This can be done with smaller investments of designed rolling units, clean up and additional coiler/winder equipments in each wire rod plant. So the production can split in wire and strip in the same plant that means bigger volume in one unit and two different markets and that will create better economic terms.

Figure 13: Additional rolling equipment

Figure 14: Additional cooling and winder equipment

61

4

The Second Step

The second step in the development of continuous casting was the “near net shape” casting technology included upcast for metal wire and horizontal/vertical thin strip/sheet casting of copper/copper alloys and the “Thin slab” casting in the steel industry. But on the other hand a lot of alloys need a rolling treatment which means "thin” slab. All this before mentioned technologies have not the biggest advantages in the future: The vertical “Thin Slab” steel technology works only for a limited amount of grades and needs a big investment, space and additional heat treatment before rolling near net shape. The copper VC, its alloy strips and wire technology has a limited potential for better productivity and greater width by the great solidification range and the difficult shell building in a moving mould and/or strand. The horizontal copper alloy production is the next target. Based on the better mould and graphite, using pure graphite moulds with inline cooling system instead of the traditional horizontal moulds will increase efficiency. The next big step was what we know as “near net shape casting”: Which is defined as „direct dimension, grade and final surface by one step cast process = one step production“. The single belt process for brass and bronzes is a special example and technology in this area. Any kind of copper alloys for sheet and strip production by the traditional way: Vertical casting cakes – preheating – hot rolling with milling and cold rolling incl. heat treatments includes a higher production than the horizontal casting of strip, milling and rolling including heat treatments which includes a lower production rate but lesser invest and working steps. This technology was first developed in the steel industry. A pilot plant was later used for test runs for different copper alloys by Wieland-Werke AG. Tin bronzes have to be during solidification absolutely stress free. The lower part of a horizontal HAZELETT machine can fulfil this parameter. It was calculated, expected and realized by a single belt to have a high production

Figure 15: Single belt casting machine, scheme

62

Figure 16: Single belt casting in operation

rate, approx. 10 to 20 t/h and strip width from 400 to 800 mm with a thickness of approx. 4 mm. The pilot plant was working with 10 to 20 t/h on basis of 400 mm u 6 to 10 mm cross section for different copper alloys. Cooling is done indirectly by the under belt. The upper surface was insulated to maintain a directed solidification. The planar level of the strip is not so important because of the following rolling of the strip. This technology is working reproducible and well for copper and micro-alloyed copper. For all copper alloys with dendritic solidification this technology does not work. The single belt technology has also economic limits by the small volumes of copper alloys in each cast/plant, which means changing the alloy to often. For the strip production in the steel and copper industry the future is vertical TWIN-ROLL, in the aluminium industry the horizontal TWIN-ROLL and/or the TWIN BELT HAZELETTcaster. The vertical TWIN-ROLL caster includes very high production rate from 40 to 70 t/h, near net shape casting with cross section from 1,200 u 12 mm2 to 2,000 u 0,7 mm2, with good surface quality. The low roughness level does not need additional milling. Now the TWIN ROLL is “state of the art” in the steel industry and also in the aluminium. Today the biggest success in the flat product market is the vertical/horizontal “TWIN ROLL”. The TWIN BELT caster for sheets is working in aluminium recycling plants. The TWIN-ROLL is a very old technology invented by Bessemer and is also a good and typical story for patents: it was patented but never producing! These Figures include the technology and process in the future of the production in the flat market. The numbers and terms in the “flat” steel industry tell of the big success and increasing volume. The copper industry has TWIN ROLL today only as pilot plant, which is a small industry in comparison to steel and aluminium. The result of this pilot plant tells the TWIN ROLL- process is really the coming technology in the flat market. Step by step, from pure copper to copper alloys and all in big volumes, depending on the possible investments, can be done for this technology by one company or by one special R+D company with different shareholders

63

Figure 17: Bessemer patent

Figure 18: Scheme TWIN ROLL vertical NA copper pilot plant

from the industry. This process will create very good economic terms and excellent quality and surface – as we have seen in the steel industry – and will also create a big chance/change for the copper industry. On the way into the future to invest in copper TWIN-ROLL-technology – R+D included – is a small investment in comparison to the steel but a big step for copper. High alloyed copper like

64

Figure 19: Scheme Castrip® Process

Figure 20: Scheme TWIN ROLL horizontal Alumina plant

brass or bronzes will have the traditional way by continuous casting vertical as cakes or horizontal as thin slab or thicker sheet. Additional remarks about tools in this technologies: Ceramic materials and valves: The ceramic material for flow control valve – especially in the copper industry – has to fit a long time corrosion and also thermo-shock resistance combined with zero penetration by the molten metal and high mechanical strength. This is the imagination

65

Figure 21: Parameters NEAR NET SHEET

Figure 22: Castrip Plant Production

of a non existing material. Materials with high contents of Al-Oxide, SiC, SiN or high temperature metal-oxides mixtures will fit the most above parameter from the material side, depending on metal and/or alloy. The other side is that all valves are working only mechanically in metal industries. Years ago it was described by INTERATOM in patent DE24970 as a channel to control and pump.

66

Figure 23: Average UCS Thickness

This technology is very well working as a pump in a closed sodium-cooling-circuit and was also good working in a pilot plant as valve for vertical casting copper. The flow rate variation was approx. 1:9, which means nearly 4 to 40 t/h – other volumes are possible by construction – in a flat channel of 5 u 80 mm and a cuprostatic pressure of approx. 300 mm – higher pressure is possible and tested too. The energy consumption was calculated and recorded lesser than 3 t/h and an investment of approx. 100.000 €/valve each additional valve for the same casting unit

Figure 24: Scheme electromagnetic pump/valve

67

Figure 25: Electromagnetic valve in operation

approx. 15.000 to 25.000 €. This kind of technology in combination with high temperature materials can vary the flow rate by joystick and/or automatic control very exactly – drop by drop – and continuously into the full variation rate. Like cast (drive) by wire without mechanical parts – especially for vertical cast – the E-valve will block every pressure coming from the ladle/ tundish. For the horizontal process this valve will give a defined metal speed = weight regulation also. Another tool is the level control by ionisation, very useful in protected and very small areas. By this limitation and in this areas laser, electromagnetic fields or other common sensors are useless. A good example for this kind of level control is the soot free technology with a “blank” metal surface and the protective area. The combination of the gas input and the sensor creates an excellent signal of the impedance by the ionised halo from the out burning gas. This signal can be used for the electronic and electric devices to control the level surface of the melt in the mould. All this tools combined in the near net shape casting technology for flat products can give any ideas about the future: The combination of electromagnetic valve with ionisation level controls blank metal surface sensor with high speed valve, drop by drop controlling without metal pressure and can bring the continuous casting technology to progress.

5

The Third Step

And what is the next and third step in and for continuous casting? The mathematic modelling is on a good way to become a common tool. The development of graphite and mould construction as “tools” is a permanent process. The simplification of materials will go on due to cost pressure. The international and national bench marks will bring the

68 industry in the situation to work better and closer together and to develop tools like booster, stirrer, magnetic valve, ionized level control, graphite and ceramic grades. In my opinion the next step has to be the mathematical modelling of alloys on the basis of Medema and/or Hückel/Debye and the quantenmechanical theory of metals. It has to be a vision for the future, that it is possible to define and design a real material by his parameter and physical data. Beforehand the big job is to collect and search the parameter and data for different temperatures. Based on this to go in a metal databank, using many different mathematical models for the creating – first in virtual – of a new metal alloy with exorbitant advantages and second out of this knowledge the real alloy can go in production. The hardware for this third step is present, the software has to be done by brainstorming in the next near future, for the target “model of metal”. This is the third but not the last step of continuous casting process and technology.

6

Conclusion

Near net shape continuous casting, vertical TWIN-ROLL, sophisticated graphite moulds and new combined “tools” are one step in the future also the cooperation between companies as the second step to reduce the economic terns The great step in the next years are the “model of metal” by the interaction between metallurgist and mathematicians to build up databases and models for better material and technology and also for lesser consumption of our worldwide sources.

7 [1] [2]

References

Handbook continuous casting, 1980 div. Patents UK, 2397, 1861 ; Bessemer US, 2 935 251 ; US, 2 946 100 ; US, 3 098 269 ; US, 3 157 921 ; ASARCO CH, 362 800 ; CH, 361 888 ; Wieland Werke AG DE, 10 2004 027 194.1 ; Norddeutsche Affinerie AG DE, 24970 ; INTERATOM JP, 5 32 96 00; HITACHI [3] P. Wolber and K.H. Boller, SGL and E. Kress, Sundwiger Messingwerke: Graphite, Personnel information [4] Dr. W. Dürrschnabel and Dr. H.R Müller, Wieland-Werke AG: single belt casting, Personnel information [5] Dr. A. Khoury, M. Porten, Dr. St. Schneider, Deutsche Gießdraht GmbH: wirerod Southwire, Personnel information [6] Dr. W.H. Müller, SMS-MEER: soot free, Personnel information [7] H. Oelmann, Dr. D. Rode, R. Frankenberg, KME: CuFe2P, Personnel information [8] Dr. M. Schwarze, SMS-MEER: contirod, “Rolling of copper flat sections with a contirod ...”and” 25 years contirod, developments, trends, new product mix” [9] R.L. Wechsler, Castrip: report at 5. June, 2005 in New York: TWIN ROLL steel [10] NOVALIS: Reference list JUMBO 3CM/3C, 2005

69 [11] [12] [13] [14]

NOVALIS/PECHINEY: JUMBO 3C & 3CM continuous casting technology NOVALIS: Newsletter Februar 2005 STAHL UND EISEN: 1891 N°11 P929 ... Pictures: Norddeutsche Affinerie AG Wieland-Werke AG SGL KME Castrip NOVALIS/Pechiney SMS-Meer

70

Continuous Strip Casting of Magnesium Alloy by a Horizontal Twin Roll Caster H.Watari 1, T. Haga 2, N. Koga 3, K. Davey 4 1

Oyama National College of Technology, Oyama, Japan Osaka Institute of Technology, Osaka, Japan 3 Nippon Institute of Technology, Saitama, Japan 4 The University of Manchester, Manchester, UK 2

1

Abstract

This paper is concerned with the development of a strip-casting technology for manufacturing magnesium alloy sheets. The aim of the work is to establish a manufacturing process and technology to facilitate the economical manufacture of high-quality magnesium sheet alloys. Magnesium alloy AZ31B was used to investigate the appropriate manufacturing conditions for use in twin-roll strip casting. Temperatures of the molten materials and roll speeds were varied to find the appropriate manufacturing conditions. The effects of manufacturing conditions on possible forming were clarified in terms of roll speeds and roll gaps between upper and lower rolls. In addition, microscopic observation of the microstructure of the finished casting was performed. It was determined that a magnesium sheet of 2.5 to 3.5mm thickness could be produced at a speed of 20m/min by a horizontal copper roll caster. Mill stiffness and a method of predicting the cast sheet’s thickness were investigated to determine the appropriate manufacturing conditions. It was also found that the cast magnesium sheet manufactured by roll-strip casting could be used for plastic forming if the appropriate wrought magnesium sheets were produced after the roll casting process.

2

Introduction

Magnesium alloys are expected to play an important role as next-generation materials, with the potential to help lighten total product weight when magnesium products are used to replace aluminum and mild steel products. The specific density of magnesium alloy is 2/3 that of aluminum and 1/4 that of iron. When alloyed, magnesium has the highest strength-to-weight ratio of all structural metals. Moreover, magnesium has received global attention from the standpoint of environmental preservation because of the ease of recycling metallic materials. The utilization of magnesium alloys has depended mainly on casting technology (e.g., thixo-forming) because of their less workable characteristics due to the crystal structure of the hexagonal close-packed lattice. Recently, demands have arisen in the automotive and electronics industries to reduce the total product weight [1]. Unfortunately, high manufacturing costs continue to be a major barrier to greatly increased magnesium alloy use. A key to solving this problem is the development of roll strip casting technology to manufacture magnesium sheet alloys economically while maintaining high quality.

71 The authors, therefore, investigated the effectiveness of twin-roll strip casting for magnesium alloys [2]. This paper describes the forming characteristics of the cast magnesium alloy sheets after being hot-rolled in a warm deep drawing test and establishes the appropriate manufacturing conditions for producing high-quality strip using a purpose-built strip-casting mill. The influences of such process parameters as materials of roll, casting temperature, and roll speed are ascertained. A simple method of predicting the sheet thickness of cast strip is introduced. A warm deep drawing test of the cast magnesium sheets after being hot rolled was performed to demonstrate the formability of the magnesium alloy sheets produced by a roll strip casting process. The microstructure of the manufactured wrought alloy sheets was microscopically observed to investigate the effects of the hot rolling and heat treatment conditions on crystal growth in the cast products.

3

Experimental

3.1

Horizontal Twin-RollCcaster and Experimental Conditions

Figure 1 illustrates the horizontal twin roll strip casting process used in the experiment. A source of molten metal feeds into the space between a pair of counter-rotating internally cooled rolls. The principle dimensions of the horizontal twin roll caster are presented in Table 1. The inclination angle of the mill in Figure 1 was set to zero degrees. Lc in Figure 1 indicates the contact length between rolls and the molten metal. Table 2 presents the experimental conditions for investigating appropriate manufacturing conditions to successfully produce magnesium alloy sheets by twin roll strip casting. Casting temperatures were varied from 630 °C to 670 °C to find the best casting conditions, as indicated in Table 2. Temperatures of the molten magnesium in the melting pot and tundish were measured by thermo-couples.

Figure 1: Schematic illustration of horizontal twin roll casting process

72 Roll casting speeds were varied from 5 m/min to 30 m/min in order to examine which roll speed was appropriate for solidifying the molten magnesium. The roll gap between the upper and lower rolls was determined to be from 2.0 mm to 3.6 mm by simple calculation results based on basic solidification theory. No shielding gases were used in the experiment. Table 1: Dimensions of roll caster and tundish Rolls Materials Upper roll (mm) Lower roll (mm) Roll speed (m/min) Inclination angle (deg ) Tundish Material

Copper, Copper alloy, Mild steel 300*150 300*150 0-150m/min (Max.) 0 Insulator, mild steel

Table 2: Experimental conditions Temperatures (°C) Roll speeds (m/min) Roll clearance (mm) Shield gas

3.2

630, 650, 670 5, 10, 15, 20, 25, 30 2.0-3.6 No shielding

Material and its Refining Process

The material used in the experiment was AZ31B. The physical properties of the material are listed in Table 3. Magnesium ingots were heated to 680 °C in a melting pot with an electric furnace. In the magnesium melting process, magnesium oxide and other suspended nonmetallic matter were removed with flux that preferentially wet the impurities and carried them to the bottom as sludge. After the refining process, the molten magnesium metal in the melting pot was carried to the strip caster and poured onto the cooling slope to manufacture magnesium strip. Table 3: Physical properties of material Density (kg/m3 103) Liquidus temperature (°C) Solidus temperature (°C) Specific heat (J/kg°C) Thermalconductivity (W/m·°C) Latent heat (kJ/kg)

3.3

1.78 630 575 1040 96 373

Hot Rolling Process after Twin Roll Strip Casting

The hot-rolling process was performed to obtain wrought magnesium alloy sheets with globular and fine microstructures to be used for plastic forming. The cast strip sheets were milled to obtain sheets with 2.0 mm thickness to remove oxide film. The cast strip was heated and rolled in the hot-rolling process. Rolling temperatures were varied from 200 °C to 300 °C. The milled sheet was rolled by several rolling pass schedules until the sheet became 0.8 mm thick. Next,

73 the 0.8 mm-thick sheet was rolled again until the sheet became 0.5mm thick. Finally, the rolled magnesium sheet was annealed at 350 °C for two hours, and cooled in an electric furnace. A 5m/min roll speed was chosen in the hot-rolling process. At 250 °C, cracks were seen during the hot-rolling process, even though the reduction was less than 10 %. A temperature over 250 °C was chosen to keep the cast products from cracking.

4

Results and Discussion

4.1

Stiffness of Roll Caster, Relation Between Roll Load and Roll Gap

Figure 2 illustrates the typical relation between roll gap and rolling load obtained by load cells attached to the mill. The result presented in Figure 2 is the case of a mill fitted with copper rolls. The horizontal axis reveals the differences between initial roll gaps and the obtained sheet thickness. It expresses pure deflection of the mill during the strip-casting process. The circles in the figure represent results for an initial roll gap of 2.0 mm, the squares for an initial roll gap of 2.8mm, and the triangles for an initial roll gap of 3.6 mm. From Figure 2, the spring constant of the mill was approximated as 10009 (N/mm). Roll loads per unit width of the cast sheets during casting were from 50 to 150 (N/mm) in the present experiment.

Figure 2: Relation between roll loads and roll gap

4.2

Thickness of Cast Sheet

Thicknesses of cast sheets were measured to investigate the casting phenomenon of magnesium alloy in the twin roll casting process. The sheet thicknesses of three forming directions were measured. Figure 3 presents an example of results for a copper roll caster. The circles indicate

74

Figure 3: Relation between roll speed and sheet thickness tween roll loads and roll gap

results for a roll gap of 2.0 mm, the squares for a roll gap of 2.8 mm, and the triangles for a roll gap of 3.6 mm. It is seen that sheet thickness gradually decreases as roll speed increases. The dotted line represents a theoretical sheet thickness obtained by simple 1-D solidification modeling for a 2.0 mm roll gap. Figure 3 indicates that a roll speed of over 20 m/min is necessary to manufacture strips less than 2 mm thick using a 2.0 mm roll gap.

4.3

Microstructure of Cast Sheet

Figures 4(a), 4(b), present photographs of the microstructures of hot-rolled sheets after the rollstrip-casting process using a copper-alloy caster. The sheet depicted in Figure 4(a) was hot rol-

Fig. 4(a): Microstructure of cast sheet hot rolled at 200°C (without annealing process)

Fig. 4(b): Microstructure of cast sheet hot rolled at 200°C (with annealing process)

75 led at 200 °C, and no annealing process was used. The photo in Figure 4(b) was annealed at 350 °C for two hours after hot rolling at 200 °C. We can see that the crystals were well homogenized with an appropriate annealing process, although the grain sizes of the crystals became larger by recrystallization.

4.4

Plastic formability of Obtained Wrought Magnesium Alloy Sheet

After the cast magnesium sheets were hot rolled, a warm deep-drawing test was performed to examine the forming characteristics of the magnesium alloy sheets produced by twin roll strip casting. The diameter of the punch was 28.8 mm. A lubricant solution was used. The limiting drawing ratio was investigated by a deep-drawing test at 250 °C. A drawing speed of 30 mm/s was chosen in the test. A limiting drawing ratio of 2.6 was obtained in the warm deep-drawing test, as indicated in Figure 5. The result presented in Figure 5 suggests that the wrought magnesium alloy sheets that were hot rolled after the strip casting process had plastic formability equivalent to that of the wrought magnesium alloy sheets manufactured by the conventional DC casting process.

Figure 5: Cup drawn in a warm deep drawing test (DR=2.6, hot rolled at 250°C, with annealing)

5

Conclusions

AZ31B magnesium alloy was cast by using a horizontal twin-roll caster. The obtained cast sheets were hot rolled, and a warm deep drawing test was performed to demonstrate the effectiveness of twin-roll strip casting of magnesium alloys. The following conclusions were obtained. 1. In the hot-rolling process, a temperature exceeding 250 °C was chosen to keep cast products from cracking. 2. An appropriate annealing temperature was effective for homogenizing the microstructure of the rolled cast sheets after the strip casting process.

76 3. The grain size of the manufactured wrought magnesium alloys sheet was less than 10 micrometers. The obtained magnesium alloy sheet exhibited an equivalent limiting drawing ratio in a warm-drawing test.

6 [1] [2] [3] [4]

References S. Yoshihara, H. Yamamoto, K. Manabe and H. Nishimura, Journal of Materials Processing Technology, Vol. 143-144, 2003, 612. F. Moll, M. Mekkaoui, S. Schumann and H. Friedrich, Proc. of the 6th Int. Conf. Magnesium Alloys and their Appln., DGM 2003, 2003, p936. H. Watari, K. Davey, M. T. Alonso Rasgado, T. Haga, Journal of Materials Processing Technology, Vol 155-156, 2004, 1662. H. Watari, R., Paisarn, N. Koga, T. Haga, K. Davey, M. T. Alonso Rasgado, Key Engineering Materials, Vols. 274-276, 2004, 379.

77

Strip Casting of Mg-Al based Alloy with Ca by Twin Roller Caster K.Matsuzkai, K.Hatsushikano, Y.Torisaka, K.Hanada and T.Shimizu National Institute of Advanced Industrial Science and Technology(AIST), Tsukuba, Japan

1

Abstract

AZ61 with 0.25wt%Ca strips were directly prepared from molten by twin roller caster. The addition of Ca enable to perform the strip casting process of Mg alloys in air without cover gas. This is expected to simplify process. The cast strip is composed of fine equiaxed grains with a grain size of 20Pm for the thickness of 1.0 mm. The mechanical properties increase with decreasing thickness, and the yield stress and ultimate tensile strength reach 160 MPa and 240 MPa, respectively, for the strip with 1 mm thickness. This is due to the grain refinement introduced by strip casting. A further improvement of mechanical properties was achieved by thermomechanical treatment. It is thought that the strip casting process is useful for production of Mg alloy sheets and the Mg alloy with a small amount of Ca is suitable for strip casting.

2

Introduction

Mg alloy is the lightest metallic constructional material with a density of 1.78 Mg/m3 and have an attractive possibility for weight saving the vehicle. At present time, most of Mg components are produced by using Die Cast process because of its good castability. On the other hand, Mg wrought alloys have superior mechanical properties compared to cast alloys. In conventional Mg sheet production process, a thin sheets is produced through many steps, because of poor formability of Mg due to hcp structure . This might lead to an increase in the cost of Mg wrought alloys. The strip casting process makes it possible to directly obtain the sheets with thickness less than one or two mm, leading to the reduction in cost and energy consumption. Furthermore, the strip casting have an effective process for refining the microstructure and reducing the segregation, leading to the improvement of mechanical properties. However, the report for the strip casting of Mg alloys has been limited [1–4]. This paper intends to produce Mg-Al based sheets with Ca by strip casting, and clarify the microstructure and mechanical properties.

3

Experimental Procedure

Pure Mg, Al Zn and Ca were used in the present study. Firstly, about 100 g of Mg-Al-Zn alloy with 2.5%Ca were melted in the graphite crucible under the argon atmosphere by induction melting. The obtained alloy ingot was inserted into a steel crucible with appropriate Mg, Al and Zn, and melted into a Mg93Al5.75Zn1Ca0.25 alloy by electric resistance furnace. The mixture of Ar and SF6 was used during the melting to protect the molten from oxidation or burning. After

78 them, the molten alloy was move to the twin roll caster and poured into a tundish between rolls. The casting was performed without a protective gas. The steel rolls with a diameter of 300 mm and a width of 50 mm was used. Roll gap was in the range of 0.5 to 4 mm and roll speed was 16 m/min. The casting temperature was in the range of 893 to 923 K. The cast strips with a thickness of 1.5 mm were subjected to warm rolling and heat treatment. The microstructure was examined by an optical microscope. Hardness was measured by a micro Vickers hardness tester and tensile properties were measured by a Shimazu Autograh material testing machine.

4

Results and Discussion

Figure 1 shows the external appearance of a as-cast Mg alloy strip with a thickness of 1.0mm . No oxidation is observed on the surface of the sheet. The sheets with a thickness of 4mm are also obtained without oxidation. In the case of AZ61 alloy, the strips with a thickness above 3mm show black or brown surface due to the oxidation. Moreover, the cover gas was needed to protect molten alloy in the tundish from burning. An addition of 0.25%Ca is effective to prevent the oxidation or burning. This protective effect of 0.25%Ca was valid for AZ31 and AZ91 alloy.

Figure 1: External appearance of as-cast Mg93Al5.75Zn1Ca0.25 with a thickness of 1.0 mm

Figure 2 shows the microstructure of center area of the strip cast Mg alloy with different thicknesses. The strip with a thickness of 4 mm is composed of dendrite and equiaxed grain with a average size of 40 μm. With decreasing thickness, the dendrite region and grain size decrease. For the 1 mm, the fine equiaxed grains with an average size of 20 Pm are obtained. This microstructure is favorable for the mechanical properties. It is reported [1–4] that strip cast AZ31or AZ91 alloys show a dendrite structure. In the present study, the alloy has a narrow freezing range and the roll speed is higher compared to that of the reported study. Therefore, the favorable microstructure is obtained. Figure 3 shows the X-ray diffraction patterns of as-cast Mg93Al5.75Zn1Ca0.25 alloy strips. The strips consist of hcp-Mg and Mg17Al12 and no peaks corresponding to Al2Ca are observed. The intensity of (002) peak increases with decreasing thickness. This suggesting that the oriented structure with a basal plane parallel to the surface of the strip is developed Mechanical properties of as-cast strips are summarized in Table 1. The yield stress(YS) and ultimate tensile strength (UTS) increases with decreasing thickness and reach 160 MPa and 240 MPa, respectively, for 1mm thickness. There is also an increase in elongation and hardness. It is notable that there is no significant difference in mechanical properties for the strip of

79

Figure 2: Microstructure of as-cast Mg93Al5.75Zn1Ca0.25 with different thicknesses, (a) 4 mm, (b) 3 mm, (c) 1 mm

Figure 3: X ray diffraction patterns of as-cast Mg93Al5.75Zn1Ca0.25 strips

1.5 mm thickness between Mg93Al5.75Zn1Ca0.25 and Mg93Al6Zn1(AZ61). This suggests that the addition of a small amount of Ca such as 0.25 % is not harmful to mechanical properties of Mg alloy. The increase in mechanical properties is due to the grain refinement, introduced by casting.

80 Table 1: Mechanical properties of as cast Mg alloy sheets produced by twin roller casting Mg93Al5.75Zn1Ca0.25

Mg93Al6Zn1(AZ61)

(1mm) (1.5mm) (3mm) (4mm) (1.5mm)

YS (MPa) 160 120 110 105 120

UTS (MPa) 240 200 175 160 210

Elongation (%) 10 7 5 4 8

Hardness (HV) 75 75 65 60 75

However, theses value are slightly low compared to those of AZ61 sheets produced by a conventional sheet production process. This is maybe due to the inhomogeneous microstructure and then the improvement of mechanical properties is expected to be achieved by a subsequent thermomechanical treatment. The strip with a thickness of 1.5 mm annealed for 10h at 673 K shows an improvement of elongation to 12 %.The subsequent rolling at 473K to 1.0mm thickness and annealing at 673 K for 1 h causes an increase in both strength and ductility and YS, UTS and elongation reach 160 MPa, 280 MPa and 18 %, respectively. A further improvement is expected by optimizing the condition of the thermomehcanical treatment.

5

Conclusion

Mg93Al5.75Zn1Ca0.25 strips can be produced by twin roll caster without cover gas. The addition of 0.25%Ca is effective to prevent the oxidation or burning. The strip with a thickness of 1mm consists of fine eqiaxed grain with a size of 20 Pm. The mechanical properties increase with decreasing thickness and YS, UTS and elongation reach 160 MPa, 240 MPa, 10 %, respectively, for the 1mm. A further improvement is achieved by the thermomechanical treatmen. The mechanical properties the Mg93Al5.75Zn1Ca0.25 strip is comparable to those of the alloy without Ca, indicating that a small amount of Ca have no significant influence on mechanical properties of Mg alloy. It is concluded that the strip casting process is useful for production of Mg alloy sheets and the Mg alloy with a small amount of Ca is suitable for strip casting.

6 [1] [2] [3] [4]

References S.S. Park,J.G.Lee,Y.S.Park and N.J.Kim, Materials Science Forum, 419–422,2003,599 B.S.You,C.D.Yim,B.S.Kam adnW.W.Park, Materials Science Forum, 488–489, 2005, 337 Y.Nakaura and K.Ohori, Materials Science Forum, 488–489,2005,419 C.Yang,P.Ding,D.Zhang,F.Pan, Materials Science Forum, 488–489, 427

81

New Strip Casting Process for Magnesium Alloys Fr.-W. Bach, M. Hepke, A. Rossberg Institute of Materials Science (IW), University of Hanover, Germany

1

Abstract

One of the main reasons for the hesitant application of magnesium sheets is their high price which is primarily caused by a lack of adequate feedstock for their production. Currently, thick, oval direct chill (DC)-billet cast slabs, which have to be milled before rolling, are used industrially. Therefore a main research aim at IW is to cast and deform magnesium near-net-shape strips in one heat in order to improve the forming capacity and simultaneously cut production costs. For this purpose a twin-roll-stand was modified to function as a Hazelett-type caster. Two variable casting belts, a melt-feeding device, an inert-gas plug furnace and a cooling system had to be adapted to a twin roll stand. These components were constructed especially for casting highly reactive magnesium melts contamination-free. Additionally, several security- and control measures were installed. First casting experiments with the standard alloy, AZ31 (MgAl3Zn1), showed the feasibility of the new plant-concept. Minor deficiencies were remedied by plant-modifications and/or changes of the construction materials. Currently, optimized parameter sets for several alternative alloys are under examination. The cast strips produced with the new caster-type were analyzed concerning their suitability to serve as rolling feedstock. The samples show a roll-able surface and a significantly finer microstructure than conventional cast products. The properties of first sheets (200 u 200 u 1 mm) made of this new feedstock have been determined. In order to evaluate the process holistically, the quality of finished magnesium sheets and the cost saving potential of the new process chain were compared to standard production techniques.

2

Introduction

Due to their high mass-specific mechanical properties, magnesium sheets offer a significant weight-saving potential in modern vehicle constructions. To increase the range of applications e.g. in the automotive industry, magnesium sheets for deep drawing purposes have to comply strict requirements regarding corrosion resistance, geometric tolerances, forming capacity, joining technology and surface quality [1]. The feasibility of automotive applications is subject to various current research projects. For example in the BMBF-funded project “ULM – Appropriate process chain for ultralight components made of magnesium sheet metal for transportation” an engine bonnet for the VW Lupo was developed as a demonstrator part. This 100% magnesium solution has a Class-A surface, shows acceptable corrosion resistance and is joined by using various techniques [2]. The range of automotive applications implies among others: closures (e.g. door panels), seat-cups, splash-boards, roof sections, convertible roof covers, instrument carriers, trunk panels, etc. In contrary to the above said, magnesium sheets have not

82 yet made the step into series application. Besides the natural aloofness of conservative design engineers, the financial aspects of the application of magnesium sheets are the main obstacles. Firstly, high purity magnesium wrought alloys are at least as expensive as aluminum. Currently the Chinese production is dramatically increasing which relieves this situation [5]. Secondly, the continuous DC-casting process of the rolling feedstock is comparatively complex and involves the usage of costly cover gases [7]. Before rolling can begin, the cast surface has to be milled, because the stationary dies cause cracks and wrinkles in the surface of the billets. Thirdly, the relatively thick slabs have to be hot-rolled due to the coarse grain sizes in the hexagonal lattice. This involves several reheating steps between the numerous rolling passes. Finally the production of the high quality parts requires heated deep drawing, corrosion protection and joining (Figure 1).

Figure 1: Process chain (from the pre-production to the finished part) [6]

The aim of this contribution is to reduce the production costs by improving the casting- and rolling process. This can be achieved by applying an appropriate continuous casting principle and thus producing better deformable feedstock.

3

Plant Concept

In the strip casting process, near-net-shape feedstock can be produced very economically. This group of casting principles is characterized by an increased productivity and a high quality-level. The comparatively rapid cooling causes a fine grain structure. If moving dies are applied, the surface can result smooth enough for rolling without pre-treatment. Two of the main caster principles shown in Figure 2 were analyzed regarding their suitability for strip-casting magnesium.

Figure 2: Twin-roll casters, single belt caster [3]

The twin-roll casters (Figure 2 left) consist of a nozzle (2) and two rolls (1). The melt solidifies directly between the cooled rolls. The orientation can be vertical or horizontal e.g., as

83

Figure 3: Hazelett-caster [3]

shown in [4]. The deficits of this method are the direct contact of the melt and rolls (surface cracks) and the difficult (inner) cooling of the rolling-cylinders. The principle of single belt casters, as pursued by [8], is shown on the right image. The melt is poured on top of a cooled belt (1) and has to be covered with protective gases (2). The problem of this caster, besides the gas consumption, is the asymmetric solidification and the surface quality of the free upper side of the strip. Since magnesium shows strong anisotropy and reactivity, problems in these directions are predictable. The Hazelett-caster (Figure 3) applies two belts (1), at least 4 rolls (2) and a tundish (3). The melt enters a closed system and the moving belts transport the strip out of the solidification zone. From the multitude of continuous casting principles, the Hazelett-type was chosen, because it offers several advantages regarding the properties of magnesium melts. Firstly, it operates with a closed die so that there is principally no need for using large amounts costly protective gases. Secondly, it functions horizontally which (in case of a failure) minimizes the amount of leakage compared to vertical (gravity) casters. Thirdly, the cooling of the belts is much simpler and more effective than an inside cooling of the rolls. Additionally, it is possible to cast very thin, near-net-shape strips (min. 1 mm). The cast skin is much smoother because the die moves along with the melt. Finally it is possible to deform the solidified strips between the main rolls which causes homogenization and grain refinement. In order to benefit from these general advantages, the dimensions and materials of the construction had to be adapted precisely to the requirements for casting magnesium. The general demands for all components in melt-contact are that they do not react with magnesium, that the melt does not stick to them and that they endure the occurring thermo-mechanical stresses. The functions which have to be fulfilled are melting under protective gas, heating to casting temperature, dosing, bide a holding period and then continuously solidify, cool, transport and deform the strip. Figure 4 gives an overview of the chosen construction. It can be seen that a plug furnace (6) was adapted above the caster. It is also possible to cast with a pressurized oven via a dosing tube from below the caster, which may be desired for safety reasons, since gravity forces the melt back into the system in case of failures. The melt feeder (8) changes the profile of the melt channel from a round tube to a rectangle. In case of a blocking in the caster (early solidification) it is possible to empty the furnace with a drain plug. This part is made of cast steel which stays bright after heating and is inert to magnesium. As lateral seals fixed steel panels (4) were utilized. Fortunately magnesium does not stick to them, thus it was possible to seal the system without using complicated chains (so called dam blocks) which move along with the cast strip. The belt adjustment (9) on the one hand suits to tighten the belts. On the other hand the cast- and the roll gap as well as the declination of the belts can be varied widely with this solution.

84

Figure 4: Magnesium specifically modified Hazelett-caster

The cooling zone is shown in detail in Figure 5. Magnesium melt enters the cast gap through a melt inlet (5). The final material choice for this part was titanium because of its low thermal

Figure 5: Cooling zone

conductivity which is needed in order to minimize heat losses from the heated into the cooled zone. The casting belts are pressed onto the seals by the pressure bridge (4). It also avoids buckling of the belts. Two water-air spray nozzles (1) cool the thin belts and hence the cast strip.

4

Experiments / Casting Parameters

Experiments were so far conducted with the standard magnesium alloy AZ31 (MgAl3Zn1). Firstly geometrical parameters such as the cast- and roll gap and the declination were varied. Secondly the balance between early solidification (blocking of the caster) and melt leakage (Mg flows though the caster) was established by variation of the holding time, melt- and caster temperatures, cooling rate and casting speed. Calculations backed by former experiments served as a staring point from which optimization was conducted. Adequate results were obtained using the parameter set shown on Table 1.

85 Table 1: Casting parameters for AZ31 Melt temperature [°C] Feeder temperature [°C] Belt speed [m/min] Holding period [s] Cooling [l/min]

Range 650 – 750 600 – 670 0–7 0–5 gas / spray: 0 – 80

Selected 690 670 2,1 1 argon: 80

With these parameters is was possible to cast 1200 mm strips of AZ31 with a thickness of 8,3 mm. The length was only limited by the melt volume in the furnace. Anyhow, some deficits were registered. The strip exits the moving die with a high temperature (approx. 450 °C) , therefore reactions with the ambient air cannot fully be avoided. A secondary spray cooling at the exit can solve this issue. Secondly, after longer operation times, magnesium tends to stick to the belts and seals. A continuous lubrication (e.g. with boron nitride) should be adapted.

5

Results and Conclusions

The strips were analyzed regarding their surface quality and their microstructure. Table two shows an overview of the average grain sizes in common magnesium rolling feedstock. Only pre-deformed material, e.g. by extrusion has a finer microstructure than the thin cast strips. Therefore strip cast material can be rolled with higher strain from the beginning on. In addition to the near-net-shape geometry, this is a large cost saving factor in contrast to billet cast feedstock. Figure 6 allows a comparison between the strips produced in the modified Hazelett caster and conventional industrial billet cast slabs. It can be seen that the strip surface (where no reaction with air occurred) is suited to enter the rolling process without milling, whereas billet casting always demands such a preparation. In conclusion it can be said that the presented laboratory scale results show the feasibility of the new process chain for magnesium sheet production via strip casting. As soon as minor defi-

Figure 6: Comparison of surface quality and microstructure

86 cits are remedied, high quality, near-net-shape feedstock can be produced very economically with the newly modified Hazelett caster. First estimations predict a cost saving potential of approximately 50 %, resulting mainly of less rolling stages, heating cycles and material losses, in comparison to billet cast feedstock. Table 2: Comparison of average grain sizes for different rolling feedstock Alloy AZ31

6 [1] [2] [3] [4] [5]

[6]

[7] [8]

Steel mould casting 500 μm

Extrusion 20 μm

Billet casting 300 μm

Strip casting 80 μm

References S. Schumann, H. Friedrich, Materials Science Forum, Vols. 419–422 (2003), pp. 51–56 P. Juchmann: Magnesium-Knetwerkstoffe für komplexe Ultraleichtbauaufgaben, Braunschweiger Symp. Faszination Karosserie, 2003, Seite 188–195 A. Stepanov, A. Neustruev, J. Zilberg, Blechherstellung aus der Schmelze, „Metallurgie“, Moskau, 1978 H. Pircher, R. Kawalla, German Patent No: DE 100 52 423 C1, Verfahren zum Erzeugen eines Magnesium-Warmbands, Patent owner: Thyssen Krupp Stahl AG, Jan. 2002 Fr.- W. Bach, M. Schaper, A. Kuhlmeyer, M. Schäperkötter, J. Weber, A. Rossberg: Moderne Blechwerkstoffe für die Automobilindustrie; ein Vergleich, Berichte aus dem IWU, Band 22, 2003, Seite 9–26 Fr.- W. Bach, M. Rodman, A. Roßberg: Improvements to the rolling- and deep-drawing process of magnesium wrought alloys, DGM conference Magnesium, Wolfsburg Dec. 2003, p. 285ff F. Pravdic, D. Leitlmeier, Leichtmetallzentrum Ranshofen, Vertical Direct Chill Casting of Magnesium, DGM conference Magnesium, Wolfsburg Dec. 2003, p. 675ff H. Palkowski L. Wondraczek: Herstellung von Magnesiumband mittels Single-BeltCaster: Grundlagen und Ergebnisse, Metall, 58. Jg. Heft 12, 2004

87

Production of Twin Roll Cast AA6016 for Automotive Applications M. Dündar1, Ö. Keles1, G. Anger2 Assan Aluminum, østanbul, Turkey AMAG Automotive GmbH, Ranshofen, Austria

1 2

1

Introduction

The growing demand for fuel-efficient vehicles to reduce energy consumption and emission stimulate development of age-hardenable aluminum alloys to meet the requirements of automotive applications. A continuous increase has been observed in share of aluminum sheet products among all aluminum parts employed in auto bodies. Inner and outer panels are typical applications for sheet products. In this regard, 6000 series aluminum alloys are the primary candidate to achieve comparable performance of conventional material, i. e. steel, that has been employed since the beginning of the century. The characteristic properties of aluminum; high strength and stiffness to weight ratio, good formability and corrosion resistance exist in AA6016. However, the cost of aluminum and its conversion cost remain the biggest impediment for its large scale use in automotive applications [1–3]. Cost effective solutions in production, recycling, or in short, effective life cycle analysis of aluminum alloys, will bring it to a position at which it can find variety of application fields, not only limited to automotive industry. Production of aluminum sheet by Twin Roll Casting (TRC) route rather than by conventional DC casting and hot mill method offers an opportunity to substantially reduce the cost barrier, which could lead to an increase in its use. Combination of low operational and investment costs to the TRC with the shorter production times and flexibility in switching from one alloy to another in casting operation stands for strong economical aspects of this production method. Superior micro structural features inherited to the material due to the solidification mechanism and its good response to down stream processes are material related issues believed to accelerate the acceptance of twin roll cast sheet in the market [4–5]. Present study investigates the performance of twin roll cast AA6016 aiming automotive applications. As-cast and processed strips were subjected to the microstructural and mechanical characterization techniques to test their compliance to specific application of automotive industry.

2

Experimental

Industrial scale coils, having average weight of 9 tons, were produced in ASSAN Aluminum, Turkey by employing 2200 mm wide SpeedCaster® at the width of 1800 mm. Chemical composition of cast strips are given in Table 1. Since this material is aimed to be used in automotive applications, one of the basic requirement is that to produce strip having defect free surface. Solidification mechanism operating at the caster roll-gap, in general, is believed to be inadequate and promote segregation behavior at the roll-strip interface while producing alloys with wide

88 solidification range. Thus, 6000 series alloys are among those that have not been produced with a consistent and reproducible quality, up to date. Ability to increase control over the casting parameters is the key factor for tailoring overall micro structural features of a strip produced with TRC. As-cast strip surface quality, in terms of surface segregations and ripples has exhibited very good performance, as far as final product quality expectations were concerned. Down stream processes of the coils were carried out in the facilities of project partner, AMAG Rolling GmbH, Austria. The strips were processed through cold rolling to 1 mm as well as up to 2,5 mm, then solution heat treatment in a continuous heat treatment furnace, water quenching and investigated after aging 7 days at room temperature. Table 1: Chemical composition of TRC AA6016 Si 1,06

Fe 0,17

Cu 0,07

Mn 0,07

Mg 0,34

Al 98,25

Through-thickness micro structural investigation of as-cast samples was conducted in two directions, namely parallel and perpendicular to the casting direction. Samples were polished with standard metallographic techniques: ground with SiC paper, polished with 3 Pm diamond and finished with colloidal silica. Macro and micro structural investigations of the as-cast samples were done by using a Zeiss Axiotech Vario model optical microscope. Micro structural features were revealed by etching 0,5% HF solution. Grain structure was observed with cross polarized light after electro polishing with Barker’s solution. Semi-quantitative analyses of micro structural constituents were conducted by using JEOL 5600 SEM equipped with Oxford EDS unit. Line and dot mapping analysis were conducted for determining chemical content of individual constituents.

3

Results and Discussion

3.1

Micro and Macro Structure of As-Cast Strip

Important characteristic of twin roll cast alloys, is the heterogeneity of the grain structure through the thickness of the cast strip. This structural heterogeneity is generally attributed to the cooling rate gradient encountered in the twin roll casting process. Very high solidification rate of initially contacting liquid metal with caster rolls generates a supersaturated aluminum matrix at the outermost layers of the strip. During the movement course of solidified metal in the roll gap, metal is forced to pass through the opening between two caster rolls, called indicated casting gauge. This process induces limited amount of plastic deformation at the shallow depth of the strip. Thus, the macroscopic structure near the surface revealed the characteristic appearance of a featureless zone at the outer fibers and gradually increasing and elongated ones at the quarter plane of the thickness (Figure 1). Obviously, alignment of the grains, especially those located at the outer fibers, reveal different grain structures in parallel and transverse to the casting direction (Figure 2 a and b)

89

Figure 1: Grain structure of as-cast strip at the surface. Note the border of featureless zone marked with the arrows

Figure 2: Grain structure of as-cast material (a) parallel, (b) transverse to the casting direction

Decrease in casting gauge has pronounced effect on the proportion of plastically deformed volume through the thickness. The angled position of the grains becomes more parallel to the casting direction while intermetallic particles, decorating the grain boundaries, are more aligned in the same direction. Center plane of the strip is occupied by more or less equiaxed grains. Like the other alloys produced with the same technique, centerline segregation at the centerplane of the strip is an unavoidable consequence of solidification mechanism of this production method [5–6]. Tubular form of individual segregates of those does not alter equiaxed grain structure of the center plane. Two distinctive segregation patterns were observed under specific casting condition. One of them was in the form of solute reach channels called centerline segregation, as was already explained and the others are dispersoids. Limited depth from the free surface of the strip is decorated with very fine intermetallic particles in the size of 1–3 Pm (Figure 3). Generally four intermetallic phases are possible: Mg2Si as the age-hardening phase, Į-AlFeSi(Mn) (Al5Fe3Si(Mn) Chinese script in a as-cast condition, Mn as substitution of Fe possible), ȕ-AlFeSi (Al5FeSi little plates which are reducing ductility and ʌ-phase (Al8FeMg3Si6 - is binding Mg) (Figure 4).

90

Figure 3: Distribution of intermetallic phase at the vicinity of the surface

Detailed investigation of segregate channels at the center plane, called center-line-segregation, has shown that they generally contain coarse eutectic cells and occasional coarse intermetallic phases. Despite of two different intermetallic particle geometry in the microstructure, their size and distribution are very uniform and consistent compared to those observed in its DC cast counterparts. While intermetallic particle size of DC cast material is larger than 6 Pm, their morphology exhibit large aspect ratios, as well. This particle geometry is more prone to align itself in deformation direction, i.e. rolling direction. It has already been known that manipulation of casting parameters might have strong influence on tailoring micro structural features to certain extend. However, in the present study, all efforts were concentrated on optimization of casting parameters to achieve minimum or no surface segregations. This approach has allowed not only for obtaining segregation and ripple free surfaces but also satisfactory productivity values in casting operation.

3.2

Micro and Macro Structure of Heat Treated Material

Twin roll cast strips of alloy AA6016 were cold rolled to 1,15 mm and then subjected to solution heat treatment. The characterization of the material was done after aging for 5 days at room temperature to T4 temper. Similar processing steps were also applied to the DC cast and hot rolled strip for fair comparison. Table 2: Chemical composition of DC AA6016 material Si 1,11

Fe 0,17

Cu 0,08

Mn 0,07

Mg 0,35

Al 98,17

The DC cast hot rolled AA6016 material which was used for comparison purpose shows a very similar chemical composition as the produced TRC 6016 material (Table 2).

91

Figure 4: Grain structure of TRC (a) and DC cast (b) material after recrystallization

Substantially small grain size was encountered through the thickness of TRC material after solution heat treatment which readily leads to re-crystallization. Contrary to the non-uniform grain size distribution of as-cast strip, re-crystallized grain structure at the final gauge is equiaxed and uniform, regardless of its position in the thickness. Their sizes are almost half of those in DC cast material (Figure 4). It was previously noted in the literature that the structure of the sheet prior to the final recrystallization, both in terms of the size and distribution of second phase particles controls the microstructure developed during recrystallization [7]. Microstructural constituents of both materials, after T4 treatment, were also investigated. Figure 5a shows distribution and morphology of intermetallic particles at the mid thickness of TRC AA6016. Compared to as-cast structure in which interdendritic area is decorated with coarser particles, along with fine ones, their sizes were found to be relatively smaller after solution heat treatment and T4 treatment.

a) b) Figure 5: Intermetallic particle size and distribution at the half thickness of the sheet

As was expected, DC cast AA6016 revealed much coarser particles that are aligned in the rolling direction. Their aspect ratio is considered to be the major factor for this alignment (Figure 5b). Different solidification characteristics during casting operation of both material determine their particle size and geometry. Coarser particles of DC cast/hot rolled strips are prone to be breaking up into two or more fragments, especially during cold rolling stage of down stream

92 operation. Corner profile of particle ends, sharp and rectangular, causes intense plastic straining in the nearby matrix and leads to void nucleation at the corners of the particles. Rolling to thinner gauges leads this damage mechanism more likely to occurs.

3.3

Mechanical and Forming Properties of Heat Treated Material

Processing route and their parameters were determined to fulfill the typical requirements of automotive industry for stamping operation of critical parts. The heat treatment was done in an industrial scale on the continuous heat treatment line (floater type furnace). T4 temper was achieved by solution heat treating at 540 °C/90 s followed by water quenching and aging 7 days at room temperature. Mechanical tests were conducted in transverse direction to the rolling direction. Mechanical properties of paint bake response were investigated on the material exposed to 205 °C for 30 min after T4. Table 3: Mechanical properties of TRC and conventional DC cast AA6016.

Yield Strength (MPa) Tensile Strength (MPa) Uniform Elongation (%) Elongatin (A50) Elongation (A80) n value (4–6 %) r value (8–12 %)

TRC T4 110 220 22 29 26 0,29 0,64

T62 230 275 9 12 – – –

DC T4 105 210 23 29 26 0,31 0,65

T62 210 260 10 13 – – –

TRC and conventional AA6016 exhibit comparable results in both tempers. T4 and T6 tempers of TRC coils has slightly higher yield and tensile strength. Regardless of testing gauge, i.e. 50 or 80 mm, almost identical total elongation values of TRC material were obtained with DC cast material. The uniform elongation is controlled by the relative strain hardening rates up to the maximum load while the extent of post-uniform elongation depends on both strain hardening and strain rate sensitivity. Similar to the material produced with conventional methods, substantial fraction of the total elongation belongs to the uniform elongation in the case of TRC material. All achieved mechanical properties with the current TRC AA6016 are in accord with the previously published data [8–10] and comply the standard European specifications for automotive sheet applications. Mechanical characterization was extended by constructing Forming Limit Diagrams (FLD) of both materials to determine limiting strains under simultaneously operating complex stress states, as encountered in industrial stamping operations (Figure 6). FLDs were constructed for their T4 tempers at which stamping is carried out. DC cast material has slightly higher limiting strains for the right side of the diagram than that of TRC. As this behaviour is assessed in the light of strain hardening exponent, that is relatively higher in DC cast material, might lead to diffuse necking by retarding local thinning and resulting premature failure. However, for the

93

Figure 6: FLD of both material at T4 temper

strain combinations of Hminor < 0 and Hmajor > 0, representing strain combinations in stretching, TRC material exhibits better performance compared to its counterpart.

4

Summary

1. Contrary to the coarse and aligned characteristics of intermetallic particles in the material produced with DC casting/hot rolling route, very fine and uniformly distributed intermetallic particles of TRC sheet have pronounced contribution to the complex straining conditions of stamping operations and failure mechanism. 2. TRC sheet shows satisfactory performance in bending process which is unavoidable joining (such as hemming) and forming method of automotive industry. 3. Fine intermetallic particles improve surface appearance in bending and overall bendability performance. 4. A phenomenon, orange peeling, that deteriorates the surface appearance due to the coarse grain structure, can be avoided with much finer grain structure of TRC sheet. 5. Not only the uniaxial tensile properties, but also complex deformation characterization technique, FLD, shows that TRC sheet has almost identical performance with that of DC cast counterpart. 6. Supersaturated outer fibers of the strip formed upon contacting with the caster rolls and associated solidification mechanism might result in different age hardening kinetics as exposed to solution heat treatment and subsequent age hardening. This phenomenon needs further detailed studies. 7. In connection with the same reasoning, influence of alloy chemistry on manipulating mechanical performance of TRC strips need further investigation, especially with changing Si content.

94

5 [1]

References

W.S., Miller, L., Zhuang, J., Bottema, P.De, Wittebrood, A., Haszler, Vieregge, 2000, Materials Science and Engineering, (A280), 37 [2] S.A. Arnold, 1993, J. Metals, 45 (6), 12 [3] Y., Muraoka and H., Miyaoka, 1993, J. Mater. Proc. Tech. (38), 655 [4] K., Sears, Automotive Engineering: Strategic Overview, 1997, (1), 55 [5] S, Ertan, M., Dundar, Y., Birol, K. Sarıoglu, C., Romanowski, 2000, Light Metals, TMS [6] M., Dundar, K., Sarıoglu, Y., Birol, A.S., Akkurt, and C., Romanowski, 2002, ed. Das, S., Automotive Alloys, TMS [7] G.B., Burger, A.K., Gupta, P.W., Jeffret, and D.J., Lloyd, 1995, Materials Characterization, 35, 23–39 [8] L., Zhuang, R.De, Haan, J., Bottema, C.T.W. Lahaye and P. De, Smet 2000, Materials Science Forum; 331–337:1309 [9] J., Hirsch, 1997, Materials Science Forum; 242:33 [10] S.M., Hirth, G.J., Marshall, S.A., Court, D.J., Lloyd , 2001, Materials Science and Engineering; A319–321: 452

95

Magnesium Upward Direct Chill Casting Fr.-W. Bach, S. Schacht, A. Rossberg Institute of Materials Science (IW), University of Hanover, Germany

1

Abstract

Since 2001, the development of a new continuous casting technology for magnesium alloys at IW has made mayor steps from laboratory to industrial scale. Compared to the first trials with 40 mm diameter billets the process has undergone many changes and enhancements towards the currently used technology for casting high quality 90 mm and 203 mm diameter billets. With proceeding development during a FP5-Project named “EuroMagUpCaster”, various advantages of the process have become apparent. These achievements do not only affect safety topics, but also the possibility to significantly increase productivity, to lower installation costs, etc. Compared to conventional vertical direct chill casting the upcasting process differs in several issues. This concerns especially the cooling conditions in the mould and the possibilities of exact thermal control of the molten metal right outside the mould. An effective secondary cooling zone reaching high cooling rates by usage of only small amounts of water-spray is another highlight of this technology. This contribution is giving a detailed insight into the actual stage of development of this innovative technology, which has the potential to become the state of the art in continuous casting of magnesium in future.

2

Introduction

In order to increase the industrial application of magnesium wrought products, high quality feedstock for rolling and extrusion has to be available at an acceptable price level and in various alloy compositions. The DC-casting technique, as for steel and aluminum, is suited to produce high volumes of such material in a continuous casting process. In several research- and industrial plants magnesium is cast in vertical (gravity) DC-casters (Figure 1, left). Unfortunately only few standard dimensions and alloys can currently be purchased. Alternative feedstock has to be ordered in high quantities which are not suited for today’s small series production. Regarding the quality of such products a standardized level has not yet been established. Especially in the case of alternative wrought alloys, the resulting microstructures and surface properties often do not fulfill the high requirements of forming processes such as rolling or extrusion. Therefore milling of the surfaces and homogenization of the material has to be implemented in the process chain. Another drawback is the security aspect of gravity casters. In the case of magnesium, small disturbances in the casting process can easily result in ruptures of the cast billets. In comparison to aluminum, these failures are more likely due to the lower strength of the just solidified material and the smaller process window, which is caused by the high thermal conductivity and the low volume specific heat capacity. Furthermore leakage of magnesium melt is much more critical since the high reactivity with oxygen leads to strong fires at atmosphere which can not easily

96

Figure 1: Comparison of VDC- and UDC-casting principles [e.g. 1, 2]

be extinguished. Often flux (salt) has to be used in order to suffocate these fires. This method is quite effective, but goes along with the emission of corrosive gases, which aggressively damage the surfaces of surround machines, etc. Therefore a system which prevents high volume melt leakages offers great advantages for the secure operation of continuous magnesium casters. At the Institute of Materials Science of the University of Hanover, the Upward Direct Chill (UDC) casting principle (Figure 1, right) was chosen for a magnesium specific adaptation. Since this casters requires an active dosing against gravity, it also offers the possibility to retain the melt in the furnace in case of billet ruptures. In the following article, the optimized plant construction and the so far examined casting experiments, which have been conducted within an EU - 5th Framework project shall be presented.

3

Plant Concept

Since the year 2000 a magnesium specific UDC-caster has been developed and optimized at the Institute of Materials Science [3]. A first laboratory scale version with a diameter of 40 mm showed the feasibility of the process. In 2001 the EU-funded “EuroMagUpCaster” project was granted. The scope of this international collaboration is to introduce this caster principle into near industrial production. For this purpose two 90 mm moulds and one 203 mm (8 inch) mould were designed and build at IW and at the project partner KME. Figure 2 illustrates the 3-D construction of the furnace with the attached mould. Within the pressurized container magnesium can be melted in an inert crucible under protective gas. When the gas pressure is increased for dosing, the melt flows upwards through a heated riser tube which is via a compensator on top of the oven lid. The mould consists of a heated titanium inlet, a water cooled copper cylinder and a graphite insert. Through this porous ring, a viscous lubricant can be pressed into the mould. It avoids sticking of the melt and closes the shrinkage gap between billet and mould, which increases the heat transfer into the copper cooler.

97 Above the mould, a secondary cooling ring, which contains several water-gas-spray nozzles is installed. The spray angle is tilted, so that water cannot drop downwards into the mould. At several points in the caster, thermocouples and pressure sensors are installed, in order to document and control the casting speed, the spray pressures, the lubrication flow and the heater temperatures.

Figure 2: Pressurized furnace with 203 mm mould

Figure 3: Assembly with feeding device [4]

On the photo of plant (Figure 3) the feeding device, which pulls the billet out of the mould, can be seen. This very rigid construction consists of 4 columns, the mould platform, the fixation of the starter block, a spindle and a belt driven motor. This first device allows pulling speeds up to 300 mm/min and a maximum billet length of 900 mm. For the near industrial application in the EuroMagUpCaster project, the consortium partner, Römer Födertechnik GmbH, has build three further plants with a max. pulling length of 2800 mm.

4

Experiments / Casting Parameters

Simulation results of the Spanish project partner, Inasmet [5], and experiences with the 40 mm mould, served as starting point in the determination of suited casting parameters. Fortunately the first set of temperatures, speeds and pressures already allowed the successful, stable production of 400 mm billets made of cpMg (99,5 %). It was soon shown that this process allows large parameter variations in order to optimize the billet quality. Due to safety reasons first experiments were conducted without secondary cooling. Table 1 shows the parameter sets for alloys with strongly different solidification behavior. In order to increase the casting speed, the secondary cooling has to be activated. First results of ongoing experiments lead to the estimation that the speed can be increased up to 30% in comparison to Tab. 1.

98 Table 1: Casting parameters for d=90 mm mould, lubrication 150–200 ml/m² surface Melt temperature [°C] Casting speed*1 [mm/min] Dosing over-pressure [mbar] Holding period [s] Primary Cooling [l/min]

AZ31 650-690 75-95 180-200 4-7 20-40

AZ80 630-670 60-80 180-200 4-7 30-40

ZEK100 670-690 80-105 180-200 4-6 30-40

*1 without secondary cooling

In Figure 4 it can be seen that a stable cooling condition in the mould is reached after approximately 200 seconds. The temperature of the titanium inlet (Ti ring) is plotted with higher resolution. Its alternating run is caused by early solidification on this part, which also results in the formation of a rough cast skin as can be seen in Figure 5. The thin rings which solidify there too early are pulled up after a certain time and the resulting cracks are closed by the following melt. This deficit shall be overcome by: a) higher melt velocity, b) stronger heating, c) better isolation and d) a changed geometry of the Ti-inlet.

Figure 4: Data for process analysis (AZ 31; without secondary cooling)

5

Results, Conclusions and Outlook

In the first stage, the billets were analyzed regarding the formation of the liquid pool and the appearance of the cast surfaces. Figure 5 allows to evaluate the results for cpMg and AZ80. Both allows were cast approx. 40°C above their liquidus temperature with strongly different speeds under otherwise equal conditions. It can be seen, that the intermetallic phases containing alloy AZ80 exhibits quite a different behavior. Still, in both cases the continuous casting was successful and the loss in the chipping is smaller than 8 mm on the diameter. The cast skin shall further be improved by a modified construction of the Ti-ring and the other measures mentioned above.

99

Figure 5: Comparison of liquid pool and billet surface (mould diameter d= 90mm)

Figure 6 gives an example for the microstructure of the cast billets. The results of the alloys AZ80 and ZEK100 were chosen, because here the fine microstructure (approx. 120 μm), the nearly globular grains (strongly reduced appearance of dendrites) and the homogeneous distribution of the precipitations (only AZ80) can clearly be observed. As main reason for this positive result, the controlled cooling conditions (high cooling rates) and the contamination free melt processing were determined. The alloy AZ31 shows a further reduced average grain diameter when the secondary cooling is activated. Results on the quantification of this parameter will be published soon to come.

Figure 6: Micrograph of AZ80 and ZEK100 (middle of the billet, d=90 mm, as cast condition)

As a next step the cast products will be tested at IW regarding their mechanical properties and by the other EuroMagUpCaster project partners with focus on their deformation behavior. Various samples have been delivered for extrusion, forging and even rolling experiments. Judging from the microstructure and small sample tests, significant improvements are being anticipated. Furthermore, the cast products of the upscaled mould (d = 203 mm) are awaited with great interest, since in case of similarly positive results the step into series application of this modern caster type comes much closer.

100

6 [1]

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

References F. Pravdic, P. Egger: Einfluss der Gießparameter auf die Qualität von AZ31-Strangussbolzen; Leichtmetallkompetenzzentrum Ranshofen, Gießereiforschung, Band. 57, 2005, S. 26–33 F. Bach, S. Schacht: Vertikaler Strangguß - ein aufstrebendes Verfahren, Werkstoffwissenschaftliche Schriftenreihe, Band 64, 2004, S. 4–10 U. Holzkamp: Entwicklung einer magnesiumgerechten Stranggusstechnologie, Dissertation, Fortschritt-Berichte VDI, Nr. 623, 2002 F. Bach, S. Schacht: New continuous casting process for Magnesium Alloys, IMA conference 2004; New Orleans, Louisiana, 9-12 May 2004 A. Landaberea, P. Pedrós, E. Anglada, I. Garmendia: Numerical Simulation of the upward continuous casting of magnesium alloys, DGM Conference, Continuous Casting 2005, Neu-Ulm

101

Spray Forming of Advanced High Strength Aluminum Alloys P. Krug, B. Commandeur PEAK Werkstoff GmbH, Velbert, Germany

1

Abstract

Most recently several new types of aluminum alloys were developed at the R&D Centre of PEAK Werkstoff GmbH, located in Velbert, near Düsseldorf. In any case, these new alloys make use of the high solidification velocity during spray forming. This leads to interesting microstructure with a high volume content of primary phases. These primary phases which can be pure Si-crystals or intermetallic phases like Mg2Si, Al3Fe are responsible for the good strength also at elevated temperatures. In the case of Mg2Si volume fraction of up to 25 % could be reached. Due to the low density of the magnesium silicide, the overall density of this alloy has been reduced down to 2,5 g/ccm, but with comparable strength to normal 2xxx series alloys. After spray forming a subsequent extrusion is mandatory to close residual porosity resulting from the atomising process. In some cases a additional heat treatment is applied to optimise properties according to the application. This presentation will show you the “making of” such alloys, their properties as well as present and potential applications.

2

Introduction

Aluminum alloys exhibit several advantageous properties or combinations of such properties like e.g. low density together with high thermal and electrical conductivity. Nevertheless, Young’s modulus, thermal expansion and wear resistance can be substantially improved by adding silicon. Unfortunately, casting of such alloys is not appropriate when the Silicon contents will exceed 22 wt.-%. Usually Silicon is limited up to 18 wt.-% used in alloys for pistons or monolitic engine blocks. Exceeding this limit leads to a brittle behaviour of the casting alloy and, therefore, there is a need for a contemporary manufacturing method. Such a method is spray forming of high Silicon-Aluminum alloys. Contents of up to 35 wt.-% can be easily obtained with in-situ formed primary silicon particles. It is worth to be noted that 100% of the silicon is precipitated primarily out of the melt and no eutectic silicon will be present after spray forming. The fine and homogeneous distribution of the silicon particles in a spray formed billet will lead to a good machinability and formability. Such alloys show superior behaviour in terms of stiffness and wear resistance. High surface quality can be installed only by turning and, thus, subsequent grinding and coating is not necessary any more as the example of a spool valve for variable camshaft timing shows.

102

3

Producing “Impossible” Alloys

Due to the high solidification velocity a variety of beneficial effects will occur. • • • •

Extended solubility Extended stability of phases Suppressed eutectic phases Enhanced nucleation of preferred phases

In addition to such metallurgical effects, spray forming offers the opportunity of injecting powders during spray forming process. Applying different conveying methods it is possible to inject • • • • • •

Spherical powders Sharp edged powders Nonmetallic powders (e.g. Carbides, Borides, Oxides) Metallic powders of same type (Overspray) Metallic powder of different type (e.g. Si-powder) Nanoscale powders

The combination of the metallurgical and the injection tools will lead to unusual materials with interesting properties. In investigating such production methods should lead at the end to a special principle of alloy design which should be claimed as “tailoring” of alloys redirected to specific customer’s requirements [1].

3.1

Metallurgical Tools

The binary system aluminum-silicon is well investigated and the influence of small additions (refining Silicon with Phosphorus or modifying the Eutectic with Sodium, Strontium or Antimony) is well established among the foundries. The influence of these small additions will diminish as solidification velocity increases and the solidification velocity will dominate the final structure and occurring phases. An interesting summary is given in [2, 3]. Starting from solidification near equilibrium condition one will expect a microstructure of a hypereutectic Al-35%Si-Alloy primary Si-Particles and Aluminum-Dendrites and Aluminum/ Silicon Eutectic. Increasing the solidification velocity a microstructure as shown in Figure 1a. will occur. The picture reveal the microstructure of a spray formed Al-35%Si alloy (DISPAL S220). The Primary Silicon becomes the dominating phase and the whole silicon content will be preciptated as small and homogenous distributed particles. Increasing the solidification velocity furthermore , i.e. by electron beam welding of the spray formed material will lead at least again to a microstructure with aluminum dendrites and a granular eutectic (Figure 1b). The coarse Silicon primary particles shown in that figure nucleated at already existing Silicon particles which did not dissolve completely during welding. It is worth noting that other alloy systems show similar behaviour. For example in Figure 2 an Al-Mg-Si alloy is given [4]. Clearly the chinese script morphology of the Mg2Si–Al-eutectic can be examined (Figure 2a). In hypereutectic alloys a irregular eutectic and cuboidal primary

103

Figure 1a: Al-35%Si (DISPAL S220) , spray formed. Silicon is completely precipitated as primary phase

Figure 1b: Electron beam weld seam in same material as figure 1a. Aluminum dendrites and fine eutectic is reintroduced to the microstructure

Mg2Si will occur (Figure 2b). Increasing solidification rate leads to an uniform microstructure with globular and fine primary Mg2Si-precipitates. This alloy system is of specific interest since the density of a 25 vol-% Mg2Si –alloy will be below 2,5 g/ccm.

Figure 2a: Al-8%Mg-3%Si, sand cast. Chinese script Mg2Si and Al-dendrites

Figure 2b: Al-9%Mg-5%Si, sand cast. Irregluar eutectic and primary Mg2Si precipitates

The Al-Mg-Si alloy shown in Figure 3 is spray formed. Although there is a certain size distribution among the Mg2Si particles the show only a weak tendency to coarsen during heat treatment which make this alloy applicable for high temperature service. A couple of alloys have been sprayed to “play around” with high volume contents of promising phases. In Figure 4 two of these experimental alloys are shown. The captions also indicate the Vicker’s hardness of the as sprayed material. Hardness as sprayed: 207 HV30.

3.2

Injection Tool

As mentioned before spray forming offeres the opportunity to inject almost every type of powder. The activities with Overspray powder reinjection led to a stable, high efficiency process. First attempts to use the same equipment for injection of ceramic particles – especially of abra-

104

Figure 3a: Al-16%Mg-8%Si, spray formed

Figure 4a: Al-30%Cu-10%Mg spray formed. Hardness as sprayed: 180 HV30

Figure 3b: Same as figure 3b, higher magnification

Figure 4b: Al-25%Cu-6%Mn spray formed

sive powders like silicon carbide – were not successful. During a public funded project an alternative method of powder transportation was elaborated. This so called “dense pack conveying” principle allows low velocities of the powder during transport into the spraying chamber. In the meantime several powders have been injected successfully up to 30 to 40 vol.-%. A interesting example is the injection of particles which will already be present in the alloy. For example, during the spray forming of an Al-35%Si-alloy additional Silicon powder was injected. At the end one will receive a multimodal distribution of Silicon particles, ex-situ and insitu formed. This method enables to influence the Si distribution when for wear resistance a small amount of coarse Silicon is required. The next step will be the injetion of nano scale powders. This should lead to supreme materials since really nanodispersion hardened alloys can be produced. Due to the insolubility and thermal stability of nano-Al2O3 in Aluminum alloys a superior hot strength and creep resistance can be implemented. There is a chance to replace the costly and time consuming mechanical alloying. Preliminary feasability studies have revealed that the injection of nanoscale powders is possible but needs further modification and optimisation of the existing equipment. A public funded project is set up within the WING programme of the Federal Ministry of Research and Technology to work out an efficient as well as an effective processing route for such nanoMMC’s.

105

Figure 5a: Al-35%Si (DISPAL S220) , Figure 5b: Same as figure 5a. with differspray formed. Addition Silicon powder was ent Si-size injected. injected

4

Summary

Spray forming is more than just another production of existing alloys. It offers unique opportunities to create supreme materials with outstanding and tailored properties. We just openend the door a little bit. Let us enter a space of nearly unlimited metallurgy.

5

Acknowledgements

The authors would like to thank the Federal Ministry of Research and Technology (BMBF), the Deutsche Forschungsgemeinschaft (DFG), Stiftung Industrieforschung and Arbeitsgemeinschaft industrieller Forschungsvereinigung (AIF) for funding research projects related to spray forming.

6 [1]

[2]

[3] [4]

References P. Krug, B. Commandeur, “Sprühkompaktieren von Aluminium-Hochleistungs-Legierungen – Pflicht und Kür“, Abschluss-Kolloquium des SFB327, Band 7, 2004, p.123–136. W.J. Boettinger, J.W. Cahn, S.R. Coriell, J.R. Manning, R.J. Schaefer, „ Application of Solidification Theory to Rapid Solidification Processing, Semi-Annual Technical Report, Springfield, Va. : NTIS, 1982 W.Kurz, D.J. Fisher, “Fundamentals in Solidification”, Transtech Publication, Switzerland, 1989 H. Hanemann, A. Schrader; Ternäre Legierungen des Aluminiums, Atlas Metallographicus Band II, Teil 2, 1952, Verlag Stahleisen m.b.H., Düsseldorf, S.128

106

A Method of VDC Hot Top Mould Design and Setting of Process Conditions I.F. Bainbridge1 & J.F. Grandfield2 Cooperative Research Centre for Cast Metals Manufacturing (CAST) 1 Division of Materials, The University of Queensland, Brisbane, Qld. 4072 Australia 2 CSIRO, Division of Manufacturing and Infrastructure Technology, Locked Bag 9, Preston, Victoria, 3072, Australia

1

Abstract

The design evolution of VDC hot top moulds has occurred largely by innovation and practical refinement. Similarly, conditions under which a particular mould design is operated are set and remain unchanged throughout the cast. During a cast process conditions within the mould may change such that control settings become inappropriate, or are at such a variance as to result in the production of scrap product. The paper presents an approach to mould design and the determination of process conditions based on an understanding of the critical parameters such as mould length, metal head and air pressure balance which control cast product quality. New knowledge of fundamental properties controlling molten metal meniscus behaviour within the mould is applied to the design of a 152mm hot top mould.

2

Introduction

Hot top moulds are used for the casting of extrusion billet. Several mould designs are available commercially and a number of in-house designs are in general use. Whilst hot top moulds have generally not been used for the casting of rolling ingot, design and operating principles now applied to rolling ingot systems are similar to those used for the hot top billet moulds. Hot top mould designs are generally innovative and practical, but they frequently incorporate features that are thought to be necessary or desirable, with many of these features evolving as attempts are made to differentiate systems competing commercially, or to improve cast product quality. The understanding of the complex interacting processes occurring within the mould to produce the cast product has not necessarily kept pace with the practical development of moulds and associated systems, including process control systems. With the advent of computer modelling, attempts are being made to simulate these complex processes and thereby refine existing mould designs and to generally improve the process as demanded by the constant competitive pressures for the production of a superior product at a lower cost. The work introduced in this paper proposes an approach to mould design and the setting of process conditions based on application of basic principles of mould heat flow and solidifying product physical properties.

107

3

Mould Design

The position of the liquid/solid metal surfaces as solidification commences within the mould is a critical factor to be considered in the mould design. The product surface structure within the mould is shown schematically in Figure 1. The important components determining the product surface quality have been identified as the meniscus, the molten metal head and the solid surface formed by the two cooling mechanisms, viz., heat extraction from the mould and heat extraction by the sub-mould cooling. The refractory overhang is also important, too short and the meniscus will not be stable and gas will bubble up through the melt. The key dimensions are (Figure 1): • A and B, the height and width of the meniscus; • C the molten metal head height; • D the shell length, i.e., the total length of the solid surface projecting into the mould from the point of impact of the sub-mould cooling; and • E the UCD. • F the refractory overhang • G the effective mould length.

Figure 1: Schematic of hot top mould, showing surfaces and important dimensions

The height A, and width B, of the meniscus may be determined by application of the model of Baker and Grandfield [1] (Figure 2) which solves the Laplace-Young equation and is dependent upon:

108 • • •

the surface tension of the molten metal J; the contact angles formed at the mould and hot top refractory surfaces, VM and VR; the pressure balance between the molten metal head height H and the gas pressure Pg.

Figure 2: Meniscus model from [1]

The height A is given by

A H  H 2 

2J  cosV M  cosV R UJ

(1)

In order to calculate the meniscus width B, one must solve equation (2) a first order ordinary differential equation for y(x) (the shape of the meniscus which together with the initial condition y(0) = 0 can be solved using a suitable numerical method such as Runge-Kutta; there is no exact integral for the function).

Gy Gx

§ ª UJ UJ 2 º · tan ¨ acos «cosV M +  H  A y  y »¸ J 2J ¼ ¹ ¬ ©

(2)

Stable meniscus sizes for three different alloys were calculated by the above model using molten metal surface tension data for various aluminium alloys available from the work of Bainbridge [2]. The approximate mould length required for a particular mould diameter may then be obtained by adding the values for the UCD, calculated from the method of Flood et al. [3] or a full 2D heat and fluid flow model (using casting speeds from standard commercial practice) to the meniscus height for a particular alloy (Table 1). A definitive mould length cannot be obtained unless accurate data is available for the mould heat flow, which in turn is dependent upon the mould gap thermal conductivity. Published data for a range of alloys and casting conditions is not presently available, but is part of an on-going research program at CAST.

109 Table 1: Calculated stable meniscus size and UCD for three different alloys, for a 152 mm diameter mould , a VR of 10° and a zero head height; plus critical mould dimensions as calculated from these two parameters (refractory plate overhang is calculated as B + 20%). Alloy

J Meniscus Size (mm) (Nm–1) A B

UCD (mm)

Effective Mould Length (mm)

1050 6063 7075

0.63 0.48 0.77

12.9 16.1 8.8

18.9 21.4 15.5

6.0 5.3 6.7

4.3 3.7 4.7

Refractory Plate Overhang (mm) 5.2 4.4 5.6

The data as derived above may then be applied to the critical dimensions of the mould. The effective mould length is the sum of the stable meniscus height and the shell length (approximated by the UCD) (A and D Figure 1), whilst the minimum refractory overhang dimension (F in Figure 1) is based on the stable meniscus width. The final mould dimensions may then be estimated by adding the effect of metal head height to the stable meniscus calculations. The above principles have been applied to the design of a mould for different alloys and operating conditions. A mould designed for the casting of one type of alloy is not commercially acceptable, hence a compromise in terms of effective mould length had to be chosen and then the critical operating conditions of casting speed and casting temperature (these affect the UCD) and metal head height (this affects A) chosen to obtain a meniscus that remained within the calculated stable size range for the particular alloy being cast. A fully instrumented version of the mould has been built and used to cast product on a pilot VDC casting unit. Results from the mould trials have shown excellent agreement with the model’s predictions. Billet quality has been shown to be totally predictable, e.g., Figure 4, showing the effect of a metal head height increase on cast surface. The smooth surface was cast within the mould design parameters, the rough surface resulted when the metal head was moved outside these parameters.

Figure 3: 6063 alloy billet cast in experimental mould

110

Figure 4: Effect of a variation in metal head height or casting speed on the effective mould length

4

Process Conditions

The models used to derive the basic dimensions of the mould include operating parameters, in particular metal head height, casting speed and molten metal temperature. A mould based on a design that incorporates these factors must obviously be operated in a manner that maintains these parameters within acceptable ranges for the mould design and alloy being cast. The mould design calculations presented above were initially based on known casting speeds, metal temperature and molten metal head height conditions used commercially. The effect of a variation in one or more of these process parameters on the critical mould dimensions is shown in Figure 4, and had to be taken into account in the final mould design and then in the manner in which the mould was operated. Note the significant effect that metal head height has on effective mould length. Present process controls generally rely on pre-set values for each of the critical operating parameters, with operator interception to change any particular value during a cast only occurring if a major problem arises. There is no real time measurement of mould conditions and the use of these conditions to control the process based on comparison of the measured values and model requirements.

5

Discussion

Existing mould designs have been derived over time, largely by empirical methods. The availability of data on critical properties of the cast product as it forms within the mould now permits mould design to be based on an approach that recognises factors critical to the formation of high

111 quality cast product. This modelling approach is still rudimentary considering the complexity of the processes occurring within the mould. The two models proposed take into account five variables for the calculation of the meniscus dimensions and eight variables for the calculation of UCD. If the latter calculation was modified to calculate the shell length, i.e., the UCD plus that solid formed due to mould cooling then at least two further variables would be added. Using the models it is possible to identify the variables that will have the most significant effect on the process and the limits to which each must be maintained to guarantee the casting of high quality product. Notwithstanding these limitations of the present model, the methodology is a significant departure from conventional mould design methods and recognises the conflicts and compromises that must be made to arrive at a commercial mould design capable of casting a range of alloys. Further, the application of the concept to the process conditions recognises that the mould was designed using specific parameters that normally form part of the operating envelope. It is suggested that this link is not commonly made in commercial operations at this time.

6

Conclusion

A method of designing a VDC casting mould based on molten metal meniscus size and UCD has been derived and the results of the initial testing of the concept suggest that the combination of the mould design and operating parameters linked to the design may result in cast product of more consistent and higher quality. The concept is a significant departure from the present empirical design methods with the potential to offer to the industry more reliable and cost saving casting systems. In addition, the concepts may be applied to the DC casting of other metals, e.g., magnesium.

7

Acknowledgement

The work was supported by the CRC for Cast Metals Manufacturing (CAST). CAST was established under, and is supported in part by the Australian Government’s Cooperative Research Centres Scheme.

8 [1]

[2]

[3]

References Baker, P.W. and J.F. Grandfield. The Role of Surface Tension Forces in Gas Pressurized VDC Casting. in Seventh Australian Asian Pacific Conference Aluminium Cast House Technology. 2001. Hobart, Tasmania, Australia: TMS, USA. p. 195–204 Bainbridge, I.:The Influence of Molten Metal Surface Properties on the Formation of Surface Defects on Vertical Direct Chill Cast Aluminium Alloy Products:Engineering, Physical Sciences and Architecture:The University of Queensland:2005 Flood, S.C., P.A. Davidson, and S. Rogers. A Scaling Analysis for the Heat Flow, Solidification & Convection in Continuous Casting of Aluminium. in Modelling of Casting, Welding and Advanced Solidification Processes VII. 1995: TMS, USA. p. 801–808

112

Continuous Casting of Non Ferrous Metal Micro Wrought Shapes J. Bast1, E. Bombach2 1

TU Bergakademie Freiberg Deutsche Solar AG Freiberg

2

1

Introduction

In the last years micro technology has developed into a key technology. An increasing need of micro-structured products is projected for the future. Information and communication technology, aircraft and aeronautics applications [1], as well as medical technology [2] and in particular the automotive industry [3] are regarded as growing markets for the use of micro-structured systems. Micro-components are not only used in mass produced parts, but also in products which are manufactured as prototypes or in small quantities. At present the mass and large-scale manufacture of micro-structured components is dominated by silicon etching techniques, lithography methods [4] and injection moulding. Since these methods are usually two-dimensional, it is complicated to produce a third dimension in such components. Furthermore, because of the high cost of equipment these methods are not suitable for small and medium quantity production of micro-structured components. That is why methods are gaining importance, that are based on conventional mechanical cutting technologies, e. g. drilling, milling, turning, grinding, micro eroding and laser machining [5]. Unfortunately, with the exception of the above mentioned injection moulding, casting methods play only a minor role for the production of micro-structured components, although they permit a wide geometric spectrum and allow the manufacture of different materials.

2

Continuous Casting

Conventional continuous casting has an interesting potential for near net shape production of wire and rods direct from the melt. Continuous casting methods are suitable for producing rods made of materials that are difficult to forge or roll and for producing products, where the casting process leads to significant cost reductions because of cross-section and dimensions of the products. Presently the smallest economically produced cross-section made by a conventional continuous casting process is above a diameter of 5 mm. With the conventional continuous casting process a metallic melt is poured into a mould. By cooling the mould walls the heat is removed from the melt, so that the metal solidifies inside the mould. The metal assumes the shape of the mould and can be pulled-out as rod material. In the mid-1980s at Chiba Institute of Technology in Japan a crystal growing process for the production of mono-crystals was developed [6]. This process makes it possible to produce very small cross-sections, with rod dimensions well below the dimensions of conventionally produced rods. In this process the mould is heated. The temperature of melt is held a little higher than the solidification temperature of the cast material. Heat is extracted from the cast product by means of a cooling device located nearby the mould

113 exit. The temperature of the mould is higher than the temperature of the rod [7], so that the heat dissipation occurs parallel to the cast direction and the nucleation of crystals on the mould surface is prevented [9]. The difference between conventional and micro-continuous casting is demonstrated in Figure 1.

Figure 1: Schematics of conventional and micro-continuous casting processes

3

Experimental Micro-Continuous Device with Horizontal Arrangement

Based on the OCC-method [6] an experimental horizontal micro-continuous device was developed in the Institute of Mechanical Engineering of the Technical University Bergakademie Freiberg [8]. See Figure 2. The metal is molten and held in the furnace. At one end of the furnace wall is an outlet pipe which also holds the mould. The pipe and the mould are heated. The mould opening corresponds to the cross-section of the rod to be cast. The height of the melt in the furnace is adjusted so, that the outlet is well filled with molten metal by metallostatic pressure. At the beginning of the casting process a starting dummy is attached directly to the melt nearby the mould exit. The dummy is cooled. The melt located in the mould loses enough heat, to solidify inside or nearby the end of the outlet while coupled to the dummy. The dummy is moved backwards by a pull-out device and the attached profile wire is pulled out of the mould. An experimental apparatus was designed and built for the production of micro-profiles. See Figure 3. The equipment consists of four main components: the melting system with mould heating, the mould, the cooling system and the pull-out system. Furthermore, the equipment is

Figure 2: Schematics of the experimental horizontal micro-continuous device

114

Figure 3: Basic design of the trial equipment for micro continuous casting of small profiled rods

Figure 4: Melting system

connected to a computer system for data collection and process control. A resistance heated muffle furnace was used for melting and holding. It contained a crucible made of steel. The pipe-like outlet holding the mould protrudes from the melting furnace. A resistance heated tube furnace surrounds the outlet and maintains its temperature. See Figure 4. The moulds are exchangeable and consist of steel or graphite. In the centre of the moulds is an opening with a cross-section equivalent to the geometry of the profiled rod. See Figure 5. The design and the surface characterization of the mould strongly influence the parameters of the casting process and on the quality of the rods. For the production of rods with a high quality the cooling system is very important. During the investigation three systems were tested: • Cooling by water flow • Cooling by water bath • Cooling by air. Because of the high specific heat of water, the turbulences when impacting the rod and the high heat transfer by cross-flow, flowing water is a very effective cooling system. Unfortunate-

115

Figure 5: Mould in the melting crucible

Figure 6: Cooling of the rod by air

ly, the direct impact of the water stream on the cast rod leads to vibration with resonance effects, which ruptured the rod and influenced the quality of the rod negatively. Because of the low distance between the cooling stream and the mould during the casting process cooling water was sprayed on the mould. This leads to sudden chilling of the mould and to a change of the cooling conditions of the system. In the investigation showed that air cooling was the best choice. See Figure 6.

4

Casting Results

The developed micro continuous casting device was used for the production of profiled rods of tin and aluminium with a cross-section below 1 mm and 900 mm length. The profiled rods had circular and more complicated cross-sections. During the investigation the following tin rods were produced: • Circular profile with a diameter of 1.0 mm, 0.9 mm, 0.5 mm and 0.3 mm with 900 mm length, • Star-like profile with 900 mm length, • Square profile with 400 mm length. Furthermore, circular profiles of aluminium and lead were produced with 400 mm length. Figure 7 displays the enlarged images of the profiled tin rods. For the examination a metallo-

Figure 7: Rods with circular (top) and star-like (bottom) cross-section

116

Figure 8: Outline of rod, geometry of the mould and micrograph of profiled cast rod

graphic specimen of the rods was embedded into resin and a micrograph was made. Figure 8 shows the outline of the cross-section, the geometry of the mould and the cross-section of the cast rod. The surface quality corresponds to the results which can be obtained with conventional manufacturing methods (micro-drilling, grinding, high-speed milling) or the die casting process. The surface roughness is 0.5 to 1.5 microns.

5

Summary

The continuous casting process permits near net shape production of rods directly from the melt and processing of a wide variety of materials. With a growing need of micro-structured components it is possible, that the micro-continuous casting process of small rods can widen the spectrum of existing processes. For the production of micro-structured rods a horizontal microcasting device was developed. With this equipment micro-structured profiles of tin and aluminium with dimension from 0.3 to 1.0 mm were produced. Furthermore, experimental and theoretical investigations about the analysis of process parameters and process control were carried out.

6

Acknowledgements

The financial support provided by the Deutsche Forschungsgemeinschaft Germany is gratefully acknowledged.

117

7 [1] [2] [3] [4] [5] [6] [7] [8] [9]

References J. Fahrenberg, Werkstattstechnik 2000, 90, 11/12, 484–486 I. Beltrami, et. al. In Tagungsband Aachen 10.-11. Juni 1999, 39– 420 T. Seubert, in Micromaterials for Automotive, Leipzig, 26. Juni 2003, 39 W. Ehrfeld, Feinwerktechnik, Mikrotechnik, Messtechnik, 1992, 100, 282–286 J. Hesselbach, et. al., Werkstattstechnik 2003, 93, 3, 119–128 A. Ohno, Metals, 1986, 38, 14–16 H. Soda, G. Motoyasu, A. McLean, A. Ohno, Advanced Materials & Processes, 1995, 4, 43–45 E. Bombach, Diss. TU Bergakademie Freiberg, April 2004 H. Soda, G. Motoyasu, A. McLean,S. D. Bagheri, D. Perovic, Cast Metals, 9, 1996, 1, S. 37–44

118

Influence of Quality of Water and Surface Roughness on Quenching Rate Jacek Król, Eckehard Specht Otto-von-Guericke-University, Magdeburg, Germany

1

Introduction

When measuring the heat transfer from a hot surface quenched by a spray, researchers typically use water that has been carefully distilled so as to remove any dissolved impurities. From experience in industrial practice it is known that the quality of the water has an influence on the quenching rate in continuous casting. The effect of additives on pool boiling heat transfer has received much more attention than their role in spray cooling. Pool boiling studies done by Najibi [1] showed that dissolved salts precipitate on the heater surface during boiling. Surfactants also have an important effect on pool boiling [2]: they promote bubble nucleation and foaming in the liquid and significantly increase heat transfer. Qiao and Chandra [4] found that dissolving a surfactant in the spray water significantly increased heat transfer. The surfactant reduced the liquid-solid contact angle and produced foaming in boiling droplets; both effects increased the solid area wetted and enhanced surface cooling. King et al. [5] observed the evaporation of salt solution droplets placed on a hot stainless steel plate. In opposite, they found that the dissolved salts reduced the vapor pressure of water and therefore decreased the droplet evaporation rate. Until now no studies are know which considered the influence of the quality of the water systematically. The roughness of the strand’s surface influences the Leidenfrost temperature and there wise the quenching rate, too. But there are no papers which quantitatively investigated the influence of the roughness. Both effects were experimentally researched using a relatively new method of measuring technique.

2

Experimental Set-Up

The measurement set-up, sketched in Fig. 1, was used to investigate atomized spray quenching. In Atomized Spray Quenching [6], the spraying water is atomized into fine droplets by compressed air and sprayed onto a hot surface. The drops partially evaporate and then move away from the superposed airflow. Thus, the vapor film is avoided as it is the case for other quenching techniques. The main component of the measurement set-up was a thin metallic sheet, electrically heated up far above the Leidenfrost temperature. This metal sheet was cooled from one side by the water spray. During heating up, the water spray was covered. When the stationary temperature was reached, the water spray has been started. On the black colored opposite side the transient surface temperature was measured using an infrared camera. Because of the small thickness (0.3 mm) and high conductivity the temperature on both sides are similar. From the given heat flow and the temperatures the local heat transfer coefficient especially in the transition range from film to nucleate boiling can be determined with a high accuracy. To correlate

119 the heat transfer with the water spray characteristics, the distribution of the drop velocity and the drop diameter of the spray were measured with a combination of 2D-Phase-Doppler-Anemometer and patternator. Internal mixing air blast atomizers were used for the water spray generation. With these nozzles, the water is mixed with compressed air inside the nozzle and ejected afterwards. The air used for atomization was supplied up to pressure of 0,5 MPa. The quality of water was regulated with different concentrations of salt and soap. With sand paper a defined roughness of the surface was adjusted.

Figure 1: Experimental set-up

Figure 2: The patternator for measuring of the impingement density

The impingement density was measured with the patternator sketched in Fig. 2. It consisted of several tubes with a diameter of 5 mm, arranged in line array parallel to the spray axis. The water amount Mw was collected by the tubes over a period of time t. The collected water was stored in small bottles. The mean impingement density was calculated from Eq. (1),

120

m s

4˜ Mw ʌ ˜ d t2 ˜ ǻt

(1)

where dt is the tube diameter. Therefore, a stagnation flow similar to the flow pattern used for quenching was measured. The heat flow caused by water spray qsp was computed from the energy balance by Eq. (2) at the investigated location, where dT/dt was measured qsp

cp ˜ U ˜ s

dT  qlos . dt

(2)

The heat loss by radiation, free convection and conduction in metal sheet was negligible. The heat transfer coefficient of the water spray Dsp was calculated from the temperature difference between the hot surface Ts and the water spray Tsp by Eq. (3)

qsp

3





D sp ˜ Ts  Tsp .

(3)

Experimental Results

The aim of the measurements of the atomized spray quenching was to find out the influence of the main parameters on heat transfer such as: diameter of the droplets, velocity of the droplets, impingement density, quality of water, surface roughness. Figure 3 depicts profiles of the impingement density, measured with the patternator in a plane 200 mm in front from the nozzle. The measurements were carried out with a constant water flow through the nozzle with different air pressure supplying the nozzle. It can be seen that the maximum of the impingement density is at the spray centre. If the water flow is constant, the maximum impingement density decreases with increasing air pressure.

Figure 3: Impingement density distribution in radial direction

121

Figure 4: Distribution of mean drop velocity and diameter

Figure 4 presents profiles of the mean volumetric diameter d30. The water flow through the nozzle is constant at 7 kg/h, and the air pressure is varied from 0.2 MPa to 0.4 MPa. The largest drops were measured in the centre of the spray. For a constant water flow through the nozzle with an increasing air pressure, the drop diameter decreases. To estimate the expansion of the spray, a qualitative profile of the impingement density is added to Fig. 4. The measured maximum velocity and mean volumetric diameter are about v = 30 m/s and d30 = 20 m, respectively in the centre of the spray. The heat transfer coefficient of atomized spray quenching is presented in Fig. 5 depending on the surface temperature for a high and a low impingement density. For both impingement densities, the heat transfer coefficient is independent from the surface temperature for the values above the Leidenfrost temperature of about 300 °C.

Figure 5: Heat transfer coefficient of atomized spray quenching vs surface temperature

The heat transfer coefficient depending on the impingement density is presented in Fig. 6 for water spray quenching and for atomized spray quenching. In atomized spray quenching the vapor film does not form like in water spray for surface temperatures above the Leidenfrost temperature. The impingement density exerts the highest influence on the heat transfer coefficient.

122

Figure 6: Heat transfer coefficient versus impingement density

For atomized spray quenching the heat transfer coefficients also increase with the air pressure. This occurs because of the changing droplet characteristics, mainly the change in the velocity. The heat transfer coefficient increases with the droplet velocity. The effect of droplet size on the heat transfer coefficient was not found to be significant in this investigation of atomized water spray. Atomized Spray Cooling leads to much higher heat transfer coefficients than water spray cooling. Figure 7 presents the mean heat transfer coefficient for atomized spray quenching depending on the surface temperature for different concentration of salt and soap. The experiments show that the amount of the dissolved salt and soap in distillated water decreases the heat transfer coefficient compared to the pure distillated water. The run of the temperature is presented in Fig. 8 for smooth and rough surface. With sand paper a defined roughness of the surface was adjusted. The roughness of the surface exerts high influence on the quenching rate. This occurs because of the larger droplet contact area for a

Figure 7: Heat transfer coefficient vs surface temperature

123

Figure 8: Run of temperature with time for smooth and rough surface

rough surface than for a smooth surface. For increasing impingement density the difference of the quenching rates between smooth and rough surface decreases.

4

Conclusions

The experiments demonstrate that the heat transfer coefficient is mainly determined by the impingement density. Atomized spray cooling has a much higher heat transfer coefficient than spray cooling. The amount of the dissolved salt and soap in distillated water decreases the heat transfer coefficient compared to the pure distillated water. Higher concentration of salt and soap increases a little the Leidenfrost temperature. Surface roughness does not influence the Leidenfrost temperature, but exerts high influence on quenching rate for the smallest impingement density.

5 [1] [2]

[3] [4] [5]

References Najibi, S.H., Muller-Steinhagen, H., Jamialahmadi, M., “Boiling and Non-Boiling Heat Transfer to Electrolyte Solutions,” Heat Transfer Engineering 1996, 17, 46–63 Hetsroni, G., Zakin, J.L., Lin, Z., “The Effect of Surfactants on Bubble Growth, Wall Thermal Patterns and Heat Transfer in Pool Boiling,” Int. J. of Heat and Mass Transfer 2001, 44, 485–497 Qiao, Y.M., Chandra, S., “Spray Cooling Enhancement by Addition of a Surfactant,” ASME J. of Heat Transfer 1998, 120, 92–98. King, M.D., Yang. J.C., 1997, “Evaporation of a Small Water Droplet Containing an Additive,” Proceedings of the 32nd National Heat Transfer Conference, Vol. 4, pp. 45–57 Puschmann, F., Specht, E., “Atomized Spray Quenching as an Alternative Quenching Method for Defined Adjustment of Heat Transfer,” Steel Research 2004, 75, 283–288

124

Electromagnetic Casting of Aluminum and Steel Billet Using Slit Mold J. Park1, M. Kim1, H. Jeong2, G. Kim2 1

Research Institute of Industrial Science and Technology, 2POSCO

1

Introduction

Many defects, such as an oscillation mark, a crack, etc., were formed in the surface of the billet into which the steel was cast through the continuous casting. Such surface defects affect seriously the productivity and workability in the continuous casting process 1–3). Surface defect causes any delay in the process, any labor loss and even any material loss. Steel industries, which have to reduce the used amount of fossil fuel, are striving to realize the hot direct rolling process without any process of reheating the billet as one of the energy saving processes. From this viewpoint, EMC technology is being studied. This technology has been developed for the moldless casting of any material with a light specific gravity and a good electric conductivity using the electromagnetic force in stead of the mold 4–5). However, the specific gravity of the steel is relatively high, and its thermal conductivity and electric conductivity are low. Besides, its casting speed is high. Therefore, it is thought that the moldless casting of the steel is almost impossible. This is why steel industries turn their attention to the study on application of EMC in the soft contact casting form using the mold. There are two EMC methods. One is to use the high frequency magnetic field of tens of kHz or higher6–11), and the other is to use the low frequency magnetic field of 60Hz to 200Hz12). According to the former method, the meniscus has a good stability but it has a difficult point that a special mold and a special power device are required. On the other hand, the latter method has a major challenge that the meniscus is unstable. This study has been conducted as a part of the study to improve the surface quality of the billet into which the steel was cast through the continuous casting by applying the EMC technology using the high frequency magnetic field to the steel. In this study, the effect of the mold shape on the billet surface shape was examined through the continuous casting experiment of the steel, and the effect of various factors on the surface quality of the billet was researched by observing the shape of the early solidified shell and measuring the meniscus shape and the mold flux consumption13,14). In the aluminum continuous casting, what has not yet been solved in aluminum continuous casting technology is a simultaneous improvement of surface quality of the cast and casting speed. In this research, an electromagnetic casting technology has been proposed that can increase casting speed as well as improve the surface and inner quality of the cast through application of electromagnetic casting technology and electromagnetic stirring technology by means of the slit mold being studied in electromagnetic casting of steel. The possibility of the proposed idea in this work has been confirmed through basic study.

125

2

Concept of EMC

Following is a brief explanation about the concept of soft contact EMC using the high frequency magnetic field. As shown in Figure1, in case the electric current is applied to the coil, the magnetic field is induced in the mold and the electric current gets to be induced by the magnetic field. This electric current generates the magnetic field and an electric current in the molten metal inside the mold through the mold segmented by slits. This induced current not only heats the molten metal but also generates the electromagnetic force in the molten metal by acting on the magnetic field. This technology using the Joule heating and the Lorentz force is just the EMC principle using the high frequency power. The electromagnetic force enlarges the meniscus curvature of the molten metal in contact with the mold, thus, improves the inflow of the mold flux and at the same time reduces the contact pressure between the shell and the mold. The meniscus is heated by the Joule heating and a thin early solidified shell gets to be formed in the lower part of the meniscus, which inhibits any hook, the root of the oscillation mark from taking place.

Figure 1: The concept of EMC using slit mold

3

EMC of Steel Billet

3.1

Lab Scale Experiment

The billet surface morphology of face are shown in Figure 2, respectively. It can be known from Figure 2 that in case the electromagnetic field was not applied, the OSM was clearly formed. It can be also known that in case (b) the coil current 670 A was applied, the face surface shape was much improved, but the OSM still remained in the corner. Further, it can be known that in case (c) the coil current 830 A was applied, both the surface shape of the face and the surface shape of the corner were greatly improved so that the OSM was formed about 0.15 mm deep or below. It was also identified that in case (d) the coil current 1100 A was applied, the respective

126 surface shapes of the face, the corner and the off-corner were all greatly improved, but grooves took place along slits in the casting direction of the billet. It is thought that the electromagnetic field acted more intensely than under any proper condition and a great magnetic pressure acted on the molten steel in the slit parts where the magnetic flux density was relatively high, and as a result, the molten steel was pushed into the mold, and thereby these grooves were formed. This mark is about 1mm deep, being deeper than the OSM in case the electromagnetic field was not applied. From this result, it can be known that any optimal electromagnetic condition is required for improvement of the billet surface.

Figure 2: Surface morphology of the billet

Figure 3: Early solidified shell of steel Billet

Since the forming position and the shape of the early solidified shell have a great effect on the surface quality of the billet, it is very important to identify the effect of the electromagnetic field on the shape of the early solidified shell. 100 g of the FeS powder was wrapped in the Al thin sheet, and it was added directly to the molten steel in the mold at the terminal stage of the casting. Figure 3 shows the shape of the early solidified shell respectively in case the electromagnetic field was not applied and in case the coil current 830 A, which was thought to be the optimal condition in this study, was applied. The black part in the photograph is the part where sulfur was intermixed, and it is also the solidified part after FeS was added. It can be known that in case (a) the electromagnetic field was not applied, the early solidified shell was thick and uneven, having a hook, the root of the OSM formed, and further that after the solidified shell was formed, the molten steel containing sulfur overflowed so that it existed even in the outer face of the solidified shell. On the other hand, in case the coil current 830 A was applied, the early solidified shell was thin and even, having no hook formed.

127 3.2

Conventional Scale Experiment

The mold plate was the copper alloy identical to the material used in POSCO’s billet casting operation. Its length was 800 mm, the size of casting section was 163u163 mm and an inside corner radius of 8 mm.

Figure 4: Mold powder consumption versus casting conditions

Figure 5: Typical appearance of billet strands

The amount of the mold powder consumption of the EMC operation is shown in Figure 4 for various casting conditions compared to the conventional casting operation. In the EMC operation, it was in the range of 110–130 % of the conventional casting. This can be explained as the work of the magnetic force in such a way that the magnetic force made the meniscus into a curved shape, and thus the mold flux would easily flow into the gap between the strand and the mold wall. Joule heat also helped in melting the powder and in keeping the molten flux at relatively high temperature compared to the conventional casting. Particularly, it is noteworthy that the mold powder consumption in case of EMC without mold oscillation was no less than that of the conventional casting. Figure 5 shows typical billet surfaces for various coil currents. In case that the coil current was 0 A, i.e. the conventional casting, billets had the OSMs of its nominal depth in the range of 0.45 r 0.15 mm. When the coil current was 500 A, the OSM disappeared a lot to its nominal depth in the range of 0.20 r 0.05 mm except the corner region. When the coil current was 1000 A, the OSM disappeared everywhere to its nominal depth in the range of 0.10r0.04 mm. When the coil current increased to 1200 A, the OSM appeared again in the form of wave mark although its nominal depth was not severe in the range of 0.18r 0.05 mm.

128

4

EMC of Aluminum Billet

In the case of general moldless EMC, there is a limit to increasing of the casting speed due to the unstable molten metal surface. Since the depth of the molten metal pool is also low due to the low casting speed, it is difficult to stirring the molten metal to make fine grains. But in the case of EMC technology proposed in this study, it is shown that casting speed can be increased since molten metal surface and mold contact through slit mold, and then the molten metal is maintained stably. Fine grain can be produced through stirring of the molten metal by installation of an electromagnetic stirring (EMS) device at the below of the EMC coil, which is due to the increased depth of the molten metal pool.

Figure 6: EMC mold system

4.1

Experimentals and Method

A schematic view for experimental devices are shown in Figure 6. The mold is a cylindrical shape with an inside diameter of 100 mm and a length of 200 mm. The slit is machined in some part of mold for longitudinal. After about 40 kg of 2024 Al alloy and A356 Al alloy were melted, respectively, and casting was made through variation of casting speed, EMC current, EMS frequency, and current. The effect of various parameters on the quality of the billet was reviewed. Moreover, the depth of the molten metal pool within the mold was measured with a 2 mm stainless steel rod during casting. The shape of the solidified pool was observed through injection of molten zinc into the molten metal within the mold at the final stage of casting. The specimen after casting was taken and microstructure was observed.

4.2

Results and Discussion

In Figure 7, a surface morphology of the 2024 billet produced by general casting is shown. Casting speed was elevated to 0.25 m/min and breakout occurred at 0.3 m/min. It can be seen

129

Figure 7: Surface appearance of the 2024 Al billet without EM

Figure 8: Surface appearance of the 2024 Al billet at EMC under coil current 1000 A

that a severe ripple mark was formed on the surface of the billet and the mark became weaker as casting speed increased. The surface morphology of the billet produced by the EMC at a coil current of 1000 A is shown in Figure 8. The ripple mark is disappeared. Moreover, casting speed was also increased up to 0.45m/min and breakout occurred at 0.5 m/min.

5

Conclusion

5.1

At the EMC of Steel Billet

The depth of the oscillation mark was decreased from about 0.6mm to 0.1 mm or below and any appropriate electromagnetic condition was required for obtaining the optimal surface shape. Electromagnetic continuous casting apparatus was suited on a billet caster of POSCO works, and commercial scale casting test was successfully performed. The major observations with the EMC operation include that the oscillation mark was improved a lot to its nominal depth of less than 0.1 mm]

5.2

At the EMC of Aluminum Billet

Casting speed as well as surface quality of the billet can be improved with electromagnetic continuous casting technology.

130 It was confirmed that surface quality can be improved and internal structure can be globular through the simultaneous application of electromagnetic casting technology and electromagnetic stirring technology.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

References S. Kumar, I.V. Samarasekera and J.K. Brimacombe: ISS Trans., 1997, June, 53 K. Kawakami, T. Kitagawa and Y. Hiratani: Tetsu-to-Hagane, 67(1981), 1190 K. Kawakami : Tetsu-to-Hagane, 74(1988), 1204 D. C. Prasso, J. W. Evans and I. J. Wilson: Metall. Mater. Trans. B, 26B(1995), 1243 Ch. Vives : Met. Tran. B, 16B(1985), 377 I. Sumi, K. Sassa and S. Asai: Tetsu-to-Hagane, 78(1992), 447 T. Li, S. Nagaya, K. Sassa and S. Asai: Metall. Mater. Trans. B, 26B(1995), 353 J. P. Park, H. T. Jeong, D. J. Sim and H. Y. Kim: Korean Ins. of Met. & mat., 36(1998), 1598 S. Itoyama, H. Tozawa, T. Mochida and K. Kurokawa : ISIJ Int., 38(1998), 461 H. Nakata, M. Kokita and K. Ebina : Tetsu-to-Hagane, 80(1994), 711 S. Furuhashi, M. Yoshida and T. Tanaka : Tetsu-to-Hagane, 84 (1998), 625 T.T oh, E. Takeuchi, M. Hojo, H. Kawai and S. Matsumura : ISIJ Int., 37(1997), 1112 H. Kim, J. Park, H. Jeong and J. Kim, ISIJ Int., Vol.42 (2002), No.2, p171–177 J. Park, H. Jeong, H. Kim and J. Kim, ISIJ Int., Vol.42 (2002), No.4, p385–391

131

Aluminium Alloy Strip Casting Using an Unequal Diameter Twin Roll Caster T. Haga1, H. Watari2, S. Kumai3 1

Osaka Institute of technology, 5-16-1, Omiya, Asahiku, Osaka city, 535-8585, Japan Oyama national collage of Technology, 771, Nakakuki, Oyama city, Tochigi, 323-0806, Japan 3 Tokyo Institute of Technology, 4259, Nagatuda, Yokohama city, Kanagawa, 226-8502, Japan 2

1

Abstract

An unequal diameter twin roll caster with a long solidification length was devised to cast aluminium alloy strip with a thickness of about 5 mm at speeds higher than 20 m/min. The characteristics of the unequal diameter twin roll caster with a long solidification length are as follows. The diameter of the lower roll is four times larger than that of the upper roll. The solidification length is long. The casting speed is high. Using the unequal diameter twin roll caster, a 4.5-mm thickness of 6111 strip could be cast at a speed of 30 m/min. Low superheat casting and semisolid casting was adopted to the unequal diameter twin roll caster with a long solidification length. The microstructure of the as-cast strip was equiaxed and spherical, not columnar. The mechanical properties of the strip rolled from roll-cast strip were almost as same as that of the strip made from cast ingot.

2

Introduction

In automobiles, decreasing weight is one of most important problems to be solved. Aluminium alloy parts are adopted instead of steel parts to reduce the weight. However, aluminium alloy parts are more expensive than steel parts. Many parts in automobiles are made from thin plate. It is desirable to economically produce aluminium alloy thin strip, and roll casting has this capability. Aluminium alloy strip can be cast directly from molten metal by a twin roll caster. The twin roll caster can eliminate the need for many processes: for example, grinding of the ingot surface, hot rolling and some degree of cold rolling. This shows that much energy and operational costs can be saved. Moreover, the need for many machines and space for machines can be reduced. The cost of installation can be saved, too. Thus, the twin roll caster has the advantage of cost saving. The twin roll caster also has the advantage of rapid solidification. The microstructure becomes fine by rapid solidification, and the mechanical properties can be improved. In summary, the twin roll caster has many advantages. However, the twin roll caster also has disadvantages: slow casting speed, limitation of usable alloys and center segregation. These disadvantages must be addressed. In the present study, an unequal diameter twin roll caster with long solidification length was devised in order to reduce the disadvantages of the roll casting of aluminium alloys. The unequal diameter twin roll caster with a long solidification length is different from the conventional twin roll caster for aluminium alloys and the conventional unequal twin roll caster. In the present study, 6111 aluminium alloy was cast into strip using an unequal

132 diameter twin roll caster with a long solidification length, and the mechanical properties were examined.

3

Experimental Conditions

The unequal diameter twin roll caster with long solidification length (UDTRCLS) used for the experiment is shown in Figure 1. Experimental conditions of roll casting by the UDTRCLS are shown in Table 1. 6 kg of 6111 aluminium was melted in the crucible by an electric furnace in air. The melt in the crucible was transferred to the UDTRCLS. The melt was poured on the lower roll through the cooling slope. The roll was rotated at desired speed (30 m/min) when the melt was poured. The special operation was not carried out at the start of the casting. The lubricant was not used on the roll surface. The melt head was maintained at 30 mm at the top of the roll. The cooling slope was used in order to decrease the temperature of the melt. The cooling slope was water cooled, and the surface was coated by BN. Casting load was 0.18 kN per unit width. This load was very small, because the hot rolling was not desired. However, this load was enough to make heat transfer between the strip and the roll better. The cast strip by the UDTRCLS was homogenized and cold rolled to 2 mm. Intermediate annealing was applied to the cold rolled strip, and then the strip was cold rolled to a thickness of 0.5 mm. T6 heat treatment was applied to the 0.5-mm thickness of cold rolled strip. Tension and180 degrees bending tests were used to examine the mechanical properties.

Figure 1: Photograph of unequal diameter twin roll caster with long solidification length in casting

133 Table 1: Experimental conditions Upper roll Lower roll Roll speed Solidification length Load of roll Specimen Melt temperature Heat treatment : T4,T6 Homogenization Intermediate annealing Cooling slope

4

Material: mild steel, Diameter: 250mm,width: 100 mm water cooling, non-lubricant Material: mild steel, Diameter: 1000mm, width: 100 mm water cooling, non-lubricant 30 m/min Upper roll: 60 mm, Lower roll: 300 mm 0.18 kN/mm (per unit width) 6111 (6 kg) 655 °C Solution: 530 °C – 4 h, Water quenching Aging: 160 °C – 6 h 540 °C – 4 h 530 °C – 2 h Material: mild steel, Length: 300 mm, Width: 100 mm Water cooling

Result and Discussion

The 6111 aluminium alloy strip could be cast continuously at a speed of 30 m/min using the UDTRCLS. The thickness of the 6111 strip was 4.5 mm. In the CTRCA, a 7-mm thickness of strip was cast at speeds slower than 5 m/min. The productivity of the UDTRCLS is three times superior to that of the CTRCA, and so the problem of low productivity in the CTRCA is solved by the UDTRCLS. The UDTRCLS strip was cast from the semisolid slurry, which contained a solid fraction of about 5 %. The semisolid slurry did not solidify around the nozzle and the dam plates, and clogging did not occur. Figure 2 shows the surfaces of the as-cast UDTRCLS strip. The upper surface of the as-cast strip is different from the lower surface. The lower surface is rich in metallic luster, and the upper surface is poor in metallic luster. When solidification occurred, the solid fraction of the semisolid slurry in the upper side of the strip was larger than the solid fraction in the lower side. Therefore, the upper side was not the same as the lower side. However, the upper side became the same after the cold rolling.

Figure 2: Surface of 6111 as cast strip

Figure 3 shows the microstructure of a cross section of the 6111 as-cast UDTRCLS strip. Figure 4 is an enlarged view of Figure 3. The microstructure of the cross section of the strip cast by the CTRCA is usually a columnar structure and is symmetric in the thickness direction. The

134 thicknesses of the solidification layers cast by the two rolls are the same. As confirmed by Figures 3 and 4, the microstructure of the strip cast by the UDTRCLS is different from the microstructure of the strip cast by the CTRCA. The microstructure is equiaxed, not columnar. The microstructure around the interface of the solidification layers is a spherical structure. This is the effect of the semisolid casting. The thicknesses of the solidification layer cast by the upper and lower roll are not the same. The solidification layer cast by the lower roll is thicker than the solidification layer cast by the upper roll, as shown in Figure 3. This effect is due to the difference of the solidification length. The solidification length of the lower roll is five times longer

Figure 3: Cross section of 6111 as cast strip

Figure 4: Enlarged view of cross section of 6111 as cast strip

135 than that of the upper roll. The interface between the solidification layers is near the upper surface. This reduces the segregation. The material around the interface toward the upper roll solidified rapidly because the thickness of the upper solidification layer is thin. This center segregation remains an important problem to be solved. This problem could be solved or the degree of segregation could be decreased. The microstructure of the cross section of the 0.5-mm-thickness strip cast by the UDTRCLS after homogenization, cold rolling and T6 heat treatment is shown in Figure 5. The nonuniformity of the microstructure is somewhat improved. The roll-cast strip is usually used after some type of treatment, such as homogenization and cold rolling. Therefore, the nonuniformity of the microstructure in the thickness direction of the as-cast strip is not a defect, if the nonuniformity can be improved by heat treatment and plastic deformation.

Figure 5: Cross section of 6111 strip of T6 condition after some treatments

Figure 6: Outer surface and cross section of 180 degrees bent strip. Thickness was 1mm and heat treatment was T4.

136 The mechanical properties of the strip cast by UDTRCLS after T6 heat treatment are shown in Table 1. The mechanical properties of the roll-cast strip are almost the same as that of the strip made from ingot. It is said that the mechanical properties of the roll-cast strip are inferior to that of the strip made from ingot. This disadvantage is lessened by the UDTRCLS. Figure 6 shows the result of the 180-degree bending test on the UDTRCLS specimen. There was no crack at the outer surface. This shows that the roll-cast 6111 strip has enough ductility for hem forming. Hem forming is an important forming application used in the panels of automobiles.

5

Conclusions

An unequal diameter twin roll caster with a long solidification length was designed and assembled. Experimental casting was performed to investigate the characteristics of this twin roll caster and the strip cast from semisolid slurry. This caster can cast a 6111 aluminium strip of 4.5mm thickness at a speed of 30 m/min. This shows that the casting speed can be increased without decreasing the thickness of the strip. A tensile test and 180-degree bending test were performed. The strip showed good mechanical properties.

6 [1] [2] [3]

References A.I.E.Nussbaum, FATA Hunter SpeedCaster Inauguration, LIGHT METAL AGE, Oct. (1997)34–38 N.Toyama, H.Aho, H.Arai, H.Yoshimura, Direct casting of stainless sheet by unequaldiametered twi roll method, Tetsu-to-Hagane, Vol.71(1986)A245-A248 D.B.Love, J.D.Nauman, Controlling the physical and mechanical properties of cast stainless steel band, TMS Proc. of an international Symposium on casting of Near Net Shape Products, Hawaii, 1988, 597–611

137

Fabrication of High Purity Copper Rod with Unidirectional Solidification Structure by Continuous Casting Using Cooled Mold Hoon Cho, Duck-young Hwang, Han-shin Choi, Shae K. Kim, Hyung-ho Jo. Korea Institute of Industrial Technology; 994-32 Dongchun-dong, Yeonsu-gu; Incheon, 406-130, Korea

1

Abstract

It is generally known that OCC (Ohno Continuous casting process)using heated mold has to be introduced to produce cast rod representing single crystal and uni-directional solidification morphology. However, in the present study, it is expected that unidirectionally solidified high-purity copper rod could be fabricated by optimization of withdrawal speed in typical continuous casting process using cooled mold. This paper discusses the production of small cross section high-purity copper rod by the process of continuous casting, and the effect of varying the parameters of withdrawal speed and casting temperature on the cast grain morphology of the rod and its drawing characteristics to fine wire. The heat transfer within the graphite mold is also discussed with respect to the withdrawal speed and casting temperature.

2

Introduction

Recently, the application of copper as a substitute of gold bonding wire has been investigated because formation of brittle intermetallic compound such as aluminide between aluminium substrate and bonding wire can be suppressed. Especially, copper being cheaper than gold and with higher conductivity and better stiffness, is a viable, cost-effective alternative[1]. In order to manufacture copper ultra fine wire for bonding wire in integrated circuit package, continuous casting process, which can produce high purity copper rod with small cross section, has to be optimized to produce cast rod without internal defects and to control microstructure orientation which can prevent wirebreaks in wiredrawing process[2]. The paper presented here is mainly aimed at investigation of influence of varying parameters of pulling speed, superheat and rod diameter on grain morphology of the casting rod and on its drawing characteristics to ultra fine wire.

3

Experimental Procedures

In the vacuum vertical continuous casting system, the graphite mold (nozzle) was attached to bottom of melting furnace and was designed to cast rod in the range of 7 to 11 mm in diameter. Especially, the top of the mold was heated by copper melts and the bottom of the mold is cooled by cooling jacket. Therefore, it is very important to understand heat transfer within the graphite mold during continuous casting process. Several thermo-couples were inserted to the mold with

138 different depth to measure temperature profile of the mold and to forecast solidification front position in the mold. In order to prevent the contamination of copper cast rod and to control heat transfer within the mold, high-purity copper rod with high conductivity was used for starting bar during continuous casting process. Table 1: Chemical composition of high purity copper used in this study (unit : ppm) Element 5N-Cu Cast-rod

Ag 0.1 0.04

Al 0.1 0.08

Cr 0.04 0.04

Fe 0.5 0.5

Mn 0.002 0.002

Ni 0.2 0.018

Pb 0.2 0.11

S 2 0.9

Si 0.5 0.5

Microstructure evolution and mechanical properties of high-purity copper casting rod were performed with respect to casting temperature, withdrawal speed and cast rod diameter. In order to investigate the variation of micro-orientation with different process parameters in continuous casting process, EBSD(Electron Back Scattering Diffraction, Oxford INCA Crystal) was used.

4

Result and Discussion

4.1

The Influence of Withdrawal Speed on Unidirection Solidification

The present work on casting 9mm diameter high-purity copper rod has shown there is an optimum withdrawal speed around 30 to 50 mm/min, which leads to small columnar grains oriented

Figure 1: Longitudinal cross sections of the Cu rod with different withdrwal speed(diameter : 9 mm, casting temp. 1200 °C)

139 at an angle to the axis of the cast bar. Figure 1 shows longitudinal cross sections of the Cu rod with different withdrawal speed when rod diameter is 9mm and casting temp. is 1200. With a withdrawal speed of up to 30 mm/min, macrostructure of the cast rod represents columnar grains, which is solidified and grown in parallel direction with continuous casting direction. In contrast, as the withdrawal speed increases, the angle between the direction of casting and the predominant crystallization direction increases[3]. Figure 2 shows pole figure of the copper rod with respect to different withdrawal speed. As shown Figure 2 (a), unidirectionally solidified copper rod with (100) direction produced by a control of solidification front position within mold through continuous casting process using a cooled mold. As the withdrawal speed increases to 100mm/min, the misorientation angle of grains increases and solidification occurs in several directions such as (100) and (110). The solid-liquid interface(Solidification front) position was measured by making a temperature profile within the mold.

Figure 2: Pole figure of the Cu rod with different pulling speed(diameter : 9mm, casting temp. 1200 °C)

Figure 3 shows schematic illustration of solidification front and crystal growth direction with different withdrawal speed(V, mm/min) when rod diameter is 9mm and casting temp. is 1200. As withdrawal speed increases, distance between top of the mold and liquid/solid inter-

Figure 3: Schematic illustration of solidification front and crystal growth direction with different withdrawal speed(V, mm/min) when rod diameter is 9 mm and casting temp. is 1200 °C

140 face(solidification front) becomes far. As shown in Figure 3 (a), the solidification front could enter the location where the mold temperature is similar with copper melts when withdrawal speed is 30mm/min. Thus, heat transfer mainly occurs according to copper rod used for starting bar. It can be mentioned that unidirectionally solidified grain result from heat transfer direction parallel with casting direction. In contrast, the solidification front could enter the location where the mold temperature is lower than copper solidus temperature when withdrawal speed is 100 mm/min. Thus heat transfer also may occurs according to perpendicular to casting direction. The heat transfer induces solidification of crystal with several growth direction such as (100) and (110) as shown in Figure 2(b).

4.2

The Influence of Casting Temperature on Unidirection Solidification

In order to investigate the effect of continuous casting temperature on unidirection solidification of Cu cast rod, the continuous casting with different casting temperature varied 1150 °C and 1200 °C was carried out. The macrostructure showing unidirection solidification is illustrated in Figure 4 and the pole figure of Cu cast rod uis shown in Figure 5.

Figure 4: Longitudinal cross sections of the Cu rod with different withdrwal speed(V) and casting temperature (T)

As shown in Figure 4 and Figure 5, Cu cast rod was solidified unidirectionally when withdrawal speed is set to be up to 30 mm/min even though the casting temperature was varied to 1150 and 1250 °C. Consequently, it is very important to control and predict the position of solidification front within continuous casting mold in order to produce Cu cast rod with unidirection solidification

141

Figure 5: EBSD measurement result of the Cu rod with different withdrawl speed (Q) and casting temperature (T)

structure. Among the continuous casting process parameters including withdrawal speed, casting temperature and cast rod diameter, it can be mentioned that withdrawal speed during continuous casting process is dominant process parameter to control the position of solidification front within the mold.

5

Conclusions

It is expected that unidirectionally solidified high-purity copper rod could be fabricated by optimization of withdrawal speed in typical continuous casting process using cooled mold. The unidirectionally solidified grain represents growth direction in (100). As the withdrawal speed increases to 100mm/min, the misorientation angle of grains increases and solidification occurs in several directions such as (110) and (100). It is very important to control and predict the position of solidification front within continuou casting mold in order to produce Cu cast rod with unidirection solidification structure. Among the continuous casting process parameters including withdrawal speed, casting temperature and cast rod diameter, it can be mentioned that withdrawal speed during continuous casting process is dominat process parameter to control the position of solidification front within the mold.

142

6 [1] [2]

[3]

References Harman G., Wire bonding in microelectronics: materials, processes, reliability, and yield.(New York : McGraw-Hill, 2nd ed., 1997), 1–11 Wilson R. et al., “Continuous casting of high purity small diameter copper rod”(Paper presented at copper ’90, refining. Fabrication, markets, Vaesteras, Sweden, 1-3 Oct. 1990) 238–244 Ohno A. et al., “ Studies pertaining to the position of the solidification front in horizontal Ohno continuous casting system”( Paper presented at AFC-8, Bangkok, Thailand, 17-20 Oct. 2003) 689–699

143

High Speed Roll Casting of Al Alloy and Mg Alloy Strips T. Haga1, H. Watari2, S. Kumai3 1

Osaka Institute of technology, 5-16-1, Omiya, Asahiku, Osaka city, 535-8585, Japan Oyama national collage of Technology, 771, Nakakuki, Oyama city, Tochigi, 323-0806, Japan 3 Tokyo Institute of Technology, 4259, Nagatuda, Yokohama city, Kanagawa, 226-8502, Japan 2

1

Abstract

A high speed twin roll caster focused on aluminum alloy and magnesium alloy was designed and assembled. Vertical type was adopted from the point of easiness of pouring of molten metal in the roll bite. Improvement of increase of cooling rate and casting speed was attained in this caster. A356 for casting was cast to investigate the characteristics of this roll caster. AZ31 magnesium alloy for forming, AM60 and AZ91 magnesium alloy for casting were cast to investigate the roll-castability of magnesium alloy. Aluminum and magnesium alloy strip thinner than 4 mm was cast at speeds from 60 m/min to 150 m/min. Use of copper roll, non-use of parting materials and reducing of the thickness of the strip were contributed to improve the cooling rate and casting speed of the strip. The microstructure of the strip cast by the high speed twin roll caster was usually duplex structure. Center of the thickness was equiaxed or spherical structure, and surface sides were short columnar or equiaxed structure. Ununiformity of the microstructure at thickness direction could be improved after cold rolling and heat treatment. This caster was useful for aluminum alloy which freezing zone was wide. Strip of A356 showed good ductility. Annealed A356 strip was not broken at 180 degrees bending test. Cold deep drawing operated to T4-strip. Warm deep drawing of AM60 and AZ91 strip could be operated, and LDR was greater than 2.0.

2

Introduction

Aluminum alloy for casting has good castability. Therefore, the roll casting of the aluminum alloy for casting is easy. However, it is said ductility of the aluminum alloy for casting is poor. Therefore, aluminum alloy for casting is not suitable for the forming. It is possible that improvement of the poor ductility of the aluminum alloy for casting by the rapid solidification using the twin roll caster. If the aluminum alloy for casting can be used for forming, the variety of the aluminum alloy for recycle can be reduced. Improvement of poor ductility of aluminum alloy for casting was investigated. Properties of the cast strip were investigated by metalography, tension test, bending test and deep drawing. Magnesium sheet is normally made from slab by hot rolling. It is difficult to do cold rolling of magnesium alloy. The hot rolling is essential for magnesium sheet. Therefore, the slab or the strip must be heated up to hot condition at every rolling pass. This makes the sheet of magnesium very expensive. If the thin strip was cast directly from molten metal, number of heating and rolling could be reduced. As the result, the cost of the magnesium alloy sheet could be reduced. The twin roll caster is suitable for the magnesium alloy, as the strip can be cast directly from molten metal. The high speed twin roll caster is better to make the price lower. AZ31 is normal-

144 ly used for sheet metal forming. AZ61 and AM60 are better than AZ31 from the point of the strength. However, it is difficult to make sheet of AZ61 and AM60 as the hot rolling of these magnesium alloys is very difficult. Therefore, roll casting of AZ61 and AM60 is very useful. In the magnesium alloy, castability at high speed roll casting was investigated. Properties of cast magnesium alloy strips were showed by the hot rolling and warm deep drawing. The hot rolling was operated on as-cast strips. Warm deep drawing was tested on the hot rolled strip.

3

Experimental Conditions

A high speed twin roll caster (HSTRC) is shown in Figure 1 and experimental condition is shown at table 1. Magnesium alloy was cast in the air without inert gas.

Figure 1: Schematic illustration of a vertical type high speed twin roll caster (HSTRC)

Table 1: Experimental conditions Roll material speed aluminum alloy Magunesium alloy superheat Cooling slope separating force Solidification length melt head

4

Material: copper, Size: diameter 300 mm, width 100 mm Lubricant: non-use, Cooling: water 60, 90, 150,180 m/min A356 AZ31, AM60, AZ91 15 °C Material: mild steel, Size: length 300 mm, width 100 mm, Inclination angle: 60 degrees, Coating: BN, Cooling: water 0.14 kN/mm 100 mm 100 mm

Roll Casting of Aluminum Alloy

Figure 2 show roll casting in operation and as cast A356. A356 could be cast to the strip continuously at speeds from 60 m/min to 150 m/min. The strip was thinner than 4.0 mm. The HSTRC was able to cast at speed 10 times higher than that of the conventional twin roll caster

145

Figure 2: Experiment of roll casting and as cast strip

for aluminum alloy (CTRCA). The strip did not stick to the roll without the lubricant. The strip cast by HSTRC was half thickness of the strip cast by the CTRCA. The strip cast by the HSTRC was thinner than that cast by the CTRCA without the hot rolling. The microstructure of the strip cast by the HSTRC is different from that of the strip cast by the CTRCA. The surface and thickness distribution were improved by the rolling. 20 % reduction was enough to improve the surface and the thickness distribution. The cause of this defect is oscillation of the meniscus at the tip of the nozzle. This defect could be improved by cold rolling. The surface of the strip cast by the CTRCA is same at both sides. The microstructure of most of the cross section was equiaxed structure or spherical structure. This is the effect of semisolid casting and rapid solidification. Microstructures of cross section of A356 strip at as cast and T6 condition are shown in Figure 3. The microstructure of as cast

Figure 3: Microstructure of cross section of A356 strip

Figure 4: Surface of A356 strip after 180 degrees bending. treatment : homogenization, cold rolling, annealing and 180 degrees bending

146

Figure 5: Cold deep drawing of T4-A356

strip was not uniform at thickness direction. This ununiformity of the microstructure was improved by cold rolling and annealing. The microstructure of T6 strip was almost uniform. Eutectic Si of T6 strip was spherical and very fine. Annealed 0.5mm thick strip could be bent at 180 degrees without crack at outer surface as shown in Figure 4. Figure 5 shows the result of cold deep drawing of T4 heat treated A356 sheet. The cold deep drawing could be operated to T4-A356 of 1.0 mm thick. These results show roll-cast A356 had good ductility.

5

Roll Casting of Magnesium Alloy

High speed roll casting of magnesium alloys was tried. The casting was operated in the air without sealed gas.The magnesium strip could be cast at the speeds up to 150 m/min. The castability became better as the Al content became greater. The castability of AZ91 was better than AZ31. The relationship between the roll speed and the strip thickness was as same as that of the aluminium alloy. The strip thickness became thinner as the roll speed became higher.The strip, which was thinner than 2.0 mm, could be cast by the high speed twin roll caster. The thickness could be reduced down to 0.5 mm only by 3 times of hot rolling. This shows that the magnesium sheet could be made at very low cost.

Figure 6: Surface of roll cast AZ31 strip at speed of 120 m/min

Figure 6 shows surface of the as-cast strip and the rolled strip. The as-cast strip had metallic luster at proper casting condition. They say that it is difficult to make the sheet of AM60 and AZ91 as hot rolling of slab is difficult. However, hot rolling of AM60 and AZ91 could be operated, and the sheet of 0.5 mm thick was obtained.

147

Figure 7: Microstructure of roll cast AZ31 strip at speed of 120 m/min

Figure 7 shows the microstructure of roll cast AZ31 strip. The microstructure of as-cast strip was not uniform at thickness direction. The microstructure of center area was spherical structure, and the near surface was the equiaxed structure. This was not the character of the magnesium alloy but the high speed twin roll caster. This non-uniformity of the microstructure became almost uniform after the hot rolling. The warm deep drawing of the roll cast magnesium strip could be operated. Result was shown in Figure 8. L.D.R. (limiting drawing ratio) was greater than 2.0. The warm deep drawing of the magnesium alloy for casting like AM60 and AZ91 was able. The warm deep drawing of AZ91 is more useful than that of AZ31, as the AZ91 is stronger than AZ31.

(a) AM60 (Pd = 40 mm, t = 0.5 mm) (b) AZ91 (Pd = 32 mm, t = 0.5 mm) Figure 8: Warm deep drawing of roll cast magnesium alloy (T = 250 °C, LDR = 2.0)

6

Conclusions

A high speed twin roll caster of vertical type was designed and assembled to cast aluminum alloy and magnesium alloy thin strips. Some devices were adopted to realize rapid solidification of the strip. Casting was operated, and ability of the high speed twin roll caster was estimated. The A356 aluminum alloy could be cast at speeds from 60 m/min to 150 m/min. The magnesium alloy strip could be cast at high speed, too. Thickness of the strip was from 1.5 mm to 3.5 mm. The microstructure of the strip was not columnar structure but equiaxed structure. It

148 became clear that the high speed twin roll caster has ability to improve the deterioration of the aluminum alloy by impurity. The 0.5 mm thick of sheets of Am60 and AZ91 magnesium alloys was obtained by hot rolling from the roll cast strip. Warm deep drawing could be operated to AZ60 and AZ91 sheet, and their L.D.R was greater than 2.0.

7 [1] [2] [3]

References M.Yun, X. Yang, D.V.Edmonds, J.D.Hunt, P.M.Thomas: Cast Met.Vol.4-2(1991),p.108 D.V.Edmonds, J.D.Hunt: Extraction Refining and Fabrication of Light Metals, CIM, Ottawa, (1991), p.257 T. Motegi: Proc. of the ICAA-6, (1998), p.297

149

Simulation / Modeling

150

151

State-of-the-Art in the Modelling of Aluminium and Copper Continuous Casting Processes J.-M. Drezet1,2, M. Gremaud2 and M. Rappaz1 1

Computational Materials Laboratory, Ecole Polytechnique Fédérale de Lausanne, IMX-STI, Station 12, CH-1015 Lausanne, Switzerland 2 Calcom-ESI SA, Parc Scientifique, CH-1015 Lausanne, Switzerland

1

Abstract

With the advent of powerful and cheap computers, modelling of solidification processes at the macroscopic scale has become a standard practice in industry, in particular in continuous casting processes. Indeed, commercial software packages are available for the modelling of heat and fluid flow, as well as for stress-strain calculations. Such software can even be used in an “inverse way” in order to deduce casting parameters (e.g., heat transfer coefficients) from measurements (e.g., temperature). Electro-magnetic stirring, which is increasingly used to refine grain structures, can also be modelled by coupling hydromagnetic and thermal aspects. In more advanced approaches, convection in the liquid, heat exchange and stress developments are coupled all together in a mixed Lagrangian-Eulerian formulation. The start-up phase, which is crucial for many continuous casting processes (e.g., DC casting of Al alloys), is another field of development for which mixed formulation is promise full. Modelling of macrosegregation is still a critical issue as it can have different origins: convection, solidification shrinkage, grain movement/sedimentation, deformation of the mushy zone. Substantial progresses have been made in this area as well. Besides macroscopic aspects, modelling of microstructure and defect formation is an active field of research, in particular for microporosity and hot tearing. The present contribution will review the state-of-the-art modelling of aluminium and copper semi-continuous or continuous casting processes, at both the industrial, more macroscopic, approach and the still more academic, microscopic, level. This review is based on the communication given by M. Rappaz at the occasion of the 40th anniversary of the R&D Centre of Hydro-Aluminium (formely VAW) in Bonn in May 2004 [1].

2

Introduction

From an academic point of view, the VDC (vertical direct chill) casting process is schematically shown in Figure 1. Although fairly simple in its principle, it involves several interplaying phenomena which finally render modelling approaches complex. At the macroscopic scale (topics 1 to 4), heat and fluid flow modelling is of course the first step in order to predict the delivery of metal, the melt pool depth, the thermal gradient, the local solidification rate, etc ... Fluid flow is also essential in determining macrosegregation, i.e., transport of solute species at the macroscopic scale, and grain structure, i.e., influence of convection on dendrite fragmentation and grain transport. Almost as important is the calculation of stress build-up and strains, since this conditions air gap formation between the ingot and the mould or bottom block, and thus heat transfer.

152 Knowing the deformation of the slabs also helps to calculate the mould shape that will minimise scalping operations, while stress assessment is a key element to predict hot cracking formation. Although fluid flow and solid deformation are governed by the same basic conservation equations (conservation of mass and momentum), their associated rheologies are so different that combined Eulerian/Lagrangian approaches are required to handle the large displacements/ small stresses of the fluid and the small displacement/large stresses of the solid. Although several softwares can now calculate in a coupled way these two aspects, the approaches usually remain “one-phase”. If such approaches are very useful in predicting the combined interaction between heat flow, fluid flow and solid deformation, they cannot address in details phenomena occurring in the mushy zone. Recently, one has seen the emergence of “two-phase” approaches in which the solid and liquid equations are averaged over a typical volume element, with appropriate exchange terms, in order to predict hot cracking tendency.

Figure 1: Schematics of the VDC casting process. Topics 1–4 correspond typically to a macroscopic approach, whereas topics 5–8 correspond to a microscopic one. Typical variations of the thermal gradient, G, of the solidification rate, v, and of G/v are shown as a function of x [1].

For a metallurgist, macroscopic entities such as cooling rate, residual stresses, etc., are fine… but not sufficient! He wants to have access to macro- and microstructures as well as to defects such as microporosity and hot cracking (topics 5 to 8 in Figure 1). While Cellular Automata or granular-type approaches in the mid-nineties were a step forward for the prediction of grain structure formation, the advent at the same time of the phase field method in the materials science community gave great hopes to solve the problem of microstructure prediction. Amazing progresses have been made, but the technique still remain delicate to use and very CPU-intensive. For the prediction of microporosity, one has seen recently the emergence of 3D computations combining solidification shrinkage, gas segregation, nucleation and growth of pores. Finally, the prediction of hot tearing has really become a …. “hot” topic since the first approach combining strains of the solid skeleton and liquid feeding was published in 1999. The present paper reviews very briefly the state of the art of solidification modelling, focusing mainly on DC casting, and will outline some of the challenges that remain. It is largely inspired by the

153 work carried out in two European research projects, EMPACT (1996-2000) and VIRCAST (2000-2004) and in the on-going project POST (Porosity Stress).

3

Macroscopic Modelling

3.1

Heat Flow

Heat extraction is of course essential in solidification modelling, not only at the process scale, but also at the level of the microstructure-defects. The thermal gradient, G, and solidification rate, v, can be used in microstructure maps [2]: the fineness of the microstructure is essentially a function of (Gv), i.e., of the cooling rate, while the type of microstructure (e.g., columnar, equiaxed) depends on G/v. Niyama’s criterion for microporosity prediction is also a function of G/v [3]. It should be pointed out that, even in the steady state regime, the solidification rate v at the liquidus position depends on position and is not equal to the casting speed, vc, as shown in Figure 1. Modelling of heat flow in DC casting is no longer a challenge from a numerical point of view, in the start-up or steady state phases. The enthalpy method with evolving/activated meshes is robust in handling the strong non-linearity associated with latent heat release. The main questions in this area are related to materials properties and boundary conditions. For both, inverse methods have become a standard practice to calibrate the calculations [4]. As an example, Figure 2 shows the heat flux deduced at the surface of a DC cast ingot as a function of the distance from the top liquid surface. One can clearly distinguish the heat flux associated with primary cooling, the air gap formation (nil flux) and the strong heat extraction associated with water cooling.

Figure 2: Computed heat flux as a function of the distance to the top liquid level for an AA5182 alloy [4]

3.2

Fluid Flow

Despite the advanced CFD (computational fluid dynamics) software on the market, fluid flow calculations remain a task of specialists, especially when solidification occurs simultaneously. Since such calculations are performed in a single domain, containing the liquid and solid pha-

154 ses, a penalty method is used to make the velocity in the solid resume to the casting speed. The viscosity can be made a function of the volume fraction of solid or, better, a drag term similar to Darcy’s equation for a porous medium can be introduced in the momentum conservation equation. On the other hand, as the laminar viscosity of metals is very low, a turbulence model is also required to get realistic fluid flow pattern and intensity. One of the challenges that remain in this field is the validation of the calculated velocity field. Although some probes were designed by Vives [5], one has very often to rely on water models to calibrate the flow [6–7]. Another challenge is the interaction between the flow and the mushy zone. Besides fragmentation of dendrites [8] and formation of feathery grains [9,10], which are both favoured by convection, it is not easy to model accurately the narrow boundary layer near the liquidus in a transient regime, since it moves over time and normally requires adaptive-evolving meshes. As an example, Figure 3 shows the velocity field calculated with Fidap in a round billet fed from the lateral side [10]. The maximum shear rate of the liquid occurs at the opposite side of the gating and favours the growth of twinned dendrites in this region.

Figure 3: Velocities and temperatures at the surface and in the vertical symmetry-plane for an AA1050 DC cast billet, injected from the lateral side, as calculated with the Fidap software. Temperatures are in °C, (Tliq = 657 °C, Tsol = 645 °C). In-flow velocity is 0.02 m/s, from [10].

3.3

Stresses and Strains

Deformations of the solid during cooling and stress build-up are important issues in DC casting of aluminium alloy [11–13]. Deformation both limits the heat exchange with the mould and bottom block, and thus the production rate, and modifies the shape of the ingot. As shown schematically in Figure 1, butt curl, butt swell and lateral faces pull-in make the final shape of the slab to deviate from a parallelepipedon. In a transverse section as well, a convex mould has to be used if a rectangular cross section at the end of solidification is desired, since pull-in at the corners is smaller than at the mid-rolling faces. Stresses are also important as they induce hot cracking or even “cold” cracking. In this area, modelling is very mature, partially thanks to the EMPACT program, and many commercial software exist. However, deformation especially of the lateral faces is strongly influenced by the inclination of the thermal gradient, and thus by the sump shape. For example, an

155 increase of the casting speed makes the liquid sump deeper and thus the thermal gradient more horizontal, in which case pull-in is increased. By the same argument, the shape of the liquid pool being influenced by convection, and therefore by the distribution bag, accurate solid deformation calculations require to couple them, not only with heat-, but also with fluid-flow simulation. This challenge is not met by all commercial software. By the same argument, compressive stresses of the mushy zone can expulse interdendritic liquid out (i.e., sponge effect [14]) thus leading to deformation-induced segregation. Tensile stresses on the contrary lead to an opening of the mushy zone, inducing either segregation (healed hot tears) or hot cracks.

3.4

Macrosegregation

Despite the effort done during the EMPACT program, macrosegregation, i.e., solute inhomogeneity at the scale of the ingot, still remains a challenge in DC casting for several reasons [15]. First, macrosegregation can be induced by solidification shrinkage (inverse segregation) and exudations, by forced or natural convection, by grain movement and sedimentation and by deformation of the solid skeleton in the mushy zone. These effects have been addressed separately, sometimes in a combined approach (e.g., shrinkage and deformation [14] or shrinkage and natural convection [16]), but never in a comprehensive model. Second, each of this topic is complex. For example, deformation-induced segregation requires to couple stress-strain calculation, including the mushy zone, with heat-, mass- and solute-transport in the liquid phase [14]. Third, some of these phenomena are very localized. For segregation induced by natural convection, the origin is the region where both the liquid velocity and the solute gradients are non-zero. It is limited to a region of maybe 1 mm thickness at most near the liquidus surface [16,17].

Figure 4: Macrosegregation results for a 2D section of an AA5182 DC cast slab, as calculated with the CalcoSOFT software. The various pictures represent, from left to right: the steady state temperature profile, the isofractions of solid showing clearly the mushy zone position, the velocity field, the streamlines and the average Mg concentration [16].

156 Nonetheless, some interesting attempts to model macrosegregation have been made. For example, a result obtained in a 2D section of a DC cast slab of an AA5182 alloy is shown in Figure 4 [16]. It has been computed with the software CalcoSOFT, considering shrinkage and natural convection. The main contribution to macrosegregation in this case was found to be due to shrinkage: it induces a negative segregation zone near the centre of the ingot, the magnitude of which corresponds to concentration profile measurements.

3.5

Coupled Heat Flow, Fluid Flow and Solid Deformation

As stated before, the thermal field induces stresses and strains in the coherent solid and convection in the liquid region. Fluid flow transports also heat, while deformation of the solid modifies the thermal contact with the bottom block or with the mould. It is therefore necessary to couple these aspects. One of the goals of the POST project [18] is to implement a mixed LagrangianEulerian representation in the 3D code ProCAST [19]. Indeed, a significant difficulty which arises when one wants to model the start-up phase of continuous casting, is the large change in size of the calculation domain in the direction of the casting velocity. This expansion cannot be accounted for by simply deforming the finite elements: such a method would lead to a severe distortion of the elements with a damaging effect on the accuracy of the calculation. A solution to this problem is provided by the use of a dynamic mesh, i.e. a mesh in which new elements are automatically added during the course of a calculation. The strategy is therefore based on the technique of element activation. All the elements exist in the mesh from the beginning of the calculation so that the total number of elements is constant all over the simulation, but the elements can be active (regular) or inactive. Although present in the mesh, inactive elements are skipped during the assembly process and therefore have no influence on the calculation. A zone made of a sufficient number of layers of inactive elements is defined in the initial mesh between the liquid distribution system and the moving part of the casting. When the material begins to move following the bottom block, the nodes in the expanding domain will move at the same speed because of the Lagrangian nature of the description in this region. The layer of inactive elements in contact with the active region is activated and the corresponding elements are stret-

Figure 5: Automatic generation of the accordion domain. The thickness of the elements is set here to a small value to make the elements visible. One layer of elements of the accordion domain is already stretched (POST project, [18]).

157 ched to account for the expansion of the domain. As soon as the thickness of this transition layer reaches a given value, the elements get fixed in size and the next layer of inactive elements is activated. The inactive elements act as a reservoir which compensates for the expanding size of the domain by progressive stretching and activation. The behavior of this set of elements mimics the unfolding of an accordion: this is why it is called below the "accordion domain" (see Figure 5). Periodic boundary conditions are used to ensure that the layers of inactive elements have no effect on the calculation, in other words, to ensure the continuity of the solution from one side of the accordion domain to the other side.

4

Microscopic Modelling

4.1

Grain Structure Formation

In continuous casting, grain structure formation has to consider the nucleation of grains, the fragmentation of dendrite arms and grain growth at a microscopic scale, the motion of the grains and their thermal history at a macroscopic scale. All these phenomena are of course influenced by the temperature field and by fluid flow. Conversely, the presence of equiaxed grains modifies the apparent viscosity of the fluid and their movement also transports heat (and solute). As these phenomena are fairly complex, a semi-coupled approach can be made: the macroscopic aspects of the process are first calculated using appropriate average conservation equations (and thermo-physical data [20]). Then, the motion and growth of equiaxed grains are calculated as a “post-processing” module, using as input the local heat extraction and the velocity field computed at the macroscopic scale. For Cu based alloys, the influence of electromagnetic stirring (EMS) on grain refinement has been studied in a Bridgman furnace by Campanella and co-workers [21]. Metallographic inspection of the specimens, temperature measurements and numerical simulations performed with CalcoSOFT [22] revealed that the efficiency of EMS is strongly dependent upon the penetration of the liquid in the mushy zone and therefore upon the position of the convection vortices with respect to the liquid front. The results were analysed on the basis of a dendrite fragmentation criterion similar to Fleming‘s criterion for local remelting of the mushy zone.

Figure 6: Schematic representation (left) of local fluid flow in the liquid, ul, and in between the dendrites, udl. The projection of udl on the thermal gradient (parallel to the z-axis), udl,z, has to be compared with the velocity of the isotherms VT . On the right, schematics of remelting (1) and fragmentation (2) of secondary dendrite arms [21].

158 The new criterion tells that local remelting occurs when the component of the fluid flow velocity along the thermal gradient becomes larger that the speed of the isotherms, as depicted in Figure 6.

4.2

Microstructure Formation

The prediction of the macro- and microstructures obtained at the end of solidification is of great interest for aluminium producers. Indeed, the mechanical properties of aluminium alloys, but also defects such as hot tearing, are strongly dependent upon these parameters. This is probably why the phase-field method has attracted the attention of so many scientists over the past decade [23–26]. Indeed, with a fixed mesh, such a model is capable in principle to predict the formation of dendrites, eutectics, peritectics, etc ... It can directly use information coming from phase diagram calculations. Nevertheless, it has some inherent difficulties. In particular, the mesh must be fine enough (typically 0.1 micron) to correctly predict the curvature and the diffusion field around a dendrite tip, thus making the time step of the explicit scheme small as well. This makes the calculations CPU intensive and limits the applications to domains of typically 1 mm in size or less, unless adaptive meshes are used.

4.3

Microporosity Formation

Although probably less important than hot tearing for DC casting, microporosity is the key defect in shape castings. While resulting from the same combined effects occurring within the mushy zone, namely solidification shrinkage and gas segregation, this defect is usually classified into shrinkage and gas porosity. Taking into account these two phenomena, a quite general module has been developed within the software CalcoSOFT [27]. Darcy’s and mass conservation equations are solved only for the mushy zone, using a dynamic and regular finite volume grid superimposed to the fixed finite element mesh used to solve the thermal exchanges within the casting. In this way, an accurate description of the shrinkage-induced pressure drop in the mushy zone can be obtained. Appropriate boundary conditions have to be set at the limit of the mushy zone, taking into account free liquid pockets (regions directly connected to ambient air), closed pockets (regions of liquid totally surrounded by the solid and/or mould) and semi-open regions of liquid (regions of liquid connected to an open region through the mushy zone). At the same time, segregation of gases such as hydrogen is calculated in each volume element. Using the pressure and temperature fields, conditions for the nucleation and growth of pores are then determined, taking into account the important overpressure associated with the curvature contribution. Knowing the amount of microporosity formed in each volume element, macroporosity in hot spots and pipe shrinkage at free surfaces can be deduced simultaneously.

4.4

Hot Crack Formation

Modelling of hot cracking is probably one of the most challenging issues in solidification. Indeed, it involves thermal stresses/strains of the coherent solid, transmission of these deformations to the partially coherent mush and interdendritic liquid feeding. For many years, very simple

159 criteria, such as the solidification interval, were used to simply address the composition-dependence of the susceptibility of an alloy to hot cracking [28]. In 1999, however, a new approach, based on Niyama’s criterion for microporosity formation [3] was derived at EPFL (so-called RDG model after the names of the authors [29]). Assuming that the strain rate component perpendicular to the thermal gradient at the roots of the dendrites is uniformly transmitted to the non-coherent mushy zone, a simple mass balance combined with Darcy’s equation allows us to calculate the pressure drop in the interdendritic liquid induced by these deformations and solidification shrinkage. If this pressure falls below a cavitation pressure, a crack is assumed to form. Due to its simplicity, such a physically-sound approach can be easily implemented as a postprocessing calculation of a deformation simulation [30]. In order to refine the RDG approach and be more rigorous in the handling of the solid and liquid interactions, the viscoplastic volume change (dilatation/densification) associated with the thermally-induced deformations in the mush, and thus to the thermally-induced “opening-up” of the dendrite network have been recently included in a compressible model of the mushy zone [31–32]. A challenge in the area of hot cracking, and to some extent of porosity formation, is the description of the gradual transformation of a continuous interdendritic liquid film to a continuous and coherent solid. In low concentration alloys, which are the most sensitive to hot cracking, this is achieved by coalescence or bridging of the primary phase [33–35]. Eventually, the numerous parameters entering the compressible model describing the mechanical behaviour of the mushy zone have been determined not only for binary Al-Cu alloys [36] but also for industrial alloys, such as the AA5182 alloy [37].

5

Conclusion

Like DC casting, the development of models is a continuous process. It evolves with the power of computer, with the increased knowledge gained from experimental observations, with the advent of new numerical techniques (phase field, cellular automata, granular methods, etc.) and with the development of new approaches (hot tearing, two phase method, etc …). Nowadays, macroscopic simulation tools are being routinely used by industry, despite their apparent high cost (hardware, software, training of people, maintenance, creation of data bases and expertise for thermo-physical and other properties, validation, …). Yet, there is not a single company which could imagine not using such tools, since this is the only way to get an increased knowledge of the process and therefore to reduce scrap, improve quality and diminish production costs. New sophisticated methods developed recently by academic partners for microstructure and defect formation will also make their way to industry in the future. This is the next logical step to characterize cast ingots from a metallurgical aspect, thus leading the way to properties prediction. The authors would like to acknowledge the scientific and financial support of the EU research projects EMPACT and VIRCAST and of the project POST.

6 [1]

References M. Rappaz: State of the art and new challenges in modelling of aluminium sheet ingot castings, in Hydro-Aluminium Bonn publications, 2004, p.14

160 [2] [3] [4]

[5] [6]

[7]

[8] [9] [10] [11] [12]

[13]

[14] [15] [16] [17]

[18] [19] [20]

W. Kurz, D. Fisher, Fundamentals of Solidification (Trans. Tech., Aedermansdorf, Switzerland, 1998) E. Niyama, T. Uchida, M. Morikawa, S. Saito, A method of shrinkage prediction and its application to steel casting practice, AFS Int. Cast Metals J. (Septembre 1982) 52 J.-M. Drezet, M. Rappaz, G.-U. Grün, M.Gremaud, Determination of thermophysical properties and boundary conditions of DC cast aluminum alloys using inverse method, Met. Trans. 31A (2000) 1627 S.A. Argyropoulos, Measuring velocity in high-temperature liquid metals: a review, Scandin. J. Metallurgy 30 (2001) 273 Dong Xu et al.: The Use of Particle Image Velocimetry in the Physical Modeling of Flow in EM or DC Casting of Aluminum: Part I. Development of the Physical Model, Mater. Trans 29B (1998)1281 J. Dong Xu, W. Kinzy Jones, Jr., J.W. Evans, The Use of Particle Image Velocimetry in the Physical Modeling of Flow in Electromagnetic or Direct-Chill Casting of Aluminum: Results of the Physical Model Including Bag Geometry, Blockage and Nozzle Placement, Met. Trans 29B (1998) 1289 A. Hellawell, J.A. Sarazin, R.S. Steube, Channel convection in partly solidified systems, Phil. Trans. R. Soc. Lond. 345A (1993) 507 S. Henry, P. Jarry, M. Rappaz, <110> Dendrite Growth in Aluminum Feathery Grains, Met. Mater. Trans. 29A (1998) 2807 S. Henry, G.-U. Gruen, M. Rappaz, Influence of Convection on Feathery Grain Formation in Aluminum Alloys, Met. Mater. Trans. 35A, pp 2495–501, (2004) J.-M. Drezet, M. Rappaz, Modeling of Ingot Distorsion during Direct Chill Casting of Aluminum Alloys, Met. Mater. Trans. 27A (1996) 321 W. Droste, J.-M. Drezet, G.-U. Grün, W. Schneider, 3D-Modeling of Ingot Geometry Development of DC-Cast Aluminum Ingots During the Start-Up Phase, in Continuous Casting, Eds: K. Ehrke and W. Schneider, (DGM, Wiley-VCH, Frankfurt, 2000) p. 175 W. Droste, G.-U. Grün, W. Schneider, J.-M. Drezet, Thermo-mechanical modeling to predict shrinkage, shape and mold openings for dc-cast rolling ingots, Light Metals 2002, Ed. W. Schneider (TMS) p.703 T. Kajitani, J.-M. Drezet, M. Rappaz, Numerical simulation of deformation-induced segregation in continuous casting of steel, Met. Mater. Trans. 32A (2001) 1479 C. Beckermann, Modelling of Macrosegregation: Applications and Future Needs, Int. Materials Rev. 47 (2002) 243 T. Jalanti, Etude et Modélisation de la Macroségrégation dans la Coulée Semi-Continue des Alliages d’Aluminium, PhD Thesis 2145, EPFL (2000) Th. Kaempfer, M. Rappaz, Modeling of Macrosegregation During Solidification Processes Using an Adaptive Domain Decomposition Method, Model. Simul. Mater. Sc. Engng 11 (2003) 575 POST Research Project, Calcom ESI, CH-1015 Lausanne, Switzerland ProCAST User Manual, Version 2005, Revision 2.1, ESI Group, 75761 Paris, France H. Combeau , B. Appolaire , B. Dussoubs and N. Houti: Prediction of the macro and microstructures of solidification in DC casting of aluminium, Aluminium (2004, vol. 80), p. 614

161 [21] Th. Campanella, Ch. Charbon and M. Rappaz: Grain refinement induced by electromagnetic stirring: a dendrite fragmentation criterion in Met. Trans. A. vol. 35 (2004), p. 3201 [22] CALCOSOFT User Manual, Version 2004, Revision 3.2, Calcom ESI, CH-1015, Lausanne, Switzerland [23] J. A. Warren and W. J. Boettinger, Prediction of Dendritic Growth and Microsegregation Patterns in a Binary Alloy Using the Phase-Field Method, Acta mater. 43 (1995) 689 [24] J. Tiaden, B. Nestler, H. J. Diepers, I. Steinbach, The multiphase-field model with an integrated concept for modelling solute diffusion, Physica D 115 (1998) 736 [25] B. Nestler, A. A. Wheeler, A multi-phase-field model of eutectic and peritectic alloys: Numerical simulation of growth structure, Physica D 138 (2000) 114 [26] W. J. Boettinger, J. A. Warren, C. Beckermann, A. Karma, Phase-Field Simulations of Solidification, Ann. Rev. Mater. Res. 32, (2002) 163 [27] Ch. Pequet, M. Gremaud, M. Rappaz, Modeling of Microporosity, Macroporosity and Pipe Shrinkage Formation during the Solidification of Alloys using a Mushy-Zone Refinement Method, Met. Mater. Trans. 33A (2002) 2095 [28] J. Campbell, Casting, Ed. O. Butterworth-Heinemann. (1991) [29] M. Rappaz, J.-M. Drezet, M. Gremaud, A New Hot-Tearing Criterion, Met. Trans. 30A (1999) 449 [30] J.-M. Drezet, M. M'Hamdi, S. Benum, D. Mortensen, H. Fjær, Hot tearing during the start-up phase of DC cast extrusion billets, in Aluminium Alloys, Eds. P.J. Gregson and S. J. Harris (2002) Part 1, p. 59 [31] C. L. Martin, M. Braccini, M. Suéry, Rheological behaviour of the mushy zone at small strains, Mat. Sc. Eng. A325 (2002) 293 [32] M. M'Hamdi, A. Mo and C. L. Martin, Two-phase modelling directed towards hot tearing formation in aluminium direct chill casting, in Met. Mater. Trans. Vol. 33A, July 2002, p. 2081 [33] M. Rappaz, A. Jacot and W. Boettinger Last stage solidification of alloys: a theoretical study of dendrite arm and grain coalescence, Met. Mater. Trans. 34A (2003), p.467 [34] V. Mathier, A. Jacot, M. Rappaz, Percolation of equiaxed grains during last stage solidification, Model. Simul. Mater. Sc. Eng. 12 (2004) 1 [35] S. Vernède and M. Rappaz: Transition of the mushy zone from continuous liquid films to a coherent solid, accepted for publication in Philosophical Magazine 2005 [36] O. Ludwig, B. Commet, J.-M. Drezet, C. L. Martin, M. Suéry, Rheological behaviour of partially solidified Al-Cu alloys: Experimental and numerical study, in Modelling of Casting, Welding and Advanced Solidification Processes X, Ed. D. Stefanescu et al (TMS, Warrendale, USA, 2003) p. 183 [37] O. Ludwig, J.-M. Drezet, P. Ménès and M. Suéry, Rheological behaviour of a commercial AA5182 aluminium alloy during solidification presented at ICASP, Stockholm 2005, and submitted to Mat. Sciences and Eng. A

162

Numerical Simulation of the Growth of Interacting, Equiaxed Dendrites with a Two-Scale Model 1

2

2

M. Jurgk , H. Emmerich , R. Siquieri 1

Max-Planck-Institut fur Physik komplexer Systeme, Dresden RWTH Aachen

2

1

Motivation

Simulating the solidification of a stable phase at the expense of a metastable phase is one of the elaborate multiscale problems of computational materials science. Starting from atomistic considerations of atom attachment over dendritic growth dynamics of a single crystal to grain growth and nally to the inuence on material properties of macroscopic cast metals the phenomenon of solidification spans 9 orders of magnitude in length scale. For an understanding of the multiscale nature of solidification paths models of dendritic growth at the scale of micrometers have played a central role. These models include the atomistic dynamics via transport and attachment coeficients. Moreover, they allow for a coupling to macroscopic transport dynamics via scaling relations for the growth velocity of the dendritic tip depending on system parameters like e.g. the surface-tension anisotropy or the strength of the undercooling. Apart from the study of dendritic growth in order to gain insight into the underlying, fundamental, physical mechanisms, one goal is to improve macroscopic application-oriented solidification simulations in their quantitative power of prediction as well. The growth dynamics of a single isolated dendrite has already been studied and understood quite well and within this contribution we will present new scaling relations for dendritic growth which account for the interaction of several dendritic crystals in a melt.

Figure 1: Left: The original problem posed by alloy solidification in a characteristic domain of sample. Right: Illustration of the homogenised two-scale model.

163

2

Model

If an alloy solidfiies and any flows in the melt are neglected, the controling transport mechanisms for the solidification process are the diffusion of heat and the diffusion of solute. The heat diffusion in alloys is usually several orders of magnitude faster than the solute diffusion. This fact motivates the introduction of two different scales in the modeling of alloy solidification: a microscopic scale for the solute diffusion and a meso/macroscopic scale for the heat diffusion. The movement of the solid-liquid interface of a dendritic crystal takes place on the microscopic scale of the slowly diffusing solute while the interaction between several of these crystals takes place on the macroscopic scale. The initial situation is shown in the left part of Figure 1 and is as follows: • Everything is in 2D. • Equiaxed dendritic nuclei are homogeneously distributed in the melt of a binary alloy. • The process is modeled by a sharp-interface model and the equations for both diffusion fields T and c have to be solved in the same, whole domain, but only in the liquid phase (one-sided model). The separation into two scales is mathematically done by the introduction of the scale parameter H, which species the ratio of the solute-diffusion scale to the heat-diffusion scale,

H

lc lT

and for the derivation of the final equations shall satisfy the condition H  1. The final homogenised two-scale model is achieved by means of asymptotic expansions in H

x· x· x· § § § TH (t , x) T0 ¨ t , x, ¸  H ˜ T1 ¨ t , x, ¸  H 2 ˜ T2 ¨ t , x, ¸  ... H¹ H¹ H¹ © © © x x x· § · § · § cH (t , x) c0 ¨ t , x, ¸  H ˜ c1 ¨ t , x, ¸  H 2 ˜ c2 ¨ t , x, ¸  ... H¹ H¹ H¹ © © © x x § · § · Xe (t , x) H ˜ X0 ¨ t , x, ¸  H 2 ˜ X0 ¨ t , x, ¸  ... H¹ H¹ © © plugging these expansions into the modeling equations and collecting all terms with H up to the second order [1]. The equations of the two-scale model and the situation after the homogenisation procedure (right part of Figure 1) are as follows: Now, there is the macroscopic, homogenised heat-diffusion equation

L· w§ ¨ T  H s ¸  K 'T cX ¹ wt ©

0 ,

(1)

164 which we complement with a boundary condition of Biot type,

wT wn

Bi TA  TB ,

(2)

B

and there is the microscopic growth problem for each nucleus:

wc  D'c wt

0

(3)

J

(4)

'cX n .

(5)

b2  cI  ceq  N  b1T L D

wc wn

I

The macroscopic problem (1)–(2) and the microscopic problems (3)–(5) are coupled in both directions: The term L/cX · wt Hs in equation (1) expresses the latent heat which is released at some point of the macroscopic temperature field due to the microscopic growth of the phase interface at this point which is modeled by the change in the local solid volume fraction Hs. In the other direction the coupling is visible by the term –b1 T in equation (4) which expresses the dependence of the microscopic phase-equilibrium concentration on the macroscopic temperature.

3

Numerical Implementation

We have modified and extended a simulation for the computation of the growth of a single dendrite under control of a single diffusion field [2]. The numerical setup is illustrated in Figure 2a: The fast diuffsing temperature field is defined on the points of a coarse grid (thick, full lines) while the slow concentration field is defined on the points of a fine grid (thin, dashed lines). One nucleus is located at each point of the coarse temperature grid. The important features of the numerical code are: • Spatial and temporal discretisation is done with finite differences. • For the temporal integration of the diffusion equations the explicit first-order scheme (Euler) is used. • The microstructure part: • The phase interface is a polygon line, which moves with higher accuracy than grid spacing, because it is discretised independently of the diffusion-field grid. (see Figure 2b) • Because a single grid causes artificial numerical anisotropy, a stack of rotated grids (usually 4) is used for the discretisation of the microscopic diffusion field. (see Figure 2c) • The measured dendrites’ properties are the tips’ radii and growth velocities.

165

a) b) Figure 2: Illustrations of the numerical setup

4

c)

Results

The scale parameter H can be related to the nuclei density U: H v1U With the computer simulation we have investigated • the dependence of the tip-growth velocity Xtip on the scale parameter H or on the inverse of the nuclei density 1/U, respectively, and • the inuence of the Biot number Bi which characterises the heat transport out of the domain on the growth-mode selection. The 1/r-dependence of Xtip is shown in Figure 3. There appears a transition between two growth modes at Hcrit = 1/Ucri = 0.0011. The existence of these two solution branches can also be seen in analytical investigations [3].

4.1

Solution branch X1n

A smaller density of crystal nuclei and thus a larger value of H means a larger scale for the heat diffusion. As long as this scale is that large, that the transport of latent heat away from the nuclei can still be accomplished completely, the solid fraction of the nuclei grows considerably in time during the evolution. As a consequence the term L / cX u wt Hs is large and dominates the heat-diffusion equation (1). The respective analysis results in kinetic dendrites with a kinetic coecient depending strongly on the crystal density (see [3]). That is why the velocities of the solution branch decrease steadily with increasing nuclei density.

4.2

Solution branch Xn2

The solution branch Xn2 appears for higher nuclei densities and thus smaller values of H. Due to the smaller scale for the heat diffusion the transport of the arising latent heat is not complete and the nuclei grow slowly. Hence the solid-fraction term L / cX · wt Hs does not dominate the heatdiffusion equation (1) and consequently in the analysis the temperature term in the GibbsThomson relation (4) does not depend on H (like it does in the case). This is why the solid-fraction term and thus the scale parameter H or the nuclei density U, respectively, does not determine

166

Figure 3: The new numerically obtained scaling relation with two solution branches

Figure 4: The inuence of the Biot number on the scaling relation

the growth rate of the nuclei for this solution branch and the growth follows the laws of capillary dendrites.

4.3

Biot Number

The analytic investigations show an inuence of the Biot number Bi on the critical value Hcrit where the transition between the solution branches appears [3]. In Figure 4 the H-dependence of

167 the tips’ growth velocity Xtip is plotted for the two Biot numbers Bi = 1 and Bi = 10. The transition is shifted to larger values of H for a smaller Biot number.

5

Outlook

The directions of our current and future work are •

to check if these new scaling relations improve the quantitative validity of multiscale approaches like e.g. in [4] by replacing the empirical metallurgical relations, • to study the two-scale model for other growth morphologies like peritectic growth, • to include melt ow as a transport mechanism in the model, and • to incorporate a phasefield model for the extension to 3D.

6 [1] [2] [3] [4]

References C. Eck, P. Knabner, S. Korotov, J. Comp. Phys. 2002, 178, 58 T. Ihle, H. Müller-Krumbhaar, Phys. Rev. E 1994, 49, 2972 H. Emmerich, M. Jurgk, R. Siquieri, Phys. Stat. Sol. (B) 2004, 241, 2128 C.Y. Wang, C. Beckermann, Metall. Mater. Trans. A 1993, 24A, 2787

168

Monte Carlo Simulation of Grain Growth in Three Dimensions D. Zöllner, P. Streitenberger Otto-von-Guericke-Universität Magdeburg, Fakultät für Naturwissenschaften, PF 4120, D-39016 Magdeburg

1

Abstract

A Monte Carlo Potts model algorithm for single-phase normal grain growth is presented, which allows one to simulate the development of the microstructure of very large grain ensembles in three dimensions. The emphasis of the present work lies on the investigation of the relaxation process. Different initial grain structures characterized by different initial size distributions are subjected to grain growth via Monte Carlo simulation. The temporal development of 3D grain structures reaches independent of the initial size distribution, after an initial period of time, a quasi-stationary self-similar coarsening regime where all scaled size distribution functions collapse to the single universal, time independent size distribution f(x). The relaxation process to this universal self-similar state is studied by following the temporal development of quantities like the average grain size, the standard deviation of the grain size distribution and topological correlations.

2

Introduction

One method to simulate grain growth is the Monte Carlo (MC) simulation based on a numerical realization of the Potts model [1]. In 1984 Anderson et. al. for the first time introduced the Potts model to simulate the grain growth kinetics and to investigate grain size distribution and topology in two dimensions [2, 3]. Especially the last years opened up growing possibilities of MC investigations of grain growth in three dimensions (cf. e.g. [4, 5, 6, 7] and the literature within).

Figure 1: Two shots of 3D grain growth (left: 0th MCS; right: 500th MCS) taken from a simulation run; the grains are shaded for a better presentation

169 The aim of the present work is the investigation of normal grain growth by studying the temporal development of quantities like the average grain size, the standard deviation of the grain size distribution, the microstructure and topological correlations. It is shown that the simulated self-similar volumetric rate of change of all grains can be fitted to an average scaled growth law that allows us to model the grain growth behavior by an analytic mean-field theory.

3

3D Monte Carlo Simulation

3.1

Monte Carlo Potts Model Algorithm

A Monte Carlo algorithm for single-phase normal grain growth in three dimensions has been implemented. Therefore the microstructure is mapped on a cubic lattice with 26 nearest neighbors and periodic boundary conditions. Each lattice point is called one Monte Carlo Unit (MCU). Each MCU has assigned an integer value representing the orientation Q of the lattice point. While in earlier simulations the size of the 3D lattice was usually limited to 100u100u100 MCU’s and therefore, those simulations were terminated to small grains we choose in the present 3D simulations the number of MCU’s up to 250 u250u250. The time unit of the simulation is called Monte Carlo Step (MCS) and is defined as N reorientation attempts where N is equal to the number of MCU’s. That means, in our present 3D simulation N = 15625000. Further details can be found in the literature [2–7, 8]. Each initial structure was subjected to grain growth 10 times for 1000 MCS. For the analysis of the simulation the data of each step were summed up over the 10 runs.

3.2

Growth Law, Scaling Regime and Size Distribution

The coarsening process (Fig. 1) has been investigated by following the temporal development of an initial Poisson Voronoi Tessellation and an initially rayleigh distributed grain structure.

Figure 2: Temporal development of the mean grain radius together with the fit of the growth law (Eq. 1): I – initial period; II – self-similar coarsening regime

170 Both 3D grain structures as they have been simulated by the MC procedure follow, after an initial period of time, the well-known growth law for the average grain radius (cf. Fig. 2)

 R ! b ˜ t ȕ  const ,

(1)

where the growth exponent of the initially rayleigh distributed grain structure is numerically given by = 0.4998. The constant b is given by b = 0.4322 corresponding to a growth parameter Ȗ k/  R ! R ! 12.6543 k is the kinetic constant of the simulation. After the initial period the grain size distribution function (GSDF), F(R, t), is self-similar, i.e., it is characterized by the scaling law

F(R,t)

g(t) ˜ f(x) .

(2)

g(t) is a time-dependent factor and f(x) is the normalized GSDF with x = R / , where R is defined as the radius of an volume equivalent sphere and is the average grain radius. All GSDF within the quasi-stationary self-similar (QSSS) coarsening regime collapse to the single universal, time independent size distribution f(x) (cf. Fig. 3).

Figure 3: Temporal development of the GSDF of the initial Voronoi Tessellation for different time steps (stars:

Figure 4: Comparison of the temporal development of the standard deviation of both initial grain structures

171 The temporal development of the GSDF is quantitatively shown through the development of the standard deviation sd, which is defined by sd

 x 2 !   x ! 2 and x n

³x

n

˜ f(x)dx .

In the QSSS coarsening regime sd stays nearly constant, while the temporal development of sd towards this state depends on the initial GSDF (Fig. 4).

3.3

Microstructure and Topological Correlations

In order to investigate the coarsening process of the microstructure to the QSSS state the temporal development of the number of neighboring grains s of grains with relative size x has been observed. In Figure 5 s vs. x is shown for the initial state (Voronoi Tessellation) and the QSSS state after 500 MCS.

a

b

Figure 5: Number of neighboring grains s of grains with size x (a – initial Voronoi Tessellation, b – QSSS state: in terms of the residual sum of squares (= sum of all squared residuals, abbreviated as rss) a quadratic relationship relating s to x yields clearly a better approximation to the simulation results than the linear fit)

The initial Voronoi Tessellation shows a certain correlation between the number of neighboring grains and their size in so far as larger grains have more neighbors than smaller grains (cf. Fig. 5a). During relaxation this size correlation becomes stronger and, within the QSSS state, can be described by a self-similar quadratic function (Fig. 5b). Because of the changes in the topology the mean number of neighboring grains or faces changes with time. For the Voronoi Tessellation has a value of = 16.7993, while the initially rayleigh distributed grain structure gives = 13.55. During the simulation both values change and yield within the QSSS state = 13.9292 respectively 13.9339. The value 13.93 is close to 14 which is the value of the mean number of faces of Kelvin’s model [9]. Lord Kelvin’s model describes a possibility to fill a 3D space completely with objects of the same size. These so called tetrakaidecahedron have each 14 faces (eight hexagonal and 6 quadratic faces) which are slightly curved.

3.4

Average Scaled Growth Law and Mean-Field Theory

Aside from the initial period of time the grain growth can be characterized by the volumetric rate of change V 1/ 3V ~RR , which depends only on the scaled grain size x and therefore is time

172

Figure 6: a – Average scaled growth law (stars: simulation data; line: Eq. 4); b – Simulated size distribution (stars) compared to analytic functions

independent and self-similar [10, 11]. In terms of mean-squared errors a quadratic function RR ( x) (cf. Fig. 6a) gives the best approximation to the simulation results. A mean-field theory based on a quadratic form of RR vs. x and the Lifshitz-Slyozov-Hillert stability conditions has been given in [8, 12, 13]. In this theory the size distribution function is given by the one-parameter family

f ( x) a ˜ x0a ˜ ea ˜

§ a ˜ x0 · x ˜ exp ¨  ¸, a2 ( x  x0 ) © x0  x ¹

x d x0 ,

(3)

where a = a(x0) is defined by the scaling requirement <x> = 1 [12, 13]. The upper cut-off x0 serves as a free adjustable parameter. The solid line in Fig. 6a shows the relation RR ( x) as it follows from a least-square fit of the theoretical expression [8]

RR ( x)

k ª§ 3 · 2 6 x0 3 x02 º ˜ «¨ 1  ˜x » ¸˜ x  J ¬«© a( x0 ) ¹ a ( x0 ) a ( x0 ) ¼»

(4)

to the simulation data of RR , yielding x0 = 4.1275 and consequently a(x0) = 16.795. The fit gives a good average description of the simulation results of RR vs. x. The resulting analytic GSDF Eq. 3 is in very good agreement with the simulated size distribution (cf. Fig. 6b). For x0 = 9/4 and consequently a(x0) = 3 Eq. 4 reduces to a linear function and Hillerts GSDF results [14], for x0 o f Eq. 3 yields Louats GSDF [15], which are both clearly in contradiction to the present simulation results (cf. Fig. 6b). It is the relatively small quadratic term in Eq. 4 that is responsible for the excellent agreement between the simulated and analytical GSDF in Fig. 6b.

173

4

Conclusions

In this paper MC simulation results on normal grain growth in three dimensions have been presented. After a short initial period of time a quasi-stationary state of the evolving grain structure is reached, which is characterized by a parabolic grain growth kinetics and a self-similar GSDF that is nearly time invariant. The local topology of the simulated grain microstructure can be characterized by a quadratic size correlation function. From the MC simulation results of 3D grain growth the average scaled growth law as a function of the scaled grain size was determined. The analytic GSDF based on a mean-field description of the simulated microstructure shows a very good agreement with the GSDF resulting from the 3D MC simulation.

5

Acknowledgements

The authors would like to thank the Professors Martin E. Glicksman, Günther Gottstein and Dimitri Molodov for useful discussions and the Deutsche Forschungsgemeinschaft for financial support under GKMM 828.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

References R.B. Potts, Proceedings of the Cambridge Philosophical Society Vol. 48, 1952, 106–109 M.P. Anderson, D.J. Srolovitz, G.S. Grest, P.S. Sahni, Acta Metallurgica Vol. 32, 1984, 783–791 D.J. Srolovitz, M.P. Anderson, P.S. Sahni, G.S. Grest, Acta Metallurgica Vol. 32, 1984, 793–802 P. Anderson, G.S. Grest, D.J. Srolovitz, Philosophical Magazine B59, 1988, 293–329 X. Song, G. Liu, Scripta Materialia Vol. 38, 1998, 1691–1696 Q. Yu, S.K. Esche, Materials Letters Vol. 57, 2003, 4622–4626 A. D. Rollett, P. Manohar, in Continuum Scale Simulation of Engineering Materials, edited by Dierk Raabe et.al., Wiley-VCH, Weinheim, 2004, 77–114 D. Zöllner, P. Streitenberger, in Proceedings of the 2nd Joint International Conference on Recrystallization and Grain Growth, Annecy, France, TTP, 2004, 1129–1136 The Kelvin Problem, edited by D. Weaire, Taylor & Francis, London, GB, 1996 W.W. Mullins, Journal of Applied Physics Vol. 59, 1986, 1341–1349 M.E. Glicksman, Philosophical Magazine Vol. 85, 2005, 3–31 P. Streitenberger: Scripta Materialia Vol. 39, 1998, 1719–1724 P. Streitenberger: in Proceedings of the 1st Joint International Conference on Recrystallization and Grain Growth, Aachen, Germany, Springer Verlag, 2001, 257–262 M. Hillert, Acta Metallurgica Vol. 13, 1965, 227–238 N.P. Louat, Acta Metallurgica. Vol. 22, 1974, 721–724

174

Thermal Conductivity of Ternary and Multi-Component Aluminum Alloys up to and above the Melting Temperature R. Brandt1, W. Bender2, G.-U. Grün2, G. Neuer1 1

Institute for Nuclear Technology and Energy Systems (IKE), University of Stuttgart, 70550 Stuttgart, Germany, E-mail: [email protected] 2 Hydro Aluminium GmbH, Bonn, Germany

1

Abstract

The determination of thermophysical properties in a broader temperature range is an essential prerequisite for proper process optimization by numerical modeling techniques, which are now commonly used in the industry. Here, results of a measurement campaign on various ternary AlSiMg and multi-component AlSiCu+ alloys using a new electric resistivity measurement unit are presented. The evaluated temperature dependent data are discussed regarding alloying element and equipment related influences.

2

Introduction

Thermophysical properties of metals in the melting range are crucial to optimize processes e.g. powder production by atomizing or casting by means of numerical models simulating solidification. Existing calculation programs model both the bulk heat transfer and the fluid dynamic process during solidification aiming to predict e.g. time dependent temperature profiles in the various states of the material: liquid, “mushy region”, solid, determine solidification rates etc. These numerical models are important tools for the better understanding of the physics of solidification and to improve quality of products, however, they are extremely sensitive against thermophysical properties. The demand and the availability of these properties is summarized by Ludwig, Quested, and Neuer [1]. Beside viscosity and surface tension the knowledge of the thermal conductivity is of particular importance. As the thermal conductivity is strongly related to microstructure of the solid material the determination during solidification is a key aspect, because the development of the microstructure highly depends on the cooling rate. Therefore the measurement technique should enable to measure the thermal conductivity in the liquid state, in the “mushy” region, and in the solid state before and after melting. Only very limited thermal conductivity results of aluminum alloys are published till now [2,3,4,5,6,7]. Only one investigation of the thermal conductivity of aluminum alloys with systematic variation of individual alloy components has been found [8]. The Si content of binary AlSi alloys had been varied between 0 - 3 -7 and 11 %Si. In the frame of the European Network Program "Microstructural Engineering by Solidification Processing" (MEBSP) a number of aluminum based alloys has been produced by Hydro Aluminium Deutschland GmbH. Specimens of the binary system AlSi, of the ternary system AlSiMg, both produced out of pure material basis, and of a AlSiCu+ alloy, representing a commercial multi-component system with fixed values for all other alloying elements but varied in

175 Si and Cu content, have been distributed to various laboratories in order to measure the density, surface tension, viscosity, thermal conductivity, and electric resistivity as functions of temperature and of chemical composition. All samples were presented in the as-cast condition.The complete list of all 21 different materials has been presented by Brandt and Neuer [9] together with the measured densities at room temperature and the electric resistivities measured at 25 °C, 50 °C, and 75 °C. Thermal conductivity results of the AlSi alloys are published by Brandt and Neuer [10] at temperatures up to 800 °C. The data evaluated for the ternary and multi-component systems shall be discussed here.

3

Measurement Technique

The electric resistivity was measured by using the four probe technique, whereby a cylindrical sample of diameter D is loaded by a direct current I. The electric resistivity is determined by measuring the voltage drop U between two probes positioned at a distance of S using the following equation:

U

'U S D2 ˜ ˜ I 4 S

(1)

Details of the measurement technique are described in [9,10]. For high accuracy measurements in the solid state the specimen is mounted on a vernier caliper with electrical insulated razor blades on the flanks. This allows measurements with variable distance S. Specimens of 3 mm to 10 mm in diameter and 100 mm to 200 mm length can be used. The whole arrangement is dipped in an oil bath to get a well defined and homogeneous temperature distribution. Thus steady state measurements up to about 100 °C are possible. Measurements in the molten state of a metal are much more difficult than measurements only on solids because the sample must be kept inside of a ceramic tube. This involves that geometric data and sample temperature are more difficult to be determined and the complicate mechanics of the electrodes may lead to increased measurement errors. Solutions of such problems are described in [10]. For the investigation of the aluminum alloys described here specimen rods of 5 mm in diameter and 100 mm long were inserted in a tubular crucible of aluminum nitride which was mounted in its vertical position in a vacuum chamber. Here the distance between the voltage probes was fixed to S = 50 mm. The vacuum chamber is surrounded by a tube furnace, which enables measurements in vacuum or in each available gas atmosphere at temperatures between 30 °C and 1600 °C. The apparatus is described in more detail in [10]. The measurements were carried out between room temperature and about 800 °C using a constant electrical current of 2 amps. Higher currents give better voltage drop signals but the high contact resistance between current electrodes and specimen causes a rapid heating of the sample edges. All measurements were performed under high vacuum conditions. A heating rate of 360 K/h was applied in the solid and in the liquid state with interrupts before each measuring point, during melting and solidification the heating/cooling rate was reduced to about 100 K/h in order to get better resolution for the melting/solidification temperatures. The thermal conductivity was calculated from the measured electric resistivity by using the Wiedemann-FranzLorenz law O /V = L T, which describes the relationship between thermal conductivity O and electrical conductivity V = 1/Ufor metals. L is the Lorenz number, which depends on the type

176 of metal and also to some extent to its temperature, and T is the absolute temperature in Kelvin. For aluminum and its alloys it has been shown that by using the theoretical Sommerfeld value of L0 = 2.445 * 10-8 V2/K2 the calculated conductivity data may be too high by about 13 % at room temperature, but with increasing temperature this deviation becomes smaller, and above 200 °C the deviation is lower than 2 % for the studied aluminum alloys [3, 10].

4

Measurements and Results

Table 1 and Table 2 compile together the sample designations, related chemical composition, density, and electric resistivity US at solidus andUL at liquidus temperatures both during heating and cooling cycles. A complete list of the electric resistivity measurements of all samples is presented in [10]. On the basis of these data the thermal conductivity values discussed below are then calculated as described in the section above. Table 1: Sample designations and some characteristic values of the investigated ALSiCu+ alloys sample No. Si [%] Cu [%] others [%] density [g/cm3]

2,730

2,747

2,765

2,710

2,755

UL [P:*cm] UL/US

AlSi-13 AlSi-15 6 6 1 3 2.5 – 3 2,748 2,783 heating cycle 563 °C 538 °C 13,17 13,38 626 °C 611 °C 31,03 31,27 2,36 2,34

545 °C 13,50 595 °C 32,16 2,38

550 °C 13,78 602 °C 32,67 2,37

536 °C 14,39 598 °C 32,84 2,28

550 °C 14,19 580 °C 33,10 2,33

531 °C 14,03 580 °C 33,36 2,38

Liquidus UL [P:*cm] Solidus US [P:*cm] UL/US

cooling cycle 616 °C 603 °C 30,71 31,29 541 °C 530 °C 16,59 18,22 1,85 1,72

597 °C 32,40 537 °C 17,89 1,81

593 °C 32,34 536 °C 18,06 1,79

586 °C 32,54 533 °C 18,71 1,74

575 °C 33,42 543 °C 19,59 1,71

568 °C 33,19 539 °C 19,06 1,74

Solidus

US [P:*cm] Liquidus

AlSi-16 AlSi-17 AlSi-18 AlSi-19 AlSi-21 9 9 9 11 11 1 2 3 1 3

The tables show that during melting the electric resistivity increases by a factor UL/US of about 2.2–2.4 for all measured AlSiCu+ and AlSiMg alloys. This agrees with the values for pure aluminum (2.21) and binary AlSi alloys (2.3–2.4) measured before [10]. Fig. 1 exemplarily shows the resulting thermal conductivity development during heating and cooling for a AlSiCu+ alloy with 9 wt% Si and varying Cu contents, while in Fig. 2 the values for fixed 3 wt% Cu and variable Si contents are displayed. It is evident that, compared to the binary AlSi, an addition of 1% Cu lowers the conductivity considerably, 2% Cu causes a light further decrease, while an addition of 3%Cu has nearly no further influence on thermal conductivity. In the case of variable Si-content (Fig. 2) the conductivity decreases continuously with increasing Si-content.

177 Table 2: Sample designations and some characteristic values of the investigated ALSiMg alloys sample No. Si [%] Mg [%] density [g/cm3] Solidus US [P:*cm] Liquidus UL [P:*cm] UL/US Liquidus UL [P:*cm] Solidus US [P:*cm] UL/US

AlSi-6 5 0,6 2,669 heating cycle 567 °C 11,71 634 °C 27,79 2,37 cooling cycle 624 °C 27,51 538 °C 12,78 2,15

AlSi-8 7 0,6 2,666

AlSi-10 9 0,6 2,659

AlSi-11 12 0,3 2,640

AlSi-12 12 0,6 2,642

569 °C 12,75 621 °C 28,83 2,26

567 °C 13,08 599 °C 29,05 2,22

569 °C 13,73 609 °C 31,61 2,30

580 °C 14,96 590 °C 31,55 2,11

609 °C 28,63 542 °C 13,98 2,05

596 °C 29,70 549 °C 15,02 1,98

576 °C 31,37 548 °C 18,23 1,72

579 °C 31,15 545 °C 16,98 1,83

Figure 1: Thermal conductivity of AlSi9Cu+ with varying Cu-content, calculated from electric resistivity measurements for heating and cooling cycles

A common factor in both figures is the fact that the values for the solid sample in the cooling cycle after solidification are remarkably lower than those values measured for these temperature levels during the heating cycle. In order to better understand the reason of this behavior, additional measurements at room temperature with the more accurate vernier caliper technique (VC) were carried out for the already processed sample AlSi-6 and compared with corresponding calibration measurements before the high temperature measurements.

178

Figure 2: Thermal conductivity of AlSiCu3+ with varying Si-content, calculated from electric resistivity measurements for heating and cooling cycles

This is shown in Fig. 3. Here, the resulting values of measurements at room temperature of the re-solidified rod for different distances and locations are compared with the results from the high temperature measurements. As it can be seen the value for the VC (S = 50 mm), which means the vernier caliper measurement over the same distance as during high temperature measurement nicely correlates with the final value of the high temperature measurement. But, the two values at a zone of 30 mm from the lower end (S = 30 mm bot.) and from the upper end (S = 30 mm top) demonstrate that there is a thermal conductivity gradient along the length of the rod after the resolidification. This inhomogeneity in the results could have a variety of reasons. A gradual change in microstructure including segregation processes due to the solidification conditions could be a possible explanation. But also the development of shrinkage porosity inside the sample caused by hydrogen segregation or insufficient volumetric compensation within the probe crucible during the end of the cooling phase would result in such strong gradual differences. For further clarification regarding the possible cause the processed sample rod AlSi-15 was cut along the longitudinal symmetry plane and porosity conditions and microstructure were evaluated. Fig. 5 clearly visualizes the most probable cause for the inhomogeneity. Quite obvious are the large differences in porosity along the rod. Most probably due to gravity the upper part of the sample shows large pores, which in the cross section area nearly seem to separate both rod parts, although the sample was still intact when removed from the ceramic tube. So, still electrical contact existed throughout the duration of the experiment. In the section taken of the mid part of the sample the porosity is definitely lowest, while in the lower part a slight increase occurs again.

179

Figure 3: Thermal conductivity for sample AlSi-15 during heating and cooling phase. Single results on the left are conductivities before (run 1) and after (run 3) resolidification, measured with vernier caliper (VC) at different sections of sample.

Figure 4: Thermal conductivity for AlSiMg alloys with varying Si-content, calculated from electric resistivity measurements for heating and cooling cycles

180 Parallel evaluation of the microstructure revealed a more or less uniform structure along the whole rod with no significant differences for an explanation of the variations in the thermal conductivity as shown in Fig. 3. Therefore, the most probable reason for the differences in the resistivity during heating and cooling cycle is the evolving porosity inside the sample during the solidification phase of the measurement sequence. This is supported by the results of the ternary AlSiMg alloy variants. As can be seen in Fig. 4, the temperature dependent thermal conductivity for the 0.6 wt% Mg and variable Si contents during heating and cooling shows minor differences between the results achieved in both cycles. This coincides nicely with the related microstructure investigation. Fig. 6 displays in a comparable way to Fig. 5 the structure along the longitudinal cross section of the sample AlSi6. It is clearly visible that here less porosity has developed and therefore a lower difference between heating and cooling cycle should be reasonable. Coming to an overall evaluation of the measurement campaign, the systematic investigation of temperature dependent thermal conductivity has led to the following general results: In case of AlSiCu+ alloy material the conductivity does not show dramatic changes for fixed Si values with growing Cu content, except for the step from the binary to the multi-component

Figure 5: Microstructure of the multicomponent sample AlSi-15 (6% Si, 3% Cu, 3% others) at upper (top), middle and lower (bottom) position of longitudinal symmetry plane

Figure 6: Microstructure of the ternary sample AlSi-6 (5% Si, 0.6% Mg) at upper (top), middle and lower (bottom) position of longitudinal symmetry plane

181 variants. Typical for this material is the rather significant increase in conductivity from room temperature to approx. 300 °C in the heating cycle, which is less visible during cooling. This does not occur in case of zero copper content, which means the binary AlSi9 alloy sample. • For the AlSiMg0.6 alloy variants a slight increase of the conductivity from room temperature to 300 °C only occurs for low Si values, for higher Si additions a continuous decrease is measured. • For AlSiCu+ and AlSiMg alloys the increase of Si always leads to a decrease in conductivity level for both, a fixed Cu and Mg content. • In those cases where strong deviations between heating and cooling cycle of the measurement could be linked to the occurrence of significant porosity material data only derived from the heating cycle should be introduced into simulation data bases.

5

Acknowledgements

The authors would like to thank the Deutsche Forschungsgemeinschaft DFG, Bonn, Germany for the financial support of the work presented here.

6 [1] [2] [3]

References

A. Ludwig, P. Quested, G. Neuer, Advanced Engineering Materials 3, 2001, 11–14 J. Auchet, J. L. Bretonnet, Rev.Int.Hautes Temper.Refract. 26, 1990,181–192 L. Binkele, M. Brunen, Thermal conductivity, Electrical Resistivity and Lorenz Function Data for Metallic Elements in the Range 273 to 1500 K, Report No. JÜL-3006, Forschungszentrum Jülich, 1994 [4] R. E. Taylor, H. Groot, et al., High Temp.-High Press. 30, 1998, 269–275 [5] H. Szelagowski, R. Taylor, High Temp.-High Press. 30, 1998, 343–350 [6] Overfelt R A, Bakhtiyarov S I, Taylor R E, 2002, High Temp.-High Press. 34 401–409 [7] J. Blumm, J. B. Henderson, L. Hagemann, High Temp.-High Press. 30, 1998,153–157 [8] M. Rappaz et al., in: Modelling of Casting, Welding and Advanced Solidification Process VII, Eds. M. Cross and J. Campbell; The Minerals, Metals & Materials Society, 1995, 449–457 [9] R. Brandt, G. Neuer, Advanced Engineering Materials 5 No. 1–2, 2003, 52–55 [10] R. Brandt, G. Neuer, presented at the ECTP 2005 (http://www.ectp.sav.sk/) to be published in High Temp.-High Press

182

FEM Simulation of Near Net Shape DC Billet of High Strength Al-Mg-Si Alloy H. Nagaumi1, Y. Takeda1, Suvanchai P.2, T. Umeda2 1

Nippon Light Metal Company Ltd., Kambara, Japan Chulalongkorn University, Bangkok, Thailand

2

1

Introduction

Aluminum casting and forging products are used for aircraft, vehicles, industrial machinery, precision machine, etc. Lighter automobiles are increasingly demanded in order to improve fuel efficiency and drivability. Many of suspension components require superior and reliable properties such as high strength and toughness. Aluminum forging, with its high strength and ductility, is a promising alternative for such suspension parts since it can result in weight reduction of up to 40 % (compared with conventional steel parts). The high deformability of aluminum forged parts is also highly suitable for maximum safety. But the higher cost of aluminum forgings has –up until now- been a barrier to fully capture the growing market for these applications. Generally, in case of the aluminum suspension that is forged using the extrusion round bar, it leads to make the manufacturing cost rise because the homogenization, pre-heating and an extrusion processes are necessary. Recently a new DC casting process of near net shaped ingot (beam blank ingot), which aimed to reduce the manufacturing cost of suspension, was proposed [1, 2]. Since the shape of the near net shaped billet is complicated than usual round billet, many troubles in the casting are originated from the ingot shape. For example, the peculiar problems such as the surface crack and the mold grasped by solidifying billet, which led to the interruption of casting, resulted from the complicate shape and casting conditions. The mechanism of these peculiar phenomena had not yet been clarified. In this study, a coupled flow, solidification and thermal deformation FEM analysis and casting experiments of near net shaped semi-continuously cast billet were carried out in order to clarify the peculiar problems. By comparing the calculated results with the experimental ones, the cause of the surface crack and the mold grasped by the solidifying billet during was clarified.

2

Casting Experiment

The material used in this study is the high-strength Al-Mg-Si alloy that was improved from AA6061, and the chemical composition of the alloy is shown in Table1. The alloy was cast by a lab-scale DC casting machine using the hot-top method. The schematic shape of mold used in this experiment is shown in Figure1, and the casting conditions were shown as follows; Casting temperature of 720 °C, Casting speed of 85–110 mm/min, Water supply of 150–200 l/min.

183

3

Analytical Method

3.1

Flow and Solidification Analysis

In this analysis, the followings were carried out; a coupled flow, solidification and thermal deformation FEM analysis. The geometry used in the study of flow and solidification analysis is as shown in Figure 1. The calculation domain includes billet, mold and bottom block, where the casting length is 500 mm. The flow and solidification analysis was calculated using a commercial solidification package, CAPFLOW. Enthalpy method was used to analyze the latent heat evolution of solidification. Considering the temperature dependency of density, specific heat and thermal conductivity, measured values, described later, were used. The physical properties used in the calculation are shown in Table 2. The thermal boundary conditions at the interface between billet/water-cooled mold, billet/cooling water and billet/bottom block are prescribed using the Newton’s law of cooling assuming h, heat transfer coefficient, as a function of temperature as shown in Table 3.

Figure 1: Schematic geometry and boundary conditions

Table 1: Chemical composition of the Al-Mg-Si alloy (in wt %) Element wt % Si 1.00

Fe 0.18

Cu 0.40

Mg 0.83

Mn 0.37

Cr 0.27

Ti 0.02

Al Bal.

Table 2: Physical properties used in the calculation

Density, kg/m3 Specific Heat, J/(kgK) Thermal Conductivity, W/(m/K)

Billet Al-Mg-Si 2680 See Figure 2 See Figure 3

Bottom Block A6061 2700 1080 150

Mold Cu 8000 385 394

184 Table 3: Relation of heat transfer coefficient and billet surface temperature

T,ºC 0 578 580 650 700

h1, W/m2 K 20 20 1000 2000 2000

T,ºC 200 300 400 582 660

h2, W/m2 K 2000 500 500 10000 20000

T,ºC 50 100 150 200 500

h3, W/m2 K 3000 10000 20000 600 4000

700

4000

h1: Heat transfer coefficient between billet and mould h2: Heat transfer coefficient between billet and bottom block in the center of the billet. h3: Heat transfer coefficient from billet and bottom block to the intruded water in the gap

3.2

Thermal Stress Analysis

In this study, thermal stress during casting was calculated by a commercial structure analysis package, ANSYS. The thermal histories obtained from the CAPFLOW were used as input data to an elasto-plasticity model which simulated the thermal stress and distortion of the billet. The calculation domain includes the billet, mold and bottom block, where mold and bottom block are assumed to be as a rigid body and do not consider their thermal distortion. In order to simulate the thermal distortion of the billet the element type was employed for a large deformation element called as SOLID 45. The effect of friction forces which happened between the billet and the bottom block and also between the billet and the mold, was considered by using the contact element called as SURFACE CONTACT 169 and 171, and the friction coefficient was set to 0.1. A multi-linear isotropic work-hardening law was determined to fit the experimental data. Von Mises criterion law determined the yield criterion. Young’s modulus, linear expansion coefficient and the work-hardening were considered as a function of the temperature and these values were obtained by experiment as shown later. In the solid-liquid coexistence region, an average linear expansion coefficient was estimated from a volumetric shrinkage percentage of 6%.

4

Thermal Properties

Figure 2 shows the specific heat of this alloy was continuously measured from room temperaº ture to 700 C. Solidus temperature (580 ºC), liquidus temperature (554 ºC) and the equilibrium solidification temperature range (74 ºC) of the alloy were obtained by specific heat curve [3]. The relation between fraction solid and temperature calculated by Gulliver-Scheil model using Thermo-Calc and specific heat curve is also shown in Figure 2. The solidification temperatures from liquidus temperature down to solid fraction of 0.8 are almost same for the two methods, but the final stage of solidification is quite different. In this study, the Gulliver-Scheil model was used to evaluate solidification and distortion behaviors. The linear expansion coefficient was continuously measured from room temperature to 600 °C using the vertical type of thermal expansion apparatus. Density at each temperature was calculated from these data. The thermal diffusivity was measured at a specified temperature using the laser flash method. Figure 3

185

Figure 2: Specific heat-temperature curve

Figure 3: Thermal conductivity and linear expansion coefficient as a function of temperature

shows the thermal conductivity which is determined from density, specific heat and thermal diffusivity, and the linear expansion coefficient, as a function of temperature. The thermal conductivity increases from room temperature to solidus temperature and it tends to decrease when the temperature exceeds solidus temperature, whereas the linear expansion coefficient increases with the rise in the temperature.

5

Tensile Testing during Solidification

In this study, tensile tests from room temperature to solid-liquid coexistence region were carried out by using the high temperature tensile test equipment with induction heating. At first, the specimen was heated up to temperature above its liquidus temperature, and then kept for a certain period of time. Next, it was cooled down to the test temperature at cooling rate 1 °C/s. By this means, a solidification structure was obtained. Finally, tensile load was applied at strain rate 10–1 s–1 and the load-displacement was recorded. Figure 4 illustrates the relationship between Young’s modulus and temperature. With the temperature rise Young’s modulus decreases and when the temperature exceeds the solidus temperature (580 °Cȉ, it rapidly decreases as shown in Figure 4. Figure 5 shows the relationship between tensile strength and temperature. The tensile strength decreases with the temperature increases. When the temperature reached 625 ºC, the

Figure 4: Young’s modulus as a function of temperature

Figure 5: Tensile strength versus temperature

186 value of tensile strength became zero, so the temperature 625 °C is ZST (Zero Strength Temperature) of this alloy at which corresponding solid fraction is 0.74.

6

Results and Discussion

6.1

Casting Experimental Results

Typical fracture occurrence position of the billet is shown in RHS of Figure 6. The cracks are well generated at D, E and F position of the billet as shown in Figure 6, and the cracks are generated on the surface and then propagate inside the billet. The surface cracks seem to initiate in the concave parts of the billet owing to constrained contraction. These cracks are different from the round billet cracks which occur in the center of the billet [4–5] and here we call the crack of the near net shaped billet as surface cracks. Figure 6 also indicates the fractured SEM images. The fractured surfaces of the shaped billet have a rupture structure where intergranular fractures with remaining liquid around interdendritic regions were observed, namely internal crack occurs in the solid/liquid coexisting (mushy) state above the solidus. Many intermetallic compounds and precipitates around grains were observed, and it is conjectured that the crack occurs easily because the intermetallic compounds and precipitates become fragile among grains.

6.2

Distortion Profile

The comparison of cast and calculated profiles of the billet with designed mold shape is shown in Figure 7. The distortion profile of cast billet agrees well with the calculated ones. Both legs were shifted inside by the solidification contraction. A neck falling phenomenon due to solidification shrinkage is observed. A comparison between calculated and cast butt-curls was made. The calculated value of but-curl (14 mm) agreed well with the cast result (13 mm). Therefore it is considered that this thermomechanical modeling in this study is appropriate from the comparison between cast and calculated distortion of the shaped billet.

Figure 6: Crack occurrence position and SEM images of the fracture surfaces of the billet

187

Figure 7: Comparison of cast and calculated profile of billet with designed mold shape directions

6.3

Figure 8: Distribution of plastic strain along X and Y the at start-up state

Mechanism of Surface Crack

Figure 8 shows the distribution of plastic strain along X and Y directions obtained from thermal distortion calculation at start-up state. It is easily understood for crack to generate at the D position since the tensile strain of the X direction concentrates in the D position as shown in Figure 8 (a). It is the reason why both legs were shifted inside and contacted with the mold by the solidification shrinkage. Consequently the reaction forces by the mold are generated and the D position receives the tensile stress. It is also likely to generate the cracks at E and F positions because the tensile strain of Y direction concentrate in the E and F positions. A neck falling phenomenon occurred and then the billet contacted to the mold due to solidification shrinkage as shown in Figure 7, so the reaction force by the mold is generated. As a result, the cracks become easy to be generated since the E and F division receive the tensile stress. Crack initiation position predicted by the thermal stress analysis reproduced well the real casting result.

6.4

Cause of the Mold Grasped

The mold grasped by the billet is a peculiar phenomenon always occurred in the shaped billet casting. When the phenomenon occurred, the billet did not moved down with bottom block because the mold was clipped by the billet. Figure 9 indicates the distribution of temperature in the cross section of the billet, (a), and the contact situation between the billet and mold, (b) and (c). Solidification is uneven at each location of the billet as shown in (a). Solidification had finished in the areas such as two legs and the neck areas, but the solidification had not yet finished in the center area. The contraction during solidification forced the two legs to shift toward the mold due to the uneven solidification. The mold, hence, is grasped by the billet and so the contact pressure between the billet and mold is generated as shown in (b), (c). It is considered that the phenomenon of the mold grasped by the billet occurs when the contact pressure becomes bigger and exceeds the drawing force applied in the bottom block.

188

Figure 9: Distribution of temperature in the billet and contact situation between billet and mold

7

Conclusions

A coupled flow and solidification and thermal deformation FEM analysis and casting experiment of near net shaped semi-continuously cast billet were carried out. The causes of the peculiar problems such as the surface crack and the mold grasped by shaped ingot were clarified. The main results are given as follows: 1. The cracks of irregular shaped billet are different from the ones of round billet. They are generated on the surface and then propagate inside the billet. The fractured surfaces of the shaped billet have a rupture structure where intergranular fractures with remaining liquid around interdendritic regions are observed, namely internal crack occurs in the solid/liquid coexisting (mushy) state above the solidus. 2. The surface crack of the near net shape billet is originated from the reaction force which arises by contacting the mold by the distortion of the solidifying billet. 3. The phenomenon of the mold grasped by the billet occurs when the contact pressure becomes bigger and exceeds the drawing force applied in the bottom block.

8 [1] [2] [3] [4] [5]

Reference M. Anderson, R. Bruski, D. Groszkiewicz and B. Wagstaff: Light Metal, 2001, 847–853 K. Sugita, E. Sagisaka, S. Ichikawa: J. P. Patent # 7-67598, July 26, 1995 H. Nagaumi: J. Jpn. Inst. Light Met. 50, 2000, 49–53 I. Farup, J. M. Drezet and M. Rappaz: Acta Meterialia, 49, 2001, 1261–1269 H. Nagaumi, T. Umeda: Journal of Light Metals, 2 , 2002,161–167

189

Numerical Simulation of DC Casting; Two Ways to Interpret the Results of a Thermo-Mechanical Model W. Boender, A. Burghardt, E. P. van Klaveren Corus RD&T, IJmuiden, The Netherlands

1

Introduction

The manufacture of aluminium products starts with the casting of ingots that are subsequently rolled to plate or sheet. Plates and sheets are used for a wide range of products, for example aeroplanes and beverage cans. The most widely used casting technique is direct chill (DC) casting, which is a semi-continuous process. Cold cracking, which is initiated by mechanical stresses at temperatures significantly below the solidus temperature, is one of the problems that impede the efficiency of DC casting. Figure 1 shows two types of cold cracks that may occur in an ingot [1]. In the 1990s, numerical simulations of the mechanical behaviour of an ingot during casting were added to the tools that are used to solve these problems [2]. Here, the interpretation of the results of such simulations to assess the likelihood of cold cracking is discussed. In a thermo-mechanical model, both the thermal and mechanical phenomena in an ingot are simulated numerically. The model developed within Corus RD&T has been successfully employed several times to improve the casting operations in the plants of Corus Aluminium. It was set up using the finite element package MSC.Marc. To illustrate how the results of a thermo-mechanical model can be interpreted, the casting of a 2000u510 mm2 ingot of an Al - 4.5% Cu alloy at a casting speed of 60 mm/min was simulated. This alloy was chosen, as it is a representative of high-strength alloys. The input of the model consists of the dimensions of the mould, the ingot and the bottom block; the process parameters for the drop; the physical properties of the ingot and the bottom block; and the boundary conditions. Because of symmetry, only a quarter of the ingot has to be simulated. Figure 2 shows a perspective drawing of this quarter. The length of the ingot was 2000 mm in the simulation, which gives a complete picture of the start-up. The bottom block was made of steel. The flow rate of the cooling water per unit length of the mould opening and the drop rate were similar to the ones Hannart et al. [2] had used for casting AA2024. Information on the physical properties of the Al - 4.5% Cu alloy and of the material of the bottom block can be found in [3]. From top to bottom, the following areas are distinguished at the vertical sides: the primary cooling zone, the air gap and the secondary cooling zone. Direct contact between ingot and mould occurs in the primary cooling zone. A falling film of water cools the ingot in the secondary cooling zone. In the air gap between these two zones, the ingot is neither in contact with the mould nor with the cooling water. At the butt of the ingot, there may either be direct contact with the bottom block; a narrow gap between them filled with air and vapour; or a wide gap filled with cooling water. Based on the local distance between the butt and the bottom block, the model determined which situation occurred. See [3] for more information on the boundary conditions applied.

190

Figure 1: Types of cold cracks

2

Figure 2: Mould, ingot and bottom block

Normal Stresses

Figure 3 shows the temperatures inside the ingot at a cast length of 1.88 m. It covers the same quarter of the ingot as shown in Figure 2. Figure 4 shows, for the same instance, the normal stresses in the x-direction, i.e. the stresses perpendicular to the casting direction and parallel to the rolling face. The picture on the left-hand side shows the inside of the ingot and the picture on the right its exterior. Tensile stresses are positive and compressive stresses are negative. There are large tensile stresses in the lower interior part of the ingot, and there are large compressive stresses at the lower part of the rolling face. The shell of the lower part of the ingot has already reached the temperature of the cooling water. However, the interior of the lower part is still hot. As the interior continues to cool down, it shrinks, which the rigid shell counteracts. The stress state in the lower part of the ingot results from this balance between tensile stresses in the interior and compressive stresses in the shell. For the y- and z-direction, similar stress distributions are obtained [3].

Figure 3: Temperatures

Figure 4: Normal stresses in x-direction

191 Only a few measurements of stresses in a billet [4, 5] or an ingot [5, 6] have been reported. The results above are in line with their findings; tensile stresses exist in the interior and there are compressive stresses below the surface. Stress is a symmetric 3u3 tensor, which contains three normal stresses and three shear stresses. In the solidified part of the ingot, none of the three normal stresses is negligible. Therefore, a real triaxial state of stress exists. A consequence of the symmetry of the stress tensor is that, at each point, three mutually perpendicular directions exist along which there are no shear stresses. These directions are called the principal stress directions, and the corresponding normal stresses are called the principal stresses U1, U2 and U3, where U1 t U2 t U3. To gain more insight into the stresses within an ingot, the principal stresses and their directions were evaluated. This approach also enables an assessment of the maximum normal stress criterion, which is assumed more applicable to brittle materials, like gray cast iron, than the other theories [7, 8]. This criterion states that failure occurs when either the maximum principal stress, U1, reaches the uniaxial tensile strength or the minimum principal stress, U3, reaches the uniaxial compressive strength [9]. Figure 5 shows the directions of the major principal stress for the ingot's interior, i.e. the wide symmetry plane, and its surface, i.e. the rolling face. The major principle stress will equal U1 if U1 is greater than the absolute value of U3, else it will equal U3. The major principal stresses are directed horizontally in the centre region. Going towards the narrow side, these tensile stresses turn more and more to the vertical direction. Griffith postulated that a crack would open up in the plane normal to the direction of maximum stress [10]. Figure 6 shows the values of the major principal stresses for the interior and the surface of the ingot. Assuming cracks only occur in regions of large tensile stresses, and that they propagate perpendicularly to the tensile stress directions, one can conclude the following. Areas with a high probability of crack formation are found about 100 mm below the surface of the narrow side or a few 100 mm above the butt. The cracks near the narrow side would then propagate following a path that resembles J-cracks. Those near the bottom block would propagate upwards like trouser cracks. The observation of the directions of the major principal stresses, therefore, led to hypotheses for the start locations and the directions of trouser and J cracks.

Figure 5: Directions major principal stress

Figure 6: Major principal stresses

192

3

Fracture Mechanics

The results of others [11, 12] for the maximum principal stress when casting AA7050 can be compared with the tensile strength that they had measured. The calculated stresses appear to be less than the tensile strength. This finding agrees with the experience of the authors. The fact that cold cracks nevertheless occur is, therefore, due to the presence of stress raisers. To determine the effect of the size of an imperfection on an ingot's mechanical behaviour, fracture mechanics was applied. An existing crack propagates only if the total energy of the system decreases [13]. Following this energy approach, the maximum size of an imperfection that can be tolerated under a certain loading can be determined. For a through-thickness crack, which is ideally sharp, in an infinite plate under remote tensile loading, a simple expression was derived:

Vc ac

KIc Sa

(1)

KIc2

S V2

where KIc is the plane strain fracture toughness, V is the tensile stress, a is half the length of the crack and the subscript c stands for critical. Cracks longer than 2ac grow unstoppably. According to Chang and Kang [11], the minimum value of KIc for AA7050 in the as-cast condition at room temperature is 8.54 MNm–3/2. The fracture toughness of AA2024, which solidified at 0.77 °C/s, is about 14.63 MNm–3/2 [14]. Using these values and equation (1), the critical crack lengths at a few tensile stresses were calculated to demonstrate the application of fracture mechanics. The selection of the stresses is based on the tensile stresses in Figure 6. Table 1 shows the results. These guesses at the critical crack lengths can be compared with sizes of defects in an ingot. For example, Gauckler et al. [15] mentioned the sizes of a few types of inclusions in aluminium. Especially oxide skins, whose maximum size may be 5 mm, may be greater than the critical crack length. Hence, they may cause cracks in areas with high stresses. Table 1: Critical crack lengths for AA7050 and AA2024 Stress Mpa 50 100 150 200 250 300

Half crack length AA7050 mm 9.29 2.32 1.03 0.58 0.37 0.26

AA2024 mm 27.25 6.81 3.03 1.70 1.09 0.76

193

4

Conclusions

Two methods were used to interpret the results of a TM-model. The application of the major principal stresses demonstrated where cracks might be initiated, and how they might propagate. The application of fracture mechanics produced guesses at the critical crack length, i.e. from which minimum size cracks would grow. Internal defects, like voids and inclusions, should be smaller than this critical size. If imperfections are distributed homogeneously in an ingot, a large volume with high tensile stresses will contain more of them than a small volume. Hence, in the former case, it is more likely that a crack will grow than in the latter. Therefore, to avoid cracks it is not only necessary to lower the stress levels but also to decrease the sizes of volumes with high stresses. The thermo-mechanical model is a tool to find ways to achieve this.

5 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

[11] [12] [13] [14] [15]

References J. Du et al. in Light Metals 1998 (Ed.: B. Welch), TMS, Warrendale, PA, USA, 1998, p. 1025 B. Hannart, F. Cialti, and R. van Schalkwijk in Light Metals 1994 (Ed.: U. Mannweiler), TMS, Warrendale, PA, USA, 1994, p. 879 W. Boender et al. in Light Metals 2004 (Ed.: A.T. Tabereaux), TMS, Warrendale, PA, USA, 2004, p. 679 A. Roth, M. Welsch, and H. Röhrig, Aluminium 1942, 24, p. 206 G. Seeger, Gießerei 1951, 38, 14, p. 325 S.A. Levy, R.E. Zinkham, and J.W. Carson in Light Metals 1974, (Ed.: H. Forberg), The Metallurgical Soc. of the AIME, New York, NY, USA, 1974,Vol. 2, p. 571 J.H. Faupel and F.E. Fisher, Engineering Design, (New York, NY, USA: John Wiley & Sons, Inc., 1981), p. 242 N.E. Dowling, Mechanical Behavior of Materials, Prentice Hall, New Jersey, USA, 1999, Chapter 7, p. 256 S. Timoshenko, Strength of Materials, D. van Nostrand Company Inc., Princeton, NJ, USA, 1958, Part II, p. 444 G.C. Sih in Mechanics of Fracture 1 - methods of analysis and solutions of crack problems (Ed.: G.C. Sih), Noordhoff, Leiden, The Netherlands, 1973, Introductory chapter, p. XXIX K-M. Chang and B. Kang, Journal of the Chinese Inst. of Engineers 1999, 22, 1, p. 27 H-M. Lu, K-M. Chang and J. Harris in Light Metals 1997 (Ed.: R. Huglen), TMS, Warrendale, PA, USA, 1997, p. 1091 M. Janssen, J. Zuidema, R.J.H. Wanhill, Fracture Mechanics, Delft University Press, Delft, The Netherlands, 2002, p. 3 R.K. Paramatmuni et al., Material Science and Engineering 2004, A 379, p. 293 L.J. Gauckler et al., J. of Metals 1985, 37, 9, p. 47

194

Continuous Casting of Hypermonotectic AlBiZn Alloys: Experimental Investigations and Numerical Simulation M. Gruber-Pretzler1, F. Mayer1, M. Wu1, J. Moiseev2, B. Tonn2, A. Ludwig1 1

University of Leoben, Leoben, Austria Clausthal University of Technology, Clausthal-Zellerfeld, Germany

2

1

Abstract

Hypermonotectic alloys separate into two melts prior to solidification. Therefore, continuous casting of hypermonotectic alloys is still a challenge that inhibits the use of these bearing materials in industry. As the secondary phase distribution has a major impact on the wear properties of the final product, it is of utmost importance to understand its formation. The motion of the secondary phase droplet during demixing of the melt is caused mainly by Marangoni transport and gravity-induced sedimentation. By a series of experiments with hypermonotectic AlBiZn alloys, a laboratory scale strip casting device has been adapted to the special behavior of these alloys. The resulting droplet distributions were experimentally investigated by varying not only the alloy composition but also relevant process parameters. Beside this investigation the formation of secondary phase distribution was also modeled with a two-phase volume averaging approach. The Bi-enriched minority liquid phase is treated as the second phase whereas the parent melt as first. The model solves the mass, momentum, species and enthalpy conservation equations for both liquids including droplet nucleation, solute redistribution and monotectic reaction.

2

Introduction

Alloys with a miscibility gap in the liquid state, especially for those with gross concentration above the monotectic point (hypermonotectics), are potential bearing materials for the automotive industry [1, 2]. As hypermonotectic alloys separate into two melts prior to solidification their continuous casting is still a challenge. Difficulties arise from the fact that the secondary phase droplets move due to Marangoni forces and gravity-induced sedimentation. After a period of less attention, the interest on the solidification of hypermonotectic alloys has increased again, because the EU has proscribed the use of Pb containing alloys. Therefore, the development of new Pb-free bearing materials is of great importance and so alloys based on Al-Bi are again of particular interest for materials research and development [3, 4]. Strips of hypermonotectic alloys are produced at the TU Clausthal on a vertical continuous casting unit at a laboratory scale. This unit insures the possibility of high cooling rates that are necessary for the casting of hypermonotectic alloys in order to cross the demixing interval very rapidly. Due to consequent prevention of minority-phase sedimentation, a fine distribution of Bi drops in the Al matrix can be produced. However, the influence of additional alloying elements, like Zn, or of differences in process parameters, like casting speed, casting temperature and cooling conditions on the distribution of the minority phase is still open and thus addressed in the present paper.

195 Theoretical descriptions of the demixing of hypermonotectic Al-based alloys were already suggested in [5–6]. A numerical treatment of this process was presented by two of the authors in 2003 [7, 8]. Their work addressed the influence of Marangoni force and of gravity-induced droplet sedimentation on the droplet distribution. However, the studies presented in [7, 8] were done for 2D square geometry which was assumed to be cooled from all sides. In the present study, a first numerical description of a strip casting process for the solidification of a hypermonotectic binary AlBi10 alloy is given. The temperature field and the distributions of droplet fraction and density are presented and discussed. This paper consists of two parts, the first is the experimental part and the second the numerical part.

3

Experimental Part

3.1

Description of the Continuous Casting Unit

The vertical continuous casting machine at the TU Clausthal consists of the mold with pouring nozzle and tundish and the mechanical-hydraulic strand withdrawal mechanism. The ready-assembled continuous casting copper mold is made out of two side frames and two spacers. The spacers confine the strand laterally and keep the mold gap at a constant width of ' = 10 mm. A steel pouring nozzle is set into a notch at the top of the mold and a tundish symmetrically on top of that. Figure 1 illustrates schematically the assembled mold ready for use. The tundish (1) and the steel pouring nozzle (2) are pre-heated up to the pouring temperature. The melt is heated up in a separate furnace also to that temperature and then poured into the tundish in order to start the casting. There are two cooling water circuits. In the first circuit the cooling water is flowing through the cooling water supply (5) towards the cooling channel (6) where the water directly impinges the strand (7). This cooling channel acts as secondary cooling, while the upper part of the mold operates as primary cooling. In each half of the mold there are a total of 17 cooling channels and associated cooling grooves. With the cooling circuit (4) an additional heat removal from the

Figure 1: Scheme of the middle cross section through the mold (detailed information in the text and see ref. [9])

Figure 2: Bi droplets distribution in an AlBi8Zn6 strand

196 mold is realized. The amount of water flowing through the two cooling water circuits can be varied depending on the required cooling intensity. With the maximum cooling possible a temperature gradient at the solidification front of G = 50 K/mm [3]. Steady-state casting conditions are found to be reached after a few seconds. Therefore, a uniform phase distribution along the length and width of the strand can be established.

3.2

Evaluation and Discussion of Results

In the primary cooling zone which extends over the first L = 23 mm of the copper mold, the melt is cooling rapidly. When the temperature falls below the binodal temperature, droplets of the minority phase start to form. From the experimental point of view, it is obvious to assume that the main nucleation happens in areas close to the walls. Further growth the nucleated droplets is then governed by diffusion. Due to the Marangoni effect, which is caused by the temperature-dependence of the interfacial tension, the continually growing particles migrate towards the middle of the strand. This favors their collision and coagulation. The Marangoni force is proportional to the temperature gradient and proportional to the square of the droplet diameter, whereas the Stokes forces increases only linearly with the diameter. The occurring droplet motion is the overlay of casting flow, Marangoni motion and sedimentation. Figure 2 gives a qualitative impression of the Bi-droplet distribution in the strand of AlBi8Zn6. The structural constituents occurring in this case were produced at a casting velocity of Vcast = 450 mm / min, a casting temperature of about Tcast = 960 °C and a cooling water flow rate of about Q = 4000 l/hour for each mold side frame. In all ternary Al-Zn-Bi alloys the edge strand zones are depleted in Bi-droplets up to a depth of 4 mm. In the Al-Bi system the monotectic composition occurs at cm = 3.4 wt.% of bismuth. This composition corresponds approximately to the contents measured in the edge zones. In opposite, the core zone of the strand is enriched in droplets and reveals around twice as much as Bi content compared to the initial composition. Figure 3 to 6 give an overview on the influence of variations in alloy composition and process parameters like casting speed, casting temperature or cooling water flow rate on the maximum and the mean Bi-droplet size. Figure 3 shows how the mean and maximum droplet size in the enriched middle zone by varying the Zn content for the two alloys, AlBi6Znx and AlBi8Znx. Figure 4 shows the influence of the casting speed on the mean and maximum droplet size for the AlBi8Zn6 alloy. It is visible that increasing Zn content increases the Bi-droplet size to a certain level for both, mean and maximum values. However, if the amount of Zn in the alloy is larger than 10 wt.% the droplet size decreases again. On the other side, high Bi amount leads to a clear increase of the droplet size. The influence of the Bi-content on the droplet size is substantially stronger than the one of Zn and higher radial components. If a casting is desired which reveal only small size maximum droplets, the optimum casting speed for an AlBi8Zn6 alloy would be around Vcast = 550 mm/min. This finding is true for a casting temperature of Tcast = 960 °C and a cooling water flow rate of Q = 4000 l/hour. A small casting speed would favor the drop growth already in the steel feeder. Evidently, a larger casting speed would lead to a deep melt pool and thus higher temperature gradients in front of the solidification front. This increases the importance of the Marangoni convection and hence influences the appearing droplet distribution.

197

Figure 3: Influence of Zn- and Bi-content on the mean and maximum Bi-droplet size in AlBi8Znx and AlBi6Znx

Figure 4: Influence of casting speed on the mean and maximum Bi-droplet in AlBi8Zn6

Figure 5: Influence of quantity of cooling water on the mean and maximum Bi-droplet in AlZn6Bi8 alloy

Figure 6: Influence of casting temperature on the mean and maximum Bi-droplet in AlZn6Bi8 alloy

Figure 5 shows the influence of the cooling water flow rate on the maximum and mean Bidroplet size. For alloys which should reveal only small size maximum droplets, the optimum cooling water flow rate is around Q = 2000 l/hour for Tcast = 960 °C and Vcast = 550 mm/min. A smaller flow rate/cooling rate favors the droplet sedimentation, whereas a higher flow rate/ cooling rate Marangoni motion is encouraged. Figure 6 shows the influence of the variation of casting temperature on the maximum and the mean Bi-droplet size. It is obvious that small casting temperatures lead to smaller temperature gradients ahead of the solidification front and thus higher viscosity of the parent melt. Therefore, the droplet motion is made difficult.

198

4

Simulation Part

4.1

Model Description

Because of the fact that in the present paper process modeling is of importance we skip a detailed model description. The reader is referred to the original papers of corresponding authors [5, 6]. Nevertheless, a brief description of most important model assumptions is given. The used two phase model considers the parent alloy as the first, L1, and the forming droplets as the second phase, L2. During the monotectic reaction the monotectic matrix is transformed directly from L1. Therefore, the solidification of monotectic matrix is modeled by increasing the L1 viscosity and releasing latent heat on reaching the monotectic temperature. The decomposed L2 droplets approaching the monotectic reaction front are modeled to be entrapped in the monotectic matrix by applying the same enlarged viscosity at or below the monotectic point. A similar approach is used by [1,6,10–12]. The model is based on solving the mass, momentum, species and enthalpy conservation equations for both liquids including droplet nucleation, solute redistribution and monotectic reaction [7]. Both Stokes and Marangoni motion is taken into account. In addition to the above mentioned model considerations the following assumptions are made: • • • •

Gravity-induced sedimentation is modeled with the Boussinesq approach. Both liquid phases are assumed to have the same viscosity. Collision and coagulation of droplets are not yet taken into account. Diffusion in a single droplet is thought to be infinite.

4.2

Geometry and Boundary Condition

For the process simulation of a binary AlBi10 alloy a casting velocity of Vcast = 828 mm/min and a casting temperature of Tcast = 792 °C is considered. Due to the geometry of the casting a 2D symmetry has been chosen for the simulation. The mold is schematically shown in Figure 1 and in a 3D view in Figure 7. Here (c) indicates the copper mold, (d) indicates the used steel pouring nozzle on the top, (e) and (f) show the primary cooling zone, where (e) indicates the part of the mold where ideal contact is assumed and (f) the lower part where already solidification shrinkage and thermal contraction of the solid shells is considered. (g) indicates the position of the secondary cooling zone. In Figure 8 the applied boundary conditions are shown. (c) gives the position of the inlet, where a velocity inlet with the casting speed Vcast is considered. A heat transfer coefficient (HTC) of h = 100 W/m²K and a temperature of Tfeed = 682 °C is considered for the steal feeder (d). (e) is divided in two parts: the upper part where h = 1000 W/m²K and Tmold = 202 °C and a lower part where h = 100 W/m²K and Tmold = 202 °C is used in order to model the start of the secondary cooling. For the rest of the secondary cooling (f) we have applied h = 10000 W/m²K and Twater = 23 °C. For the outlet (g), outflow is considered. A grid of 15960 cells and 16523 nodes is used. The applied material properties are the same as used in [7]. Tab. 1 gives the values for the used thermodynamic properties of the system. As initial conditions, we start with hot melt (Tinit = 792 °C) at rest (Vinit = 0 m/min). Cooling and inflow are then resulting in an acceleration

199 of the melt in the strand, the formation of a solidifying shell and nucleation and growth of Bidroplets.

4.3

Results and Discussion

Figure 9 shows three properties in the upper part of the casting, namely temperature field, volume fraction and density of the Bi-droplets 44s after switching on flow and cooling. The pictures are overlaid with two isothermal lines: binodal (Tb = 786 °C) and monotectic (Tm = 657 °C). Table 1: Thermodynamic phase diagram used for the simulation Monotectic temperature Monotectic concentration L2 monotectic concentration Critical Temperature Melting point of Al Melting point of Bi Gross concentration Slope of liquidus at c0 Partitioning coefficient

Tm cm cL2 Tc TfA TfB c0 m k

930 K 0.47 at.% 83.4 at.% 1310 K 933 K 543 K 1.415 at.% 148.1 K per at.% 51.72

657 °C 3.526 wt.% 97.493 wt.% 1037 °C 660 °C 270 °C 10 wt.% 20.42 °C per wt.% 9.55

The temperature field displayed in Figure 9a shows the temperature evolution from the upper border of the steel feeder (c) and to the end of the copper mold (d and e). According to the experimental conditions, a gap between casting and copper mold is considered 15 mm below the steel feeder (e). This gap is indicated in Figure 9 by the white gap between mold and casting. The secondary cooling zone starts with that gap and moves on fare below the copper mold. t = 10 s after switching on cooling und inflow, the temperature distribution reaches steadystate. The complicated boundary conditions applied for the considered strip caster lead to a temperature distribution (Figure 9 a) in the solidifying strand that has a strong temperature gradient

Figure 7: Casting Mold (described in the text)

Figure 8: Geometry and boundary conditions (described in the text)

200 in the upper part of the copper mold (d). As the Marangoni force is proportional to the temperature gradient, the droplets which form along the wall are forced to move towards the centre of the strip (Figure 9 b). The movement of the droplets towards the centre of the strand is overlaid by Stokes motion and is therefore at the origin of two small vortices which occur at the wall (f). Figure 9c shows the density, n, of the Bi-droplets. The black curve overlaid is the value of n along the strand at a depth of 25 mm (n = 0 is at the center of the figure). Due to the small temperature gradient in the steel feeder nucleation appears visible slightly below the binodal temperature (Figure 9c). It can be seen that the nucleation of the droplets is somewhat oscillatory in nature (g). At present, the origin of these oscillations is unclear. Zhoa [5] mentioned that oscillations appear in numerical calculations for describing the microstructure evolution in hypermonotectic alloys. The reasons therefore are changes in the supersaturation of the liquid caused by Marangoni and Stokes motion. Detail studies on that issue are ongoing.

5

Concluding Remarks

Figure 9: c steel pouring nozzle, d (ideal contact) and e start of the secondary cooling (air gap). (a) Temperature field with a strongest gradient in d; (b) volume fraction: f indicates the position of high Marangoni motion towards the centre of the cast; (c) density of droplets: g indicates oscillatory in the droplet density. The black curve overlaid is the value of droplet density along the strand at a depth of 25 mm.

First results of investigations with ternary Al-Bi-Zn alloys have been presented in the first part of the paper, intended mainly to examine the influence of additional alloying elements, like Zn, and various process parameters. The results intend to show the following facts: • Zn content variation has minor effect on the maximum and the mean Bi-droplet size compared to Bi-content variations.

201 • For the alloy AlBi8Zn6 the casting conditions which would result in the smallest maximum droplet size for Tcast = 960 °C have been found to be: a casting velocity of Vcast = 550mm/min combined with a cooling water flow rate of Q = 2000 l/hour. For lower casting temperature the maximum droplet sizes can even been reduced. • Further investigations are planed to evaluate the influence of different grain refiners, especially based on Al-Ti-C and Al-Ti-B, on the fineness of the minority phase distribution. In the second part of the paper preliminary simulation results for a strip casting of AlBi10 have been presented. The complexity of the demixing process of hypermonotectic alloys together with the different boundary conditions particular for the considered caster make the qualitative interpretation of the results difficult. However, the following statements can be made: • The temperature field shows the highest gradient in the upper part of the copper mold. • The Marangoni force increases with increasing temperature gradient which causes the motion of the Bi-droplets towards the centre of the strip. • Oscillating droplet nucleation appears spatially in the strand. At present the reason for that is not clear. Maybe changing supersaturation in the liquid caused by Marangoni motion and Stokes sedimentation are of importance. Detail studies on that issue are ongoing.

6

Acknowledgements

This work was financially supported by the ESA-MAP project ‘Solidification Morphologies of Monotectic Alloys-MONOPHAS’.

7 [1] [2]

References

L. Ratke, S. Diefenbach, Mater. Sci. Eng. 15, 1995, p. 263 B. Predel et al, Decomposition of Alloys : The Early Stages, ed. Walter H. U., Ashhy M. F., Berlin: Springer, p. 517 [3] J. Moiseev, S. Vogelgesang, H. Zak, H. Palkowski, B Tonn, Metall 58, 2004, p. 289 [4] J. Moiseev, H. Zak, H. Palkowski, B. Tonn, Aluminium 81, 2005, p. 92 [5] J. Zhao, L. Ratke, Scripta Mater. 39, 1998, p.181 [6] J. Zhao, L. Ratke, Z. Metallkunde, 89, 1998, p. 241 [7] M.Wu, A. Ludwig, L. Ratke, Modell. Simul. Mater. Sci. Eng. 11, 2003, p. 755 [8] M.Wu, A. Ludwig, L. Ratke, Metall. Mater. Trans. 34A, 2003, p. 3009 [9] B. Prinz, DE Patent 40 03 018 A1, 1991 [10] J. Alkemper, L. Ratke, Z. Metall. 85, 1994, p. 365 [11] S. Diefenbach, Ph. D. Thesis Ruhr-Univercity Bochum, 1993 [12] L. Ratke et al., Materials and Fluid Under Low Gravity, ed L. Ratke et al. Berlin: Springer, 1995, p. 115

202

Numerical Simulation of the Upward Continuous Casting of Magnesium Alloys A. Landaberea1, P. Pedrós1, E. Anglada1 and I. Garmendia1 1

INASMET Foundation, San Sebastian, Spain

1

Abstract

The continuous casting of magnesium alloys in vertical upward direction is a novel technology which can be employed for the production of semi-finished materials circumventing the main disadvantages of using conventional casting processes since the risks of burning and explosion are practically eliminated. The present investigation deals with the simulation of the upward continuous casting of round billets of magnesium alloys. The equations for the flow field with heat transfer are numerically solved by a finite volume method and the solidification is accounted via an enthalpy-porosity formulation where the mushy region is modeled as a pseudo porous medium. The obtained temperature distribution is then used as input for a thermo-mechanical analysis to determine the stress field in the billet during the casting process. Several configurations have been simulated and comparison of computed results with available experimental data is provided.

2

Introduction

Continuous casting is the key technology for cost effective manufacturing of semi-finished materials for further processing to obtain high quality final products and has been the state of the art in steel and aluminum processing for decades. However, due to the specific properties of magnesium alloys, i.e. tendency to burn of the melt when exposed to air, the danger of explosion when the melts come into contact with water and the low strengths at temperatures above 300 ºC, the application of conventional gravity continuous casting techniques with magnesium alloys is only possible assuming high risks. In order to prevent these difficulties, the Upward Direct Chill (UDC) casting principle has been proposed at IW – University of Hanover for magnesium alloys where the continuous casting is operated vertically against the gravity. Initial works have been further developed within the EU – 5th Framework funded “EuroMagUpCaster” project (see [1] for further information). Over the years, several mathematical models have been developed to help to understand the processes occurring during continuous casting most of them focused on steel and aluminium alloys (see for example [2, 3]). As part of the “EuroMagUpCaster” project, a model has been developed to study the UDC process which is described in the present article.

203

3

Model Description

The model calculates the temperature distribution and the associated thermal stresses in the billet during the upward casting process. To do this, a combination of commercial together with self developed codes has been adopted. Thus, the commercial Computational Fluid Dynamics (CFD) code FLUENT [4] has been used to calculate the flow field, temperature distribution and solidification of the billet. FLUENT is based on the Finite Volume Method (FVM) and uses an enthalpy-porosity method for modelling the solidification. The thermo-mechanical analysis has been performed with MSC.Marc [5] which is a general purpose code based on the Finite Element Method (FEM) with well recognized capabilities for nonlinear problems. A series of translation codes have been developed to pass the mesh and thermal solution from FLUENT to MSC. Marc so the temperature history at each node is applied as thermal load in the stress analysis.

3.1

Computational Domain

Due to the geometry and configuration of the casting process, the problem can be considered as axially symmetric and simplified to a two dimensional formulation. The computational domain for the thermal problem is sketched in Figure 1 and includes both the mold and the billet. In the results presented in this work, the mold has not been considered for the mechanical calculations. FLUENT considers the horizontal axis as the axial direction and the vertical as the radial. According to this, the gravity vector has been specified to act in the negative horizontal direction. The domain has been meshed with quadrilateral cells in FLUENT which convert to linear quadrilateral elements in MSC.Marc

Figure 1: Computational domain

3.2

Boundary Conditions

Regarding the flow and thermal problem, at the inlet section a uniform velocity and temperature are imposed according to the casting process parameters. In the mold, convection to the primary water cooling loop is assumed with a heat transfer coefficient of h = 4500 W/m2/K and a water temperature of 20 ºC. At the external surfaces of the mold natural convection to ambient air (h = 10 W/m2/K, T = 20 ºC) is applied. After the mold, a combination of natural convection

204 (h = 10 W/m2/K) and radiation (H = 0.8) to ambient air (T = 20 ºC) is applied at the surface of the billet. In order to include the effect of gap formation between the billet and the mold due to shrinkage during solidification and further cooling a contact resistance Rc has been applied at the interface between the billet surface and the mold. This contact resistance is assumed to vary with the liquid volume fraction being zero (perfect contact) when the billet is in liquid state and Rc when the billet has solidified. The value of the contact resistance has been modified until the results of the model have been adjusted to the experimental data provided by IW. At the outlet section, a mass flow balance is applied. Afterwards, the temperatures computed with FLUENT are used as input to MSC.Marc for the stress analysis where an essentially unconstrained billet is assumed since, together with the symmetry conditions along the axis, only the longitudinal displacement of the node in the axis located at the outlet section has been restricted.

3.3

Material Properties

The model uses thermally dependent physical properties of the materials for both the thermal and mechanical problems. The properties have been obtained mainly from [6, 7]. The model assumes a elasto-plastic behavior of the billet material during solidification. It must be pointed out that the mechanical properties above 300 ºC are not available so they have been extrapolated to nearly null values at solidus temperature. In addition, the thermal expansion coefficient for temperatures above the liquidus temperature has been assumed to be zero to account for the liquid phase in the model.

4

Results

The model has been applied to the stationary continuous casting of 90 mm diameter billets of pure magnesium and AZ80 alloy. Results are presented hereafter. 4.1

Pure Magnesium

The computed temperature and streamline distributions for pure magnesium at a casting speed of 95 mm/min are presented in Figure 2. As expected, the billet cools down in the mold due to the action of the water cooling and some reheating occurs at the surface of the billet after the mold due to the heat stored in the central region and to the lack of a secondary cooling unit. The temperature at the surface of the billet just after the mold is around 550 ºC. A large recirculation generates after the step of the melt distributor due to the abrupt change of section. Figure 3 compares the shape of the solidification front obtained in casting trials with that predicted by the simulation after final adjustment of the model showing a good correlation. The model predicts that the billet solidifies before the melt reaches the mold insert. This is in agreement with the surface cracks generated on the billet (see Figure 3) which can be attributed to the stepwise filling and solidification of material in this zone. As can be observed in the temperature distribution, the corner region between the melt distributor and the mold insert is refrigerated by the water cooling so a better insulation between these parts should be introduced to avoid

205

Figure 2: Pure magnesium. Contours of a) temperature (ºC) and b) streamlines

Figure 3: Pure magnesium. Experimental and computed solidification front and billet surface quality

the creation of the surface cracks. Several simulations have been performed with different casting conditions to try to solve this problem and modifications to the casting assembly have been proposed which are currently under preparation. The contours of circumferential stresses developed in the billet together with the deformed shape are given in Figure 4. The billet exhibits a contraction in length and in diameter due to the solidification and further cooling. In the figure the residual stresses in a cross section in the billet away from the mold is also included. As can be seen, compressive residual stresses appear at the outer zone which turn into traction inside the billet and again into compression near the axis. Some oscillations are originated in the distribution due to numerical difficulties in the solution coming from the highly non linear properties of the material and the very low introduced resistance at high temperatures. Figure 5 shows the radial displacement at the billet surface in the mold region. Initially, a dilatation of the billet is predicted which is not constrained since no contact algorithm has been included in the calculations. Later, the billet shrinks due to the solidification and further cooling with a radial contraction of 0.113 mm at the outlet of the mold which gives an overview of the size of the gap between the mold and the billet surface.

206

Figure 4: Pure magnesium. Contours of circumferential stress (Pa) in the billet and radial distribution of circumferential stress at a section away from the mold

Figure 5: Pure magnesium. Radial displacement (m) of the billet surface in the mold region

4.2

AZ80

Figure 6 shows the temperature distribution and streamlines computed for AZ80 alloy at a casting speed of 60 mm/min. The temperature contours follow a similar distribution as in the case

207 of pure magnesium with cooling in the mold region and some reheating after the mold. The temperature at the surface of the billet when it leaves the mold is around 400 ºC, lower than in the previous case mainly due to the lower casting speed so the material has more time to cool inside the mold. In addition, the recirculation zone after the step of the melt distributor has almost disappeared also due to the reduced velocity. The calculated liquid volume fraction is given in Figure 7 where the wider mushy zone of the alloy can be clearly observed. The simulation is compared with the solidification front obtained in casting trials showing also a good agreement. Again, the solidification takes place before the melt reaches the mold insert leading to the generation of the surface cracks observed in the trials (see also Figure 7). Also, from the results of the simulation, it can be said that the material presents a coherency temperature and develops some strength in the mushy zone before reaching the solidus temperature.

Figure 6: AZ80. Contours of a) temperature (ºC) and b) streamlines

Figure 7: AZ80. Experimental and computed solidification front and billet surface quality

208

5

Conclusions

A mathematical model has been developed to analyze the fluid flow, heat transfer and stress generation during the upward continuous casting of magnesium alloys. The predictions have been found to be in good agreement with available experimental data. The model has shown that, for the selected casting conditions, the solidification occurs before the melt reaches the mold insert leading to the formation of cracks at the surface of the billet. The main reason for this effect is the low efficiency of the thermal insulation between the water cooling and the melt distributor. Through a series of simulations, several modifications have been proposed to the casting assembly to overcome this problem and to obtain higher quality billets. The model has proved to be an efficient tool to help in the design and optimization of the casting process. It has been also applied to other magnesium alloys and to an upscaled mold assembly with d = 203 mm for which preliminary casting parameters have been defined.

6 [1] [2] [3] [4] [5] [6] [7]

References Fr.-W. Bach, S. Schacht, A. Rossberg, DGM Conference, Continuous Casting 2005, NeuUlm J.E. Kelly, K.P. Michalek, T.G. O’Connor, B.G. Thomas, J.A. Dantzig, Metall. Trans. A 1988, 19A, 2589–2602 J.M. Drezet, M. Rappaz, Metall. Mater. Trans. A 1996, 27A, 3214–3225 FLUENT´s Manual version 6.2, Fluent Inc., 2005. MSC.Marc 2005, User’s Manual, MSC.Software Corporation, 2005 R.D. Pehlke, A. Jeyarajan, H. Wada, Summary of thermal properties for casting alloys and mold materials, University of Michigan, 1982 Magnesium and Magnesium alloys (Ed.: M.M. Avedesian, H. Baker), ASM International, USA, 1999

209

New Possibilities in the Simulation of Continuous Casting Processes with WinCast-Conti H. Ricken1, C. Honsel2 1 Technische Universität München, Institute of Metal Forming and Casting (utg), Walther-Meißner-Straße, D-85747 Garching 2 RWP GmbH, Am Münsterwald 11, D-52159 Roetgen

1

Abstract

The casting-simulation WinCast was adapted to the continuous casting process. The thermal and mechanical equations are coupled and solved for every time step. Thus the material flow and the shrinkage of the billet can be predicted. The interactive mesh-generator enables the user to change geometry very freely. Boundary conditions and process parameters may be varied in a wide range. This freedom in variation enables the user to change parameters and add special geometrical features to existing or planned system configurations. Special material properties may be integrated into the model. The Institute of Metal Forming and Casting(utg) at the Technische Universität München has developed a test facility to predict heat transfer coefficients against gap length in the mould. Now a lot of parameter studies with the real system can be replaced by simulations. The ability of the user to improve the model is a major step to upgrade production facilities.

Figure 1: Structure of the WinCast-Conti software package

2

WinCast

It was in 2002 when the Institute of Metal Forming and Casting at the Technische Universität München began continuous casting simulations with the RWP software package WinCast.

210 At this time was difficult to calculate the transient heat transfer between strand, mould and cooler. Especially the complex phenomenon the of the gap-building between mould and cooler was yet to be solved. The advantage of freedom in modelling with pentaeder elements and the convenience in setting various boundary conditions revealed the foundations to create a powerful simulation for the modelling of continuous casting processes.

Figure 2: Thermal results

Figure 3: Stress and distortion

211 These attributes were the most important properties of the programm package giving the impulse to develop a special continuous casting module. In 2005 the transport of material during the process and the coupling of thermal and stress calculation were implemented. The transport of matter is calculated not only for molten metal and strand, but also for the cooling water. With the new module there is a direct coupling between the thermal results, the resulting geometrical deviation of the whole system and the changing heat transfer caused by gaps opening and closing. For calculating the complete thermal behaviour of the mould it is also essential to determine not only the cooling and solidification of the melt, but also the heating of the cooling water. With the new WinCast-Conti module it becomes possible to calculate the temperature of the effluent water in dependency on the influent water and the current heat transfer through the

Figure 4. Section of upper and lower coolant pipe

mould. This is a major step to calculate the thermal behaviour and heat conductivity of mould and cooler in an early stage of design. So with the heat balance between strand and coolant the complete thermal behaviour of the heat exchanger can be calculated. With the ease of temperature measurement in cooling water as well as in melt and strand surface outside the mould the consistency between simulation and reality can be proved. The graphic representation of the results provides a comprehensive insight to the process. It is possible to select materials or sections to evaluate special regions of interest. The distortion is represented with a user-defined superelevation. So the user is able to improve the shape and the mounting of the parts relevant to the heat transfer process. It is also possible to export values to a chart for further analysis.

212 At present, utg operates a test facility to quantify heat-transfer coefficients between the different materials used in the heat exchange system. The configuration of this test facility is similar to a horizontal continuous casting machine. In the middle of this system there are uniaxial and quasi-statical conditions of heat transmission. Different contact conditions and gaps are tested as well as special classifications of material. This new database increases the correlation between the new simulation and the reality in industrial casting machines.

213

Modeling Continuous Casting of Metal Wire Rods Chang Hung-Ju1, Hwang Weng-seng1, Chao Long-sun2, Pan Wensen3, Lai Yi Lin3 1

Department of Material Science and Engineering, National Cheng Kung University, Tainan, Taiwan Department of Engineering Science, National Cheng Kung University, Tainan, Taiwan 3 Metal Industries Research & Development Center 2

1

Introduction

Because the electrical industrial is developing fast, the requirements for the metal wire rod’s quality and size become rigid especial for the metal wire of electrical packaging, which is made of metal wire rod. Continuous casting is applied in the production of metal wire rod. The process is similar to that of the steel continuous casting, but the casting speed is slower than steel’s. For controlling the quality, it is necessary to build an analysis model to understand the phenomena of heat transfer and solidification in the continuous casting of metal wire rod. The modeling of continuous casting systems is a problem of a great mathematical and industrial significance. Continuous casting involves many complex physical phenomena and up to now no model can include all of the phenomena at once. In recent years, heat transfer models for continuous casting have been developed [1–3], but most of them studied the continuous casting of steel. This paper is to develop the heat transfer model of continuous casting for producing metal wire rods. The numerical method is the finite difference method. The effective specific heat/enthalpy method is used to handle the release of latent heat. Casting speed and cooling rate are very important variables for the continuous casting and the simulations for different working variables can investigate their effects on the heat transfer behavior of the continuous casting. Simulations can give us important output data, such as temperature distribution and the profile of solidified shell. It can help the industrial to shorten the experiment time and seek the best operation condition.

2

Mathematical Model and Numerical Method

The physical model of the continuous casting machine of metal rods is shown in Fig. 1. The metal in the crucible is in liquid state due to the heat from the heater around the crucible. The liquid metal enters the graphite mold from the bottom of the crucible. The graphite mold is surrounded by the cooling system. Because of the heat extraction of the cooling system, the liquid metal becomes solidified soon after entering the graphite mold. The liquid metal and solid casting move down together with constant casting speed. To build the mathematical model, the following basic assumptions are made firstly: 1. 2. 3. 4. 5.

The system is in the steady state. The system is the axial symmetry. The natural convention effect is ignored since the diameter of the metal rod is small. The casting speed is constant. The thermal conductivity and heat capacity of metal are temperature-dependent.

214

Figure 1: The schematic diagram of continuous casting

The concerned area of this study is the region of fast temperature variation and solid/liquid transformation after the liquid metal enters the graphite mold. Since the length of the graphite mold is much larger than its inside diameter, the metal temperature does not have any significant variation before leaving the graphite mold. Consequently, the graphite mold and the metal inside it are the computing domain. According to the previous statements and assumptions, the governing equation (energy equation) can be written as

U C p eff V

wT wz

1 w § wT · w § wT · ¨k r ¸ ¨k ¸ r wr © wr ¹ w z © wr ¹

(1)

where r is the density, V is the casting speed, k is the thermal conductivity. C peff is the effective specific heat, including the effect of latent heat. The equation is used in the metal and in the graphite mold, the convection term U C peff V wT/wz does not exist in the energy equation. In this study, because the nonlinear effect of the latent heat makes the iterative computations difficult to converge, the transient equation is solved instead to ease off the nonlinear effect. In solving this equation, the effective specific heat/enthalpy method is used to deal with the latent heat. The boundary conditions of the computing domain are given as follows: 1. At the top of metal, the temperature is fixed and is the temperature of the liquid metal in the crucible. 2. At the bottom of metal, there is no temperature variation, i.e., no temperature gradient. 3. At the centerline of metal, it is symmetric. 4. At the interface of metal and graphite, there exists a contact resistance and the heat flux is continuous. The expression of heat flux can be given by qmetal = qgraphite = hmg(Tmetal–Tgraphite)

(2

215 where hmg is the effective heat transfer coefficient whose inverse is the contact resistance. qmetal, qgraphite, Tmetal and Tgraphite are the heat fluxes and temperatures of metal and graphite at the interface. 5. At the outer surface of the graphite mold, an effective convective boundary condition is applied. q = heff(T – Tref)

(3)

where q and T are the heat flux and temperature at the outer surface. heff is the effective heat transfer coefficient which is used to handle the effect of the cooling system. Tref is the reference temperature and here is the water temperature of the cooling system. The effective specific heat/enthalpy method is the combination of the effective specific heat method and the enthalpy method. It owns the advantages of these two methods: fast convergence and good accuracy. At one time step, the relationship of temperature, effective specific heat and enthalpy can be written as

C p eff

en 1  en T n 1  T n

(4)

The numerical method is the finite difference method. The backward difference is used for the time derivative and the centered difference is for the spatial derivatives. The control volume method is utilized to derive the difference equations for both the governing equation and boundary conditions. Because the difference equation is nonlinear, the iteration method is used to solve the difference equations.

3

Results and Discussions

3.1

Verification of Numerical Analysis of Continuous Casting

In the literature, many numerical simulations of continuous casting only put the metal in the computing domain. This study is based on the experiments from Metal Industries Research and Development Centre (MIRDC) in Taiwan. The temperature data available are from the measured points on the surface of the graphite mold. Accordingly, to compare the computing temperatures with the measured ones, it is necessary to include the graphite mold and the metal in the computing domain. In this research, the interface between the metal and the graphite mold is regarded as non-perfect contact, which is handled by applying Eq. (2). According to the MIRDC’s experiments, the working conditions are (1) the metal material is copper, (2) the metal temperature at the inlet of the graphite mold is 1473 K, (3) the height, inside diameter and thickness of the graphite mold are 23.5 cm, 0.4 cm and 0.9 cm, (4) the casting speed is 0.667 cm/sec. The effective heat transfer coefficient (heff) in Eq. (3) is 2700 W/m2.which is estimated by using the experiment data (inlet and outlet temperatures of water and flow rate) from the cooling system. The computing temperature distribution is shown in Fig. 2 and Fig. 3 is the enlarged one near the inlet of the graphite mold. From these two figures, it can be found that the primary temperature variation is along the axial direction and the variation in the radial direction is small. To verify the feasibility of the numerical simulation, the computing results are compared with the experimental ones. Since the temperatures in the metal are not easy to measure, there

216 measured points were set on the outer surface of the graphite mold. Their locations are 1, 6 and 12 cm from the top of the graphite mold. Table 1 shows the computed and measured temperatures of these three locations. The last measured point was arranged at the outlet of the graphite mold. The computed outlet temperature of the metal is 133, while the measured one is 150. From the comparison of these four temperatures, the computed temperatures are close to the measured ones and therefore the feasibility of the numerical simulation can be verified. Table 1: Comparison of computing and measured temperatures

Graphite surface 1(cm) Graphite surface 6 (cm) Graphite surface 12 (cm) Metal outlet

Numerical Solution Temperature() 125 77 42 133

Figure 2: Computing temperature distribution of the metal and graphite mold

3.2

Experiment Solution Temperature () 119 52 36 150

Figure 3: Enlarged figure of the temperature distribution near the inlet of the graphite mold

Comparisons of Casting Speed

The working conditions set up in this section are similar to those in the last section, besides casting speed. The casting speed which this research used are 0.4 cm/sec, 0.6 cm/sec and 0.8 cm/sec. The temperature distributions along the central axis for different casting speeds are shown in Fig. 4. From this figure, the temperature gradient increases with the casting speed. When the speed is faster, the duration time of metal in the graphite mold is shorter. This makes the amount of heat taken by the cooling system less and therefore the metal temperature goes down slower.

3.3

Comparisons of Interfacial Heat Transfer Coefficient

The effective interface heat transfer coefficient can be decided according to the cooling water’s temperature and flow rate. The working condition in the first section is that the casting speed is 0.6 cm/sec and the effective interface heat transfer coefficient is 2700 W/m2K. The situation

217

Figure 4: Temperature distributions along the axial direction for different casting speeds

Figure 5: Temperature distributions for different effective heat transfer coefficient

that the effective interface heat transfer coefficient is raised or reduced is equivalent to that of raising or reducing the cooling rate. To study the effect of effective heat transfer coefficient, another two heat transfer coefficients 1000 W/m2K and 4700 W/m2K were chosen. The temperature distributions along the central axis for different heat transfer coefficients are shown in Fig. 5. From the figure, it can be clearly found that the outlet temperature of the metal decreases as the effective heat transfer coefficient increases, however, in the upper part of the graphite mold, the temperature curves are very close to one another.

3.4

Comparisons of Materials

The control variable this section is the casting material. The chosen materials are gold, silver and copper. The thermal properties of the metals are shown in Table.2. Temperature distributions along the axial direction for these three metals are shown in Fig. 6. The temperature decrease of gold along the axial direction is the smallest since it has the smallest thermal conductivity

Figure 6: Temperature distributions along the axial direction for different casting materials

218 and heat capacitances. The silver’s temperature gradient is smaller than either gold or copper, because the thermal conductivity of silver is the highest. Table 2: Thermal properties of different metals Material

Au

Ag

Cu

Melting point, CpS, CpL, Latent heat, Thermal conductivity (W/m K)

1064.4 0.163 0.166 62.76 317.9

961.9 0.262 0.310 104.2 356

1084.8 0.386 0.494 205 334

4

Conclusions

To investigate the heat transfer behavior of the continuous casting of metal rods, a mathematical model is built in this paper. The numerical method is the finite difference method. In the beginning of the study, the computing temperatures are consistent with the measured ones and thus the feasibility of the numerical simulation can be verified. From the computing results, it can be found that the temperature variation in the metal is primarily along the axial direction and most of the temperature decrease occurs in the upper of the graphite mold. As the casting speed increases, the temperature gradient increases. When the effective heat transfer coefficient increases, the outlet temperature of the metal decreases, however, in the upper part of the graphite mold, the temperature curves are very close to one another. Among gold, silver and copper, the temperature decrease of gold along the axial direction and the temperature gradient of silver are the smallest.

5 [1] [2] [3] [4] [5] [6] [7] [8]

References M. Bamberger and B. Prinz: Mathematical modeling of the temperature field in continuous casting. Z. Metall. Vol. 77(1986), No. 4, p. 234–38 S. Louhenkilpi: Simulation and Control of Heat Transfer in Continuous Casting of Steel, Doctoral Thesis, Helsinki University of Technology, Finland, 1995 S. Louhenkilpi, E. Laitinen, R. Nieminen: Real-Time Simulation of Heat Transfer in Continuous Casting.Met. Trans. B. Vol. 24B (1993), August, p. 685–693 J. S. Hsiao, "An Efficient Algorithm for Finite Difference Analysis of Heat Transfer with Melting and Solidification," Numerical Heat Transfer, Vol. 8, pp. 653–666, 1985 J. A. Dantzig, “Modeling Liquid-Solid Phase Changes with Melt Convection”, International Journal of Numerical Methods in Engineering V28, n8, pp. 1769–1785, August 1989 J. P. Holman, “ Heat transfer,” 9th edition, Mcgraw-Hill company, New York, USA, 2002 M. Uoti, M. Immonen, and K. Harkki, “Theoretical and Experimental Study of Vertical Continuous Casting of Copper.” R. Wilson, “A Practical Approach to Continuous Casting of Copper-Based Alloys and Precious Metals,” IOM Communication Ltd., London, UK, 2000

219

Modeling of Macrosegregations in Continuous Casting of Sn-Bronze M. Gruber-Pretzler, F. Mayer, M. Wu, A. Ludwig Christian-Doppler Laboratory for Multiphase Modeling of Metallurgical Processes, Department of Metallurgy, University of Leoben, Franz-Josef-Str. 18, A - 8700 Leoben, Austria

1

Abstract

Macrosegregations in DC casting may be caused by various reasons. Besides sedimentation and flotation of equiaxed grains, feeding flow, thermal and solutal buoyancy driven flow, and inlet flow are potential mechanisms for the formation of macrosegregations. However, the relative importance of these phenomena is difficult to estimate. With a two-phase volume averaging model the formation of macrosegregations for a DC casting of a Sn-bronze is simulated. By comparing the results of four different cases, it turns out that for the conditions applied, feeding flow creates the strongest macrosegregations, namely positive segregations at the casting surface and negative segregations in the center. On the other hand, negative surface segregations form if only thermal buoyancy flow is considered.

2

Introduction

For almost all practical solidification processes inhomogeneous distributions of alloy elements at the scale of the whole casting, known as macrosegregations, are found. In order to predict the formation of those undesired ‘defects’ numerical methods have been developed intensively during the last years [1–12]. In a recent publication of two of the present authors [13] the columnar-to-equiaxed transition has been modeled with a three-phase volume averaging approach, where the motion of grains, the melt flow caused by shrinkage and thermo-solutal buoyancy, and the growth of a columnar front were considered. In the present paper we have used only the columnar solidification part of the above mentioned three-phase approach [13]. In this part, the permeable mushy zone is thought to be composed of cylindrical ‘dendrites’ with a given primary dendrite arm spacing, O1. Feeding flow, through the mushy zone, as well as thermal and solutal buoyancy flow are accounted for. In order to investigate the role of feeding flow as well as of thermal and of solutal buoyancy flow in the formation of macrosegregations in DC casting of Sn-bronze four different simulations were compared: (i) considering only feeding flow; (ii) only thermal buoyancy flow; (iii) only solutal buoyancy flow; and (iv) without considering feeding flow and thermo-solutal convection.

220

3

Simulation Description

3.1

Model Description and Assumptions

Here, only a short outline of the used columnar solidification model is given. For more details the reader is refereed to the original publications [11–13]. The model considers two phases, the melt and the growing columnar dendrites. For these two phases the conservation equations of mass, species and enthalpy are considered. In addition, the momentum conservation equation for the melt is solved. The main assumptions of the model can be gathered as follows: • The thermodynamic for the binary CuSn system is considered by assuming a simplified phase diagram. A constant redistribution coefficient, k, and a constant liquidus slope, m, was used. At the peritectic temperature the solid fraction reaches already 95–98%. Therefore, and because of the fact that the model for the peritectic reaction is still under development, we assume that the remaining liquid solidifies over a small and arbitrary temperature interval. • Nucleation and growth of equiaxed grains are ignored. • Columnar dendrites are thought to start growing at the mold wall as soon as the temperature drops below liquidus. • The columnar dendrites are approximated by growing cylinders. • A shell-wise growth driven by diffusion around the cylinder is assumed. • Mechanical interaction between the mush and the flow is calculated via Darcy’s law and the Blake-Kozeny permeability model. • Corresponding source terms to account for feeding flow and thermo-solutal buoyancy driven flow is introduced.

3.2

Geometry Information and Boundary Conditions

For the process simulation a casting velocity of Vcast = 1.92 mm / s and a casting temperature of T0 = 1389 K is considered. Due to the cylindrical shape of the mold an axis symmetric simulation has been chosen. The mold is schematically shown in Figure 1a where (c) gives the position of the nozzle, (d) indicates the free surface on the top, (e) shows the upper part of the mold which is assumed to be insulating, (f) shows the lower part of the graphite mold which is surrounded (g) by a copper–steel mold including a water cooling. In Figure 1b the boundary conditions are shown. (c) gives the position of the inlet, where a pressure inlet is considered. A heat transfer coefficient (HTC) of h = 50 W/m²K and a nozzle temperature of TSEN = 1292 K is considered for the submerged entry nozzle (SEN) region. For (d) HTC and temperature have a value of h = 50 W/m²K and Tsurface = 325 K. For (e) almost ideal insulation is assumed with h = 10 W/m²K and Tmold = 1292 K. (f) has h = 3000 W/m²K and Tmold = 550 K and (g) h = 1000 W/m²K and Twater = 300 K. A constant velocity Vcast is taken for the outlet (h). For the nozzle and for the free surface a slip condition is used. The mold wall is assumed to move with the casting velocity. Here, a slip condition for the liquid phase and a non-slip condition for the columnar phase have been chosen. A grid of 9016 cells and 9296 nodes is used. The chosen conditions are the same as used in [14]. As initial conditions, we start with hot melt (Tinit = 1292 K) in rest (Vinit = 0 m/s). The presented results are taken when steady state were reached.

221

Figure 1: (a) Continuous casting: c nozzle; d free surface; e graphite mold with isolation; f graphite and copper mold; g steel mold with water cooling. (b) Grid and interfaces for boundary conditions (details given in the text).

To estimate the relative importance of feeding flow, thermal and solutal buoyancy flow, and inlet flow the simulation results for four different cases are compared: • Case A: No thermal and solutal buoyancy flow and no feeding flow is considered. We have used equal densities for the liquid and the solid. • Case B: Here only solutal buoyancy flow is considered. The solutal expansion coefficient has been chosen to be Ec = 0.11 wt.%–1. • Case C: Here only thermal buoyancy flow is considered. The thermal expansion coefficient has been chosen to be bt = 8.6 · 10–5 K–1. • Case D: For this simulation only feeding flow is considered. So we assumed the liquid and solid densities to be independent of temperature and concentration, but different. During solidification the higher solid density leads then to a shrinkage-induced feeding flow.

4

Results and Discussion

The temperature distributions of the four different cases are shown in Figure 2. In Figure 3 the velocity fields for Cases A and C are shown. In Figure 4 the mixture concentration defined as

cmix

cl ˜ fl ˜ Ul  cs ˜ f s ˜ Us fl ˜ Ul  f s ˜ Us

(1)

is compared, again for the four different cases. Here, cl and cs stand for the concentration, fl and fs for the volume fraction and Ul and U2 for the densities of liquid and solid. In all figures the liquidus (TL = 1289 K), the solidus (TS = 1230 K) and the temperature where the solidification is thought to be completed (TP = 1072 K) is shown by iso-lines. Within the columnar mushy zone, located between TL and TP, the volume fraction of solid varies from 0 to 1. As we considered the mush to be permeable, melt-flow occurs through the mush. However, the flow velocities in the mush are much smaller compared to the inlet flow. The incoming melt reveals a velocity as high as Vin = 25 mm/s. This large value is a consequent of the constant outlet velocity, Vcast, and the

222

Figure 2: Steady-state temperature distributions (in K) for the four different cases: (a) without feeding and thermo-solutal buoyancy flow (Case A); (b) only solutal buoyancy flow (Case B); (c) only thermal buoyancy flow (Case C) and (d) only feeding flow (Case D)

overall mass conservation. The inlet ‘jets’ hit the mold close to the region where first solid forms, bend inwards and form corresponding vortices (Figure 3). Due to the assumed permeability of the mush, these vortices penetrate into the mushy zone only moderately. Nevertheless, some solute which has segregated close to the dendrite tips is removed from these ‘tip’ regions and replaced by not segregated ‘fresh’ melt from the jet vortices. The washing away of segregated melt from the ‘jet-mush’ interaction zones (labeled with I in Figure 4) leads to a somewhat higher solute concentration in the bulk of the casting. As a result, weak negative macrosegregations form in those ‘jet-mush’ interaction zones and a hardly noticeable positive macrosegregation forms in the rest of the strand (labeled with II in Figure 4). This ‘jet-mush’ interaction phenomenon takes place in all considered cases. Its strength depends on the strength of the ‘jet-

Figure 3: Velocity field at the inlet region (scaled in m / s) (a) for the case without feeding and thermo-solutal buoyancy flow (Case A) and (b) the case where only thermal buoyancy flow is considered (Case C)

223

Figure 4: Steady-state distributions of the mixture concentration for the four different cases: (a) without feeding and thermo-solutal buoyancy flow (Case A); (b) only solutal buoyancy flow (Case B); (c) only thermal buoyancy flow (Case C) and (d) only feeding flow (Case D). Bright green represents the initial alloy concentration, yellow positive and blue negative macrosegregations. Flow patterns are also indicated by arrows.

mush’ interaction. If a higher mush porosity would have been chosen or if the jet would have met the mush more intensively, the amount of washed away segregated melt would have been higher and with that the resulting negative surface macrosegregations and positive ‘bulk’ macrosegregation would have been stronger. In Case A the described ‘jet-mush’ interaction phenomenon is not overlaid by other phenomena. Here, the flow from the jets and the flow caused by the withdrawal of the strand are the only machanisms present. However, even in this case the weak negative surface macrosegregations and hardly any positive ‘bulk’ macrosegregation can be seen with the color scale applied in Figure 4. This scale was chosen to optimally depict the differences in macrosegregations for the four considered cases. By comparing the temperature distributions (Figure 2), the mixture concentration distributions (Figure4) and the flow pattern for Case A (Figure3a) and B (not shown) it becomes obvious that for the considered situation solutal buoyancy driven flow is negligible. However, thermal buoyancy driven flow seems to be of more importance (see Figure 2c, 3b, 4c). The melt close to the mold is cooled down and thus increases its density, which in turn results in a downwards acceleration and with that in a strengthening of the vortices (Figure 3b). The acceleration by the denser and cooler melt and the strengthening of the vortices has three important consequences. Firstly the vortices are moved upwards; secondly cooled down and thus denser melt leaves the vortices and flows further downwards and thirdly melt from deep down in the pool is pressed/sucked upwards in the centre of the middle axis of the cylindrical casting (labeled with III in Figure 4c). This upwards flow is the reason for the fact that with thermal buoyancy flow the liquidus temperature in the middle of the casting is located somewhat closer to the SEN compared to the case without this flow (see Figure 2). It is worth to mention that the ‘jet-mush’ interaction is strengthened by the thermal buoyancy flow and thus more segregated melt is washed out which leads to slightly more pronounced negative surface macrosegregations (Figure 4c). The washed out segregated melt is accumulated

224 mainly in the mush in the centre of the casting, which is the reason for the weak positive macrosegregations appearing there (labeled with II in Figure 4c). Up to now, the surface macrosegregations are predicted to be negative and macrosegregations in the casting center are predicted to be more or less positive. However, if shrinkage-induced feeding flow is included and thermal buoyancy driven flow is ignored, the surface macrosegregations turned out to be positive and the macrosegregations in the center strongly negative, so just the opposite! Feeding flow is always directed from the dendrite tip towards its roots and thus carries segregated into the mush. Since the early work of Flemings in 1967 [15, 16], this phenomenon is known to result in a positive macrosegregation at the surface of a casting, the so-called inverse segregation1. Exactly this happens in Case D (labeled with IV in Figure 4d). At the centre of the cylindrical casting the dendrite tips approaches each other and so they form a ring which is going to be closed while solidification proceeds. On the other hand, a relatively large mush area at the center is solidifying and thus a large amount of melt is needed to feed the corresponding shrinkage. This melt is sucked into the solidifying mush via the closing ‘ring of dendrite tips’ and so a relative strong downwards feeding flow occurs in the center of the casting (labeled with V in Figure 4d, see ref. [14]). With that the solidifying dendrites are fed with less- or no-segregated ‘fresh’ melt from the melt pool and therefore a negative macrosegregation in the center region of the casting forms. In addition, heat is carried with the downwards feeding flow which leads to the somewhat lower position of the isotherms, especially at the casting center (Figure 1d). Again, the predicted macrosegregations depend on the permeability of the mush. If larger mush porosity would have been chosen, feeding through the mush would have been easier and the center downwards flow would have been significantly broader. This would have caused a broader negative center macrosegregation which, however, would not have been so strong.

5

Conclusions

The formation of macrosegregations in DC casting of a columnar solidifying Sn bronze is caused by the ‘inlet jet-mush’ interaction, the thermal buoyancy driven flow and the shrinkageinduced feeding flow. Conclusions gained from the study of the different macrosegregation formation mechanisms can be gathered as follows: • The ‘jet-mush’ interaction leads to weak negative macrosegregations on the surface and hardly any positive macrosegregations in the bulk of the casting. • For the investigated situation solutal buoyancy driven flow is negligible. • With thermal buoyancy driven flow the ‘jet-mush’ interaction is stronger, and thus the negative surface macrosegregations and the positive bulk macrosegregation is more pronounced. • Shrinkage-induced feeding flow causes macrosegregations which are opposite to those which forms with and without thermal buoyancy driven flow. Namely, positive surface macrosegregation and a strong negative center macrosegregation.

1 Nowadays, inverse (surface) macrosegregation is known to be also caused by an interdendritic flow towards the surface accompanied with reheating and remelting during the local formation of a gap between the casting and the mold.

225 • The strength of the macrosegregations is thought to be dependent on the considered mush permeability. Here, further investigations are highly needed. Although, we did not present predictions on macrosegregations for the case with both thermal buoyancy driven flow and shrinkage-induced feeding flow, it is obvious from Figure 4 that for the considered conditions the effect of feeding flow will dominate the formation of macrosegregations. However, for different conditions the opposite might be true.

6

Acknowledgements

This work was financially supported by the Austria Christian-Doppler Society (CDG) and by the Wieland-Werke in Germany for which the authors kindly acknowledge.

7 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]

References M. Rappaz, Int. Mater. Rev. 34, 1989, p. 93 J. Ni, C. Beckermann, Metall. Trans. 22B, 1991, p. 349 C. Beckermann, R. Viskanta, Appl. Mech. Rev. 46, 1993, p. 1 J. Ni, F.P. Incropera, Inter. J. Heat Mass Transfer 38, 1995, p. 1271 J. Ni, F.P. Incropera, Inter. J. Heat Mass Transfer 38, 1995, p. 1285 C.Y. Wang, C. Beckermann, Metall. Mater. Trans. 27A, 1996, p. 2754 C.Y. Wang, C. Beckermann, Metall. Mater. Trans. 27A, 1996, p. 2765 C.Y. Wang, C. Beckermann, Metall. Mater. Trans. 27A, 1996, p. 2784 C. Beckermann, JOM 49, 1997, p. 13 A.V. Reddy, C. Beckermann, Metall. Mater. Trans. 28B, 1997, p. 479 A. Ludwig, M. Wu, Metall. Mater. Trans. 33A, 2002, p. 3673 M. Wu, A. Ludwig, A. Bührig-Polaczek, M. Fehlbier, P.R. Sahm, Inter. J. Heat Mass Transfer 46, 2003, p. 2819 M. Wu, A. Ludwig, Metall. Mater. Trans., 2005, submitted A. Ludwig, M. Gruber-Pretzler, F. Mayer, M. Wu, Mat. Sci. Eng. 2005 in print M.C. Femings, G.E. Nereo, Trans. Metall. Society AIME 239, 1967, pp. 1449–1460 M.C. Femings, R. Mehrabian, G.E. Nereo, Trans. Metall. Society AIME 242, 1967, pp. 41–49

226

Continuous Casting Simulation: From Solidification and Fluid Flow to the Calculation of Grain Structures R. Eberle Wieland-Werke-AG, 89079 Ulm, Germany

1

Abstract

Nowadays the simulation of the continuous casting process is an important tool for optimizing production in the foundry. Simulation tools become more and more powerful. The possibilities start from simple calculations of the solidification and the fluid flow in the melt. The next step is the consideration of mechanical stresses and the combination of thermal and mechanical calculations called thermo mechanical coupling. An even more complex task will be the calculation of segregation effects or grain structure. A selection of the above mentioned tasks is presented and the results are illustrated, analysed and compared with plant trials. Thereby the growing demand for an excellent knowledge of the casting process with increasing complexity of the problem will be demonstrated. It is important to estimate the validity of the models for an optimization of the casting process. The prediction power of the different models will be pointed out. Furthermore difficulties and limits of this ability are discussed.

2

Introduction

The continuous casting process represents a very complex system with a lot of process parameters. Often, the parameters are correlated and usually the kind of correlation between them is unknown. These facts make it very difficult to optimize a continuous casting process by plant trials only. For a systematic proceeding one has to make a series of plant trials with the consequence of high costs. This problem can be solved by using simulation techniques to understand the interaction between the different process parameters and to find optimized parameter combinations to be tested in a few plant trials. In this paper typical examples are treated starting from simple tasks like solidification and fluid flow to more complex ones like macrosegregation. The possibilities to influence and to optimize fluid flow are illustrated in Chapter 3. Chapter 4 deals with the calculation of gap formation between the mould and the strand. For this purpose an iterative procedure is used. Finally in Chapter 5 results of macrosegregation calculations are presented.

3

Optimizing Fluid Flow in a Rectangular Shaped Mould

The kind of fluid flow has a large influence on shell formation, heat distribution in the liquid pool and the surface quality of the cast strand. With an unfavourable fluid flow shell formation

227 could be disturbed. Additionally slag from the top of the casting could be drawn into the melt and lead to a bad surface quality. Thus, it is important to optimize fluid flow in the mould.

3.1

The Status of the Process at the Moment

First a typical process is analyzed with Magmasoft®. Several steps have to be done before the calculation can be run. The geometry of the mould and the slab together with the nozzle and an inlet has to be constructed. Then the geometry is meshed. In a third step the process data are programmed. The different material properties for the mould and the slab are defined, parameters like the amount and the temperature of the water cooling are specified and different boundary conditions are determined. Casting velocity and calculation parameters terminate the collection of process data. A typical geometry is shown in Figure 1. The melt enters the mould through a T-shaped nozzle and leaves in horizontal direction towards the narrow side of the mould. The velocity at the outlet of the nozzle is determined through casting velocity and outlet area of the nozzle. In Figure 2 results of fluid flow after a calculation time of 150 seconds are shown. The fluid flow can be visualized by tracer particles.

Figure 1: Geometry of mould and slab

Figure 2:. Corresponding fluid flow for T-shaped nozzle

The melt leaves the nozzle towards the narrow side of the mould. At this side it is guided to the corner of the mould and flows back to the middle of the broad side. While flowing back the melt cools down and descents, preferential in the middle where two streamlines from opposite directions meet. Thus, in the middle of the broad side a zone with descending melt forms. There it is probable that some slag is transported from the surface to the solidification zone and is included in a region near the surface of the slab.

3.2

Solutions Avoiding Descending Zones

According to the statements above it is advisable to avoid zones with a strong descent. For this purpose, nozzles of different shape and orientation were tested with the simulation. As one ex-

228

Figure 3: Fluid flow results with an inclined nozzle

ample a nozzle was inclined towards the broad side of the mould. The results from the fluid flow calculation are shown in Figure 3. The streamlines are reflected at the broad and narrow sides of the mould so that the melt moves in an ellipse-like way. There is no descending zone any more and the melt is distributed homogeneously. This configuration should avoid slag to be included into the slab.

4

Gap Formation Between the Mould and the Strand

Considering the continuous casting process, the behaviour between the strand and the mould is an often discussed question. In the upper region of the mould the liquid metal fully wets the mould. Then the strand starts to solidify and shrinks. This leads to the formation of a gap between strand and mould which lowers the heat transfer. The solidified shell reheats, the strand expands towards the mould and the heat transfer increases again. With unstable casting parameters this process can occur several times until the strand leaves the mould. It is a very interesting task to describe this behaviour within a simulation model.

4.1

Model for the Heat Transfer

First of all we need a model to describe the heat transfer between the strand and the mould. For the casting process some kind of parting compound is used which forms a film on the mould surface. At first there is no gap formation and the film fills up the space between strand and mould completely. Then the heat transfer coefficient (HTC) D is determined by the heat transmission through the film. If dfilm is the thickness of the film and Ofilm its thermal conductivity we get a HTC of

229

D

O film

(1)

d film

With a larger gap the film cannot fill the space completely and an air gap forms. Then the complete HTC composes of three different parts. Through the air gap the heat is conducted by radiation and the heat conductivity of the air. Finally it has to pass the film. Posing all these mechanisms together we get

D

1 1

D rad 

O air d air



1

O film

(2)

d film

Drad describes the radiation, Oair and Ofilm represent the heat conductivity, dair and dfilm the thickness of the air gap and the film respectively. In Figure 4 the HTC depending on the thickness of the gap is illustrated.

Figure 4: Model for the heat transfer coefficient with a maximum film thickness of 1 mm

At a thickness up to 1 mm where the gap is fully filled with the parting compound, the HTC is large, determined by the width of the film and its thermal conductivity. If the width increases above 1 mm, an air gap forms. Then the HTC drops rapidly to low values determined by radiation and the heat conductivity of air.

4.2

Iterative Calculation of Gap Formation

With this model the gap formation for a known casting process can be calculated. Because there is a coupling between the heat transfer and the gap formation an iterative procedure is chosen. Starting with an arbitrary HTC between strand and mould the solidification and the shrinking of the strand are calculated. From the results of the shrinking a new characteristics of the HTC can be determined. This HTC is then used for the calculation of solidification and shrinking in

230 the next loop. That procedure is repeated until a stable configuration for HTC and shrinking is achieved. This proceeding will be exemplary shown by the continuous casting of billets. As starting conditions a constant HTC of 1000 W/m²K and a maximum film thickness of 1 mm were chosen. After some iterations the HTC gets stable. The starting value and the final result of the HTC are shown in Figure 5.

Figure 5: HTC for the beginning and final results after the iteration procedure

At the top of the casting there is a region with a high HTC. Here is the liquid pool and the shell starts to solidify. After the strand gets a solid shell, shrinkage begins and the HTC is distinctly lower. The reduced HTC in this area causes a reheating of the shell due to the cooling of the core. Thus the strand approaches the mould. At the end of the mould a secondary cooling becomes important, shrinkage is increased (see Figure 6) and a further reduction in the heat transfer is caused.

Figure 6: Shrinkage of the strand after the final iteration

231 The development of the air gap is regarded in Figure 6. In principle it is vice versa to the behaviour of the HTC. Near the top of the casting there is no air gap. The liquid pool and the mushy shell completely touch the mould. After solidification a small gap of 0,2 mm develops which remains nearly constant. The air gap decreases in this case at 80 % of the mould length by the repulsing of the strand to the mould. Although there are some simplifications in the model, this is a suitable approach to describe the dynamic behaviour of a strand in the mould.

5

Calculation of Macrosegregation in CuSn Alloys (Bronzes)

A well known effect for the continuous casting of bronze is the so-called inverse macrosegregation. This means that during the casting process tin enriches in the outer areas of the strand and at the same time decreases in the core. To optimize the quality of a cast product it is important to achieve a tin distribution over the whole strand which is as homogeneous as possible. Again simulation provides the possibility to predict element distribution within the strand and optimize process parameters.

5.1

The Model

In this example the tin concentration for the continuous casting of slabs was calculated with Calcosoft-2D. For this purpose a cross section along the narrow side of a rectangular geometry was chosen. The relative movement between the solid and the liquid phases is important for the calculation of macrosegregation. For the model two main effects are taken into account. These are the natural or forced convection and the shrinkage due to solidification. Therefore, first of all solidification and fluid flow have to be calculated. Furthermore additional material data are required to perform the calculation. It is necessary to know several parameters from a phase diagram like melting point of the base material, eutectic temperature, liquidus slope and eutectic concentration.

5.2

Results for Solidification, Fluid Flow and Macrosegregation

Feeding all this into the calculation results for solidification, fluid flow and tin distribution along the slab were obtained. These results are shown in Figure 7. All pictures above represent a cross section through a slab. On the right side it is cooled down by a mould. The edge on the left represents the middle of the slab. There a symmetric boundary condition was applied. The left part of Figure 7 shows the fraction solid. The change in colour from white to blue describes the solidification front. In the middle part the fluid flow is illustrated. It is assumed that the whole top of the casting represents the inlet. Below the surface the melt flows towards the cooled side and descends along the solidification front. A small part of the melt moves back towards the corner. Below this main flow a curl forms turning in clockwise direction.

232

Figure 7: Results for fraction solid, fluid flow and macrosegregation

The last part on the right side shows the tin distribution in the slab. The contour of the solidification front, which plays an important role for macrosegregation, can clearly be recognized when compared with the fraction solid. Within the solid region the tin distribution is fixed. Caused by macrosegregation tin distribution differs by –0,14 to +0,26 % from the nominal concentration. As indicated by the line in Figure 7 a profile of the tin distribution has been extracted which is shown in Figure 8.

Figure 8: Final tin distribution in the solid region of the slab

233 The horizontal line represents the nominal tin content. In the core of the slab the tin content is lowered. A few centimetres further the nominal concentration is reached again. Then it stays constant just slightly above this value and increases close to the surface by 0.25 %. This behaviour arises from inverse macrosegregation often observed in continuous casting.

6

Summary

In the preceding three chapters a brief survey of some topics which can be treated with simulation was given. It could be shown that simulation gives a lot of hints which effects during the process could occur. Nevertheless these results strongly depend on the quality of the material data used for the simulation. Especially concerning macrosegregation for the calculations often some material parameters are not known. Thus, the results must be regarded critically and have to be carefully tested in plant trials.

234

Mould Temperature Fields during Continuous Casting of DHP-Copper M. Mäkinen1, M. Uoti2 1 2

Helsinki University of Technology, Espoo, Finland Outokumpu Copper R&D, Pori (Finland)

1

Introduction

In continuous casting processes the mould is the most complex and critical part, which controls initial solidification and surface quality. The amount of heat transfer and its uniformity affect on the solidification front and liquid pool depth. This front is also influenced by other factors, such as casting speed, mould design and superheat. During casting, mould may distort due to the steep thermal gradients and air gap formed between strand and mould due to shrinkage of casting. This air gap formation starts, as solidifying metal gets strong enough to withstand the metallostatic pressure. This air gap has been determined to be the greatest single resistance to heat removal. Oscillation marks, which are formed during every mould oscillation cycle, are small depressions on the surface of the cast metal. They may cause cracking and they can act as a small air gap in the mould area [1–5]. As higher volume and better quality of production is needed, heat flux between strand and mould is investigated by measuring the inflow and outflow temperatures and the throughflow volume of the mould cooling water. Thermal variations are measured via embedded thermocouples. The aim of mould thermal monitoring is to predict breakouts and casting quality, to develop casting practise for new grades and casting powders and to determine the influences of different casting variables on mould heat transfer, heat flux profiles, shell growth rate in the mould and temperature uniformity in the mould [6–9]. In this research during continuous slab casting of deoxidized high phosphorus (DHP) copper, the temperature field in the mould plates were measured via embedded thermocouples to get more precise information how operating parameters, such as casting speed, superheat, cooling water amount, oscillation of the mould and slab size affect on heat transfer, quality and productivity. Mould plate temperatures at steady state were also compared between different locations in the mould due to changing temperature field and its fluctuations during operation. Due to this temperature fluctuation, there is also average temperature and standard deviation counted for investigated thermocouples. The adjustable trial mould of our vertical casting machine consists of copper mould and a graphite die.

2

Experimental Procedures

In the industrial scale experiment we have cast DHP-copper (Deoxidized High Phosphorus, 99.9 wt-% Cu) slabs with adjustable trial mould. The depth of liquid pool was measured by steel rod. Casting speed, superheat, cooling water amount and temperatures, oscillation of the mould and slab size were selected as casting parameters. While comparing differences of one casting

235 parameter, others are kept constant unless anything else has been mentioned. These selected casting conditions are presented in Table 1. Table 1: Casting conditions Casting speed Mould size Mould Casting temperatures / TS Secondary cooling Oscillation Oscillation amplitude

145–225 mm/min 120*1085 mm2, 155*1025 mm2, 200*850 mm2 Water cooled Cu mould with graphite linings 1140–1150 °C /1084 °C Narrow face 13–15 m3/h Wide face 27–53 m3/h 50–100 rpm 20 mm

Between the strand and the copper mould there are graphite linings, which act as a lubricant and molten copper is covered by carbon powder to prevent oxidation. Therefore no mould flux is needed. In order to evaluate heat transfer and temperature field in the water-cooled adjustable copper mould, there were 18 thermocouples embedded in both wide copper plates at several locations. The uppermost thermocouples are 16 cm below meniscus and following rows are 10 cm from each other. In this research, only six of them were used due to the scope of this article. These thermocouples locations in wide copper plates with three different mould sizes are illustrated in Figure 1 a. Figure 1 b shows mould copper plate temperatures and standard deviation for different thermocouples during casting of DHP-copper. The standard deviation of thermocouple temperatures particularizes how stable casting conditions are. For the convenience of the caster, different temperatures are shown in different coloured cells. Mould size in this experiment has been 155u1025 mm2 and casting speed was 210 mm/min.

Figure 1: a) A top and a side view of three different mould sizes with embedded thermocouples, b) Mould temperature field and standard deviation of thermocouples during casting with mould size 155*1025 mm2 and casting speed 210 mm/min.

3

Results

The temperature field in the mould wall during casting process varies a lot as can be seen from Figures 2 a–b. The width of the mould has also some effect on this phenomenon. Temperatures are slightly higher using wider mould. One reason for this is that with constant casting speed li-

236

Figure 2: Measured temperature as a function of time for mould size 155*1025 mm2 and 200*850 mm2, respectively. Casting speed has been 180 mm/min in both cases.

quid flow into the mould is greater, and more heat has to be extracted. Highest temperatures on the mould wall have been gauged usually from thermocouple 6 b. This is due to liquid flow and so called dog bone effect. This dog bone effect means that the slab is a bit thicker on the narrow faces than in the centre of the wide face. Broken lines represent wider mould (mould size 200*850 mm2) and it is this size where thermocouples 6 b and 6 c are closer to the mould narrow face (see Figure 1a.), where the liquid flow hits. Temperature fluctuation is more visible from thermocouples located near the narrow face (thermocouples 6 b and c in Figure 2 b.) especially with wider mould, but temperature fluctuation decreases in the middle of the mould wall (Figure 2 b.). The measured depth of liquid pool was few centimetres higher with wider mould as can be expected. Average temperatures and standard deviation for these measurements are illustrated on Table 2. Table 2: Average temperature fields and standard deviations for selected thermocouples with mould sizes 155u1025 mm2 and 200u850 mm2. Casting speed was 180 mm/min. 155*1025 mm2

N4a Av. temp. °C 123.6 St. Dev. 2.7

200*850 mm2

W4a Av. temp. °C 133.9 St. Dev. 1.2

N4b

N4c

N5b

N6b

N6c

99.5 1.9

61.0 1.2

118.1 1.3

126.6 1.8

89.2 1.7

W4b

W4c

W5b

W6b

W6c

101.7 1.8

58.4 1.1

122.9 2.4

134.0 2.2

100.9 2.8

Casting speed is always optimised to the best yield and quality and it plays vital role for uninterrupted production. Due to possible shell breakout, the evaluation of temperature fields and especially temperature fluctuations during changing casting speed is essential. In Table 3 average temperatures and standard deviation for three different casting speeds are depicted. As results show, increasing casting speed increases average temperatures of thermocouples. Standard deviation, on the other hand, decreases which means that temperature field changes less than before. The former is due to that more hot metal is coming into the mould as well as growing shell thickness is decreased. The latter is due to better contact between strand and mould. Here mould size is 120u1085 mm2 and amount of mould cooling water is 25 % less than results shown in Figure 2 and Table 2.

237 Table 3: Average temperatures and standard deviation for three different casting speeds. 145 mm/min 160 mm/min 170 mm/min

Av. temp. °C St. Dev Av. temp. °C St. Dev Av. temp. °C St. Dev

4a 105.0 17.9 111.0 6.8 112.7 1.1

4b 112.8 15.0 133.1 5.8 138.0 1.8

4c 83.6 5.2 94.7 4.1 96.3 3.4

5b 109.0 12.8 125.2 5.0 123.9 2.2

6b 104.7 10.3 120.6 5.8 129.9 1.9

6c 83.5 3.6 91.6 3.6 94.0 2.3

Temperature fluctuation during low casting speed is due to air gap formation between mould and strand that is due to slab local shrinkage. This air gap decreases heat removal from slab surface and affects on solidification. After increasing casting speed by 15 %, the liquid pool depth was also increased approximately by 20 %. After a certain value, where liquid pool depth was about 10 centimetres below mould there were no measured changes. The strong secondary cooling did strongly influence the temperatures at this point. The effect of casting speed on grain structures can be seen from cross section samples in Figures 3 a–b. Increasing casting speed changed orientation of columnar grains upwards and near the narrow face they turned a bit away from the edge of the slab. The direction of heat extraction is perpendicular to the sump profile. With higher casting speed the direction of heat flow will turn more vertical direction and the cooling effect of narrow face at the local point is smaller.

Figure 3: Grain structures on cross section samples where casting speed was a) 190 mm/min, b) 220 mm/min

Superheat affects on the solidification structure as well as midway and centreline crack formation. This has been investigated earlier with round DHP-copper billet [10, 11]. By examining the effect of superheat to the response of thermocouples, the finding was pretty obvious: decreasing superheat approximately 10 °C decreases average temperature field. Also standard deviation seemed to decrease, except thermocouple 4 b. With thicker mould and lower casting speed temperature fluctuation was a bit more with thermocouples 4 b and 4 c. Obviously strand is thicker when less superheat is used. To avoid clogging of the entry nozzle, superheat can not be reduced too much. The air gap formation is not strongly affected by superheat. It should be limited by other casting parameters such as casting speed and the amount and temperature of mould cooling water. Amount of mould cooling water as well as temperature of inflowing cooling water has particular significance to the air gap formation, the shell growth and the solidification structure. These have been also confirmed for round DHP-copper billet [10]. It was visible that increasing the amount of mould cooling water decreased temperatures of thermocouples that were embed-

238 ded in the mould plates. Thermocouple 4a is an exception, where temperatures are even higher compared to the situation where there was less mould cooling water. The effect of the amount of mould cooling water was more visible looking at standard deviation values. They were higher with greater amount of cooling water. After a certain amount of mould cooling water air gap start to form between strand and mould. This has been noticed from the surface of the slab that is more or less wavy. The surface of the slab always consists of some amount of small ripples on it. These ripples are oscillation marks and they act as small air gaps. Oscillation marks are formed on every cycle where mould is moving up and down i.e. oscillating. Oscillation frequency has been noted to affect on depth of those marks as well as to the distance between them [12]. Within our earlier studies it was noted that by increasing oscillation frequency from 50 rpm up to 100 rpm the average depth of oscillation mark decreased approximately from 0.45 mm to 0.15 mm resulting in a very smooth surface. The response of thermocouples for different oscillation frequencies was not so clear. From 50 rpm up to 90 rpm average temperatures increased while standard deviation decreased. By increasing oscillation frequency up to 100 rpm situation changed a bit; average temperatures decreased and standard deviation increased. The former phenomenon was due to better contact between strand and mould with smaller oscillation marks.

4

Summary

Mould temperatures and its deviations during continuous casting of deoxidized high phosphorus (DHP) copper slab for different casting speed, superheat, amount of mould cooling water, oscillation and slab sizes were determined. With wider mould average temperatures were higher than with narrow one as well as temperatures standard deviation near the narrow face of mould. The wider mould was narrower than the other ones, which meant that thermocouples are located more close to the narrow face and therefore average temperatures were higher. Liquid pool depth was increased applying wider mould. Increasing casting speed increases average temperatures whereas standard deviation is decreased. Former is due to that more liquid metal is coming into the mould and thickness of strand is decreased, whereas latter is influenced more via decreased air gap between mould and strand. Liquid pool depth increased approximately by 20 % when casting speed was increased by 15 %. After a certain value there were no measurable changes in the liquid pool depth because of heavy secondary cooling. After a certain point the amount of mould cooling water has effect on temperature fluctuation as air gap has started to form. Oscillation of the mould has been noted to influence to the surface quality as every oscillation cycle leaves small ripple on to the surface of the slab. These oscillation marks act as small air gaps. Increasing oscillation frequency decreases the depth of these marks and average temperatures are higher with less deviation of temperatures. Liquid flow can be verified with thermocouples as can be seen especially from thermocouple 6b where liquid flow hits. It was also possible to see if submerged entry nozzle was starting to clog or was misaligned, when left and right hand sides were compared to each others.

239

5 [1]

References

J. K. Park, B. G. Thomas, I. V. Samarasekera, U. S. Yoon, Thermal and mechanical behavior of copper molds during thin-slab casting (I): plant trial and mathematical modeling, Metal. Mater. Trans. B, 2002, Vol. 33B, pp. 425–436 [2] J. M. Rodriquez, A. Esteva, S. Meza, A note on the control of the solidification front in the continuous casting of copper tubes, J. Mater. Proc. Tech., 1999, 96, pp. 42–47 [3] Ozgu, M. R., Continuous Caster Instrumentation: State-of-the-Art Review, Canadian Metallurgical Quarterly, 1996, Vol. 35, No. 3, pp. 199–233 [4] J. Elfsberg, Oscillation mark formation in continuous casting processes, Licentiate thesis, Stockholm, Sweden (2003) [5] J. Kron, Air gap formation and hot tearing in solidification processing of Al- and Cu-base alloys, Doctoral thesis, Stockholm, Sweden (2004) [6] A. S. Normanton, P. N. Hewitt, N. S. Hunter, D. Scoones, B. Harris, Mould Thermal Monitoring: a Window on the Mould, Ironmaking and Steelmaking, 2004, Vol. 31, No. 5, pp. 357–363 [7] M. W. Nichols, Measuring variation in heat transfer in a slab casting mold using embedded thermocouples, 1990-PTD Conference Proceedings, pp. 45–52 [8] M. Yamamoto, T. Mizuguchi, Mold temperature measurements during semi-continuous slab casting of copper alloys and influence of operating variables, Proc. Merton Flemings Symp. Solidification and Materials Processing, TMS, 2001, pp. 407–410 [9] G. Xia, H. P. Narzt, Ch. Fürst, K. Mörwald, J. Moertl, P. Reisinger, L. Linderberger, Investigation of mould thermal behaviour by means of mould instrumentation, Ironmaking and Steelmaking, 2004, Vol. 31, No. 5, pp. 364–370 [10] M. Mäkinen, Effect of superheat and cooling on grain structure and crack formation in the continuous casting of DHP-copper, Licentiate’s thesis, Helsinki, Finland, (2004) [11] M. Mäkinen, M. Uoti, The effect of superheat on micro- and macrosegregation and crack formation in the continuous casting of low-alloyed copper, Will be published by Trans. Tech. Publications in a special issue of Materials Science Forum in spring 2005 [12] M. Mäkinen, The effect of oscillation on the surface of copper slab, TKK-report 2003, pp. 25

240

Simulation of Heat Transfer and Solidification in Continuous Casting of Copper Alloys and the Effect of Fluid Flow S. Vapalahti, S. Louhenkilpi, M. Mäkinen, P. Väyrynen Helsinki University of Technology, Laboratory of Metallurgy, Espoo, Finland

1

Abstract

A lot of mathematical models have been developed to simulate heat transfer and solidification in continuous casting. In these models, it is usually assumed that the strand is withdrawn through the machine with a constant casting speed and the convective heat transfer generated by the fluid flow is taken into account by using an effective thermal conductivity method. At Helsinki University of Technology, this kind of heat transfer model is developed (TEMPSIMU3D). It consists of two separate models, the mould and the strand model. These two models are coupled together using a so called gap heat transfer coefficient as a function of the strand surface temperature. A dynamic version, DYN3D, is also developed. The required material data for the simulations are calculated using an in-house model called CASBOA. It calculates the material data for copper binary alloys. These models are presented in this paper. Coupled models calculate the turbulent fluid flow, heat transfer and solidification simultaneously. These kinds of models are generally subjected to convergence difficulties and long computing times. In the present study, a commercial FLOW-3D package is used to make coupled simulations to investigate the effects of fluid flow on heat transfer and shell growth to estimate the justification of the effective thermal conductivity method.

2

Introduction

Continuous casting involves many physical phenomena. The main phenomena are: fluid flow, heat transfer and solidification. Most of the thermal models developed are not calculating the fluid flow at all. In these models, it is assumed that the strand (solid and liquid) is withdrawn through the machine with a constant velocity field (= casting speed). The convective heat transfer generated by the fluid flow is taken into account by using an effective thermal heat conductivity method. The equation to be solved now is the basic partial differential equation of heat conduction including the removal of latent heat of solidification and other phase transformations. These kinds of models are increasingly being used to simulate continuous casting processes. For steels, mostly one- or two-dimensional models have been developed and successfully used but in the case of copper, the model must be three-dimensional. Today, more attention has been paid on developing three-dimensional steady state and three-dimensional real-time heat transfer models. In this paper, a three-dimensional steady state heat transfer model called TEMPSIMU3D and a three-dimensional real-time heat transfer model called DYN3D are presented. The thermophysical material data for the models are calculated using in-house model CASBOA. These are also shortly presented in the paper.

241 Coupled models calculate the fluid flow, heat transfer and solidification simultaneously. The coupled models include three main equations: the energy equation for temperature, the NavierStokes equation for velocities (momentum) and the mass equation for mass. If the flow is turbulent, a turbulent model must be added. The most commonly used turbulence model is so-called k-H turbulence. This increases again the complexity of the system, because two additional equations must be solved. One example of the coupled calculation is presented and the effects of fluid flow on heat transfer and shell growth is discussed.

2

TEMPSIMU3D and DYN3D - Heat Transfer Models for Continuous Casting

TEMPSIMU3D and DYN3D are 3-dimensional heat transfer models for continuous casting. The first is for steady state and second for transient casting conditions with a capability to be used in real-time. The model equations include the transient term but in the case of steady state model, this term is neglected. The models consist of two separate modules: the mould model and the strand model. They are running iteratively and so called gap heat transfer coefficient as a function of the strand surface temperature is used to couple them. Models are based on finite difference method where upwind scheme is used. The strand model simulates the strand temperature field and the temperature related data three-dimensionally using Eq. 1 (the term UwH/wt is neglected in steady state model).

U

wH wH v wt wz

w wT w wT w wT ( keff ) (keff )  (keff ) wx wx wy w y wz wz

(1)

Here H is the enthalpy, t is time, keff is the effective thermal conductivity, T is temperature, U is density and Q is the actual casting speed. Enthalpy function H includes also the effect of all phase transformations. The convective heat transfer in the liquid pool due to the liquid flow is described with the effective thermal conductivity according to Eq. 2.

keff

k ˜ f S  A ˜ k ˜ (1  f S )

(2)

A is constant and fs is solid fraction in the mushy zone. The boundary condition for the strand in the mould is given by Eq. 3.

k

wT wn

hgap (T  Text )

(3)

Here hgap is the gap heat transfer coefficient, T is the strand surface temperature and Text is the external temperature, in this case the temperature of the mould surface. Eq. 3 is also used as the boundary equation for the hot side of the mould, where hgap is the same gap heat transfer coefficient as for the strand model but Text is now the temperature of the strand surface, which is calculated by the strand model. For the colder side, similar equation is used, but Text is the temperature of the cooling water and h is a constant. It is assumed that there is no heat flux through the top and the bottom surfaces of the mould.

242 The mould surface temperature is calculated by the mould model using Eg. 4 (the term

UcwT / wt is neglected in steady state model).

Uc

wT wt

k

w 2T w x2

k

w 2T w y2

k

w 2T wz 2

(4)

The model allows defining three different material layers in the mould. This is especially important in the case of continuous casting of copper where there is usually at least a copper jacket with a graphite die. In the secondary cooling zone after the mould the model has been divided into calculation domains between a pair of support rolls. This domain has been divided into four different cooling regions that are: roller contact area, pre-nozzle area, spraying area, after spray and pool water area (post-nozzle area). Roller contact area indicates heat conduction from strand to roll. Roller heat transfer depends on the roll contact length in casting direction as well as the roll cooling mode. Eq. 5 is the boundary equation for the roll cooling.

k

wT wn

h(T  Text )

(5)

Pre-nozzle and post-nozzle areas account for the indirectly cooled space between roller pairs. The boundary equation for these areas is defined in Eq. 6.

k

wT wn

4 h (T  Text )  HV (T 4  Tair )

(6)

Here h is the heat transfer coefficient for air convection, which is the minimum value, used for all exposed surfaces in the machine. Text is the external temperature, H is the emissivity, and Tair is the air temperature. For post-nozzle area heat transfer coefficient for flowing water is used. This takes into account the heat transfer to the flowing water on the strand after nozzles. In the spraying area, the strand is usually cooled by a spray of water or water-air mixture. Nozzle parameters like air and water flow have an effect on the cooling efficiency. The mechanism of cooling can be subdivided into heat transfer by radiation and convection. The radiation heat flux parameters, emissivity and the air temperature are the same as defined for the pre-nozzle area. The heat transfer coefficient for the water spraying area is defined by the Eq. 7.

h

a ˜ W b ˜ c(T )

(7)

Here h is the water spray heat transfer coefficient, W the water flow rate, and a, b and c(T) are parameters. The parameter c(T) is a function of temperature and it takes into account the effect of the strand surface temperature on the heat transfer coefficient (Leidenfrost effect).

243

3

CASBOA Software for the Calculation of Material Data

To obtain reliable results from the heat transfer simulations, accurate thermophysical material data are needed. Usually these data are obtained from literature but very seldom all required data is found. The use of inaccurate material data can lead to considerable errors in calculations as presented in [1]. Typical data needed are the density, the thermal conductivity, the specific heat, the phase transformation temperatures and the corresponding latent heat with the information how the latent heat is released during the phase transformations. CASBOA (Copper Alloys Solidification for Binary One-solid-phase Alloys) is a thermodynamic-kinetic solidification model for the simulation of solidification phenomena of copper alloys [2]. Currently the calculation can be done for pure copper and for 14 binary copper alloys for Cu-X, where X is Ag, Al, Cr, Fe, Mg, Mn, Ni, P, Si, Sn, Te, Ti, Zn or Zr. Multicomponent version is under development.

4

Results

The parameter A in the Eq. 2 generally has a value between 1 and 8. The effective thermal conductivity method also assumes that fluid flow increases liquid pool heat transport isotropically, which is not the case, because fluid flow only contributes to heat transport in the flow direction. In this work the effect of the value of parameter A was studied in a copper billet casting machine. The study was carried out using a commercial CFD package FLOW-3D. Both fully cou-

Figure 1: Liquidus and solidus isotherms for flow, A = 1 and A = 7 respectively

244 pled and effective thermal conductivity method calculations were performed in 2D cylindrical coordinates. In the coupled calculations RNG (renormalized group theory) -model is used as the turbulence model. Test calculations were carried out using A = 1, and A = 7 in Eq. (2) and these results were compared with the fully coupled calculations. The results are presented in the Fig. 1. and Fig 2.

Figure 2. Solid fraction profiles at the centre of the strand and next to the mould wall respectively

As can be seen, the parameter A has influence on the results, but more on the liquidus isotherm than on the solidus isotherm. This means that accurate results within the liquid pool can not be obtained. From the results it is possible to see a trend that if a constant value is used, too high a value leads generally to a thinner shell in the mould and to shorter pool length. Too small a value leads to thicker shell in the mould and too long a pool length, correspondingly. The parameter A has a minor influence on the solidus isotherm or on the temperature of the solid strand if the value is not badly over or under estimated as can be seen from the results. To over estimate the value to 7 leads to shell thickness to decrease by 15 centimeters near the mould wall. Models using an effective thermal conductivity method can be applied to study the temperatures in the solid shell and related data such as the shell growth and the location of the liquid pool end position. The value of A depends strongly on the geometry of caster, casting speed and it is different in the mould than in the lower part of the machine.

245

5

Conclusions

At the Helsinki University of Technology, 3-dimensional steady state (TEMPSIMU3D) and dynamic heat transfer model (DYN3D) for continuous casting are developed. The models consist of two separate modules: the mould model and the strand model. They are running iteratively and so called gap heat transfer coefficient as a function of the strand surface temperature is used to couple them. The models are validated by industrial measurements. Models using an effective thermal conductivity method can be applied to study the temperatures in the solid shell and related data such as the shell growth and the location of the liquid pool end position, but the temperatures in the liquid pool cannot be calculated very accurately.

6 [1]

[2]

References S. Louhenkilpi, M. Uoti, H. Kytönen and S. Vapalahti, Effect of Thermophysical Material Data on Heat Transfer in Continuous Casting, Modeling of Casting, Welding and Advanced Solidification Processes X, Destin, Florida, USA, May 25-30, 2003, p. 733–740 J. Miettinen, Simulation of Solidification and Calculation of Thermophysical Properties in Binary FCC Copper Alloys-Revised version, Helsinki University of Technology Publications in Materials Science and Metallurgy, TKK-MK-118, 2001

246

247

Micro / Macro Structure

248

249

Spray Forming and Post Processing of Superalloy Rings V. Uhlenwinkel1, N. Ellendt1, M. Walter2, J. Tockner3 1

University Bremen, Bremen, Germany Böhler Edelstahl, Kapfenberg, Austria 3 Böhler Schmiedetechnik, Kapfenberg, Austria 2

1

Introduction

Superalloys are mainly used in airplane engines and power plants (for gas turbines) due to their superior properties at high temperatures relative to conventional alloys. The powder metallurgy route is the standard manufacturing process for high-performance aircraft engine components. However, high cost, risk of contamination, and high oxygen content are disadvantages of this route. Spray forming as an alternative production route was investigated in the past using different superalloys (IN100, IN718, Rene 95, Rene 88DT) [1-8]. Nitrogen or argon was used as the atomization gas. With nitrogen a lower level of residual porosity in the as-sprayed material was achieved because the nitrogen is not inert and reacts with some alloying elements creating nitrides or carbonitrides [1], which affect the precipitation behaviour and finally lead to poor mechanical properties. Therefore argon is preferred as the atomization gas even though the use of argon leads to a higher porosity. It is important to minimize the argon porosity because the initial porosity can affect the material properties of the final product. Therefore, this paper is focused on the optimisation of porosity in spray formed superalloys rings and to improve the understanding of the process.

2

Experimental Procedure

In principle, the spray forming unit used is shown in fig.1. The molten metal (IN718 or U720) was superheated (app. 150 K above liquidus temperature) in the crucible by induction. Subsequently, the melt (150 kg) was pulled into the preheated tundish and left the nozzle as a cylindrical melt stream at the bottom of the tundish. This molten metal stream was atomized with a high velocity gas flow generated by a scanning free fall atomizer (frequency 15 Hz). The substrate was preheated to app. 1150 °C by an induction coil inside the rotating substrate (rotation frequency 1.2 Hz). To optimize the porosity more than 30 spray runs were carried out varying the parameters shown in Table 1. During the spray run several process parameters were recorded. The surface temperature of the substrate and the deposit were measured using a two colour pyrometer (company: Bartec, type: R 2520 - 80). Density measurements were carried out with cubes (app. 1000 mm3) using the buoyancy method (DIN EN 6018) in order to calculate the porosity. Additionally, porosity data was obtained using image analysis of micrographs, which gave a better spatial resolution. Both measurements gave similar results despite the first method being based on a volume and the second on an area.

250

Figure 1: Sketch of the spray forming plant

Table 1: Variation of process parameters Atomizer gas Nozzle diameter Melt flow Gas pressure Gas flow GMR Scanning angle

Argon, nitrogen 6.0–6.5 800–1570 0.27 to 0.43 580–900 0.43 to 1.01 ±8.5 to ±10.5

mm kg/h MPa kg/h °

The as-sprayed rings were machined, divided into four parts and post processed following three different routes: a) Hot isostatic pressing (HIP) b) Forging c) Hot isostatic pressing and forging to eliminate the residual porosity. The process parameters are summarized in Table 2. Table 2: HIP and Forging parameter of the thermomechanically processed quarter-rings Alloy HIP temperature HIP pressure HIP Soaking time Forging temperature Soaking time Ram speed No. of heats

°C MPa H °C H mm/s –

IN718 1140 100 3/6 1000 2 10 1

U720 1140 100 3/6 1100 2 10 2

251

3

Results and Discussion

The use of nitrogen as atomization gas led to pick-up of app. 500 ppm nitrogen in IN718. Even if it would be possible to get lower porosity with nitrogen this was not followed up because the high nitrogen content was not acceptable. Also it was observed that process parameters must be changed if argon is used instead of nitrogen because the cooling effect of argon is much lower (more details are reported in [10]). The scanning of the atomizer led to a typical ring shape displayed in fig. 2. The picture on the left side shows a ring in the as-sprayed condition and on the right side a cross-section is shown. High porosity areas at the left and right margins were not investigated. The core of the rings exhibited a low porosity (see fig. 3). The success of the optimising process is demonstrated by fig. 4 which shows the decrease of the average porosity together with the standard deviation. The average porosity includes measurements across the whole thickness of the ring. Step by step the porosity and the standard deviation was reduced by appropriate selection of spray parameters and finally the average porosity of the argon sprayed IN718 was app. 0.75 vol.%.

Figure 2: (a) Spray-formed ring (IN718) and (b) cross-section

Figure 3: Porosity distribution in an IN718 ring (run 341, buoyancy method)

252

Figure 4: Optimization of porosity in argon sprayed IN718 rings (image analysis)

The boundary conditions of the droplets and the deposit during the impact are responsible for the as-sprayed porosity. Therefore a correlation between process parameters (e.g. Melt flow, gas flow, …) and porosity is not very meaningful. Investigations with other base metals have already indicated that the deposit surface temperature during the spraying process is the key parameter concerning the porosity [11]. Therefore, the measured deposit surface temperature was plotted against the local porosity in fig. 5. A low deposit surface temperature caused a high porosity (cold porosity) and a minimum porosity was observed in the temperature range between 1240 and 1270 °C. The deviation of the porosity at a constant surface temperature may be in-

Figure 5: Dependency between local porosity (buoyancy method) and deposit surface temperature of spray-formed rings (IN718 and U720 using argon or nitrogen), melt flow is given in the legend

253 duced by other parameters but this could not be differentiated. Finally, it can be established that the deposit surface temperature is the most important parameter to control the porosity. Reproducibility of the process is a major need for industrial application. During the whole spray run of approximately 10 minutes the surface temperature was kept constant. This is demonstrated in fig. 6 which shows the surface temperature for different spray runs. The substrate temperature is monitored from the beginning of the operation. It decreases due to the turn on of the atomization gas. The temperature increased rapidly after the spray starts and is nearly kept constant in a range of 1250 to 1260 °C, which finally led to a uniform porosity distribution versus the thickness of the ring.

Figure 6: Reproducibility of substrate and deposit surface temperature during different spray runs with IN718

To exemplify, this porosity distribution versus ring thickness is shown in fig. 7 in comparison to some data from literature [9]. The melt was IN718 and the atomizer gas was argon. The porosity close to the substrate is completely different due to the preheating of the substrate. In this study, cold porosity in the vicinity of the substrate was inhibited totally. The area with a low amount of porosity after [9] was still double as high as the porosity achieved in this investigation. This can make a large difference for the final material properties after post processing. To improve the material properties the porosity must be closed completely. Forging with the parameters mentioned in table 2 was not sufficient to close the porosity. In contrast, hot isostatic pressing (HIP) led to a successful reduction of porosity for the IN718. The as-sprayed sample before HIP had an average porosity of 0.85 Vol.%, after HIP of three hours it was nearly zero and after 6 hours it was totally dense (see fig. 8). At the moment it is not clear what happened to the argon in the gas pores. A few ppm of argon were measured in different samples, but this did not affect the tensile properties (UTS, 0.2%YS, elongation) as reported in [12]. The U720 was HIPed as well using the same parameters, but the densification of this material was not enough to eliminate the porosity completely. The reason for this result is not clear and new experiments are prepared with varying HIP parameters.

254

Figure 7: Local porosity in as-sprayed IN718 rings from [9](marked as Whi-96) and this study (run 354), atomizer gas is argon

Figure 8: Porosity of spray formed IN718 before and after HIP (parameter s. table 2, run 307)

4

Conclusion and Outlook

Superalloy rings IN718 and U720 with a maximum diameter of 500 mm were spray formed successfully. Using nitrogen for atomization led to an increase in nitrogen content (up to 400500 ppm) which was not acceptable. Therefore, argon was used for the optimization process. After several optimization steps an average porosity of 0.75 vol.% was achieved. The porosity was uniformly distributed versus the thickness of the ring and cold porosity in the vicinity of the substrate was avoided due to the preheating of the substrate. A strong correlation between the deposit surface temperature and the porosity with an optimal range between 1240 and 1270 °C has been established. The reproducibility of the deposit surface temperature was extremely

255 good which is important into achieving low porosity. The porosity of the IN718 could be closed by HIP with moderate parameters. The same HIP parameters were not sufficient to close the porosity of the U720. New experiments are in progress to investigate the influence of the HIP parameter to close the porosity. Due to the high impact of the porosity in superalloys for aero engine applications, further investigations are necessary. In the future, a more detailed characterization of the porosity in terms of size and shape are planned. Moreover, there is an open question about the habitation of the argon which is measured in the HIPed material but cannot be seen optically.

5

Acknowledgement

The authors are grateful to the European Union which supported the work under the Growth program with the project number G4RD-CT 2002-00762

6 [1] [2] [3] [4] [5]

References

M.G. Benz, et al., Proc. ICSF-II (1993), Swansea, UK, 171–181 Zhang et al., Proc. SDMA 2000, Bremen, Germany(2000), S. 161–170 M.D. Barratt, A.L. Dowson, M.H. Jacobs, Mat. Sci.& Eng. A 383(2004) 69–77 W.D. Cai, E.J. Lavernia, Mat. Sci.& Eng. A226-228 (1997), 8–12 M.K. Hedges, A.P. Newbery, P.S. Grant, Proc. SDMA 2000, Bremen, Germany, 379–394 [6] R.M. Jones, et al., Proc. ICSF III (1996), Cardiff, UK, 71–78 [7] R.S. Minisandram, et al., Mat. Sci.& Eng. A326 (2002) S. 184–193 [8] S.D. Cai, J. Smugeresky, E.J. Lavernia, Mat. Sci. & Eng., A241 (1998), p. 60–71 [9] E.D. Whitton, P.S. Grant, D. Bryant, Proc. ICSF III, 1996, Cardiff, UK, p. 89–99 [10] V. Uhlenwinkel, R. Attwater, L. Achelis, M. Walter, Proc. PM 2004, Wien, Austria, Vol.5, S. 33-38 [11] V. Uhlenwinkel, M. Buchholz, R. Tinscher, A. Schulz, J. Fischer, R. Schröder, Proc. 4th Int. Conf. On Spray Forming, 11.-13. Sept. 99, Baltimore, USA [12] O. Caballero, D. Fournier, W. Smarsly, Proceedings of Superalloys 718, 625, 706 and Various Derivatives, 02-05 Oct. 2005, Pittsburgh, PA, USA

256

Macro- and Microstructure of Spray-Formed Tin-Bronze D. V. Kudashov, H.R. Müller, R. Zauter Wieland-Werke AG, Ulm

1

Abstract

The first step in the production of metallic materials is melting and casting. Macrostructure, microstructure, and segregation are influenced by the production process. For special materials with strong tendency to segregation powder metallurgical processes are suitable methods with comparatively high production costs. Less sophisticated systems can be cast in permanent moulds or by continuous casting, which is considerably cheaper than powder metallurgy. Spray forming is a process between these extremes considering both aspects. For example the production of Nb3Sn-super-conductors needs homogeneous material with minimized segregation. Spray-formed bronzes with tin content up to 17 % fulfil these demands. Macro- and microstructure of spray-formed tin-bronze is compared with conventionally cast alloys. Possible defects like porosity, the phenomenon inverse segregation, and micro-segregation are presented.

2

Introduction

Copper tin alloys, bronze, are one of the eldest technical alloys invented by mankind. Figure 1 shows a phase diagram of copper tin [1]. In the liquid phase tin is completely soluble in copper. In the solid state the maximum solubility of tin in copper is 15.8 % at 520 °C. With decrease in temperature the solubility is decreasing which does not have practical consequences because of the low diffusion coefficient of tin inside the copper matrix the equilibrium content of soluted tin is only achievable with extreme long annealing periods. During solidification of bronze, there are many factors disturbing the equilibrium state. Therefore, a technical two phase diagram has been found to describe suitable phases (dotted line in Figure 1). According to this diagram copper tin alloys with more than 6 % tin in the solid state at room temperature contain -Phase and -Phase. This microstructure is due to the slow transformation kinetics. With long annealing periods it is possible to achieve the equilibrium microstructure. This annealing process is not able to remove the inverse segregation which appears after continuous casting. Its effect is a higher tin concentration at the edge of the billet than in the centre which does not correspond to the phase diagram. Due to the inverse segregation, the further cold and hot working of copper tin alloys is restricted to bronzes with tin content less than 8 %. With the spray forming process bronzes with tin contents up to 17 % Sn are produced which can be further hot and cold formed without any prior homogenization process. The microstructure of spray-formed high tin bronze discussed in the following chapters. It is shown that the optimum microstructures are only achievable in a small range of process parameters during spray forming.

257

Figure 1: Phase diagram Cu-Sn [1]

3

Spray Forming

3.1

Process

The Wieland-Werke AG produces billets of high-tin bronzes by spray forming (scheme s. Figure 2). The melt is prepared in vacuum furnace with additional stirring coil. Thus the melt is homogeneous and poor of soluted gas. The melt jet is then dispersed to small droplets by an inert gas. The average droplet diameter is about 60 μm. Comparing to the metal powder production, the mass flow rate is significantly higher, namely up to 35 kg/min with a single atomizer and up to 70 kg/min with a twin atomizer. Growth direction is vertical. The Wieland spray forming device and process are described in detail in [2].

3.2

Influence of Process Parameters on Quality of the Spray Formed Billet

The spray forming process is specified by numerous parameters, which influence shape, microand macrostructure, porosity, segregation of the billet. Additionally they interact in different ways. For example the diameter of the casting nozzle and the melt level in the tundish define the melt flow rate (Torricelli formula). Figure 2 outlines the most important process parameters. The preset parameters are defined by the design of the plant or can not be changed during the run. In principle, the control parameters can be changed, but practically only the gas mass flow rate and the withdrawal speed w are controlled. At the Wieland plant the melt temperature and the gas to metal flow rate ratio (GMR) are automatically controlled. Normally scan frequency

258

Figure 2: Spray forming process parameters

and substrate rotation are kept on a default value. The operator adjusts only the withdrawal speed to keep the billet diameter constant. This is necessary, if the metal flow rate or the compacting rate changes. The parameter GMR influences the atomization (droplet size and velocity distribution in the spray cone), the compacting rate (ratio of atomized melt mass to deposited mass), the cooling of the droplets and the cooling of the deposit. In the following the effect on structure of high alloyed tin bronze billets is discussed.

3.3

Porosity

A higher GMR causes the reduction of the mean particle diameter in the spray cone. Therefore the fraction of solidified particles is increased and the heat entry into the deposit is reduced. At the time when the particles hit the target only a small percentage of particles is still liquid. This percentage of particles is too small to fill up all gaps between the solid material with liquid which causes pores. So with increasing GMR porosity does increase (compare Figure 3a, b).

Figure 3: CuSn13,5 a)GMR 0,53, porosity: 3,6 %, b) GMR=0,47, porosity: 1,2 %

259

Figure 4: Influence of Ti-content on porosity in CuSn15,5 at constant gas-metal-ratio G/M = 0,5; a) without Ti; b) with Ti

Another important factor is the chemical composition of the melt. Figure. 4a shows the structure of CuSn15 with some pores. At similar spray conditions, but with 0.2 % Ti the structure is nearly 100 % dense (Figure. 4b). Influence of reactive elements on the porosity are analysed in detail in [4]. So it seems to be very easy to avoid porosity by decreasing gas-flow rate and reactive elements – if there would not be segregation.

Figure 5: Sn-concentration in the cross section of CuSn12 billets spray-formed with different gas flow rates G in m3/s and constant melt flow rate

3.4

Segregation

The driving force for the macro-segregation phenomenon is assumed to be the centrifugal force of the rotation of the billet during the spray forming process [4]. Figure 5 demonstrates the Sn – concentration pattern in the cross section of CuSn12 sprayformed billets. The lower curve shows the segregation across the diameter at a low, and the upper at a high gas flow rate at the same melt flow rate. The amplitude of tin content variations across the cross section is decreased by increased GMR. On the other hand the increase of GMR causes an enhanced microsegregation. This microsegregation is present in the form of -Phase agglomeration (Figure 6a). In the microscopic view of a microetched billet cross section they are visible in shape of concentric circles inside the spray-formed material. In the longitudinal section the circles turn out to be 3 dimensional dome

260 shaped agglomerations (Figure 6b). This dome shape obviously is the image of the surface of the billet in process at a certain time during the spray forming process. By optimizing the process parameters all these micro-segregation phenomena are avoidable.

a)

b)

Figure 6: Spray-formed 13,5 %-tin bronze with microsegregation a) microstructure b) macrostructure

4

Differences from Conventional Casting Process

4.1

Macrostructure

Figure 7a and 7b illustrate the macrostructures of spray-formed and cast bronzes. While sprayformed bronze shows fine homogeneous grains in the longitudinal direction (Figure 7 a), the grains in continuous cast bronze are coarse and inhomogeneous (Figure 7, b).

4.2

Microstructure

The difference in microstructure between cast and spray-formed bronze with 15.5 % tin is illustrated in Figure 8. The permanent mould cast shows a dendritic structure. Between the dendrites the tin-rich -phase is enriched (Figure 8, a). The high fraction (about 30%) of this brittle, low melting phase prevents cold and hot forming. In the spray-formed structure (Figure 8, b) the

a)

b)

Figure 7: a) spray-formed 13,5 %-tin bronze with fine and homogeneous grain structure, grain size about 60 μm, b) continuous cast 12 %-tin bronze with course and inhomogeneous grains [3]

261

a)

b)

Figure 8: Microstructure of 15.5 %-tin bronze produced by a) spray forming and b) die casting. [3]

fraction of -phase amounts to approx. 5 % and particles are not connected. This structure allows cold and hot forming.

4.3

Segregation

In continuous casting of high tin containing copper alloys the so called inverse segregation is well known. In early days of the continuous casting development various publications dealt with this phenomenon [6, 7, 8, 9]. Sucking by volume contraction and metallostatic pressure are most likely explanations. The strong influence of metallostatic pressure was impressively demonstrated by Ohm [10] Figure 9 shows a plot of tin-concentration in dependence of the location inside a continuous cast bronze billet with 8 % tin and with a diameter of 7 inch. The tin concentration at the billet surface is considerably higher than in the center. The tin concentration differences reach values up to 9 %. This macrosegregation is not removable by heat treatment [1]. During spray forming process hydrostatic pressure does not exist. If spray conditions are optimized, macrosegregation is reduced to a concentration difference of about 1 wt-% inside the billet. Figure 5 shows a concentration plot of a spray-formed 12%-tin bronze at optimized spray

Figure 9: Tin concentration gradient in continuous cast 8 %-tin bronze billet, diameter 178 mm, cross section [11]

262 parameters. Macrosegregation is limited to 1 wt-% tin. The driving force for this small segregation phenomenon is assumed to be the centrifugal force of the rotation of the billet during the spray forming process [5].

5

Application

5.1

Superconductor

Superconducting wire [2, 3] with magnetic field strength up to 22 Tesla are usually low temperature superconductors which run at liquid helium temperature. The super-conducting material is Nb3Sn, produced by the bronze method. This type of magnet is applied in Nuclear Magnetic Resonance (NMR) spectrometers for analytical applications and in magnets for nuclear physics. Nb3Sn is a brittle inter metallic phase, and so it is not a metallic material which can be produced and brought into the shape of a wire by normal industrial melting, forming and machining processes. It must be produced by co-working of the metals Niobium and bronze until the compound has the shape of a wire. Finally a diffusion annealing process generates the Nb3Sn super-conducting phase. The bronze serves on one hand as carrier for tin, on the other hand as heat and current carrying material in case of super-conductivity break-down. The manufacturers of superconducting wire need a good workability of the tin-carrying bronze. By means of forming processes the tin is brought as near as possible to the niobium. The final heat treatment lets tin diffuse to the niobium and transfers both into the Nb3Sn phase. In order to show the long route from material production to the final product, the production process of superconducting wire is illustrated in Figure 10 in six separate steps: 1. Spray forming of bronze billet 2. Extrusion and cold drawing of bronze tubes and rods 3. Assembling bronze tubes and niobium rods

Figure 10: Production process of bronze for super-conducting Nb3Sn wire (schema) [2, 3]

263 4. Co-working of Niobium and bronze (extrusion, cold-drawing, heat treatment) into the shape of a wire 5. coiling of wire into final shape inside the magnet 6. Generation of Nb3Sn by final annealing process. During this process a multi-filamentary composite material is generated. Figure 11 shows a cross section through a stabilized type of superconductor wire in a production stage cold-drawing of the niobium-bronze-composite (production step 4). At a first glance 55 hexagonal grey dots surrounded by bronze are visible. Each dot consists of 85 niobium filaments which also are surrounded by a bronze matrix. The dimensions of niobium filaments and bronze channels finally reach the size of 5 to 20 μm. The final production step is diffusion heat treatment.

Figure 11: Stabilized superconductor with 55 x 85 = 4675 niobium filaments in Cu-Sn-matrix, wire diameter 1 mm (by courtesy of EAS/Hanau)

5.2

Machinable Materials for Connectors

Between the standard machinable copper alloys and the machinable, high-strength but expensive beryllium copper CuBe2Pb exists a gap in term of mechanical properties and in price, which can be closed by alloys of the system Cu-Sn. Wieland-Alloy CuSn13,5Pb shows an unusual combination of high yield strength (about 900MPa ) and low modulus of elasticity (85 GPa). This combination of properties allows employing the alloy for the spring application (connector). Good free cutting characteristics are achieved by the addition of Pb. Due to the spray forming process a fine and homogeneous lead distribution is obtained. Free cutting is substantially improved in comparison to the lead free alloy CuSn13.5. The lead containing alloy exhibits a more favourable chip shape (see. Figure12). Mechanical properties are not influenced by the lead addition.

264

a)

b)

Figure 12: Different shape and length of machining chips of high tin bronze a) without lead, b) with lead [12]

6

Conclusion

The spray forming process is specified by numerous parameters, which influence the properties of the materials. The influence of GMR on micro- and macrostructure, porosity and segregation of high alloyed tin bronze billets is discussed. The process parameters have to be set very precisely to get minimized segregation on one hand and limited porosity on the other hand. Spray-formed high-tin bronzes have various advantages compared to cast high-tin bronzes. The grain structure of spray-formed material is fine and homogeneous. Macrosegregation is minimized to 1 % tin concentration difference across the billet. Owing to these advantages the spray-formed bronze with tin content up to 17 % can be hot extruded and cold drawn without any prior homogenization. Spray forming is a modern production process which is able to shift the classical border between wrought and cast alloys to considerably higher tin contents.

7 [1]

References

K. Dies, Kupfer und Kupferlegierungen in der Technik, Springer-Verlag, Berlin/Heidelberg/New York, 1967, 514 [2] H. R. Müller, R. Zauter, Erzmetall, 2003, 56, Nr. 11, 643–650 [3] R. Zauter, K. Ohla, H.R. Müller, J. Maier, Intern. Conf. On Spray Deposition and Melt Atomization, SDMA 2003, Bremen – Germany, 2003, 5-113–5-122 [4] H. R. Müller, K. Ohla, R. Zauter, M. Ebner, Materials Science and Engineering A, 2004, 383, 78–86 [5] H.R. Müller, S. Hansmann, K. Ohla, Intern. Conf. On Spray Deposition and Melt Atomization, SDMA 2000, Bremen – Germany, 2000, 205–218 [6] H. Voßkühler, Z. Metallkunde, 40, 1949, 8, 305–311 [7] W. Roth, Z. Metallkunde, 40, 1949, 12, 445–460 [8] H. Kästner, Z. Metallkunde, 41, 1950, 8, 193–205 [9] H. Kästner, Z. Metallkunde, 41, 1950, 8, 247–254 [10] L. Ohm, Metall 43, 1989, 4, 520–524 [11] DKI - Deutsches Kupferinstitut, Berlin – Düsseldorf (1965), 14 [12] K. Ohla, H. R. Müller, M. Keppeler, A. Bögel, Metall, 55, 2001, 4, 213–215

265

Influence of the Crystallization Conditions on the Microstructure and Mechanical Properties of TiAl- and Ti3Al-Based Alloys B. A. Greenberg, N. V. Kazantseva, A.E. Volkov, Yu. N. Akshentsev Institute of Metal Physics, Ural Division, Russian Academy of Sciences, Ekaterinburg, Russia

1

Introduction

Titanium aluminides are of practical interest alloys providing unique combination of physical and mechanical properties. The TiAl-based alloys are characterized by good strength at temperatures up to 650 °C (a600 MPa), but poor room-temperature plasticity (a1–2 %). The latter property limits their commercial applications. The efforts aimed at the improvement of the plasticity of these alloys include both the design of new alloys and new methods of their production. The optimum combination of properties can be attained by the formation of a specific structure, such as a fully lamellar two-phase TiAl/Ti3Al structure with controlled content of J and D2 phases [1–2]. However, the ultimate strength, plasticity, and fracture behavior of such alloys are very sensitive to the orientation and microstructure of lamellae. In many works the properties were found to be anisotropic, depending on the interlamellar spacing, domain size, which in turn depended on the crystallization condition of and chemical content of the alloys [3–5]. This work presents a comparative study of influence of the crystallization conditions on the microstructure and mechanical properties of the TiAl- and Ti3Al- based alloys with the additions of V, Nb, and Mo. All of these alloys were prepared by the special method of pulsed volume pressing (PVP) [6].

2

Experimental Procedure

The alloys in the PVP method were melted in a vacuum of 10 Pa by an electric arc formed between the initial sample and a consumable electrode of the same composition. The melt was overheated by about 10–15 degrees above the melting temperature (this instant was detected by an optical sensor), then the copper (or steel) mold was pressed to the sample, and simultaneously the melt mirror was subjected to gas (Ar) pressure pulses. This imparts a high velocity (5–20 m/s) to the melt, which rapidly fills in the mold. In addition, a punch installed in the bottom part of the mold moves toward the melt and presses it, providing additional pressure on the mold. At the instant of casting, the melt undergoes a vibration at a frequency of 25–50 Hz. The Ti-48.%Al-1.%V alloy was prepared in a form of rods of 8 mm in diameter by various regimes such as casting into copper and steel molds under an additional pressure and casting into a steel mold under normal pressure. All of others alloys were prepared by casting into copper mold under additional pressure. X-ray diffraction examination was performed using a DRON-3 diffractometer (Cu KD radiation). The microstructure was examined using a Neophot-2 optical microscope and JEM-200CX electron microscope. Mechanical tests of samples 3u3u4,5 mm in size were performed in air at

266 room temperature by compression using an INSTRON machine at strain rate of 0,05 mm/min. The pores in the samples were studied using the ultrasound method.

3

Results and Discussion

3.1

Ti-48 at.% Al-1 at.%V

According to X-ray diffraction data the alloy prepared by all the regimes used consists of two ordered phases such as TiAl (J) and Ti3Al (D2). The alloy that was cast in a copper mold under pressure has a very fine lamellar structure. The grains in this sample are rather uniform, and their average size is 40 Pm. The TEM study shows a fine lamellar structure with a high dislocation density in the J lamellae. The D2-phase is present as fine lamellae at twin boundaries in the J-phase. The width of J and D2 lamellae is 0.06–0.2 and 0.02–0.06 Pm, respectively (Figure 1a–b). The alloy that was cast into the steel mold without pressure consists of grains both containing lamellae and free from them. The grains in this alloy are less uniform in size, which is

Figure 1: Microstructure of the Ti-48%Al-1%V alloy, TEM: (a)–(b)- cast into a copper mold under pressure, (a) dark-field image in the (1-11) J-phase reflection, (b)- electron-diffraction pattern for (a), zone axis [10-1]J II [11–20]D2; (c) cast into a steel mold without pressure, dark-field image in the (111)J reflection; (d) cast into a steel mold under pressure, dark-field image in the (1-10)D2 reflection

267 also 40 Pm on average. The alternating J and D2 lamellae are 0.05–0.4 and 0.05–0.1 Pm wide, respectively (Figure 1c). Like the alloy cast into the copper mold, the alloy cast into steel mold under pressure has a lamellar structure (fig1d). The grains are widely differing in size (from 20 to 200 Pm). Both J and D2 lamellae are wider (0.2–0.5 and 0.03–0.1 Pm).

3.2

Ti-46 at.% Al-1,3 at.% V

According to X-ray diffraction data this alloy, as the previous one, consists on two phase: J and D2. Parameters of the crystal lattices are: in D2 phase – a = 0.5726 r 0.0005 nm, c = 0.4615 r 0.0003 nm; in J phase – a = 0.3998 r 0.0001 nm, c = 0.4024 r 0.0003 nm. The structure of this alloy consists on the grain with different form. Some part of the sample contents the elongated grains. The average size of the grains is about 60 Pm. The lamellar structure forms in the middle of the grains; J and D2 lamellae are | 0.3 and 0.08 Pm wide, respectively (Figure 2a).

3.3

Ti-45 at.% Al-1at.% V

X-ray diffractogram of this alloy has also the lines of two phases: J and D2. Parameters of the crystal lattices are: for D2 phase a = 0.57086r0.0005 nm, c = 0.45626r0.0003 nm; for J phase a = 0.3982r0.0001 nm, c = 0.4117r0.0003 nm. Zone of the elongated grains occupies the most part of the sample. The central part of the grains has thin lamellae: for J | 0,2 Pm and for D2 | 0,03 Pm in wide (Figure 2b)

Figure 2: Microstructure of the alloys, TEM: (a) Ti-46%Al-1,3%V, dark-field image in the (0-12) D2-phase reflection, (b) Ti-45%Al-1%V, dark-field image in the (002) J reflection

3.4

Ti-34 at.% Al-1,6 at.% Nb-0,5 at.% Mo-0,3 at.% Cr

According to the equilibrium diagram this alloy must have D2 single phase content. However, the addition of the beta stabilized elements and high rate of cooling allowed serve high temperature E0. The diffractograms of this alloy content the lines of two phases: E0 (2) and D2 (D19). Pa-

268 rameters of the crystal lattices are: for D2 phase a = 0.574r0.0005 nm, c = 0.459 r 0.0003 nm; for E0 phase a = 0.319 r 0.0001 nm, c = 0.4117r 0.0003 nm. The structure of this alloy consists on the grains with Widmanstatten structure inside of them. The sizes of grain are from 50 up to 200 Pm. It is found that little grains form in the different places of the sample, not only in the middle of the sample as it can be seen in the case of usual cast ingot. The samples don’t have a texture or a dendrite structure characteristic. On the TEM pictures one can see the short and thick D2 plates and thin E0 plates between them. Many dislocations are seen inside of the D2 plates. We also found the small particles of Z-phase (B82) inside of the E0 plates (Figure 3a–c). After the additional aging at 900 °C–5 h., the Z-phase reflects disappeared and quantity of the dislocations were reduced (Figure 3d). The techniques of the preparation of titanium aluminides with lamellar structure determine the grain size, the orientation of lamellae and their thickness. The alloys prepared in the laboratory by zone melting and having oriented lamellar structure (polysynthetically twinned crystals, PST) have very high mechanical properties at a grain size varying between 25 and 50 Pm. The strength of such alloys reaches 1100 MPa [2]. The maximum strength of PST TiAl alloys is

Figure 3: Microstructure of the Ti-34%Al-1,6%Nb-0,5%Mo-0,3%Cr alloy, TEM: (a) bright-field image; (b) dark-field image in the (011)Z reflection; (c) electron-diffraction pattern for (b), zone axis [100]Z; (d) microstructure of the alloy after aging 900 °C–5 h, bright-field image

269 1800 MPa at plasticity of 23 %, the J interlamellar spacing in the alloy is 1.4 Pm. In turn, the strength of polycrystalline TiAl samples prepared by complex thermomechanical treatment and having a nonoriented lamellar structure is as small as 600 MPa, and plasticity of the alloy is about 6 % [7]. Compared to the “pure” PST TiAl alloys, the PST alloys containing vanadium exhibit a higher plasticity upon both tension and compression, the plasticity of the Ti48,4at.%Al-0,6at.%V alloy upon compression reaches 28 %. The vanadium-containing alloys are characterized by a more uniform distribution of D2 lamellae and by a thickness of J lamellae of 0.2–2 Pm [8]. Table 1: Strength characteristics of the alloys prepared by PVP method. N 1

2 3 4

Alloy Ti-48 at.% Al-1 at.% V a) casting into copper mold under pressure b) casting into steel mold under pressure c) casting into steel mold without pressure Ti-46 % Al-1,3 % V. Ti-45 at.% Al-1at.% V Ti-34 at.% Al-1,6 at.% Nb-0,5 at.% Mo-0,3 at.% Cr a) casting into copper mold under pressure b) + aging 900°C–5 h

H %

Vb, MPa

V0,2, MPa

32,6 26,5 28,6 18 17

1088 1041 972 1030 1184

685 734 568 844 1007

10 19

1920 1700

1314 872

So, according to the experimental results obtained in this work we could figure out some of the external factors that responsible to the successful structure of alloys with high mechanical properties: 1. Vibration and addition pressing under crystallization are beneficial for small grain structure, which form by the dendrite breakage. 2. The rough and very cold mold, lowed temperature of casting, and convection flows are also beneficial for small uniform equiaxed grains formation. 3. The cooling rate varying in the mold with different heat capacities (0.385 cal/g K for copper and 0.12 cal/g K for steel) also substantially affects the grain size, the uniformity of its distribution and the lamella thickness. In the Ti3Al-base alloy high rate of cooling allows serve “soft” plates of E0 phase.

4

Acknowledgments

This study was supported financially by the Program "National technological basis", grant 47/04/768-2004, and the Russian Fund for Basic Research Ural, grant 04-03-96008.

270

5 [1] [2] [3] [4] [5] [6] [7]

[8]

References Kim Young-Won, JOM. 1989, 41, 24–30 S. Naka, M. Tomas, and T. Khan, Mat. Sci. and Technol. 1992, 8, 291–298 M. Yamaguchi and Y. Umakoshi, Progress in Materials Science. 1990, 34, 60–73 N.V. Kazantseva, A.E. Volkov, B.A. Greenberg, A.A. Popov, V.V. Yurovskikh, Phys. Met. Metallogr. 2001, 91, 173–178 B.A. Greenberg, N.V. Kazantseva, A.E. Volkov, etallography and Temperature treating of the Materials. 2005 (in print) A.E. Volkov, A.V.Frolov, V.N.Boiko, RU Patent 2 106 226, 1996 T.Maeda, M.Hosomi, M.Okada. in The Proc. Symp. Sponsored by the Structural Materials Division (SMD) of TMS (Ed.: Young-Won Kim et.al), Warrendale, Pa, 1995, 771–778 K.F.Yao, H.Iniu, K.Kishida, and M.Yamaguchi, Acta Metall.Mater. 1995,43, 1075–1086

271

Effects of Process Parameters on the Characteristics of the Billet Sump and Related Defect Formation during DC Casting of Aluminum Alloys D. G. Eskin1, L. Katgerman2 1

Netherlands Institute for Metals Research, Delft, The Netherlands Delft University of Technology, Dept. Materials Science and Engineering, Delft, The Netherlands

2

1

Introduction

There are three main process parameters that one can change upon direct-chill casting, i.e. casting speed, water flow rate, and melt temperature. Water flow rate should typically assure efficient cooling (nucleate boiling conditions), and its increase above this sufficient level does not affect much the structure and defect formation [1] The other two process parameters are extremely important, e.g. in occurrence of casting defects [1, 2, 3, 4, 5]. Casting speed and casting temperature are also known to affect the structure formation during solidification. This is because of their influence on cooling conditions, melt flow and geometry of liquid and semi-liquid parts of the billet [1, 3, 6]. The increase in casting speed results in proportional deepening of the liquid sump, increasing of the mushy zone thickness, and in overall acceleration of solidification [3]. The melt temperature, however, has not received much attention from the viewpoint of its influence on the structure and defect formation in DC cast billets. In a very good review of DC casting [7], the authors mention the melt temperature only once and then among other parameters that can also affect the structure and quality of a billet. As for the casting defects, two are most frequently encountered in practice, i.e. macrosegregation and hot tearing. Macrosegregation patterns generally depend on the distribution coefficient k and are obviously linked to the morphology of the forming solid phase, permeability of mushy zone, magnitude of solidification shrinkage, ratio between the shrinkage velocity and the direction and velocity of melt flow, level of solute rejection to the melt, and movement of the solid phase in the liquid and slurry regions [1, 8]. Macrosegregation is known to increase with the casting speed [1, 3]. There are controversial data on the effect of melt temperature on macrosegregation [9, 10]. Shestakov [11] showed analytically that the increase in melt temperature narrowed the transition region in the billet and accelerated the solidification (local solidification time is shortened). The most important casting parameter that affects hot tearing is also the casting speed, and the main reason for that is believed to be connected to greater temperature gradients and, therefore larger thermal strains [3, 5, 12]. The optimum casting speed is a compromise between productivity, alloy composition, billet size, and quality (structure and defects). The effects of melt temperature on hot tearing was previously studied only for shape castings [13]. It was shown that higher melt temperature results in higher hot tearing susceptibility. However, direct applicability of these results to direct-chill casting is unclear. The analysis of available literature sources reveals, therefore, gaps in published experimental and modeling information. Only few experimental data are reported on the interrelation be-

272 tween the sump and mush dimensions, and melt flow patterns on one side and the macrosegregation and hot tearing on the other side. The aim of this paper is to report the results of systematic examinations of billets of binary Al–Cu alloys cast under different process conditions. Effects of casting speed and melt temperatures on macrosegregation and hot tearing are examined and correlated to computer simulated solidification patterns.

2

Experimental and Numerical Procedures

A series of experiments was performed in a pilot direct-chill casting installation in Delft University of Technology. Round billets 195 mm in diameter were cast in a 200-mm hot-top mould. Binary Al–Cu alloys containing 1 to 5% Cu were prepared using 99.85% pure aluminum and Al–47.7% Cu master alloy. Two types of experiments were performed, i.e. upon stationary and transient stages of casting. In the first case, the casting parameters were changed stepwise with producing at least 200 mm of a billet at each casting regime. In the second case, the casting speed was ramped up and then down at a constant rate, and the structure was examined at different casting speeds during the ramping. The detailed description of experimental equipment and experiments can be found elsewhere [1, 4, 5] The structure of the billets was examined in an optical microscope Neophot 30. The chemical composition was measured using a spark spectrometer SpectroMax. The hot-cracking susceptibility was estimated as the ratio of the billet cross-section affected by cracks to the total billet cross-section. The flow patterns and sump profiles were assessed numerically using the Flow3D and CFX software [1, 4] and, the case of sump profile, also with the MSC.Marc software [5]. Measurements of temperatures in the hot top and sump of the billet and measurements of the sump depth with a rod during DC casting experiments were used for validation of the numerical results.

3

Results and Discussion

3.1

Macrosegregation and Floating Grains

The magnitude of macrosegregation is clearly affected by the casting speed as shown in Fig. 1a. Casting speed also greatly influences the geometry of the billet sump as demonstrated in Fig. 1b, c. One can notice a distinct correspondence between the extent of macrosegregation in the radial direction of the billet and the width of the transition region in the billet. This correlation is illustrated in Fig. 1d. One can see that the relative copper concentration normalized to the vertical distance between solidus and liquidus isotherms is not dependent on the radial position in the billet cross-section, except for the subsurface region where additional mechanisms of macrosegregation, e.g. shrinkage-driven flows, may act complementary to the thermo-solutal convection [1].

273

Figure 1: Effect of casting speed under steady-state casting conditions on macrosegregation (a) and parameters of the sump (b, c) in the Al–4.3% Cu billet; and (d) macrosegregation of copper (as in a) normalized to the width of the transition region (as in c). Numbers in b relate to the centre of the billet: 1, calculated total depth of the sump; 1’, measured total depth of the sump; 2, calculated depth of the liquid pool; and 3, calculated distance between liquidus and solidus isotherms.

Figure 2: (a) Effect of melt temperature in the furnace on the macrosegregation and (b) the flow pattern at a melt temperature at the inlet to the hot top of 725 °C (corresponds to the melt temperature in the furnace 760 °C) in the Al–2.8% Cu billet cast at 200 mm/min (transient casting stage). Water flow rate 150 l/min.

274 The melt temperature does not affect much the segregation pattern in the bulk of a billet, which is in line with the opinion of Tarapore [9] and Reese [10]. The subsurface segregation, however becomes more pronounced on increasing the melt temperature as illustrated in Fig. 2a. Computer simulations show that the transition zone in the center of the billet becomes narrower with increasing the melt temperature whereas the transition zone at the periphery of the billet expands [4]. Therefore the correlation between the dimensions of the transition region and the extent of macrosegregation shown in Fig. 1d holds here as well. In addition, the computer simulation of flow patterns at high melt temperatures show strong currents (of solute-enriched liquid) directed towards the billet surface, Fig. 2b. The appearance of floating grains is frequently linked to macrosegregation patterns, assuming that solute-depleted grains brought to the center of the billet by melt flow contribute to the negative segregation [2]. Floating grains are characterized by a coarser internal structure indicative of their longer solidification time. Recently, based on the numerically assessed flow patterns in the slurry part of the billet, we suggested a mechanism of floating grains formation [1]. This mechanism implies that floating grains are formed in the upper part of the slurry region

Figure 3: Effects of casting speed and melt temperature of the distribution of floating grains in the horizontal cross-section (a, b) and on the calculated flow patterns in the Al–2.8% Cu billet (c–f): a, c, e, casting speed 100 mm/min, steady-state stage; b, 200 mm/min, transient stage; and d, f, 200 mm/min, steady-state stage. Melt temperatures are (c, d) 700 °C and (e, f) 760 °C. Water flow rate is 150 l/min. Temperatures are the melt temperatures in the furnace.

275 and closer to the periphery of the billet, and then are transported by the flow to the central part of the billet. The effective cooling solidification time for these grains is thus much longer than it can be presumed just from the vertical dimension of the transition zone. The effects of casting speed and melt temperature on the distribution and amount of floating grains are illustrated in Fig. 3. In the range of low casting speeds (Fig. 3a), „floating” grains spread across the entire cross-section of the billet at a low melt temperature (700 °C) and are confined to the central part of the billet at higher melt temperatures. Their amount decreases with increasing melt temperature, which agrees well with earlier predictions [9]. In this casting regime, the increase in melt temperature results in diminishing of the region with stagnant flow in the central part of the billet where the „floating” grains may grow (Fig 3c, e), with a larger fraction of „floating” grains going upwards in the slurry zone and, apparently, remelting. In the range of high casting speeds, the distribution of „floating” grains is similar at any given melt temperature, with the maximum fraction in the central part of the billet, Fig. 3b. However, the total amount of these coarse grains increases with melt temperature. We can suggest that deepening of the sump and more severe currents in the vicinity of the mushy zone with increasing the melt temperature and the casting speed as shown in Fig. 3d, f create more possibilities for „floating” grains to form, grow and settle.

3.2

Hot Tearing

Our experimental observations show that susceptibility to hot tearing during direct-chill casting depends on the composition, increases with increasing the casting speed, and decreases with increasing the melt temperature [4, 5]. Hot tearing susceptibility of an alloy depends on the dimensions and properties of the socalled vulnerable solidification range that is confined between the isotherms of rigidity and the solidus [14], or in the case of the billet – on the dimensions of the mushy zone. The increasing casting speed results in widening of the transition region, especially in the centre of the billet, where hot cracks are typically observed (see Fig. 2b, c), and in higher tensile strains concentrated in the central part of the billet [5]. Computer simulations show that on increasing the melt temperature, however, the mushy zone either retains its dimensions in the centre of the billet or becomes thinner, meaning that under the same casting conditions the alloy spends the same or less time in the vulnerable range [4]. At the same time, the liquid bath (distance from the melt surface to the liquidus) and the sump as a whole become deeper with increasing melt superheat, resulting in a larger metallostatic pressure on the mushy zone and, therefore better feeding of the melt to potential cracks, Fig. 3e, f. This is revealed in the structure as „healed” cracks – relatively long paths spread along several grain boundaries and filled with eutectic, the amount of the latter also increasing with the melt temperature [4].

4

Conclusions

Experimental and numerical studies of DC casting show the direct correlation between the degree of macrosegregation, hot tearing susceptibility, and amount of floating grains, on one hand, and the vertical distance between the liquidus and solidus isotherms and melt flow patterns in the sump of the billet, on the other hand. Macrosegregation and hot tearing susceptibility increa-

276 se with the casting speed. The increased casting temperature, however only affects the subsurface segregation and causes less hot tears.

5

Acknowledgements

This work is done within the framework of the research program of the Netherlands Institute for Metals Research (www.nimr.nl), Projects MP 97014 and MC 02134. Contribution of J. Zuidema, Jr., V.I. Savran, and Q. Du to computer simulations is highly appreciated.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

References D.G. Eskin, J. Zuidema, Jr., V.I. Savran, L. Katgerman, Mater. Sci. Eng. A 2004, 384, 232–244 E.F. Emley, Intern. Met. Rev. 1976, Review 206, 75–115 V.A. Livanov, R.M. Gabidullin, V.S. Shepilov, Nepreryvnoe lit’e alyuminievykh splavov (DC Casting of Aluminium Alloys), Metallurgiya, Moscow, 1977, p. 168 D.G. Eskin, V.I. Savran, L. Katgerman, Metall. Mater. Trans. A 2005, 36A (in press) Suyitno, D.G. Eskin, V.I. Savran, L. Katgerman, Metall. Mater. Trans. A 2004, 35A, 3551–3561 H. Nagaumi, Sci. Technol. Adv. Mater. 2001, 2, 49–57 J.F. Grandfield, P.T. McGlade, Mater. Forum 1996, 20, 29–51 M.C. Flemings, Solidification Processing, McGraw-Hill, New York, 1974, p. 364 E.D. Tarapore in Light Metals 1989 (Ed.: P.G. Campbell), The Minerals, Metals and Materials Society, Warrendale, USA, 1989, p. 875 J.M. Reese, Metall. Mater. Trans. B 1997, 28B, 491–499 A.D. Shestakov, Izv. Ross. Akad. Nauk, Metally 1996, 6, pp. 130–138 M. M’Hamdi, A. Mo, C.L. Martin, Metall. Mater. Trans. A 2002, 33A, 2081–2093 I.I. Novikov, Goryachelomkost tsvetnykh metallov i splavov (Hot Shortness of Non-Ferrous Metals and Alloys), Nauka, Moscow, 1966, p. 299 D.G. Eskin, Suyitno, L. Katgerman, Progr. Mater. Sci. 2004, 49, 629–711

277

Effect of Casting Speed and Grain Refining on Macrosegregation of a DC Cast 6061 Aluminum Alloy R. Kumar Nadella1, D. Eskin1 and L. Katgerman2 1 2

Netherlands Institute for Metals Research, Delft, The Netherlands Delft University of Technology, Department of Materials Science and Technology, Delft, The Netherlands

1

Introduction

Direct Chill (DC) casting remains a major processing route to produce large aluminum ingots, which are used for downstream processing. During this process, grain refining is normally employed to achieve finer equiaxed grains. However, the occurrence of defects such as hot cracking, macrosegregation etc. needs to be properly understood from the quality point of view. Process parameters such as casting speed, water flow rate and melt temperature can be tailored to get a sound product. Macrosegregation, which is the non-homogeneous distribution of alloying elements over a large length scale, needs to be minimized as it affects the properties of the finished product. This is because during subsequent homogenization, while microsegregation is reduced, macrosegregation remains essentially unaffected. The nature of solidification profile during DC casting together with convective flows in the ingot sump is responsible for this phenomenon [1]. Further, the degree of macrosegregation can be such that the composition in certain regions across the thickness of the ingot may be outside the registered limits established for the alloy. In addition to the casting process parameters, grain refining has an important effect on macrosegregation. Despite the large amount of literature and excellent reviews on DC casting [1,2], there has been little systematic work concerning the effect of grain refining on macrosegregation. This is particularly evident in case of commercial Al alloys, which contain more than one alloying element.

2

Experimental Methods

Commercial-scale DC casting experiments were conducted in a pilot DC casting installation at the Delft University of Technology, the details of which can be found elsewhere [3]. Experiments were carried using an AA 6061 with the composition (in wt.%): 0.51 Si, 0.23 Cu, 0.95 Mg, 0.12 Fe, 0.007 Mn, 0.076 Cr, 0.016 Zn with and without grain refining. Different casting speeds were used in these experiments while maintaining a constant water flow rate (170 l/min) and casting temperature (715 °C). The sump depth is measured by a digital length meter. DC cast billets of 192 mm in diameter obtained from a 200 mm round hot top mould are longitudinally sectioned in the center and samples of approximately 20 mm wide and 20 mm high were cut at suitable locations in the horizontal cross section of the billet. Care was taken to ensure that the sampling represents the steady state conditions during DC casting. These rectangular bars were analyzed by a spark spectrum analyzer across the billet diameter for the composition variation. Measurements were taken in all 4 sides at regular intervals of approx. 10 mm, and the

278 average values are reported. The absolute error in these measurements is ± 0.05 wt. %. Concurrent microstructural observations were carried out close to the surface and in the central portion of the billet. Samples anodized in a 3% HBF4 water solution were observed under cross-polarized light to check the grain size.

3

Results and Discussion

The process chart for a typical DC cast experiment along with an illustration of an AA 6061 billet is shown in Figure 1. To obtain a grain-refined ingot, a known quantity (2 kg/tonne) of grainrefiner in the form of Al-5Ti-1B was added to the liquid metal in the furnace just prior to commencing the DC casting. The average Ti concentrations were about 0.01 wt.%.

Figure 1: Typical process chart for the DC casting experiment of AA 6061

To study the macrosegregation patterns in this alloy, the composition profiles of the major alloying elements (Mg and Si) are plotted across the whole diameter (192 mm). The general trend for all the observed patterns is a negative segregation in the center and minor positive segregation at the mid-radius. Further, strong segregation zones are noticed close to the surface. The effect of grain refinement on the macrosegregation profiles is shown in Figure 2 for a casting speed of 8 cm/min. It can be seen that grain refining does not seem to have a considerable influence on the composition profiles. However, at higher casting speeds, it is observed (Figure 3) that both grain refined and non-grain refined samples exhibited higher segregation levels, particularly in the central portion of the billet. For the grain-refined ingot, significant Ti enrichment is observed in the central portion, which is again found to be higher at increasing casting speed. Figure 4 illustrates this at a casting speed of 12 cm/min. From the above results so far, it appears that the casting speed has major influence on macrosegregation compared to the grain refining Microstructural observations (Figure 5) near the surface and at the center indicate significant grain refining with Al-5Ti-1B addition. The grain size is refined from 400 Pm to around 75 Pm. Irrespective of the grain refinement, grains with larger dendrite arm spacing are seen in the cent-

279

Figure 2: Macrosegregation profiles for AA 6061 at a casting speed of 8 cm/min (a) Non-grain refined and (b) grain refined [' Mg and Si]

er. Further, for the grain-refined ingot, the grain size is coarser in the center compared to the surface. In general, the movement of inter-dendritic liquid (solute-rich) in the mushy zone and the transport of solid grains (solute-lean) account for the segregation patterns observed in DC casting [1]. In addition to the natural convection of solute-rich liquid, the transport of solute-lean solid phase from the periphery of the billet to the center leads to the negative centerline segregation. Isothermal dendrites formed early in the solidification process are detached and carried by the strong natural convection currents into the molten metal pool [4]. They grow isothermally at a temperature close to the alloy liquidus. It is generally accepted these isothermal dendrites with coarse cells observed in the center of the billet increase the severity of negative segregation [57]. The central part of the billet thus shows duplex structure with a mixture of grains exhibiting finer and coarser internal structure. Our observation of grains with larger dendrite arm spacing in the center of the billet present case (Figure 5b,d) is in qualitative agreement with the above statement. The effect of casting speed on macrosegregation of various Al alloys is well known [3, 4]. The main parameters that influence structure formation and macrosegregation during DC casting are the sizes of the transition region along with the flow pattern in the slurry region of the billet [3]. Increase in casting speed leads to the widening of the transition region, especially in the central portion of the billet. In the present work, the sump depth measurements indeed showed a great variation (a from 25 mm to 68 mm) as the casting speed increased from 8 cm/min to 12 cm/min. Quite the same depths were recorded with grain-refining ingots. This

280

Figure 3: Compositional variation of Ti across the grain refined billet cast at 12 cm/min

means that the increase in the severity of segregation with higher casting speed can be directly related to the depth of the sump in both cases. As expected, the surface region of the DC cast material is characterized by a strong segregation zones (in the present case either positive or negative depending on how close the measurements are made in relation to the surface). This is related to the shrinkage driven flow of soluterich liquid [4] With respect to the effect of grain refining on macrosegregation, however, there are conflicting reports in the literature. Finn et al [8] who showed that grain refining produced positive centerline segregation due to the improved permeability of the mushy zone. Opposite trends are observed by Lesoult et al [9] who showed that grain refining causes more severe centerline seg-

Figure 4: Compositional variation of Ti across the grain refined billet cast at 12 cm/min

281

Figure 5: Optical microstructures of non-grain refined and grain refined billets in the center (b, d) and close to the surface (a, c). [Arrows in (b, d) show “floating grains”]

regation. Greater centerline depletion of Mg is observed in AA5182 with grain refining and connected to the formation and transport of isothermal dendrites (or “floating grains”) in the grain-refined ingot [10,11]. Gariepy and Caron [5] studied the effect of various methods and grain refiner on macrosegregation. They found that liquid flow, which worked against natural convection, and smaller amounts of grain refiner, reduced macrosegregation. In the present case, the balancing effect of the increased permeability of the mushy zone along the increased amounts of isothermal dendrites in the central zone may lead to a nearly unchanged macrosegregation patterns in the grain refined and non-grain refined samples. Further work, however, needs to be directed towards the quantification of these floating grains. Also it would be interesting to examine the effect of Ti concentrations (i.e. amount of grain refiner) on structure and macrosegregation.

4

Conclusions

Direct chill casting experiments with and without grain refining at different casting speeds were conducted on an AA 6061. Significant structural refinement is observed. The concentration profiles for major alloying elements (Mg and Si) showed a negative segregation in the center and close to the surface. The severity of segregation increases with higher casting speed both in non-

282 grain refined and in grain-refining billet. On the other hand, grain refining does not seem to have any considerable effect with respect to macrosegregation.

5

Acknowledgements

This work is done within the framework of the research program of the Netherlands Institute for Metals Research, Project MC4.02134.

6 [1] [2] [3]

References

J. F. Grandfield and P.T.McGlade, Materials Forum, 1996, 20, 29–51 E. F. Emley, Int. Met. Rev., 1976, 206, 75–115 D. G. Eskin, J. Zuidema, Jr., V. I. Savran and L. Katgerman, Mat. Sci. Engg, 2004, A384, 232–244 [4] H. Yu and D .A. Granger, Fundamentals of alloy solidification applied to industrial processes, Proc. NASA symposium, 1984, p. 157 [5] B. Gariepy and Y. Caron, Light Metals 1991, The Minerals, Metals and Materials Society, Warrendale, PA, p. 961 [6] R. C. Dorward and D. J. Beerntsen in Light Metals (Ed.: C. M. Bickert), The Minerals, Metals and Materials Society, Warrendale, 1990, p.919 [7] Men G. Chu and John E. Jacoby in Light Metals (Ed.: C. M. Bickert), The Minerals, Metals and Materials Society, Warrendale, 1990, p.925 [8] T. L. Finn, M .G. Chu, W. D. Bennon, Micro/Macro Scale Phenomena in solidification, ASME, New York, 1992, p. 17 [9] G. Lesoult, V. Albert, B. Appolaire, H. Combeau, D. Dalouz, A. Joly, C. Stomp, G. U. Grün and P. Jarry, Sci. Tech. Adv. Mat, 2001, 2, 285–291 [10] A. Joly, G. U. Grün, D. Daloz, H. Combeau and G. Lesoult, Mat. Sci. Forum, 2000, 329-330, 111–120 [11] A. M. Glenn, S. P. Russo and P. J. K. Paterson, Met. Mater. Trans A., 2003, 34A 1513–1523

283

Effect of Melt Flow on Macrostructure and Macrosegregation of an Al–4.5% Cu Alloy A. N. Turchin1, D. G. Eskin1, L. Katgerman2 1 2

Netherlands Institute for Metals Research, Delft, the Netherlands Delft University of Technology, Department of Material Science and Engineering, Delft, the Netherlands

1

Abstract

One of the major problems in production of large ingots and billets is macrosegregation. The occurrence of this defect is associated with melt flow in the liquid sump of billets or ingots. In this paper the effect of melt flow on macrostructure and macrosegregation of an aluminum alloy is investigated using the electromagnetic pump, which allows one to organize the controlled melt flow along the solidification front. The experiments are performed on an Al–4.5% Cu alloy in a wide range of melt flow velocities and temperatures. Computer simulations of the experiment with a CFD software are used in the interpretation of the formation of macrosegregation in the presence of melt flow. The results show strong influence of melt flow on the macrostructure and macrosegregation. The concentration profile in the sample solidified without flow shows negative macrosegregation in the center part. With increasing melt flow the positive macrosegregation can be observed. The interpretation of the structure formation in the presence of melt flow is discussed.

2

Introduction

Melt flow involved in all casting processes has significant consequences for structure and macrosegregation evolution during solidification. The published studies show that the melt flow influences the size and distribution of grains, affects grain morphology [1–6], e. g., causes the columnar to equiaxed transition (CET) [4, 5] and the appearance of feathery grains in aluminum alloys [6]. Macrosegregation is a major problem during the production of large ingots and billets where it appears due to compositional differences in liquid and solid. It is believed to be controlled by melt flow in the liquid sump of a casting. Many experimental and numerical studies have been done in order to understand and to predict the macrosegregation development [7–10]. The detailed description of causes of solute movement due to fluid flows in casting processes can be found elsewhere [11]. It was shown that the segregation during casting is a result of several types of flow in the transition zone: shrinkage flow during the cooling [12], buoyancy induced flows due to thermal and solutal convection [13], and forced convection leading to movement of grains [14]. However, it was found in some works that flow can not only promote the segregation but also almost eliminate it [15]. This study is aimed to investigate the effects of controlled linear melt flow on macrostructure and macrosegregation formation in an Al–4.5% Cu alloy using specially dedicated electromagnetic pump.

284

3

Experimental Technique and Computer Simulations

An experimental set-up consists of an electromagnetic pump, control, melt-guiding and data acquisition systems [16]. A „flow-through” melt-guiding system was designed in this work in order to obtain a constant unidirectional bulk flow along solidification front with the control of the flow rate. Working on magneto-hydrodynamic principle, the pump creates a linear melt flow that goes through a rectangular launder with dimensions of 800 u 60 u 70 mm into the crucible placed on digital scales. Solidification occurs under conditions of constant melt flow along the solidification front on the bottom side of a water-cooled bronze chill built in the bottom surface of the launder. The chill has a rectangular shape 34 mm u 110 mm in plane section with the inner cavity about 10 mm u 100 mm. The water-cooled chill is designed to reproduce the linear melt flow motion along the solidification front. The objectives of experiments were to study the effect of flow velocity and melt temperature on macrostructure and macrosegregation of an Al– 4.5% Cu alloy1. An experimental Al–Cu alloy with a chemical composition of 4.40% Cu, 0.20% Si, and 0.12% Fe was prepared using 99.95% pure aluminium and an Al–47.7% Cu master alloy. Temperatures in the liquid bath of the pump and at the entrance to the launder were controlled and measured during experiments. Melt velocity is controlled by a stepwise change of voltage at the electromagnet. The weight per unit time is initially measured by the digital scales Mettler Toledo. The linear melt flow velocity was then recalculated from the values of weight per time. The samples were obtained at melt temperatures in the liquid bath between 695–745 °C and in a velocity range of ~0.05–0.60 m/s. After experiments, samples were sliced in the middle section in the longitudinal direction for the examination of macrostructure. After cutting in the longitudinal direction, samples were polished, etched with 45 ml HCl, 15 ml HNO3, 15 ml HF and 25 ml H2O solution (Tucker's reagent) for approximately 25 seconds in order to investigate the macrostructure of the whole section. The second half of each sample was used for measurements of chemical composition along longitudinal direction (flow direction) close to the chill surface. Measurements were taken each 20 mm using a spark spectrum analyzer SpectroMax. The absolute error for measured copper concentration was ±0.01 wt %. Computer simulations of the effects of process parameters, such as melt flow velocity and temperature, under conditions of constant melt flow including solidification were performed using the Flow-3D commercial software (Version 9.0). A two-dimensional flow of a molten aluminum alloy Al–4.5% Cu was assumed in a launder. A computational block consisted of two obstacles: one plate and a built-in chill with the cavity reproducing the geometry used in experiment. Thermophysical parameters for an Al–4.5% Cu alloy used in the simulations are described elsewhere [ȑȔ]. The computational time was 30 s. The grid consisted of structured uniform 2D mesh: 35840 cells in 640 u 56 square. The melt flows from the left side of the computational block with the initial temperature 700 °C at various velocities specified as the lefthand boundary conditions. The interface between obstacles and flowing melt is characterized by a heat transfer coefficient of 150 W/m2K with the temperature of obstacles 373 K. The initial condition of void state is specified as 298.15 K. When liquid reaches the cavity the solidifica-

1

All alloy compositions are in wt%

285 tion occurs. The heat transfer coefficient between the bottom of block and flowing melt is 1150 W/m2K [19]. The right-hand boundary condition is specified as outflow.

4

Results

Macrostructure examination of samples obtained under different conditions shows that melt flow dramatically changes the morphology and distribution of grains. While the macrostructure obtained under „no-flow” condition generally consists of equiaxed grains and small fraction of columnar grains at the bottom of the sample (Fig 1 a), the columnar and feathery grains oriented towards incoming flow appear in the whole volume of samples by applying relatively slow flow velocities (up to 0.10 m/s) (Fig. 1 b, c). The zone of equiaxed grains develops in the lower part of a sample obtained at 0.10 m/s. This zone expands and the grains become finer with further increasing of flow velocity (Fig. 1 d, e). The increasing of melt temperature up to 745 °C promotes the formation of columnar and feathery grains within the given velocity range. The effects of flow velocity and melt temperature on macrosegregation pattern are demonstrated in Fig. 2. Figure 2 a shows the strong influence of flow velocity on the extent of macrosegregation. The profile of sample obtained under „no-flow” conditions exhibits so-called negative macrosegregation with minimum in the center of the sample. The negative profile remains after applying slow melt flow (~ 0.03 m/s). On further increasing of flow velocity the positive profile of macrosegregation with the shift of maximum copper concentration towards the incoming flow can be observed for velocities of 0.05 m/s and 0.10 m/s. Finally, at a velocity of 0.30 m/s the pronounced positive segregation profile in the center of the sample is obtained.

Figure 1: Macrostructure of the Al–4.5% Cu alloy obtained at 700 oC and at various melt velocities of (a) without flow, (b) 0.03 m/s, (c) 0.05 m/s, (d) 0.10 m/s, and (e) 0.30 m/s; bottom side–chill surface, melt flow direction– from left to right

286

Figure 2: Macrosegregation profiles for copper concentration (Xi – Xnominal)/ Xnominal in the longitudinal section of the samples obtained at 700 °C and at different flow velocities (a) and at velocity of 0.03 m/s (b) and different melt temperatures; melt flow direction–from left to right

The effect of melt temperature at a constant velocity of 0.03 m/s on macrosegregation patterns is shown in Fig. 2 b. Increase in melt temperature has a little effect on the segregation profile with a tendency to produce larger segregation.

5

Discussion

The experimental results show that the parameters of unidirectional flow such as velocity and temperature have significant effects on macrostructure, namely: (1) solidification in flowing melt results in the formation of columnar and feathery grains in comparison with „no-flow” conditions where macrostructure exhibits equiaxed grains; (2) columnar grains are oriented towards incoming flow; and (3) increasing of melt velocity results in CET and grain refinement. Numerous works have been performed on effects of bulk [4, 5, 10, 15] and unidirectional flows [17, 18] on macrostructure development during solidification. Data produced by computer simulations (Fig. 3 a) show that the macrostructure composed of columnar grains is formed under conditions of high thermal gradient and low undercooling that agrees well with earlier observations [17, 18]; and development of equiaxed grains in the lower part of samples (Fig. 1 d, e) is associated with strong melt recirculation and, consequently, specific thermal and concentration fields in this region [4, 5, 10] (Fig. 3 b). To the best of our knowledge, there are no papers where the macrosegregation pattern is analyzed in samples obtained under conditions of constant unidirectional melt flow. The formation of negative segregation profile in the sample obtained under „no-flow” conditions can be explained by the motion of liquid melt (Fig. 4 a). The transport of solute-enriched liquid from the center of the sample to the periphery and stagnant region in the center of the sample with a

287

Figure 3: Temperature field in the chill region at an initial melt temperature of 700 °C and a melt flow of 0.05 m / s after 30 s (a) and velocity pattern at 700 °C and 0.30 m/s after 20 s (b); bulk melt flow direction–from left to right

potential for settling down solute-poor grains contribute to the development of negative segregation in the center of the sample. This trend remains throughout solidification with increasing of solid fraction. In spite of the fact that the macrostructure entirely changes to columnar at a relatively slow velocity of 0.03 m/s, the profile does not alter and exhibits again the negative segregation. It can be a result of almost identical velocity pattern as during „no-flow” conditions which tend to enrich the periphery regions. It is known that applying of melt stirring in the liquid sump of the billet during DC casting process leads to the formation of normal, namely, positive segregation [20]. The same tendency of profile changing is also observed in the present work by applying the strong fluid flow (from 0.20 m/s). This fact may be a result of appearance of strong flows in the opposite direction to

Figure 4: Velocity pattern at 700 °C under „no-flow” conditions after 3 s, arrows in the central part are artificially enlarged (a) and in the central part of the chill region at the flow velocity of 0.30 m/s after 15 s (b); bulk melt flow direction–from left to right

288 the initial flow and, thus, may increase melt stirring in the vertical section of the sample, as can be seen in Fig. 4 b. In that way the flow influences the transfer of the solute to the center from periphery of the sample. The formation of zone of equiaxed grains apparently formed by fragmentation in the right bottom corner of samples further contributes to the solute depletion of this region (Fig. 2 a). A transient situation is observed in a velocity range between 0.03 m/s and 0.10 m/s. Increasing of melt temperatures enlarges the positive segregation at a flow velocity of 0.03 m/s (Fig. 2 b). The growth of coarse columnar grains and, consequently, the formation of large interdendritic channels with increasing of melt temperature can promote stronger solute transfer within the interdendritic region.

6

Conclusions

Effects of parameters of unidirectional flow such as melt temperature and flow velocity were investigated. It is found that the solidification in flowing melt results in the formation of columnar grains deflected towards incoming flow and promotes CET and grain refinement in recirculation zone. Moreover, the melt flow changes macrosegregation pattern from negative to pronounced positive, when strong flows are applied.

7

Acknowledgements

This work is done within the framework of the research program of the Netherlands Institute for Metals Research (www.nimr.nl), Project MC 4.02134.

8 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

References D. Daloz, H. Combeau, S. Sterdjella, B. Commet, Ph. Jarry, Aluminium 2004, 6, 603–608 R. S. Rerko, H. C. de Groh, III, C. Beckermann, Mater. Sci. Eng. A 2003, 347, 186ñ197 K. Kubota, K. Murakami, T. Okamoto, Mater. Sci. Eng. A 1986, 79, 67–77 S. R. Chang, J. M. Kim, C. P. Hong, ISIJ International 2001, 41, 738–747 W. D. Griffits, D. G. McCartney, Mater. Sci. Eng. A 1996, 216, 47–60 S. Henry, G.-U. Gruen, M. Rappaz, Metall. Mater. Trans. A 2004, 35A, 2495–2501 D. G. Eskin, J. Zuidema, V. I. Savran, L. Katgerman, Mater. Sci. Eng. A 2004, 384, 232–244 J. P. Gu, C. Beckermann, Metall. Mater. Trans. A 1999, 30A, 1357–1366 Y. Yang, Q. Zhang, Y. He, Z. Hu, Sci. Techol. Adv. Mater. 2001, 2, 271–275 W. D. Griffits, D. G. McCartney, Mater. Sci. Eng. A 1997, 222, 140–148 C. Beckermann, Int. Mater. Rev. 2002, 47, 243–261 M. J. M. Krane, F. P. Incopera, Metall. Mater. Trans. A 1995, 26A, 2329–2339 A. V. Reddy, C. Beckermann, Metall. Mater. Trans. B 1997, 28B, 479–489 C. J. Vreeman, F. P. Incropera, Int. J. Heat Mass Transfer 2003, 43, 687–704

289 [15] B. Zhang, J. Cui, G. Lu, Mater. Sci. Eng. A 2003, 355, 325–330 [16] A.N. Turchin, D.G. Eskin, L. Katgerman in Light Metals 2005 (Ed.: H. Kvande), The Metals, Minerals and Materials Society, Warrendale, USA, 2005, p.1025 [17] K. Murakami, T. Fujiyama, A. Koike, T. Okamoto, Acta Mater. 1983, 31, 1425–1432 [18] L. L. Rishel, PhD Thesis, University of Pittsburgh, USA, 1993 [19] D. R. Poirier, G. H. Geiger, Transport Phenomena in Materials Processing, The Mineral, Metals and Materials Society, Warrendale, USA, 1994 [20] V. I. Dobatkin, N. A. Anoshkin, Mater. Sci. Eng. A 1999, 263, 224–229

290

Quenching Study on the Solidification of Aluminum Alloys D. Ruvalcaba1, D. Eskin1, L. Katgerman2, J. Kiersch2 1

Netherlands Institute of Metals Research, Delft, The Netherlands Delft University of Technology, Delft, The Netherlands

2

1

Abstract

Real-time observation of the microstructure during solidification may show in-situ evolution of the morphology in aluminum alloys. However, existing techniques are expensive, time consuming and not readily available. Therefore, the metallographic technique coupled with quenching of the microstructure during solidification is still considered as an alternative method. The major problem in employing this procedure is overestimation of solid fraction as compared with existing solidification models, e.g. lever rule and Gulliver–Scheil models. The present research is focused on studying the evolution of the microstructure during solidification employing the quenching technique in order to understand the overestimation of solid fraction and its behavior during quenching. Binary aluminum alloys (Al–3 wt% Si and Al–7 wt% Cu) were melted, solidified at cooling rates between ~0.03 and ~0.3 K/sec and quenched at different temperatures within the semi-solid region. Two quenching rates were achieved, i.e. ~50 and ~100 K/sec. The solid fraction evolution was measured by image analysis. Solute profiles over cross-sectioned dendrites were obtained by EPMA. Finally, the results were compared with lever rule and Gulliver–Scheil calculations.

2

Introduction

It is important to understand the development of the microstructure during solidification of aluminium alloys since the final properties and defects depend on this development. Commonly found defects such as porosity, hot tearing and microsegregation depend on the solid fraction; while permeability and dendrite coherency, which determine the properties of the macroscopic mushy zone, depend on the development of the microscopic mushy zone (e.g. on microsegregation) [1, 2]. Nowadays, methods such as: calorimetric techniques [3], optical-intensity measurements [4] and X-ray microtomography [5] may provide some information about the microstructure development. Nevertheless, some of these methods do not give information about the morphological changes while others such as X-ray microtomography are not readily available and are still under development. On the other hand, the quenching technique has been widely used as an alternative method to reveal how the microstructure develops during solidification of alloys. This technique needs, however, to be completely understood in order to produce reliable results, since the quenched interface develops instabilities that cause overestimation of solid fraction in aluminium alloys and modifies the original shape of the interface [1, 3, 6–9]. Quenching increases the rate of heat extraction and in turn undercooling at the interface. It is assumed that undercooling promotes the formation of equiaxed grains and a refined structure

291 ahead of the interface, as a result, there is a distinction between the microstructure before (i.e. columnar/coarse dendrites) and after (i.e. equiaxed/fine dendrites) quenching [1]. If quenching rate is high enough, the microstructure can be frozen and the distinction between solid and liquid is only determined by dendrites surrounded by the existing liquid at the moment of quenching. It has been found that overestimation of solid fraction is still present, even if interface instabilities formed by quenching are considered as part of the liquid phase. It has been shown that the overestimation becomes less if the two phase sample is held isothermally and if quenching takes place close to the eutectic reaction [6, 8, 9]. Accordingly, the microstructure is not properly characterized with regard to the original solidification conditions. Notably, it is important to understand the limitations of the quenching technique in order to employ it adequately when studying the microstructure development during solidification.

3

Experimental Procedure

Two hypoeutectic aluminium alloys Al3 wt% Si and Al7 wt% Cu were prepared from commercially pure ingots (i.e. 99.99 wt% Al, 99.99 wt% Si and Al47.7 wt% Cu). Then, the alloys were cast in small rods which were then cut as 15 mm Ø × 30 mm high and 11 mm Ø × 14 mm high samples. Two different crucibles were employed, an alumina crucible for the 15 mm Ø × 30 mm high samples and a graphite crucible for the 11 mm Ø × 14 mm high samples. The crucible with the sample was placed in a holder inside a vertical furnace. Then, a thermocouple was carefully placed in a drilled hole in the middle of the sample. After this, the sample was kept in the furnace until reaching a temperature of ~680 °C and then carefully stirred by tapping the crucible for less than one minute. Then, the sample was again heated up until its temperature reaches ~700 °C. After ~2 minutes the furnace was turned off and the sample was cooled down inside the furnace. Finally, the sample was quenched when a selected temperature was reached. Temperatures in the semi-solid interval (i.e. between the temperature of liquidus TL and the temperature of eutectic TE) were selected from the equilibrium diagram. The accuracy of temperature measurements is ~ ±1 °C. Once the samples were quenched, they were cut, grinded and polished. Then the samples were analysed by optical microscopy, where the solid fraction was calculated by image analysis with an accuracy of ~ ±0.02 of solid fraction. The -Al phase was considered as solid and the eutectic and instabilities were considered as the liquid phase at the moment of quenching, assuming the formation of instabilities happening after quenching. Finally, line scan measurements were performed in order to analyse the solute distribution in dendrites (i.e. microsegregation) by Electron Probe Microanalysis (EPMA). The cooling rates were calculated from the cooling curves recorded during solidification. First, the cooling rate before quenching for each sample is calculated from:

T

TL  Tq t L  tq

,

(1)

where T is the cooling rate, TL is the temperature of liquidus, Tq is the temperature of quenching, tL is the time at which TL was reached and tq is the time at which Tq was reached.

292 Then the quenching rate was calculated from:

Q

Tq  T500 tq  t500

,

(2)

where Q is the quenching rate, T500 is the temperature at which the sample is assumed to be completely solidified which is 500 °C and t500 is the time at which this temperature was reached.

4

Results and Discussion

The quenching rates achieved were: ~50 K/s for the 15 mm Ø × 30 mm high samples and ~100 K/s for the 11 mm Ø × 14 mm high samples. Table 1 shows the selected quenching temperatures and cooling rates for the different samples. It can be noticed that most of the samples show close initial cooling rates prior to quenching, therefore the analysis can allow comparison between the samples. Also, the cooling rates prior to quenching are very low, therefore the results can be compared with the two solidification models which consider the solidification extremes under normal solidification conditions. These models are equilibrium (i.e. lever rule) and non-equilibrium solidification (i.e. GuilliverScheil), which are calculated using ThermoCalc Software. Figure 1 demonstrates microstructures of the two alloys quenched at 50 K/s. This Figure shows that the formation of interface instabilities is less pronounced when quenching occurred at temperatures close to the eutectic reaction. The same behaviour was found at 100 K/s, however, at this rate finer instabilities are found at higher temperatures. Figure 2 shows that in the case of the Al3 wt% Si alloy, the overestimation of solid fraction compared to the lever rule and Scheil is less at 100 K/s at temperatures above 615 °C. The sample quenched at 50 K/s and at

Figure 1: Micrographs of samples quenched at 50 K/s for: a) Al3wt% Si at Tq = 634 °C; b) Al 3wt% Si at Tq = 578 °C; c) Al7wt% Cu at Tq = 637 °C; d) Al7wt% Cu at Tq = 551 °C

293 634 °C did not show a homogeneous distribution of phases which may be due to the size of the sample, therefore, this measurement of solid fraction may not be the one that was present close to the thermocouple tip. At temperatures below 615 °C, there is no overestimation of solid fraction, which may mean that quenching was successful in freezing the microstructure. Also, no instabilities were found in these samples quenched at low temperatures. Only the sample quenched at 50 K/s and at 578 °C showed almost complete solidification exhibiting some coarse silicon particles, attesting for the beginning of the eutectic reaction (Figure 1b). Table 1: Selected temperatures (Tq) and cooling rates ( T ) prior to quenching for Al3 wt% Si and Al7 wt% Cu alloys Al-3 wt% Si Samples Quenched at: Quenching temper- Q ~50 K/s Q ~100 K/s ature Tq [°C] T [K/s] T [K/s] 634 0.05 0.03 627 0.06 0.08 616 0.08 0.10 595 0.13 0.19 582 0.13 0.19 578 0.13 0.19

Al-7 wt% Cu Samples Quenched at: Quenching temper- Q ~50 K/s Q ~100 K/s ature Tq [°C]  T [K/s] T [K/s] 637 0.04 0.06 631 0.06 0.06 620 0.06 0.09 595 0.13 0.19 583 0.13 0.19 551 0.17 0.30

In the case of the Al7 wt% Cu alloy, the samples showed to be more susceptible to quenching (Figure 1c). Overestimation was only found in samples quenched at 631 °C and 620 °C. All the samples showed no difference in behaviour with the quenching rate (Figure 3). The lowertemperature samples showed good agreement with Scheil approximation, which might mean that back diffusion did not take place. Line-scan measurements were performed over several dendrites in Al3 wt% Si samples by EPMA analysis. Figure 4 shows the concentration profile of silicon in a dendrite arm for the Al3 wt% Si alloy quenched at 100 K/s at 627 °C. It is noticed a change in the inclination of the slope when it reaches CS = 0.87 wt% Si. After this, the slope changes again at the solid content at the quenched interface i.e. CSq =1.95 wt% Si. This change of slope takes place at the instability neck where the instability initiates at the moment of quenching. Several dendrites showed the same behaviour with a change of slope at the onset of the instability. The solute content CS is close to that calculated from Scheil approximation (i.e. ~0.87 wt% Si) for a solid fraction of

Figure 2: Calculated and experimental solidification paths for the Al3 wt% Si alloy

294

Figure 3: Calculated and experimental solidification paths for the Al7 wt% Cu alloy

0.48 and a temperature Tq = 627° C which is the temperature of quenching of the sample. However, the experimentally determined solid fraction for this temperature was ~ 0.75. The overestimation is quite high as compared with Scheil. This overestimation may be due to the existence of large instabilities which formed at the very beginning of quenching and which may have had some time to grow and develop without leaving trace. The observation of microsegregation patterns and the corresponding changes in the slope of concentration profiles may provide a link between the experimental observed structures and the morphology of the solid phase prior to quenching. Finally, the samples quenched at 595 and 582 °C showed a solid fraction and CS falling between equilibrium and Scheil, which may mean that the microstructure was frozen successfully without the formation of instabilities. Quenching efficiency may depend on the characteristics of the solid-liquid interface. It is assumed that at the beginning of solidification (i.e. low solid fractions), the solid and liquid diffusion boundary layers have a large solute pile-up which provides a suitable place for nucleation. The formation of instabilities in the liquid pool may be due to detachment of dendrite branches that nucleate at the interface and that are displaced away from the interface to allow the nucleation of more instabilities at the same active interface. The formation of instabilities may be overcome by mixing the liquid, breaking the boundary layers, or by giving enough time for piled-up solute atoms to diffuse away and distribute homogeneously in the liquid (i.e. as achieved by isothermal holding).

Figure 4: Concentration profile of Si for the Al3 wt% Si alloy quenched at 100 K/s and at Tq = 627 °C showing the solute content at the interface before quenching CS and the solute content at the quenched interface CSq

295

5

Conclusions

The study of the development of solid phase at high temperatures in the solidification range implies some difficulties when employing the quenching technique. Line-scan measurements that show a noticeable change in the solute profile may help in reconstructing the interface at the moment of quenching. By understanding the formation and kinetics of instabilities produced by quenching, it may be possible to uncover and describe the microstructure that forms before quenching.

6

Acknowledgements

This work is performed within the framework of the scientific research program of the Netherlands Institute of Metals Research (www.nimr.nl), Project MC4.02134b. The authors would like to thank W. G. Sloof and K. Kwakernaak from the Department of Materials Science and Engineering of TUDelft for their assistance on EPMA analysis.

7 [1] [2] [3] [4] [5]

[6] [7] [8] [9]

References Ø Nielsen, S. O. Olsen, Transactions of the American Foundryman’s Society 2002, 110, paper 02-096 J. Thevik, A. Mo, J. Heat Mass Transfer 1996, 40, 2055–2065 D. Larouche, C. Larouche, M. Bouchard, Acta Mater. 2003, 51, 2161–2170 S. Steinback, L. Ratke. Scr. Mater. 2004, 50, 1135–1138 L. Salvo, M. Pana, M. Suery, M. DiMichiel, Ø. Nielsen, D. Bernard, in Proceedings of the 2nd International Light Metals Technology Conference (Ed.: H. Kaufmann) 2005 (in press) O. Pompe, M, Rettermayr, J. Crystal Growth 1998, 192, 300–306 S. W. Chen, C. C. Huang, Acta Metall. 1996, 44, 1955–1965 M. Rettenmayr, O. Pompe, J. Crystal Growth 1997, 173, 182–188 E. Tzimas and A. Zavaliangos, J. Mater. Sci. 2000, 35, 5319–5329

296

Numerical Study of the Influence of an Applied Electrical Potential on the Solidification of a Binary Metal Alloy P.A. Nikrityuk, K. Eckert, R. Grundmann Institute for Aerospace Engineering Dresden University of Technology, D-01062 Dresden, Germany

1

Abstract

In this work we study numerically the influence of a homogeneous electrical field on the fluid and heat transfer phenomena at macroscale and mesoscale during unidirectional solidification of a binary metal alloy. The numerical results showed that a pulse electric discharging applied perpendicularly to the solidification front leads to a much stronger Joule heating of the liquid phase in comparison to the solid phase. It was found that on the mesoscopic scale the electric current density is not homogeneous due to the complex shape of the dendrite and the difference between electrical conductivities of the solid and liquid phases. This inhomogeneity of the electrical current density in the mushy zone leads to the increase of the Joule heating of the dendrite in comparison to the interdendritic liquid and induces a pinch force (electromagnetic Lorentz force). The main features of the resulting convection in the interdendritic liquid are discussed.

2

Introduction

Control of solidification of metal alloys is one of the most demanding problems in the electromagnetic processing of materials. One of the innovative methods of such a control is the pulse electric discharging (PED). This method allows the modification of the microstructure during solidification [1–4]. The main feature of PED consists in a series of electric impulses passing through the solidifying melt. Due to the Joule heating caused by passing of an electric current, the temperature of the melt can increase. In the case of an inhomogeneous electrical current the interaction of the current with its own magnetic field produces a Lorentz force. This phenomenon, the so called pinch effect, received recently considerable attention in magnetohydrodynamics [5–8]. The pioneering work of the study of the influence of the direct electric current passing through the solidified melt were performed by Mirsa [1]. It was shown experimentally that the direct electric potential changes the nucleation and growth processes of the solid. But the mechanism of modification of the grains size were not understood. Nakada and coworkers [2] studied experimentally the influence of PED on the solidification structure of Sn15wt%Pb alloy. The electric discharging was carried out parallel to the solidification front by means of two cylindrical electrodes located along side wall of the cavity. The electric current was non-homogeneous. It was shown that solidification structures were modified from large grains with dendrites to finer grains with globular dendrites by means of pulse electric discharging with a capacitor bank. It was proposed that the Lorentz force (pinch force) induced at the moment of discharge is respon-

297

Figure 1: Schematic description of the geometry: axisymmetric cylindrical cavity on the macroscale (a) and columnar dendrite of parabolic shape on the mesoscale (b)

sible for the break of dendrites into globular fragments due to high shear stress. But no numerical simulations were carried out to support this hypothesis. To sum up, our understanding of the complex interaction between electrical current and solidified melt is far from being complete. On of the reasons for the still rather empirical applications of PED in unidirectional solidification of metals and alloys is the lack of detailed knowledge of the main mechanisms which are responsible for the grain size modification under the influence of electric current pulse. Motivated by this fact, this paper presents the first numerical study of the influence of PED on the heat and momentum transfer during directional solidification of Sn15wt%Pb alloy. In this paper we show the details about the development of the electro-vortex flow on mesoscale produced by interaction between the electric current passing through the melt and its own magnetic field. Furthermore we analyze the influence of the duration of the electric current pulse on the cooling curves.

3

Problem Formulation

To study the influence of the direct current applied during unidirectional solidification of a metal alloy on the macroscale heat transfer we consider a cylinder with non-conducting side walls of the height H0 = 0.075 m and the radius R0 = 0.025 m filled with the superheated alloy Sn15wt%Pb, see Fig. 1a. Between the bottom and the top of the cylinder an electric potential is applied. Thus a homogeneous electric current flows through the liquid and solid phases perpendicular to the solid front. Due to the homogeneity of the electric current there is no Lorentz force produced by the interaction of the current and its own magnetic field inside of the cavity on the macroscale [4]. The top and side walls of the cavity are thermally insulated, while the bottom is cooled at a rate governed by the instantaneous wall temperature, TW(t), and a uniform and constant overall heat transfer coefficient, D :

qW  t D ˜ TW  t  Tc ,

(1)

298 where TC is the temperature of the cooling media. In this work the values for D and TC were set to 10 W m–2 K–1 and 300 K, respectively to represent an intermediate velocity of solidification of about VS | 3 · 10–4 m s–1. To study the physical processes in the mushy zone on the mesoscale we consider a columnar dendrite of a paraboloid shape without resolving the complex morphology of the its boundary, see Fig. 1b. The paraboloid shape is an accepted approximation of the dendrite, c.f. the famous work of Ivantsov [9]. The size of the domain considered is Rd = 10–4 m which approximates typical dendrite arm spacing for VS | 3 · 10–4 m s–1 [10]. The center of cylindrical coordinate system lies on the axis of the symmetry of the dendrite, and is moving with the velocity corresponding to the solidification velocity VS. Thus we have the dendrite which is flowed around the liquid phase which velocity equals to VS, see Fig. 1b.

3.1

Electromagnetic Field Calculation

To calculate the electric current density we use Ohm’s law:

& & & & j V E u uB ,





(2)

& where V is & the electrical conductivity of the mixture of solid and liquid phases, u is the velocity vector, B is the magnetic induction vector. By mesoscopic consideration of solidification V varies stepwise between solid and liquid phases. In the case of macroscopic consideration of the V variation a linear interpolation can be used:

V

V l H  V s 1  H ,

(3)

where H is the volume fraction of liquid. In the liquid phase, corresponding to H = 1, the electrical conductivity equals to Vl. In the solid phase & corresponding to H = 0, the electrical conductivity equals to Vs. The electric field intensity E is & (4) E ’M . Here M is the electric potential. To derive the electric potential we use the continuity condition of the electric current: & ’˜ j

0.

(5)

Inserting eqs. (2) and (4) into eq. (5) written in cylindrical coordinates (r, T, z), we have:

1 w§ wM · w § wM · ¨ rV ¸  ¨V ¸ wr ¹ wz © wz ¹ r wr ©



1 w w  r V u z BT  V ur BT . wz r wr

In this work we consider the axisymmetric case, thus

wM wT

0.

For the calculation of azimuthal magnetic field, BT, we use Biot-Savart’s law

(6)

299

& j

& ’u B

1

P0

(7)

where P0 = 4S · 10–7 H/m is the vacuum magnetic permeability. If there are no external magnetic fields we have only an azimuthal component of the magnetic induction, BT, given by:

BT

P0 r

R

˜ ³ r jz dr

(8)

0

For the better understanding of the dynamics of the electrical filed parameters during solidification we simplify eq. (6) under the following conditions:

& 1. The electrical current density j is homogeneous. 2. During solidification there are only solid and liquid phases, i.e. no mushy zone exists. 3. The side walls of the cavity are isolated. In this case eq. (6) has an analytical solution:

M s l

M1  M 0 AV  As  1 AV  As  1  1

(9)

where Ms–l is the electric potential on the boundary solid-liquid, M0 and M1 are the electric potentials on bottom and top, respectively, AS = H0 / HS and As = Vs / Vl. Thus eq. (9) allows us to calculate electric field intensities, ES, El and Joule heating terms V s Es2 , V l El2 in solid and liquid phases, respectively. In this work we use eq. (9) for the validation of a solution of the eq. (6), see Section 3.

3.2

Macro-Energy Transport

In this study we restrict ourselves to the hypereutectic alloy Sn15wt%Pb which has the advantage of an initially stable stratification with respect to both the thermal and the solutal density change during unidirectional solidification. Thus without forced convection, the UDS of Sn15wt%Pb is not affected by thermosolutal convection. Furthermore shrinkage-driven flow is negligible. Since the homogeneous electric field does not induced convection the energy transport equation has the following form [10]:

U

& w wH  V E2 c p T ’ ˜  O ’T  ' H U  wt wt

(10)

where O = Ol H + OS (1 – H) and cp = cpl H + cps (1 – H). The volume fraction of liquid H is calculated from the relation [11]:

300

H

T  Ts Tl  Ts

(11)

where TS = 183 °C is the solidus temperature, Tl = 218.5 °C is the liquidus temperature. The last term in the eq. (10) is the Joule heat. To further simplify the problem we evaluate the characteristic time scales. Namely the typical solidification time for O(VS) | 10–4 ms–1 and O(H0) | 10–2 m is O (102 s) . The characteristic time for solute diffusion on the macroscale is given by O ('z2 / Dl) | 103 s, where 'z is the size of the control volume (CV) ('z = H0 / 100) and Dl is the diffusion coefficient (Dl = 1.5 · 10–9 m2 s). These different orders of magnitudes justify the neglect of solute mass transport on the macroscale. Table 1: Physical properties of Sn15wt%Pb alloy Thermal conductivity O, W m–1 K–1 Specific heat cp, J kg–1 K–1 Molecular viscosity P, N s m–2 Latent heat 'H, J kg–1 Electric conductivity V, A V–1 m–1

Solid

Liquid

57.99 210.85 – – 7.48 106

26.2 233.8 1,873 10–3 54140 1.8 106

Material properties of Sn15wt%Pb alloy were calculated from a linear dependence on the mass concentration of its component and are given in the Tab 1. The material properties of pure Pb and Sn were taken from [12–14]. In this work we assume that the densities of the solid and liquid phases are identical and equal to 7889 kg m–3.

3.3

Fluid Flow on the Mesoscale

On the microscale the electric current density is inhomogeneous due to the difference between Vs and Vl, and complex form of the dendrites. The interaction between the electric current and its own magnetic field produces the pinch force. This force induces a forced convection in the mushy zone. To be able to capture the main flow structure taking place in the mushy zone we present a simplified model of the mesoscopic fluid flow. This model includes the Navier-Stokes equations decoupled from heat and mass transport in the liquid and solid phases. Here we assume that the dendrite is imbedded inside the heterogeneous fictious domain in which we globally solve the fluid dynamics problem. The corresponding N-S equations are based on the porous media theory, introducing the permeability relative to each phase [15]: & (12) ’˜u 0

U

& & & wu   U u’ u wt

& & u & ’p  P ’ 2u   FL K

(13)

where . is permeability constant which prescribe immersed boundary conditions. This value related to each phase is defined by

301

K

­ f, if H 1 ® ¯0, if H 0

(14)

& The Lorentz force FL have the following form: & & & FL j u B

(15)

Using eq. (2) the radial and axial projections of this force have the form:



FLr

V  Ez BT  ur BT2

FLz

V Er BT  u z BT2







(16) (17)

To justify the neglect of the heat and mass transport we compare the characteristic time scales for heat, mass and momentum transports with respect to the size of the control volume ('z = Rd / 100). In particular the characteristic time for the Joule heat transfer in the liquid phase is O (cp U ' z2 / O ) | 10–7 s. The characteristic time for the solute diffusion is given by O (' z2 / Dl) | 10–3 s and the characteristic time for the viscous diffusion is O (U ' z2 /P) | 10–5 s. This evaluation shows that the mass transport is the slowest process, while temperature diffusion is the fastest one. Thus, a temperature perturbation produced by Joule heating is dissipated faster than fluid flow appears. Thus we assume that the liquid and the solid phases have the same temperature. In order to obtain first insights into the fluid flow induced by the pinch force we neglect the mass transport of the solute. However, to get the real insights into the transport process taking place on the mesoscopic scale during the PED it is necessary to consider momentum, energy and solute transports coupled with each other. This task is computationally demanding and will be done in future work.

4

Numerical Scheme and Code Validation

The set of eqs. (6), (10), (13), (14) has been discretized by a finite-volume finite-difference based method. The time derivatives are discretized by a three-time-level scheme. The convection terms are discretized by a central difference second order scheme with deferred acorrection [15]. The system of linear equations is solved by using Stone’s strongly implicit procedure (SIP). SIMPLE algorithm with collocated-variables arrangement was used to calculate the pressure and the velocities. Rhie and Chow stabilization scheme was used for the stabilization of pressure-velocity coupling. More details about the coupling algorithm can be found in [16]. Time marching with fixed time step was used. For every time step the outer iterations were stopped if residual of energy equation is less than 10–4 and less than 10–13 for pressure and momentum equations. Several grid-convergence and time-step-convergence tests were preformed to define proper grids and time steps leading to grid and time-step independent solutions. For the macro-scale energy transport simulations we used 20 u 70 grid, where first and second numbers correspond to the numbers of CV in the radial and axial directions, respectively, and a time step of 1 sec for diffusion controlled solidification with a PED duration of 30 sec. For the oscil-

302 lating PED with a period 'toff = 1 s between switch on and off, we used 20 u 200 grid with a time step of 0.2 s. For the mesoscale simulations we used a structured non-uniform 350 u 350 grid to calculate the electric filed parameters and fluid flow for the case VS « 10–4 ms–1. The time step used was set to 10–4 sec. For the calculation of fluid flow for the case VS = 10–4 ms–1 we used a structured uniform grid 250 u 250 and the time step of 10–5 s. To validate the code we model the electro-vortex flow induced by an inhomogeneous direct electric current flowing between the sidewalls in a rectangular cavity. We consider the 2D case when the thickness of the cavity is much more less than both height and width. This geometry is a simplified variant of the liquid metal current limiter (LMCL) investigated experimentally by Cramer et al. [7]. The 2D scheme of the device is shown in Fig. 2a. For the numerical simulations we used Cartesian coordinates (x, y). Fig 2b shows the spatial distribution of the velocity vectors induced by the Lorentz force. It can be seen that two large quasi symmetric vortices are generated. This is in good qualitatively agreement with the experiment [7]. The vortices induced are the product of the interaction between the electric current densities jx, jy and its own magnetic field Bz shown in Fig. 2c. The validation of the solution of eq. (6) is done in the Section 4 by the comparison of the analytical solution (9) with the numerical one.

Figure 2: Code validations: scheme of the domain (a), vector plot of velocity scaled with 5 · 10–4 m/s (b) and contour plot of magnetic induction Bz scaled with P0V(M1 – M0) · 0.5 (c). Here we used M0 = 0 V, M1 = 5 · 10–4 V, V = 106, U = 6000, P = 2 · 10–3, L0 = 35 · 10–3 m, rc = 5 · 10–3 m, 70 u 140 grid.

5

Results

The first series of numerical simulations is devoted to the study of the macro-energy transport during the solidification of the alloy under the influence of a pulsed electric current. Three cases were simulated. The first is the diffusion controlled UDS without PED, the second and third one concern the UDS with PED and voltages 'M = 0.05 V and 'M = 0.1 V, which was initialized after 50 sec of solidification and stopped after 80 sec of solidification. Fig. 3a shows the cooling curves obtained at the positions z = 0.02 m and z = 0.065 m from the bottom. It can be seen that during application of PED the cooling rate in the solid and liquid phases decreases. For 'M = 0.1 the Joule heat is so large that the liquid phase is heated during the PED. To understand the im-

303 pact of the PED on the spatial behaviour of the temperature we plot in the Fig. 4 the axial profiles of the temperature obtained at 70 sec. In comparison to the first case the temperature gradient on the boundary between the liquid and the mushy zone is increased. To investigate the influence of PED duration on the cooling curves we depict in the Fig. 3b the comparison of T(t) gained for steady current and periodically switched on and off with the period 'toff = 1 s. It is clearly seen that in the case of an oscillating PED the Joule heating effect decreases. To understand the increase of the temperature in the liquid phase we plot in the Fig. 5 the axial profiles of the nondimensional electric potential calculated by means of eq. (6) at t = 70 s. For comparison we depict the analytical solution for M in the case of a lacking mushy zone, i.e. the boundary between solid and liquid phases lies at Tl. For a better understanding of the Figure we plot additionally the profile of the liquid fraction H at that time. It can be seen that due to the higher electrical conductivity of the solid phase in comparison to the liquid we have a leap change in gradient of electric potential, in other words in Ez. Thus the electric field intensity in the solid phase is less than that in the liquid phase. As a result the Joule heating, V Ez2 , in the liquid is higher than in the solid phase, see Fig. 6. This figure shows the dependency of the Joule heating on the ratio HS / H0, calculated by means of eq. (9).

Figure 3: Predicted cooling curves: comparison of cooling curves gained for different voltage 'M (a) and comparison of cooling curves for z = 0.02 V gained steady current and periodically switched on–off current (b). The period 'toff was set to 1 sec. Here z is the distance from the chill.

Figure 4: Predicted axial profiles of the temperature obtained for different voltages at t = 70 sec

304

Figure 5: Comparison of the numerically and analytically calculated electric potential scaled with Ml at t = 70 sec. Here H is the volume fraction of the liquid Vs / Vl = 4.16

The primary interest of the next series of calculations is the prediction of the evolution of the flow field induced by PED in the mushy zone and to study the distribution of electric potential, electric current and Joule heat on the mesoscopic scale, see Fig. 1b. For the calculation of electric parameters we used the following boundary conditions: Between the bottom and the top of domain we set the potential difference 'M, taken from the macro-simulations, equal to 'M = 5·10–4 V corresponding to a electric field intensity in the liquid phase of Ez = –5 Vm–1. Furthermore no current flows through the side walls in radial direction. Fig. 7 shows the predicted spatial distribution of the electric potential and electric current density vectors. It can be seen that due to the difference in electrical conductivities of the solid and the liquid phases the current density in the dendrite is higher than the current in the liquid. As a result the Joule heat increases in the dendrite in comparison to the liquid, see Fig. 8a. The Joule heat has the maximum value on the tip of the dendrite which is explained by maximum curvature of the surface in that place.

Figure 6: Analytically predicted dependence of the Joule heating term on the height of the solid phase HS scaled with the height of the cavity H0. Here V’ = Vl,s / Vl, E’ = EzH0 / 'M, Vs / Vl = 4.16.

305

Figure 7: Predicted distribution of electric parameters: nondimensional electric potential M / 'M(a) and vector plot of electric current density scaled with Vl 'M / Rd (b). Here Rd = 10–4 m, 'M = 5 · 10–4 V.

To study the fluid flow pattern induced by the pinch force we consider two cases. The first case corresponds to a very small solidification velocity VS = 10–4 ms–1. On referring to Sec. 2 this allows us to use the no-slip boundary conditions on the top, bottom and side wall of the domain, see Fig. 1b. The velocity vector plot, see Fig 8b, displays a toroidal vortex rotating in clockwise direction. Thus the interdendritic fluid flow washes the dendrite from the bottom to the tip, and in that way it will modify the solute boundary of the dendrite. As a result the shape of the dendrite will be changed. The second case is devoted to the growth of dendrite with an intermediate velocity VS = 1 · 10–4 ms–1. Thus, a considerable relative flow around the dendrite occurs. In this case, the free-slip condition was used on the side wall, on the bottom: the velocity was set to uz = VS and on the top outflow boundary condition was used. The velocity vectors plots calculated for PED with Ez = –1 Vm–1 and Ez = –5 Vm–1 are displayed in Fig. 9. It can be seen for the induced velocities smaller than the solidification velocity there is no change in the flow around the dendrite, see Fig. 9a. But if u zmax ! Vs a toroidal vortex rotating in clockwise direction appears near the tip of the dendrite, see Fig 9b. This vortex produces the upward flow which will transport the solute rejected by the dendrite to the upper part of the mushy zone. We suppose that this may lead to a larger constitutional undercooling.

Figure 8: Predicted spatial distribution of the Joule heating term scaled with Vl ('M / Rd)2 (a) and meridional velocity scaled with 2.5 · 10–4 ms–1 (b). Here 'M = 5 · 10–4 V.

306

Figure 9: Predicted spatial distribution of meridional velocity scaled with VS = 10–4 m s–1 for Ez = –1 Vm–1 (a) and meridional velocity scaled with 2.5 · 10–4 ms–1 for Ez = –5 Vm–1 (b)

To calculate the time required to establish a fully developed electro-vortical flow we introduce in analogy to [17] the volume-averaged meridional flow velocity as follows:

U rz

2 Rd2 H d

H d Rd

³ ³r 0

ur2  u z2 drdz

(18)

0

Its time history is presented in Fig. 10. The calculations were performed for two values of Ez. We found that the time of the flow establishment for all two Ez is of order O(10–3 s). This is a very promising result since the application of PED with a comparable pulse frequency would both avoid the Joule heating and reduce the consumptive power of the PED devise. We argue that the fluid flow induced by the pinch force in the mushy zone can cause a mechanical fragmentation of dendrite. Probably the change of the solute concentration on the dendrite surface caused by the convection causes the constitutional fragmentation. Thus we suppose that the main mechanism of the grain refinement by the application of PED is related to the hydrodynamics of the turbulent regime which responsible for the refining of the grains by analog with works [17, 18,19]. Summing up the results of simulations on macroscopic and mesoscopic scales, we are faced with a paradoxical situation. On the macroscale there is no convection due to the homogeneous electric current but on the mesoscale due the complex shape of the dendrite and Vs / Vl z 1, toroidal vortices appear in the interdendritic liquid which may induce macroscale convection. This assumption needs detailed consideration in the future work. We note that in the case of inhomogeneous electric current on the macroscale a pinch force appears on both macroscale and mesoscale.

6

Summary

The results of the numerical simulations of the heat transport on the macroscale showed that a pulse electric discharging applied perpendicularly to the solidification front leads to a much stronger heating of the liquid phase in comparison to the solid phase (the heating is caused by

307

Figure 10: Time history of the volume-averaged meridional velocity for different axial electric field intensities Ez in the liquid phase for VS = 10–4 m s–1

the Joule heating effect). We could show that a shorter duration of the PED pulses decrease the Joule heating of the melt. The numerical studies on the mesoscopic scale revealed that both due to the complex shape of dendrite and difference in electric conductivities between solid and liquid phases the Joule heating in the dendrite is increased in comparison to the heating in the interdendritic liquid. The Joule heating reaches its maximal value on the dendrite tip. The inhomogeneity of the electrical current density in the mushy zone induces a electromagnetic Lorentz force (pinch force). This force induces a toroidal vortex near the dendrite tip. It was shown that for the domain with side length of 10–4 m the time required for the flow to be established has the order of magnitude 10–3 s for Ez of order of O(10) Vm–1.

7

Acknowledgements

The authors are grateful to Dr. M. Peric for the source code of the Navier-Stokes solver. We thank Armin Heinze for stimulating discussions. Financial support by the Deutsche Forschungsgemeinschaft (SFB609, B2) is gratefully acknowledged.

8 [1] [2] [3] [4] [5]

References A.K. Mirsa. Metallurgical Transactions 17A, 1986, 358–360 M. Nakada, Yu, Shiohara, M.C. Flemings. ISIJ International 30, 1990, 27–33 A. Prodhan, C.S. Sivaramakrishnan, A.K. Chakrabarti. Metallurgical and Materials Transactions 32B, 2001, 372–378 M. Gao, G.H. He, F. Yang, J.D. Guo, Z.X. Yuan, B.L. Zhou. Materials Science and Engineering A337, 2002, 110–114 V. Bojarevics, T. Freibergs, E.V. Shilova, E.V. Shcherbinin. In Electrically Induces Vortical, Kluwer Academic Piblishers, Dordrecht, 1988, p. 400

308 [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]

P.A. Davidson. An Introduction to Magnetohydrodynamics. Cambridge University Press, Cambridge, 2001, p. 431 A. Cramer, G. Gerbeth, P. Terhoeven, A. Krätzschmar. Materials and Manufacturing Processes 19, 2004, 665–678 I. Kolesnichenko, S. Khripchenko, D. Buchenau, G. Gerbeth. Magnetohydrodynamics 41, 2005, 39–51 G.P. Ivantsov, Growth of Crystals, Consultants Bureau, NY, 1958, 1 , p. 76 D.M. Stefanescu. Science and Engineering of Casting Solidification. Kluwer Academic/ Plenum Publishers, New York, 2002, p.342 S. Chang and D.M Stefanescu. Metall. Mater. Trans. A, 27A, 1996, 2708–2721 The Goodfellow website: www.goodfellow.com. Metals, polymers, ceramics and other materials Verein Deutscher Ingeniere (VDI) – Wärmeatlas Berechnungsblätter für den Wärmeübertragung, 9 Auflage, Springer-Verlag, Berlin, 2002 T. Iida, R.I.L. Guthrie. The Physical Properties of Liquid Metals. Clarendon Press, Oxford, 1988, p. 288 K. Khadra, P. Angot, S. Parneix, J. Caltagirone. Int. J. Numerical Methods in Fluids 34, 2000, 651–684 J.H. Ferziger, M. Peric. Computational Methods for Fluid Dynamics, 3nd ed., Springer, Berlin, 2002, p. 423 P.A. Nikrityuk, M. Ungarish, K. Eckert, R. Grundmann. Phys. Fluids 17, 2005, 067101. A. Vogel. Metal. Science 12, 1978, 576–578 B. Willers, S. Eckert, U. Michel, I. Haase, G. Zouhar. Mater. Sci. Eng. A., 2005, in press S. Eckert, B. Willers, P.A. Nikrityuk, K. Eckert, U. Michel, G. Zouhar. Mater. Sci. Eng. A., 2005, accepted for publication

309

Microstructure and Strain Distribution Influence on Failure Properties in Eutectic AlNi, AlFe Alloys P. Olaru1; G. Gottstein2; A. Pineau3 1

INAV-S.A.-Bucharest Romania IMM-RWTH-Aachen-Germany 3 ENSM-Paris-France 2

1

Abstract

Eutectic AlNi, AlFe, alloys exhibit plastic strains to failure (usually in the range of 1%-5%), that those of conventional structural alloys. We have developed a technique to measure strains at the scale of the microstructure and have used this method to assess the variation in failure properties with microstructure. This method is capable of using the grayscale information in the image of a gridded sample to obtain sub-pixel marker displacement, and can therefore accurately determine small strain values. Microstructures that exhibit large variation in local strain distribution tend to have higher variability in tensile properties, particularly tensile ductility, compared to microstructures that accumulate strain more uniformly. Orientation and morphology of lamellar plates in lamellar colonies play, also, a role in influencing the distribution of strain. Local grain orientation, phase distribution and segregation are factors influencing the strain distribution, and therefore the properties of these materials.

2

Introduction

The combination of specific stiffness and good oxidation resistance at intermediate temperatures can provide significant weight saving for certain components. One of the obstacles to the application of eutectic AlNi, AlFe alloys components is the relatively low ductility of these materials in tension. Because, this property is not clear explicitly dealt with in component design, some degree of damage tolerance and ductility in generally required so that stress concentrations can be blunted and minor levels of damage do not produce immediate failure. Some scientist [1], have shown that a plastic strain to failure of 2.8%, is sufficient to blunt relatively large stress concentration in eutectic AlNi,AlFe alloys, and many of the currently eutectic alloys being considered have average strains to failure that meet this relatively low requirement [2]. Much investigated alloys samples, exhibit a large variations in failure stress and strain, but the results of some samples do not reach the desired 2.8%, [3]. The aim of this work is to examine the influence of microstructure on variations in strength and ductility in eutectic AlNi, AlFe alloys. This work describes the investigations in INAV-S.A. Bucharest and laboratories F.S.I.M.-U.P.B.

310

3

Experimental Procedure

Master ingots, with chemical composition given in Table1, were elaborated in vacuum induction furnace by melting high-purity metals under argon atmosphere and were casting into graphite rods [4]. These ingots were machined into small- bars and placed in high purity aluminum crucible of 57/47 u 10–3, outside/inner diameter [4]. Thin samples of eutectic alloys, were elaborated by directional solidification in XTAL-VAR 97 installation [4], at a constant growth of RQl and a constant temperature gradient at solid/liquid interface of Gl. Table 1: Chemical composition of eutectic alloys No Eut 1 Eut 2

Alloy Al-5.7%Ni Al-1.7%Fe

Ni [at%] 5.7 –

Fe [at%] – 1.7

Al [at%] Balance Balance

In order, to eliminate the nucleation and growth of mis-orientated grains along the specimen length, all thin samples were grown from seeds defined structure and crystallographic orientation [4]. Sheet samples with an olmet geometry and a gage volume of 160 mm3 were prepared from each alloy. The surfaces of olmet samples were prepared by a low stress grinding procedure followed by hand treatment through 680 grit. The samples were electro polished at temperature (–30qC), to obtain a surface mirror finish and avoid hydride formation [5]. The polished samples were gridded by evaporating gold through a 1400 line per inch titan mesh and then tested to failure in Kammrath-Weiss in-situ tensile stage in Philips SEM 515-EDS. It was selected the microscope magnification around 92 grains and 2300 to 3000 markers [Eut1-120x; Eut2-120x] were present in each acquired 706 by 468 pixel image. Tested was interrupted at regular load intervals to allow for image collection. The image resolution of the technique depends on a number of experimental parameters including a higher resolution, signal to noise ratio, and the number of grayscales present, captured.

4

Results

The results obtained for samples with chemical compositions in Table 1, are shown in Table 2: Table 2: Results of tensile testing Alloy Eut-1 Eut-2

Samples test 7 9

Plastic strain to failure 1.9-2.4 % (2.6%) 1.2-1.8% (2.0%)

Peak strain before to failure ~5.5% ~3.2%

The value for mapped samples is indicated in parentheses. The surface displacement mapping technique is an accurate technique for measuring in plane displacements of less than a pixel, which allows the local strains in a sample to be measured with a high level of accuracy. Stress-strain curves obtained by averaging strains at each load interval for both eutectic samples are shown in Figure 1.

311

Figure 1: Stress – strain curve from strain mapping, calculated for displacement mapped samples

These curve appear in a good shape, despite being measured over only 0.67 mm2 of the surface area of the Eut-1, and 0.58 mm2 of the Eut-2. Like the first conclusion, the results become more consistent as the level of strain in sample increases, and that there is considerable variation at low strain levels. The error bars used in the graph represent the variation in the strains in the four quadrants of the analyzed area. The large variation in strain at each load level indicates that these areas are too small to provide an accurate measurement of average strain for the entire sample. This local variation in strain increased as the load and average strain increased for both eutectic alloys. For measuring small strain during in-situ tensile tests in SEM -Microscopy, we introduced a method, which use pattern recognition algorithms to locate markers on a gridded sample before and after deformation. The distortions in the grid are used to calculate the strains at each marker. Gold grids with a mesh size of 20Pm were evaporated onto the samples though titanium grids affixed with adhesive tape. Electron backscatter images of the samples taken during testing were analyzed using a set of interactive routines written in the Operative Data Language (ODL). After all displacement has been determinate in this fashion, a polynomial fit maps the locations of a marker and its surrounding markers in the distorted image to the corresponding in the reference image. The strains of the sample surface are then simply the coefficients of the polynomial expansion, and they are then contour plotted to show a map of the strains in the sample .The strain maps can be overlaid on the original image or plotted separately. Some strain maps obtained by this method can be seen in (Figure 2). This technique is to capture the development of strain in our sample, and the progression can be followed as the load increase. His magnitudes of the strains are similar to those predicted by a continuum model, but the shape of the strain contours depends on the microstructure of the sample, in main manner.

312

Figure 2: Strain maps by pattern recognition algorithm method to locate markers on a gridded sample before and after deformation

5

Conclusions

This method of measurements useful for isolating the features associated with enhanced local strain. The distribution of these local strain concentrated regions also apparently results in volume effects, where samples with smaller strained volumes higher strength in the absence of extrinsic flaws. The surface displacement mapping technique is essentially a two-dimensional technique, so no information about out of plane displacements can be obtained. As a finale conclusion, an apparent fracture origin may not be the actual fracture origin, since the real origin may be subsurface.

313

6 [1] [2] [3] [4] [5]

References G.Gottstein and L.S.Shvindlerman, Grain Boundary Migration in Metals, 1999, CRC Series in Materials Science and Technology, CRC Press Florida-U.S.A A.D.Rollet and D.Raabe , Comp. Mater. Sci., 2001, 21, 69 C.Maurice, “Proceeding of the first joint international Conference, ReX &GG1, Springer Berlin, 2001, p.123 P.C.Olaru, “Proceeding of the first joint international Conference, ReX &GG1, Springer Berlin, 2001 M.J.Blackburn, J.C.Williams , Transactions of the Metallurgical Society of AIME, 239 (1967), 287

314

Dendrite Coarsening and Embrittlement in Continuously Cast Tin Bronzes T. Virtanen Tampere University of Technology, Tampere

1

Introduction

Tin bronzes (CuSnP) are widely used in different applications due to their good physical and chemical properties. They have excellent durability and resistance to plastic deformation i.e. high strength and toughness. Good mechanical properties are combined with competitive thermal and electrical conductivities. Long solidification range and low solute redistribution coefficient makes tin bronzes prone to several problems during solidification in casting process. The main problem in upcasting is a defect called ‘tin sweat’ in which tin and phosphorous segregate between dendrites and further onto strand surface forming brittle eutectoid (D+G+Cu3P). Also coarsening of dendrites is closely related to the phenomenon. Coarsening means a change in dendritic length scale by disappearance of smaller dendrite arms in favor of larger arms. It determines the local chemical composition and therefore it greatly influences the properties [1]. It can also significantly influence the time necessary for homogenizing heat treatment after casting [5]. Coarsening rate depends strongly on the fraction of solid, time spent in the solid/liquid region, dendrite arm spacing and the temperature [3, 4, 5]. Large solidification range results in coarser dendrite branches [6]. Exudations are believed to result from separation of the casting from the chill surface. During air gap stage, the shell reheats thereby permitting solute enriched (lower melting point) alloy to exude from interdendritic locations and/or grain boundaries, and to cover the shell surface and contact the mold wall [7]. Remelting leads to a considerable increase in permeability [8], while metallostatic pressure, casting temperature and heat transfer coefficient have also pronounced effect on the amount of surface segregation [9, 10]. CAS2 software has been developed by Miettinen for simulating phase changes and solute distribution during solidification of binary copper alloys containing several alloying elements [11].Very important feature of the CAS2 model is that it takes the cooling rate into account. Depending on alloy composition, cooling rate and dendrite arm diameter, the package determines the stable phases (liquid, fcc, bcc, compounds) and their fractions and compositions as a function of temperature. The aim of this work is to investigate the conditions leading to formation of undesired structure in continuous casting of tin bronzes.

2

Experiments

Coarsening of dendrites and surface segregation phenomenon were studied experimentally by upcasting samples having different tin and phosphorous contents and varying the cooling conditions. Cooling condition was varied by changing casting speed and by adding external distur-

315 bance in cooling system. Tensile tests were carried out and the samples were examined with optical microscope. The amount of eutectoid and coarse dendrites were evaluated and classified visually and compared to the elongation. The sample structures were divided into four different categories (0–3), highest number presenting the worst-case change in dendritic length scale or appearance of surface segregation. The method is explained elsewhere in more detail [12]. Due to the visual nature of the used method, trends presenting coarsening and segregation tendency are indicative.

3

Results and Discussion

CAS2 package was used to estimate several features related to solidification of tin bronzes. Also an extensive casting campaign was carried out. In this paper the typical features of the results are presented and the findings discussed.

3.1

CAS2 Model – Interdendritic Sn Composition

Due to segregation and formation of non-equilibrium phases, interdendritic tin composition is of great interest. It can be seen in fig. 1 that tin segregates strongly between dendrites during solidification. Segregation increases with increasing cooling rate. With the highest cooling rate of 1000 qC/s Sn composition of the last liquid drop is around 20 wt.% and yet the final Sn composition in the interdendritic region remains low, at about 5 wt.%. It can be concluded that 1 wt.% tin concentration is not high enough for peritectic bcc phase and further brittle (D+G) eutectoid to form at any of the given cooling rates. Therefore the final cast structure is solid solution. In figure 1b) it is shown that increase in nominal tin composition from 1 wt.% to 8 wt.% results in great difference in the interdendritic composition during solidification. It can be seen that cooling rate deviating from infinitely slow equilibrium cooling rate does not have major impact on the segregation behavior. Even if the cooling rate is very slow, compared e.g. to typi-

Figure 1: Interdendritic Sn composition versus temperature for different cooling rates. a) 1 wt.% and b) 8 wt.% Sn

316 cal upcasting process, only at 1 qC/s, tin segregates as much as in the case of very fast cooling rate at 1000 qC/s. Regardless of the cooling rate peritectic bcc phase, E, is formed and further transformed to (D+G) eutectoid, which can be seen in sudden increase in curves at temperature close to 520 qC. Curves show the maximum Sn concentration between dendrites, not the average. Cooling rate has more pronounced impact on secondary dendrite arm spacing (DAS) [13, 14], i.e. dendritic length scale, than on segregation behavior.

3.2

Coarsening of Dendrites and Surface Segregation

Variation of DAS in cross section of a single sample is very small in the normal upcast structure. In some samples containing 1 wt.% tin almost cellular growth morphology is found close to the surface when cooling condition is disturbed. Cellular growth can be an indication of very slow growth velocity [15], which in turn means that heat removal has either been slow or has dropped suddenly. It can be seen in Figure 2 that in case of severe coarsening of dendrites and surface segregation, dendrite morphology changes remarkably between the surface and the center of the sample. In Figure 2a coarse dendritic structure close to the surface is shown. Also continuous films of brittle eutectoid between dendrites can be found. In Figure 1b, structure appears to be much more homogeneous in the center, which is very typical for every sample having coarsened dendrites and surface segregation. At the slowest casting speed, 0.2 m/min, coarsening close to the surface is found in all tin compositions (Figure 3a). Slow casting speed results in very rapid development of rigid solid shell which in turn results in formation of an air gap. Heat extraction is reduced significantly and latent heat removal deteriorates. As a result first solidified dendrites get coarse. Faster casting speed improves structure to a certain extent by making it more homogeneous i.e. less coarsening and surface segregation. The amount of coarsened dendrites also increases when the tin content is increased. This may be due to wider mushy zone and shrinkage. When the length of mushy zone is increased the time with both solid and liquid coexist is also increases. Small changes in air gap results in significant variation in the length of the mushy zone liquid [16]. The increase of the phospho-

Figure 2: Sample (Sn 10 wt.% P 0.01 wt.%) in different locations. a) Dendritic structure close to the surface is very coarse and has continuous films of brittle eutectoid between dendrites and on the surface whereas dendrites in the centre b) are significantly more homogenous. Scale bar 0.1 mm

317 rous content does not increase the coarsening tendency in the normal cooling condition. The more the alloy contains tin the more of it is in form of (D+G) eutectoid. Increasing length of the mushy zone provides long open channels between dendrites in which solute rich liquid can flow. At the lowest tin content no surface segregation is observed (Figure 3b), not even at the lowest casting speed. Reason for this behavior is the fact that no brittle eutectoid is formed at any cooling rate as is indicated in Figure 1a.

Figure 3: a) Coarsening severity of dendrites in samples versus tin composition and casting speed. b) Appearance of surface segregation versus tin composition and casting speed

It is clear that the disturbance and slow casting speed together strengthen uneven air gap formation (Ÿ uneven heat removal/-shell thickness Ÿ–coarsening). Structure close to the surface becomes very inhomogeneous also with the smallest tin composition as seen in Figure 4a.

Figure 4: a) Coarsening severity of dendrites in samples with disturbed cooling condition and b) surface segregation in samples with disturbed cooling condition

It can be concluded that when disturbance exists, coarsening occurs regardless of the tin composition or casting speed. Coarsening depends strongly on the severity of external disturbance: the more cooling is uneven along the cast wire, the more coarsening occurs. The difference in the appearance of surface segregation between normal and disturbed cooling condition is not big (Figure 3a and 4b), even though some increase can be seen in disturbed cooling condition. Cooling rate has both direct and indirect influence on formation of the sur-

318 face segregation. If the cooling rate is slow, dendrites get coarser than in the normal cast structure. In the area of coarse dendrites solute enriched liquid forms long continuous films and pools of brittle eutectoid. 3.3

Elongation

At lower tin contents the difference in elongation between different casting speeds and normal (5a) and disturbed cooling condition (5b) is not distinct. Visible surface segregation of brittle eutectoid is not formed in samples containing 1 wt.% tin, even if cooling condition is disturbed. In higher tin compositions the elongation is first improving with increasing casting speed and after reaching maximum value it starts to reduce again. Pieche [17] has realized that 45q angle between the casting and crystal growth direction gives the best structure for cold-rolling deformation. Increasing tin content improves also the strength.

a)

b)

Figure 5: Measured elongation of as-cast samples versus casting speed and different tin compositions, a) normal cooling condition and b) disturbed cooling condition

Disturbed cooling condition leads to distinct reduction in elongation in samples containing 8 or 10 wt.% Tin. The phenomenon is clear for slower casting speeds. Faster casting speed results in better and longer contact between solidifying shell and the mould leading to even cooling of wire. It can be seen as more homogeneous and smaller dendrite arm spacing along the wire. Also the oscillation of wire surface temperature ceases with increasing casting speed [18]. Addition of phosphorous increases the amount of surface segregation: The higher the phosphorous content, the more of it reacts with copper forming copper phosphides. Eutectoid (D+G+Cu3P) has very low melting point i.e. it remains quite long in liquid state.

4

Conclusions

Coarsening of dendrites occurs when heat removal drops suddenly after the formation of solid shell. This change in heat extraction can be caused either by slow casting speed or external di-

319 sturbance in cooling system. Even if the increase in alloying elements favor coarsening, cooling rate determines the final dendrite arm spacing. Formation of the surface segregation is influenced by all process variables - directly on tin and phosphorous content and indirectly on cooling rate. Change in the dendritic length scale is required to obtain severe surface segregation. Coarsening is not sufficient condition alone to significantly reduce the elongation. Presence of the brittle constituent is also required. The cooling rate has more pronounced impact on dendrite arm spacing than on the segregation behavior. It was shown that in spite of the cooling rate, tin segregates in samples containing 8 or 10 wt.% Tin between the dendrites with the same final composition leading to formation of (D+G) eutectoid. At low tin compositions brittle eutectoid does not form even with very high cooling rate

5 [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

[12] [13] [14] [15] [16] [17] [18]

References L.M. Coluzzi-Mizenko, M.E. Glicksman, R.N. Smith, JOM, 1994, 51–55 V. Rontó, A. Roósz, Int. J. Metals Res. 13(2001), 337–342 T.Z. Kattamis et al., Transactions of AIME, 239(1967), 1504–1511 M. Chen, T.Z. Kattamis, Materials Science and Engineering, A247(1998), 239–247 K.P. Young, D.H. Kirkwood, Metallurgical Transactions A, 6A(1975), 197–205 Suyitno, D.G. Eskin V.I. Savran, L. Katgerman, Metallurgical and Materials Transaction 30A(2004), 3551–3561 M. Morishita et al., Light Metals 2000, 657–662 H.J. Thevik, Modelling of Casting, Welding and Advanced Solidification Processes VII 1995, 557–564 P. Reisener, S. Engler, Metall 2(1995), 116–122 E. Haug , A. Mo, H.J. Thevik, Int. J. Heat Mass Transfer, 38(1995)9, 1553–1563 J. Miettinen, CAS2 - Solidification analysis package for binary copper alloys, User manual of DOS version 2.0.0. Report TKK-MK-162, Helsinki University of Technology Publications in Materials Science and Metallurgy, Espoo (2004), p.25 T. Virtanen, Effect of process variables on embrittlement and formation of coarsened dendrites and surface segregation in upcast tin bronzes, Doctoral Thesis, to be published M.A. Martorano, J.D.T. Capocchi, Int. J. Cast Metals Res., 13(2000), 49–57 E.A. Kumoto, R.O. Alhadeff, M.A. Martorano, Materials Science and Technology, 18(2002), 1001–1006 W.F. Savage, 1980: Welding in the world No5/6 p.8 F. Kaempffer, F. Weinberg , Metallurgical Transactions, 2(1971), 2477–2483 Piesche R.G., Copper and its alloys, Proceedings of an International Conference, London, 1970, 85–91 Kotipelto A., Numerical and experimental analysis of heat transfer in the continuous casting process of copper, Doctoral Thesis, Tampere University of Technology, 2002, p.138

320

Continuous Casting of Tin Containing Alloys and their Transformation V. Lebreton, F. Sadi, Y. Bienvenu Centre des matériaux P.M.Fourt, Ecole Nationale Supérieure des Mines de Paris, UMR/CNRS 7633, B.P.87, F91003 Evry Cedex

1

Abstract

This paper deals with binary and ternary tin bronze processing. A first issue is the refinement of the as cast microstructure to facilitate homogeneization: vertical casting is usually associated with interdendritic spacings of about 100 micrometers while thin slab casting (horizontal usually) reduces the length scale for microsegregation to below 20 micrometers. This is crucial since the homogeneization time at a given temperature in the solutioning range varies with the square of that length scale. Microsegregation modelling is used to account for observations and microanalyses and to guide process evolutions. The sequence of phase transformation in ternary CuNi-Sn bronzes is presented to optimize the annealing and to reach a compromise strength / ductility / conductivity.

2

Introduction

Electrical and thermal conductivities together with corrosion resistance and mechanical strength make copper alloys attractive for a broad range of engineering applications (electric, telecommunication and automotive industries). To reach such properties by conventional processes, each parameter of the processing route (temperatures, times...) must be optimized in function of the chemical composition of the copper based alloys. Generally, three principal stages can be distinguished: first is the as-cast microstructure and the difficulties associated with the solidification heritage for the rest of the shaping process; the second stage consists in the possibility for the alloys to homogenize and the last processing stage is the influence of hot or cold deformation and annealing treatments to obtain the best compromise between electrical and mechanical properties. To obtain this compromise, one must know the influence of the microstructure at all scale on these properties during aging after homogenization. Consequently, the sequence of phase transformation in solid state is a paramount parameter, which can be studied by D.T.A, S.E.M, T.E.M, and resistivity measurements. The cases of Cu-Sn and Cu-Ni-Sn alloys are very interesting because they illustrate the solidification problems, put in light the importance of the microstructure on electrical and mechanical properties and the requirement to establish in great detail the sequence of phase transformations. The first part of this paper deals with the refinement of the as cast microstructure to facilitate homogenization in Cu-Sn and Cu-Ni-Sn alloys. The sequence of phase transformation in ternary Cu-Ni-Sn is presented to optimize the annealing and to reach a compromise strength / ductility / conductivity.

321

3

The Cu-Sn and Cu-Ni-Sn Alloys in As-Cast Condition

The Cu-Sn and Cu-Ni-Sn alloys are generally processed by melting in an electric induction furnace. Some additions such as phosphorus, manganese or charcoal are introduced to minimize the formation of oxides. The Cu-Sn binary phase diagram illustrates the difficulties associated with solidification processing for the rest of the shaping process: a succession of peritectic reaction involving brittle intermetallics, a broad solidification interval prone to macrosegregation or to incipient melting in reheating tin containing intermetallics in copper alloys are frequently responsible for brittleness. In the as-cast state and beyond 4 wt% in tin, the D+E structure of tin bronze transforms during cooling in D+Jand D+G(Figure 1). Both J and G are brittle.

Figure 1: Cu-Sn phase diagram [1]

Figure 2 : Macrostructure and microstructure in Cu-6wt%Sn alloy [2]

The microstructure of a bronze containing 6 wt% tin produced with a horizontal continuous casting process (section 40 cm x 1.9 cm) presents a columnar zones along the thermal gradient, and the inclination of the grains on the axis observed is associated with the speed of the process. The micrographs show two kinds of segregation : • A macrosegregation of tin appears at the surface of the metal. This phenomenon is bound to the high rates cooling at the interface metal/mould and to stresses exercised on the metal by extraction rolls. • A microsegregation at the dendritic scale with the G phase forming in the majority the interdendritic zone (figure 2). The measurements of the secondary dendrite arm spacing achieved on Cu-6 wt%Sn and Cu9 wt%Sn as function of cooling rate permitted to identify the empiric law [3]:

O = 320 u (dT/dt)

0.58

O: interdendritic spacings, dT/dt: cooling rate

322 The interdendritic spacings are in the range 10–15 μm in the case of alloys produced with a horizontal continuous casting, 100–150 μm with a vertical continuous casting and some micrometers with the Osprey Spray£ technology. Similar observations have been achieved on the Cu-15 wt%Ni-8 wt%Sn produced with a vertical continuous casting process. Although the macrosegregation has not been observed, the enrichment of the interdendritic spaces in tin element is present (Figure 3).

Figure 3: X-ray microanalysis on Cu-15 wt%Ni-8 wt%Sn in the as-cast condition

The segregation of Sn in interdendritic spaces formed during the solidification embrittles the material considerably. These phases are crack initiation sites as shown in figure 4. A similar result is observed in the case of the Nordic Gold eurocoin monetary alloy in which the content in tin doesn't exceed 1% in mass (Fig.5). The formation of Sn-rich phases is often considered to be responsible for the hot cracking of copper based alloys. However the optimal mechanical features of these one are bound directly to the content of tin. Thus, the capacity of these alloys to homogenize becomes important. Isopleth sections of the Cu-Ni-Sn ternary equilibrium phase diagram present a single phase domain D with a c.f.c structure. Its extent is a function of the nickel and tin contents. The increase of the concentration of nickel has the effect of decreasing the solubility of tin in copper enlarging the two phase domain D+J to the detriment of the single phase domain.

Figure 4: Chrysocale Cu-Sn3 wt%-Zn9 wt% in as-cast condition : Sn-rich phase is a crack initiation site

Figure 5: Macrosegregation of Sn in Nordic Gold in as-cast condition

323

Figure 6: An isopleth at 15wt% Ni for the CuNi-Sn ternary phase diagram [4]

Figure 7: An isopleth at 7wt% Ni for the Cu-NiSn ternary phase diagram [5]

The difficulties associated with solidification processing for the rest of the shaping process can be avoided if the alloys can be homogenized to presents a single phase at high temperatures. Another solution is to change radically the process (spray deposition process) but it is more expensive and the productivity is low.

4

Phase Transformation in Solid State: Correlation Between Microstructure and Mechanical Properties

The optimization of the mechanical properties of the copper based alloys cannot be done without knowing the sequences of phase transformations in the solid state. That means to know the influence of each transformation on these properties. In the case of certain ternary alloys such as Cu-Ni-Sn, the mechanism of strengthening is more complex and depends strongly on the phase transformations sequence during aging which can be described as following: According to this sequence, there are two stages of strengthening : the first is bound to the formation of a chemical modulated structure characterized by an alternance of tin poor and tin rich zones. This microstructure has for origin the spinodal decomposition which preserves the crystalline structure of the matrix. The variation of stress during this stage follows the empirical equation established by B. Ditchek et al. in 1978 and true later by Ph. Herman and D.G. Moris for Cu-Ni15 wt%-Sn8 wt% alloy [6, 7]: 'V = 0.41 · A · K · Y A: amplitude of spinodal modulation, K = 1/a · Ga/Gc: differential size misfit, Y : elastic constante

324 With extended aging times this equation is no longer verified. Indeed, according to the phase transformations sequence in the solid state, the J’-D022 metastable phase coherent with the matrix D forms from the modulated microstructure. Thus, a second stage of strengthening can be observed. Figure 8 shows the variation of the ultimate tensile strength with aging temperature and aging time from a homogeneized state. The comparison between these curves and the T.T.T. diagram for Cu-15 wt%Ni-8 wt%Sn established by Zhao and Notis illustrates well the formation of the metastable phase J’-D022. This precipitation leads to an increment in yield stress following the equation according to Labush [8] and Janson et al.[9]: 'V = 3.7 · Pef · H4/3 f 2/3· (U /b)1/3

Pef: shear modulus, H misfit, f: final volume fraction of J, Uparticle diameter, b: modulus of Burgers ve ctor In the case of Cu-9 wt%Ni-2 wt%Sn and Cu-9 wt%Ni-6 wt%Sn this law has been verified.

Figure 8: Effects of ageing time and temperature on the ultimate tensile strength in Cu-15 wt%Ni-8 wt%Sn alloy

Figure 9: T.T.T. diagram for the Cu-15 wt%Ni-8 wt%Sn alloy [4]

For the Cu-15Ni-8Sn alloy, the precipitation kinetics are fast since a treatment at 450 °C during 15 minutes is sufficient to produce the modulated structure and the phase J’-D022. After 15 minutes, a decrease in mechanical resistance is observed due to the appearance of the equilibrium phase J-D03. The latter phase is born at grain boundaries and growths into the matrix with time. For aging temperatures included in the range [250–400 °C] the L12 phase appears before J-D03 with a discontinuous precipitation D/J morphology. The mechanical response to J-formation is a loss of ultimate tensile strength first, then the ultimate tensile strength increases again. This phenomenon is due to the increasing fraction precipitates of the metastable phase L12. Nevertheless, the mechanical resistance doesn’t seem to reach those associated with a treatment leading only to the J’-D022 phase.

325

Figure 10: J-D03 is born to grain boundaries. The J'-D022 is observed in the matrix (bright field, T.E.M)

5

Figure 11: Discontinuous precipitation in Cu-15 wt%Ni-8 wt%Sn (S.E.M)

Conclusion

Tin containing copper alloys often present macrosegregations and microsegregations due to the segregation of tin during solidification. To eliminate these problems a homogenization treatment is required. This treatment is shorter for fine solidification structures (horizontal casting). Another solution is to change radically the process (spray deposition process) but it is more expensive and the productivity is low. The ternary alloys of the system Cu-Ni-Sn in the copper rich corner present complex sequences of transformations of phases in the solid state complex resulting of a competition between thermodynamic phase stability and kinetic of precipitation. If the J’-D022 phase is often mentioned like the hardening phase of the ternary alloys Cu-Ni-Sn by its degree of coherency with the matrix, its action is accentuated by contribution of the internal stress generated by the spinodal decomposition that occurs before.

6 [1] [2] [3] [4] [5] [6] [7] [8] [9]

References T.B. Massalski, Binary phase diagrams, 1986, 2 F. Sadi, Y. Bienvenu, F. Bacou, R. Bailly, Matériaux 2002 T.F. Bower, M.R. Randlett, “Solidification structure of copper alloys ingots”, Metals Handbook, 9th, 9, ASM, 638–645 J.-C. Zhao, M.R. Notis, Acta Mater. 1998, 12, 4203–4218 J.-C. Zhao, M.R. Notis, Scripta Materiala. 1998, 11, 1509–1516 PH. Hermann, D.G.Morris, Metallurgical and Materials Transactions A. 1994, 25A, 1403–1412 M. Kato, T. Mori and L. H. Schwartz, Acta Metall., 1980, 28, 285 R. Labusch, Physica Status Solidi, 1970, 41, 619 B. Janson and A. Melander, Scripta Metall., 1978, 12, 497

39

Suplier Session

40

329

Horizontal Casting Technology for Copper Products W. Müller, P. Schneider SMS Meer GmbH, Mönchengladbach, Germany

1

Introduction

In companies manufacturing copper tubes new investments might become necessary in the foundry due to various reasons: 1. Replacement of existing vertical continuous or semi continuous casting plants 2. Increase of casting capacity 3. Producing copper billets inhouse instead of purchasing For the selection of a suitable casting machine different process types are available. The vertical semi continuous casting requires low investment and operator’s skill with the disadvantages of re-melting of head and tail, mostly manual operation and transportation, frequent starting risk and high probability of unstable production conditions with reduced quality. These disadvantages are overcome by a vertical fully continuous casting machine, which however, requires a higher investment. As an alternative the horizontal casting process could be considered with the advantages of a continuous operation and a lower investment than the vertical fully continuous caster, especially regarding foundation, intermediate transportation of raw material or cast billets and labor costs. Furthermore the space requirements are significantly lower compared to typical vertical casting plants. A further alternative might be the horizontal casting of tube shells with the subsequent processing on a high reduction rolling mill. This is an attractive option in those cases where an existing extrusion press is working already at the limit or where a complete new production line is in consideration.

2

Horizontal Continuous Casting of Copper Billets

Although the horizontal casting process is well established for casting brass billets for more than 30 years, its application for producing copper billets is still an exception compared to the vertical casting process. However one new horizontal casting plant for copper billets has been delivered by SMS Meer Technica in 2005 and recently started the production. The typical layout of such a horizontal billet caster for copper is shown in Picture 1 with channel type induction melting furnace, holding furnace for superheating and buffering of the liquid metal, attached cooler / graphite mould, cooling water distribution unit, secondary cooling water station, withdrawal unit, movable saw and run out conveyer with storage table. A video clip may demonstrate the function of the plant. The one-strand horizontal casting plant for copper billets of a diameter of 300 mm has a typical casting speed of 110–140 mm/min resulting to an output of 4.2–5.3 t/h. At an operating

330

Figure 1: Horizontal copper billet caster

time of 5.400 h/year a total casting capacity of 22.700–28.600 tons per year is expected. Multiple strand casters of up to 3 strands are also available today. Experience has shown that a pre-condition for a reliable and quality-wise good production is the careful desoxydation of the copper melt and the avoidance of any oxygen pick-up in order to reduce the slag amount occurring with phosphor addition to a minimum. Residual oxygen content will also reduce the lifetime of the graphite die, which can be extended to up to five days with a content of oxygen below 10 ppm. The cast copper billets show a very smooth surface and an internal structure as typically obtained by horizontal casting (Picture 2). Since nearly all copper billets are further processed by an extrusion press to copper tubes, the question always comes up about the behavior of horizontal cast billets compared to vertical cast billets in the extrusion press. Concern is the eccentricity of the extruded copper shells, especially in the high ratio extrusion, where the eccentricity of the shell cannot significantly be improved during the subsequent drawing process. Extrusion trials have been carried out with vertically and horizontally cast billets of the same outside diameter (305 mm) in a new extrusion press. The eccentricity of most extruded tubes of size 65 u 3.1 were in ranges below 5 % independent of the type of casting. It therefore appears that the horizontally cast billets show the same quality as the vertically cast. However, it must be taken into account, that these trials were only based on a few billets and the results need to be confirmed in larger test series. It should be mentioned that our sister company SMS Eumuco has developed a new press, where the billet is kept in the center by a new centering system, which reduces the eccentricity

331

Figure 2: Internal structure of horizontally cast copper billet

Figure 3: Eccentricity of aluminum tubes in a 45 MN indirect extrusion press

in a reliable way. Levels of eccentricities mostly below 3 % were reached by extruding aluminum alloy No. 5052 as shown in Picture 3. Increased experience in horizontal continuous casting of copper billets combined with new developments in the extrusion process open possibilities to again consider the horizontal casting as a real alternative to vertical casting. The total investment of a horizontal caster is significantly lower than for a vertical fully continuous caster and total operating costs are lower than for a semi continuous caster.

332

3

Tube Shell Casting

In tube manufacturing an alternative to the conventional process route billet casting and extrusion is the directube® process, where tube shells are cast and further processed on a high reduction planetary rolling mill before drawn to final sizes. Picture 4 shows the general layout of the complete directube® plant with charging equipment for cathodes and return scrap, two melting furnaces, horizontal continuous casting line, which basically consists of a pressure controlled holding furnace with coolers /graphite moulds, withdrawal unit, movable saw, transportation equipment and surface milling machine. The rolling line consists of a loading section, mandrel thrust block arrangement, shell feeding section, planetary rolling mill, movable shear and inline coiling unit.

Figure 4: directube® process

The shown plant was realized in 2000. The horizontal continuous casting line is designed for four-strand casting of shells with an outer diameter of 100 mm and a wall thickness of 25 mm. Based on the number of strands an annual capacity of copper tubes between 10.000–25.000 t/y can be reached. This capacity corresponds to a casting speed of 375–450 mm/min. One important factor for an excellent casting quality is the low content of oxygen. Special melting furnaces with a siphon spout in combination with a closed transfer, launder, installed at the tilting axis of the melting furnace, avoid the pick-up of oxygen. A further key component within the casting plant is the three-chamber holding furnace with the pressure control as shown in Picture 5 with its various operation modes. Section 1 shows the furnace filled with copper at the start of a casting campaign indicating the casting chamber, inductor, pressure chamber together with filling chamber. By increasing the pressure in the pressure chamber during production the metal level in the casting chamber always remains constant and the constant metallostatic pressure in the mould ensures stable casting conditions and results in an uniform product quality.

333

Figure 5: Tube shell casting with pressure controlled holding furnace

Section 2 represents the status just before refilling with metal from the melting furnace. The fresh melt will be nearly completely absorbed in the pressure chamber and therefore turbulences and splashes of slag into the mould with a negative effect on quality are avoided during the refilling operation. At the end of a production campaign before changing the graphite die of the cooler, the pressure controlled holding furnace is emptied in the pressure chamber to the minimum level and also in the caster chamber to just above the cooler as represented in Section 3. Then the pressure is released and the melt level drops below the cooler (Section 4), thus avoiding any metal loss and enabling a rapid change of the cooler/die assembly of a multi-strand casting unit which are all together mounted on a common adapter. The two plugs at the rear end of the holding furnace are placed for emergency and completely emptying of the holding furnace. During all operation stages the metal content in the holding furnace can be monitored with a weighing device in the furnace base frame. This patented method allows automation of the melt transfer from the upstream melting furnace. The advantages of this furnace for production and quality are: 1. Constant metallostatic pressure above the mould at all times, irrespective of the filling status of the furnace therefore stable casting conditions and product quality. 2. No slag formation of the metal by using nitrogen in the pressure chamber reduction of slag inclusions in the cast product. 3. Avoidance of bath turbulence during recharging reduction of the risk of break-outs, increase of production safety. 4. Quick mould change, as the metal level drops below the mould due to release of pressure reduction of down time and no metal loss.

334 Depending on the required production rate for tube shells the horizontal casting plant is equipped for one up to four strands, which are cut to approximately 18 m single length and further processed to tubes with coil weights of 1.000 kg. The grain size of the produced shells can be influenced by the proper adjustment of cooling water flow and graphite design in order to achieve grain size as desired for the final product. Picture 6 shows the micro section of tube shells with large and with small grains. Table 1: Production Cost Overview for different copper tube production processes General Plant Data Finish tube capacity Total output (consusmption figures) Net Production hours per year Yield factor

High Ratio 60.000 1,36 5760 80

Low Ratio 60.000 1,36 5760 80

directube 60.000 1,23 5760 80

People People People People

38 0 0 38 32.500

35 0 0 35 32.500

42 0 0 42 32.500

Investment costs Machinery and equipment Tools, spare and wear parts Foundations, erection and commissioning Infrastructure, production halls, bay cranes, etc. Properties and office buildings Additional investments Sum Investments

Mio. Euro Mio. Euro Mio. Euro Mio. Euro Mio. Euro Mio. Euro Mio. Euro

26,10 0,75 4,70 0,00 0,00 1,31 32,85

27,00 0,65 4,86 0,00 0,00 1,35 33,86

26,50 0,77 4,14 0,00 0,00 1,33 32,73

Consumption figures Electricity Cooling water Natural gas Protective gas Compressed air Lubricants and other consumables

MWh/year m³/year nm³/year nm³/year m³/year Euro/year

57.466 2.039.040 5.899.034 338.498 676.800 65.250

56.319 2.240.640 5.880.012 336.838 641.520 67.500

52.697 2.027.520 4.423.357 586.363 747.000 66.250

Production costs Machinery, foundation, erection, tools, spare parts Euro/ton Other investments (long depreciation period) Euro/ton Additional costs (rental fees, maintenance, etc.) Euro/ton Electricity, cooling water, natural gas, etc. Euro/ton Tools Euro/ton Intermediate storage costs for copper materials Euro/ton Personnel Euro/ton

68,43 0,00 17,47 109,19 20,20 23,92 20,74

70,55 0,00 17,96 108,47 17,00 25,24 19,19

68,19 0,00 17,69 95,05 10,42 19,38 22,82

Sum production costs, without copper

259,96

258,40

233,54

Personnel Operators (shift personnel) Other shift personnel, foreman, technician Adminastrative and laboratory personnel Sum required personnel Average income per year ( /year/person)

T/year Factor h/year %

Euro/ton

335

Figure 6: Macrosections of tube shells of DHP copper, 100/50 mm (small grain sizes and large grain sizes)

By bypassing the extrusion press of the conventional process route the tube shell casting offers advantages regarding production costs as shown in table 1, where the production costs of the high ratio and low ratio extrusion are compared with the directube® process. Basis for the sample calculation is the production of 60.000 t/y of sanitary tubes. In order to have a comparable basis, three directube® lines have to be considered. Cost advantages for media, tools and intermediate storage result to lowest production costs for the directube® process.

4

Conclusion

New concepts and developments in continuous casting of copper for the production of copper tubes contribute to the demand of the industry for alternative cost saving process routes for higher product quality and better productivity. Horizontal continuous casting plants for copper billet and copper tube shell production mirror this most important requirement of copper tube makers for medium production capacities.

336

Horizontal Direct Chill (HDC) Casting of Aluminium – the HE Universal Caster F. Niedermair, H. Zeillinger Hertwich Engineering, Braunau, Austria

1

Introduction of “The Universal Caster”

HDC casting has well earned its place in modern Al-casthouses, and is still gaining momentum. Hertwich Engineering has successfully commissioned some 50 Horizontal Direct Chill Casting Plants (HDCs) world wide to date. Todays generation of HDC-casting machines is one of the most versatile pieces of equipment, which may be employed to produce any of the following: • T-bar • Foundry ingot • Busbar and anode rod • Extrusion billet • Forging stock • SSM-feedstock • HDC for magnesium ingots etc. • This is why the HERTWICH horizontal caster really deserves the attribute “universal caster”. Let’s look at some of the most commonly produced shapes.

2

T-Bar and Foundry Ingot

Over the past few years especially the mass producers of remelt product have discovered the Horizontal Caster to fulfil their demanding needs in terms of product quality and process control. Large scale production of high quality foundry ingot has been shifted from ingot belts to HDC. T-bar/Ingot casting on VDC casting machines have lost ground to the over the years developed HDC casting process. Figure 1 shows T-bars produced on the HDC casting machine. The VDC process has the following drawbacks compared to the HDC process: • Higher costs of the VDC caster, especially due to greater building height required, necessity of overhead crane and foundation for the casting pit. • The semi-continuous character of VDC-Casting results in lower productivity. A great amount of set-up work per drop is required, which is rather labour intensive, whereas with the HDC, continuous production runs of 3 to 20 days are common. For T-bar production only one to two operators per shift are needed (two operators for cast start and stop). • On VDC plants sawing is not integrated in the process, so that an additional sawing station plus operator is required. HDC casting employs an automatic flying saw, which cuts the Tbars to length without disturbing the casting process.

337

Figure 1: T-bars produced on Hertwich HDC

Pouring of metal into open moulds causes dross: Sows, pigs and ingots were traditionally produced employing the open mould technology. Although this technology was improved over many years, dross formation and inclusions are still unavoidable. Due to cascading, turbulence occurs when filling the mould. So a relatively big unprotected surface area is offered to the atmosphere for oxidation. The dross formation is mainly ruled by the metal temperature, pouring height and pouring rate. Values achieved during production of pure aluminium sows are shown in table 1. Table 1: Dross formation during production of pure aluminium sows Pouring height [m]

Temperature [°C]

approx. 0,2 to 0,3

700–770 > 800 approx. 750 approx. 850 to 900

approx. 0,6 to 1,0

Dross formation [kg per ton of poured metal] 0,2–0,4 0,3–0,6 2,5– 5–7

The HDC process is almost free of dross formation. It results in savings due to avoided metal losses and in inclusion-free products. On the HDC the metal flows smoothly, protected by an undisturbed oxide layer via launder and tundish to the mould. Thus leaving no chance for oxides and other impurities to get into the product. The HDC cast T-bar and foundry ingot are chilled at least ten times faster than sows and pigs. This ensures a fine and uniform grain structure as well as a uniform analysis throughout

338 the cast product. A further step ahead in the production of remelt products in terms of quality is the combination of the HDC process together with an Inline Degasser and Ceramic Foam Filter (CFF). Both items are needed to obtain T-bars and foundry ingot free from porosity and inclusions. Summary of advantages of HDC products vs. sow and open mould ingots • • • • • • •

low hydrogen, extremely low oxide inclusions fine uniform grain structure consistency and uniformity of alloying elements’ content and distribution no gravity segregation no cracks and shrink holes and no water inclusions consistent dimension, straightness, weight smooth surface, easy for stacking and strapping, compact bundles

2.1

Design Details

The most common HDC caster has a width of 3.000 mm. The 3 m wide caster may produce up to 13 tonnes per hour T-bar or 8 t/h foundry ingot. At present HE supplies the largest HDC caster built to date. It is designed to cast 4 strands T-bar 850 u 300 mm or 24 strands ingot 106 u 106 mm simultaneously. Continuous production is 17 t/h for T-Bar and 12 t/h for ingot.

Figure 2: HDC Casting of foundry ingots

339

Figure 3: Typical layout of the Universal Caster

A wide range of alloys can be produced, for instance ranging from pure aluminium to 12% silicon and up to 5 % magnesium. (Ref. Figure 2) Each product has a dedicated exit route, downstream of the flying saw (Ref Figure 3). Provided all exit systems are installed, a product change can be undertaken within one shift, by changing to a different tundish/mould set and loading the new applicable cast recipe on the PC. The fully continuous HDC process ideally lends itself for automation. This advantage has been well exploited by Hertwich Engineering. All downstream equipment is fully on-line with the casting machine and no additional personnel is required. Sawing, weighing, hard stamping, ink marking, labelling, stacking and strapping is carried out fully automatic. (Ref. Figure 4)

Figure 4: Foundry ingot automatic marking, stacking

340 During past years the plants were improved consistently and now feature automatic cast starts and stops as well as automatic tundish adjustment. The HDC plant is controlled by the Hertwich PCPLC system, which offers an error-manager system and a menu-type casting recipes. Besides, all important plant parameters are monitored, controlled and stored and are available to a clients host PC for further processing or storage.

3

Busbar and Anode Rod

In the primary aluminium business, a HDC plant is often initially purchased for producing busbar for potline construction, but designed to allow later conversion to froundry products. In phase two the busbar caster is then typically turned into a T-bar or foundry ingot caster to produce high quality remelt products for foundries, by adding the relevant handling equipment. Excellent electrical conductivity of busbar, good surface finish and straightness are achieved. The inline flying saw cuts continuously-cast busbars to exact finished length. Busbar lengths of 0,6 m to 20 m can be cast, no additional cutting operation is required and the number of welding joints is greatly reduced. Furthermore, HDC busbar production cost is significantly lower than VDC. Figure 5 shows the production of busbar on a HDC-Casting Machine.

Figure 5: Production of busbar on Hertwich HDC casting machine

4

Extrusion Billet, Forging Stock, SSM Feedstock

4.1

Extrusion Billet

To operate their own remelt plant poses a big challenge to most extrusion firms. Hertwich Engineering has developed a remelt plant to meet the limited capacity requirement of a typical extruder. The “Compact Type Remelt Plant” represents an economically very interesting concept, particularly in the range 4.000 to 20.000 tons per year, which comprises an integrated, conti-

341

Figure 6: Schematic of a Two Chamber Furnace and HDC Caster

nuously operated, fully automated system starting with the handling of scrap and ending with delivery of cut to length homogenized extrusion billets (logs). Within recent years Hertwich Engineering has supplied some 20 such plants to extruders worldwide. Heart of the CTRP is a HE HDC billet caster. Figure 6 shows the schematic of a Two Chamber Furnace and HDC Caster. The Horizontal Caster is the key machine in the effective HERTWICH Compact Type Remelt Plant (Figure 7) In-house generated extrusion scrap can be charged by a dedicated charging machine into the Two Chamber Melting and Casting Furnace. The stationary furnace consist of a melting and a holding chamber. Applying the submersion melting process allows remelting of profile scrap at a metal loss of less than 0,5 %. Primary metal and clean scrap from the market may be remelted as well. For contaminated scrap, like painted profiles, HE offers a 3 Chamber Furnace. This furnace evaporates and combusts the hydrocarbons from the paint prior to melting. Thereby additional metal loss is avoided, increasing the thermal efficiency and destroying harmful compounds like dioxins etc. Through a tap hole in the holding chamber, the metal flows via an Inline Degasser

Figure 7: Layout of a CTRP for production of extrusion billets from clean and contaminated scrap

342 and CFF to the Horizontal DC Casting machine. Melt flow from tundish to mould is close to the bottom of the tundish, hence no oxides floating on the surface of the melt may ever get into the mould. The mould is of short water cooled design with integrated lubrication for continuous casting. Size of casting machine and number of moulds determine the production rate. Depending on the size of machine and billet diameter typical casting rates range from 1000 kgs/hour to 3000 kgs/hour. However, smaller or larger machines are available. A billet diameter change is done fast and easy by changing the set of tundish and moulds. Figure 8 shows the HDC caster during billet production. Casting campaigns last several days, for instance a casting cycle starts on Monday morning and stops on Saturday morning. Cut to length billets can automatically be pin stamped on the cut face. A flying saw integrated to the HDC casting machine automatically cuts the continuously cast strands into billets of the required length. The max. cut length of billets is ruled by exit conveyor length, usually 6 m or 7 m long billets are produced. After Sawing, billets are directly fed into a HE Continuous Homogenizing Furnace for heat treatment.

Figure 8: Horizontal Caster with Flying Saw, Billet Production

4.2

Forging Stock (Spaghetti Production)

Traditionally forging stock is produced by VDC casting of regular billets followed by extruding to the required diameter. On the HE HDC Caster forging stock can be directly produced in dia-

343 meters ranging from 25 mm to 125 mm, followed by scalping. The main advantages of HDC produced forging stock are: no extrusion grain texture and substantially reduced conversion cost.

4.3

SSM Feedstock

SSM feedstock is readily produced on the HDC, by fitting an electromagnetic stirrer around the mould. This concept has been developed and successfully tested some years ago.

5

Metal Cleanliness

It must be stated that proper upstream operation and equipment for adequate metal cleanliness is a prerequisite for producing a top quality product and achieving economical cast durations, 3 to 20 days. Particular attention should be paid to the following: • sufficient furnace capacity • proper furnace operation: incl. skimming, fluxing, alloying, sufficient settling time, furnace change over (leaving a heel) • efficient degassing and –refining • adequate twin box type ceramic foam filter • rod feeding for modifying and grain refining. As a service HE offers to consult on this matter individually.

6

Conclusion

The Universal Caster from Hertwich Engineering has become a familiar sight in cast houses and extrusion shops. Its versatility, the low investment costs involved, the high quality of the product and the lean operating labour required make this plant unique. The evident trend is away from ordinary sow – or open mould ingot casting and towards the superior HDC process - a clear step forward due to quality conscious customer demand.

344

Aluminium-Semi-Continuous Casting Technic, State of the Art G. J. Brockmann Maerz-Gautschi Industrieofenanlagen GmbH., Tägerwilen, Switzerland

1

Introduction

The first patent concerning semi-continuous casting of aluminium billets, applicated by the former VAW/Germany, took place at Germany in 1936 (Figure 1). Since then many different ideas for designing casting devices followed. Finally the development focussed on the idea of lowering a casting table. To lower a casting table, the plate may be moved with a thread spindle drive, a rope system or with a hydraulic cylinder. Here the development focussed on using a hydraulic cylinder. To guide the casting table with high accuracy, usually a guiding systems with rails at the pit walls was installed. Due to problems caused by metal particles in the cooling water, sticking between bearing shoe and rail, the guiding system is now mostly implemented into the casting cylinder. The size of casting machines varies from 2 t to about 120 t metal weight per cast. The number of moulds depends on the billet or slab size and may be anywhere in the range from 1 to 144 pieces. Typical Mould No.’s for billet moulds are 12, 24, 48, 72 and 96 to match a good mould pattern. The cast length varies in the range of about 3’000 mm up to 7’000 mm, in single cases even up to more than 10’000 mm. The billet sizes are between 3” and 14”. A typical billet diameter is 7”, but also diameters up to 28” are of interest for forging ingots. The geometry of the cast product may be round (extrusion billets) or rectangular (rolling slabs) or even show an elliptical or t-formed shape. Many modern casting machines allow the cast of billets and slabs with the same machine. All wrought aluminium alloys are VDC-castable. About 70% of casted alloys is AA 6xxx, but also 1xxx, 3xxx and 7xxx are of interest. New demands of the automotive industry request casting of tin-containing alloys (tin is exchanging lead in the futures). Moulds for casting such special alloys are under development. In the following description, typical design properties of state-of-the-art aluminium-semicontinuous casting machines – or better said: Vertical Direct Chilled (VDC) Casting Machines will be discussed.

2

Casting Machine Layout

Main components of a state-of the-art casting machine (Figure 2) are: • • • • •

Casting pit (equipped with access ladder, tubing and pumps) Base frame with pit shield (top of casting pit, mostly with integrated evaporation) Pit covering platform Safety devices Mould table frame (mobile water frame, movable and/or tiltable)

345

Figure 1: First VDC Patent

• • • • • • • • •

Casting cylinder (single action, hydraulic, mostly internally guided) Platen (mostly with torque limiter) Starting head base (roof or lattice design) Starting heads (for billets) or blocks (for slabs) Anti-tilt grid guard Mould table (carrying billet or slab moulds) Mould System (Hot Top, Airsol Veil©, Airglide© and similar) Launder system (for side or central mould feeding) Hydraulic power unit (separately for casting cylinder and actuators)

346

Figure 2: VDC Components

• • • •

Cooling water system (with cooling tower, tanks, filters and oil separator) Mould oil supply system Mould air supply and control system (if required) Switch board and control desk (mostly with visualisation)

The components will be discussed in detail as follows:

3

VDC Components

3.1

Casting Pit

The casting pit is looking like a big tank, mostly with walls from concrete, sometimes from steel or partly with an internal steel cladding, submerged under the plants ground level. The dimensions depend from the size of the casting table and the desired billet or slab length. A typical shape of a 25 t VDC may be about 4’000 x 2’500 mm square and 7’000 mm depth, the resulting volume is about 70 m3 then. Smaller pits have an extension at one side in the nearground area to get access to the place under the casting table, bigger pits have a separate parallel pump pit with a door between both pits at the ground. The pit is equipped with cooling water

347 pumps with tubing and an access ladder. In the pit’s ground is usually a center bore for receiving the casting cylinder. The bore’s length is another time about the same like the pit’s depth. The bore is inside cladded with a steel tube with an inner diameter like the cylinder’s outer diameter plus spacing.

3.2

Base Frame

The base frame is placed at the top of the casting pit, incorporated into the cast house foundation and made from mild steel profiles, corrosion-proof painted. The base frame includes the pit illumination with waterproof halogen lamps. Furthermore, the base frame has slots incorporated for the connected steam extraction device. An axial fan, installed into the steam collection system, guides the extracted steam through galvanized ductwork releasing the steam through a stack into the atmosphere outside the building. On the base frame a pit shield is mounted guiding splashing cooling water and liquid metal into the casting pit.

3.3

Pit Covering Platform

A pit covering platform closes the open pit area during billet or slab removal. So the operator can handle the billets or slabs without danger. For small casting machines it can be moved by hand, bigger machines get a motor-driven platform, secured by proximity switched, connected to an interlock system.

3.4

Safety Devices

To prevent operators from falling into the open pit, fences or chains according to international safety rules are implemented. Such fences can be actuated manually or hydraulically.

3.5

Mould Table Frame

The mould table frame incorporates the cooling water system and is made from rectangular hollow cross section from stainless steel. A painting is not necessary. The water connections to the external system are done by automatic couplings or bearing inlets. The connections to the moulds are realized by flexible hoses. The design may be done in three different ways to meet the operators need: • Movable: the table is rolling on rails and can move side wards • Tiltable: the table is fixed with bearings at one side and can be tilted-up by means of hydraulic cylinders • Movable/tiltable: the table drives on rails into a tilt station beside the caster and can be tiltedup there for easy maintenance work

348 Two double acting hydraulic cylinders tilt the frame from operation to upright position, max. 85o to prevent internal pollution of the moulds.

3.6

Casting Cylinder

State of the art is an internal guide of the ram to assure a minimum rotation of the cylinder rod over the whole length of the stroke (Figure 3.). The maximum rod rotation is typically ± 4’ arc. The casting cylinder is mostly a hydraulic single action ram type cylinder, secured to a heavy steel plate, which is incorporated into the casting pit foundation. The cylinder must be designed for maximum rigidity with excentric loads on the platen. The ram may be fabricated from stainless steel or from mild steel with a multi metal-oxide coating to guarantee a smooth, low-friction lifting and lowering motion. The ram is solid up to a diameter of about 580 mm, bigger diameters are designed as tubes with a 100 mm wall thickness, filled with steel balls after erection.

Figure 3: Casting Cylinder

3.7

Platen

The platen is the mounting platform for the starter head base and fabricated of mild steel rectangular hollow cross sections with a coating approved for molten aluminum service. Stainless steel flat bars, machined after welding onto the framework, offer accurate mounting surfaces for the starter head base. The platen carries magnetic end switches for an accurate positioning of the platen in the pit. Between casting cylinder and platen, a torque limiter, is installed. This safety device protects the internal guiding of the casting cylinder in case that a piece of metal is jamming between platen and casting pit wall, causing a high torque on the system. A shear pin is breaking then, releasing the platen for free rotation to prevent overload of the internal guiding system.

349 3.8

Starting Head Base

There are two different designs for the starting head base: • Roof top box design (mostly used for slab casting) • Lattice design (mostly used for billet casting) The Base is made from mild steel with a special coating approved for molten aluminium service. The welded heavy-duty steel structure is precision machined as complete unit on both the top and bottom surfaces to ensure proper alignment between the starting heads and the moulds. The starting heads are mounted on supports with self aligning devices.

3.9

Starting Heads

The starting heads for billets are made from steel, aluminium or steel with a special shaped aluminium top. They are movable and will be centred while beeing lifted up into the moulds by means of ribs or three-ball self centring devices (Figure 4.). For rolling slab casting the starting blocks are normally made from aluminium. The alignment is done with centring pins. The centring is done by fixing the position of the moulds initially by hand.

3.10

Anti Tilt Grid Guard

A steel frame prevents the toppling of the billets once the mould table has been removed. It is mounted in the pit between the mould table and the starter head base.

Figure 4: Starting Head (right: inserted)

350 3.11

Mould Table

The mould table accepts the moulds and all piping for water, oil and air (if necessary). It may be part of the moulds themselves (AirslipTM) or only support platform for separate moulds (Airsol Veil® and Airglide®). Mould table with mounted moulds and starting head base with installed starting heads are to be handled as aligned units.

3.12

Mould System

• Hot Top: A feeder top ring from refractory material (Hot Top) reduces thetemperature drop during casting to improve the billet’s properties; approved technology for simple demands. Actual versions of this technology incorporate a continuous lubrication instead of grease coating of the bearing surface. • Airsol Veil®: Hot Top mould with a modern air/oil mist supply forming a cushion of parting agent in front of the bearing area to reduce the cooling rate to form a thinner shell zone of the billet (Figure 5.). • Airglide®: Design similar to Airsol Veil®; an additional solid graphite ring in the bearing area provides a non-wetting area, reduces therefore the oil consumption and minimizes the billet shell zone (Figure 6.). • Others: F.e. AirslipTM incorporates a semi-permeable graphite ring for direct air and oil supply. The effect on the billet properties is similar to the Airglide® technology.

3.13

Launder System

The launder distributes the molten metal to the connected moulds. It is a welded steel construction with insulated pre-cast refractory sections; important is the minimizing of the temperature

Figure 5. Airsol Veil® Mould

351

Figure 6: Airglide® Mould

drop between launder inlet and the final mould furthest away from the inlet. Launder systems are equipped with preheating devices to keep the system dry between the casts. 3.14

Hydraulic Power Unit

The hydraulic power unit comprises oil reservoir, pumps, filters and valves. The system is necessary for lifting-up the platen with mounted parts and for actuating f.e. the hydraulic cylinders for tilting the mould table frame. Dropping the platen is done by its own weight and controlled by an electrical precision valve to enable a casting speed in the range of about 20 to 300 mm/min. It is recommendable to operate the system with a flame retardant hydraulic fluid. Additionally an emergency platen lowering valve for manual operation in case of an electrical failure is installed in reach of the operator. 3.15

Cooling Water System

The cooling water passing through the moulds is collected in the casting pit. The maximum water level is simply controlled by overflow, the minimum water level of 1’000 mm above pit ground according to safety rules is controlled by a floating switch. Various water levels in the pit can be controlled by means of pressure sensors or several floating switches at the pit’s wall, controlling the drain pumps. An automatic back-flush filter is recommendable as well as an oil separator to keep the cooling water clean. Separated hot and cold well cooling water tanks together with a cooling tower outside allow individually adapted cooling water temperatures by means of a mixing valve. An additional pulsation valve, commonly used for casting rolling slabs,

352 improves the cooling effect. Finally an insulated emergency over-head water tank is necessary in case of an electrical failure to finalize a cast according to safety rules.

3.16

Mould Air and Oil Supply System

For moulds operated with parting agents a separated oil supply system must be installed, using mostly biodegradable vegetable oils. Since semi-permeable graphite rings tends to plugging, mineral oil must be used then. The mould air must be clean and dry and is supplied by an air compressor with air dryer.

3.17

Switch Board and Control Desk

Switch board and control desk allow to operate the casting machine in a convenient and safe matter. The control desk is IPC and PLC supported and usually equipped with a visualization of all caster functions. Connections to the plant’s Ethernet network allow improved functions (Scada). Operation, recipe administration, alarm logging and trend representation are to be handled in a user friendly way.

4

Product Demands

Geometry, dimensions and tolerances of the cast products are to be kept within narrow limits according to agreed rules like DIN EN standards and others. The surface quality depends on alloy, mould technology and casting parameters. The shell zone is depending from the mould design and has to be mechanically removed by milling/turning afterwards – if necessary. The structure of billets and slabs may be improved by a following homogenization.

5

Economical/Ecological Aspects

Investments and caster operation costs have to be minimized. All parts should be robust with a long lifetime and easily exchangeable in case of damage. With view to ecological demands no serious pollution should be caused by operating a VDC casting machine.

6

Development

The continuous improvement process in designing a VDC casting machine must be focussed at: • Cost reduction of equipment and operation • Improvement of product properties • Extension of the technology to f.e. new alloys (i.e. tin alloys)

353

OCP Crucible Monitoring System in Long-Term Tests W. Schmitz, F. Donsbach, Otto Junker GmbH, Simmerath / Germany

H. Hoff Lios Technology GmbH, Cologne / Germany

1

Introduction

One key design feature distinguishing induction furnaces from other heating equipment is the fairly thin ceramic lining between the live water-cooled copper conductor and the molten metal bath. Depending on the furnace size, the thickness of this ceramic lining varies between 10 and 15 cm; it diminishes noticeably as a result of wear or crucible erosion. Inductor insulating materials such as insulating varnish and bandages are heat resistant up to about 150–200 °C. If overheating occurs at this point, the insulation may become damaged or even electrically conductive, resulting in interturn short-circuiting of the coil. The coil repair effort required in this case will render the furnace inoperative for several days, even if a spare coil is on hand. In the worst case, which has rarely been documented but is nevertheless a possibility, the melt may penetrate all the way through to the water-cooled coil with all attendant risks of a furnace breakthrough and ultimately, a steam explosion. These considerations, together with furnace users’ economically motivated demands for a maximum service life of the ceramic furnace lining, call for a technology which permits a “isual” inspection of the gap between the ceramic furnace lining and the induction coil. Addressing this requirement, various technical solutions for monitoring the crucible have been proposed and implemented in the past.

2

Overview of Conventional Crucible Monitoring Systems

The most important of these is the classic earth leakage monitoring system. In this technology, a d.c. or a.c. voltage of a defined, fairly low frequency is applied to the induction coil and the system measures the current flow to earth. For this purpose the molten metal bath must be earthed via an earthing rod in the bottom of the crucible. This earth fault monitoring system, although by now a standard feature on virtually all induction furnaces, has a number of disadvantages. For one, it is not selective, i.e., defined tests and disconnection steps must be carried out whenever an earth leakage is detected so as to determine whether the fault has occurred between the coil and the molten metal or in any other part of the equipment, such as in the switchgear or even in the water recooling system. Another disadvantage of this earth leakage monitoring method is that in the event of infiltration or penetration of molten metal to the coil, evidence of this condition will be not be obtained until fairly late. As a result, the furnace must be emptied quickly if a current flow between the melt and the coil is detected. In any event, minor damage to the coil may have occurred already. Despite these drawbacks, earth leakage protection systems remain an essential safety tool in the operation of coreless induction furnaces and are un-

354 likely to be supplanted by more recent measuring technology. Earth leakage monitoring will always remain in use as an ultimate safety feature since naturally, the entire furnace including the switchgear and water recooling system must be continuously checked for earth faults. The use of thermocouples between the hot-face lining and the coil levelling mix, as well as in the furnace bottom, is another technique employed. However, this method can yield only spot measurements (at least if the cost and effort involved is to be kept within reasonable limits) and is therefore not capable of monitoring the entire crucible. In the past, wire netting in various geometrical configurations has been placed between the coil levelling mix and the hot face lining. The idea is to detect an electrical continuity between an advancing tip of molten metal and the net. One particular disadvantage of this method is that it provides no trend indication, i.e., no advance warning is given. Moreover, it is virtually impossible to identify the fault location, and numerous spurious faults will be detected since the net, as an electrical measuring device, is subject to many stray voltage and current effects in the magnetic field of the induction furnace. A further process in industrial use1), 2) relies on the use of sensor grids comprising an array of metallic electrodes in a comb-type configuration. These electrodes are used to measure the electrical resistance of the ceramic lining. As this resistance is temperature-related, it is possible to infer the temperature in specific crucible segments. Fault locations can thus be identified in relation to the furnace circumference, and an advance warning functionality is obtained. However, given the resistance-based nature of the measurement, the system must be adapted to the specific ceramic material used. The readings obtained will be affected by any change in the composition of the ceramic lining, moisture effects, and furnace-induced magnetic interferences influencing the electrical resistance measurement.

3

OCP Optical Coil Protection System

OCP (Optical Coil Protection System) stands for a latest-generation temperature measurement and monitoring technology which relies on fibre-optic sensors. Given their properties, such sensors are perfectly suited for interference-free monitoring of the crucible on induction melting furnaces. Figure 1 shows the typical crucible structure of a coreless induction furnace, with the OCP sensor cable firmly grouted into the furnace's permanent lining which is installed directly on the coil. Based on an optical fibre, the system utilizes a quantum-mechanical effect, the so-called RAMAN effect, for temperature measurement. Laser light of a suitable wavelength and modulation frequency is injected into the optical fibre. This laser light scatters on the bonding electrons of the solid state structure over the full fibre length and is detected as a backscatter spectrum. This spectrum contains the RAMAN lines, the intensity of which is a function of vibration levels in the solid state fibre structure, which in turn depend on temperature. A new, patented ’optical radar’ technique makes it possible to detect these lines locally and to measure an exact, high-resolution temperature profile through the optical fibre. Thus, OCP is a unique crucible monitoring system which enables us for the first time ever to determine the temperature field in the induction furnace irrespective of refractory type and design. By selecting a radial resolution of 60 measuring points, it is possible to represent the temperature curve in the manner of the familiar analogue clockface (Figure 2).

355

Figure 1: View of a typical crucible structure, showing the permanent furnace lining with the OCP sensor cable embedded (4)

Figure 2: OCP System monitor screen

By adopting an appropriate configuration of the sensor grids, the crucible can be vertically divided into several regions, although only a single optical fibre is used in all cases. Points of particularly high temperature, e.g., due to infiltration, erosion or cracking in the crucible, can thus be accurately localized and checked for potential hazards to the coil insulation. The sensor cable is currently rated for a maximum continuous operating temperature of about 260 °C,

356 which far exceeds the maximum temperature resistance of the coil insulation. The measuring system has a range of several kilometres, and a position resolution of 27 cm (related to the straight length of the optical fibre). The temperature resolution of this system is better than 1 K.

4

Design and Installation of the OCP Sensor Cable

The core of the OCP sensor cable, first of all, consists of a commercially available hightemperature glass fibre of the type commonly used in telecommunications. For mechanical protection, this fibre is enclosed in a stainless steel tube measuring 1.2 mm in diameter. The tube, in turn, is coated with a high-temperature insulating compound. The overall diameter of the sensor cable is 5 mm. In order to provide the fullest possible crucible sensor coverage in the direct vicinity of the coil, it is desirable to have a maximum length of sensor cable in the furnace. This is achieved by placing the sensor cable on the inside of the coil in meandering curves, using ceramic deflector elements, while taking into account the cable's minimum bending radius. Figure 3 illustrates this layout on a 6-tonne steel melting furnace. Here we have two meandering cable layers, one for the upper and one for the lower crucible region. As mentioned earlier, the cost and effort involved in providing extra layers is minimal. Once the cable is installed in this manner, the usual former is placed in the coil and a permanent lining made of high temperature resistant corundum concrete is cast, with the sensor cable thus embedded therein. For new equipment and coil overhauls, this can be done in the workshop. In the case of retrofits and under special circumstances, this work can even be carried out locally in the foundry. Since the measurement is taken at the same end at which the laser light is introduced, it will suffice, in the most basic case, to bring only one cable end out of the furnace. However, it has become standard practice to bring out both ends of each meander layer (Figure 4), which are

Figure 3: Arrangement of the OCP sensor cable on the coil of a 6-tonne induction furnace or melting steel

357

Figure 4: Lead-in bushing of an OCP sensor cable on the coil

then connected to a terminal box mounted on the furnace body. This has certain benefits at the commissioning stage and also facilitates diagnosis. Moreover, in the unlikely event of a fibre breaking inside the furnace due to inappropriate mechanical loads, it is thus possible to conduct the measurement in the reverse fibre direction, thus ensuring that the whole furnace can still be monitored. In a next step, the intended measuring end of the sensor cable is connected to an optical fibre transfer cable armoured to foundry standards. This cable is run to the location of the evaluator and connected to one of its ports. The evaluator consists of the actual measuring device in a 19” configuration plus a PC for visualization and evaluation. Both are habitually mounted in an appropriate panel or cabinet. Alternatively, the display and evaluation functions can be assigned to the visualization system of the melting furnace where this is technically feasible.

5

Display of Measured Temperature Data

The main screen of the OCP visualization system is shown in Figure 2. It displays a schematic top-down view of two furnaces, each comprising two meander layers, any one time. If more than two furnaces are monitored, the user can freely select which of these should be displayed on the left and right-hand side, respectively. As a general rule, the temperature curves for all individual meander layers are initially rendered in a screen window. For a less cluttered view, individual layers can be suppressed. This has been done in Figure 2, where the left image shows the current temperature distribution in the upper layer of Furnace 1 while the image on the right gives the current temperature distribution in the bottom layer of Furnace 2. As shown, each of these temperature profiles can be rendered in a polar or linear view. It is also possible to display a “relative” mode, i.e., a profile generated in relation to a given historical (reference) profile. In

358 this case the user will see the current temperature deviation from the selected reference profile (offset view). By selecting a playback function and entering a date in the respective window, past temperature profiles can be viewed at any time. It is also possible to show temperature profiles in an animated or “video” mode between a user-defined start and end point. In this animated mode, the temperature profiles can be rendered on the basis of absolute or offset values or in a “maximum” mode. In the latter case the system examines where maximum average temperatures have occured during a selectable time period (e.g., a shift or crucible campaign). The temperature profiles associated with such events are then shown successively in animated form. The software can also display the temperature curve at individual measuring points, or the average temperature, over time.

6

Evaluation of Temperature Measurements

Figure 5 shows the screen for entering alarm parameters. For every parameter, the user can enter a warning threshold and an alarm trigger threshold (possibly resulting in a furnace shutdown). The individual alarm criteria are “temperature”, “deviation from average”, “temperature change” and “uniformity”. At a more detailed level, these criteria are monitored as follows:

6.1

Temperature

The temperatures at one or more measuring points are monitored for overruns exceeding these preset thresholds.

Figure 5: OCP screen template for the definition of alarm thresholds

359 6.2

Deviation from Average

The measurements from which the temperature profile is plotted are initally processed into an average value representing the mean temperature in the respective zone. The system then checks whether the temperature at one or more measuring points deviates from this mean value by more than the preset threshold.

6.3

Temperature Change

Here the system determines whether a time-related temperature gradient, defined as a threshold, is exceeded at one or more measuring points. The unit in which this threshold is set in the various input fields is °C/min.

6.4

Uniformity

The uniformity parameter is largely identical with the “deviation from average” criterion, except that the averaging step and check for threshold overruns is not carried out in a single step over the entire circumference of the furnace. Instead, the system initially examines a radial sector (“pie wedge”) whose thickness is defined by the user in angular degrees (q) in the window marked “step”. This sector is then analysed in the same way as for the “deviation from average” criterion. In an interative process, this evaluation window is then advanced in a clockwise direction one measuring point at a time. The analysis is continued until the evaluation window has covered the entire circumference of the furnace. This is a valuable criterion when it comes to distinguishing local flaws from large erosion areas. In the OCP visualization system, warning and alarm messages are indicated by appropriate symbols in the temperature profiles (refer to Figure 2). The symbols can be suppressed to obtain a less cluttered view, although the underlying alarm functions remain active.

7

Practical Operating Examples

OCP systems are now successfully in use in coreless induction furnaces for melting copper alloys, aluminium alloys, cast iron and steel. In the following part of this paper I will examine three examplary cases in which crucible wear and premature crucible failures were detected in a timely and accurate manner.

7.1

2.5-tonne Vacuum-Type Coreless Induction Furnace for Melting Copper Pre-Alloys

This particular furnace is run in three-shifts to produce copper-iron pre-alloys. Such alloys pose exacting demands on the crucible material due to their aggressive chemical characteristics and fluidity. The tapping temperature is in the region of 1500 °C. The furnace is usually operated with a ready-made crucible consisting of refractory concrete. The space between this crucible

360 and the coil is backfilled with a dry ramming compound. A normal crucible campaign lasts about 2 weeks, depending largely on the degree of sintering of the backfill mix. If sintering pro-

Figure 6: Temperature profiles on a 2.5-tonne vacuum-type induction furnace at the start (left) and end (right) of a trouble-free crucible campaign, shown in absolute temperature (a) and offset display mode

361

Figure 7: Temperature profiles obtained after crucible cracking with resultant infiltration, again shown in absolute temperature (left) and offset display mode

pagated too far towards the coil the crucible will be very difficult to break out; moreover, there will be an increased likelihood of molten metal penetrating all the way to the coil in the event of a crack formation in the crucible. The degree of backfill mix sintering grows over the crucible campaign, causing the thermal conductivity of the backfilling material to increase progressively. As a result, the temperature in the permanent furnace lining and hence, the temperature local to the OCP sensor cable, will rise steadily. Figure 6a shows the temperature profile at the start (left) and near the end (right) of a normal crucible campaign in which no apparent local crucible failure occurred. Figure 6b gives the same data rendered in offset mode. Once a critical maximum temperature had been identified over several crucible campaigns, the OCP system was used as an indicator to identify the need for a scheduled re-lining. Figure 7 shows the situation for a crucible that had been in use for a week, i.e., half the normal crucible campaign. The graph on the left plots the absolute temperatures; its right-hand side counterpart gives an offset view of the same development. Over a span of a few charges, overtemperatures increasing from one charge to the next were identified in the 5 o’clock position, and the system eventually generated alarms of the “deviation from average” type. The crucible was broken out, and a crack was found at this point which had allowed the melt to infiltrate the ramming compound.

7.2

6-tonne Coreless Induction Furnace for Melting Stainless Steel

Figure 8 shows the temperature profile measured in the lower regions of a 6-tonne induction furnace for steel near the end of a crucible campaign. The “deviation from average” alarm mes-

362 sages (left) indicate general erosion towards the furnace spout. The “uniformity alarm” messages (right) point to the formation of caverns at four points. Measurements conducted on this

Figure 8: Temperature profile in the lower regions of a 6-tonne induction furnace for steel near the end of a crucible campaign. The alarm messages are of the “deviation from average” (left) and “uniformity” type.

Figure 9: Result of crucible measurements, showing an integral top-down view of the crucible. Uniform premature wear in the direction of the spout (12 o’clock) and local erosion in the 2, 5, 8 and 11 o’clock positions are readily apparent.

363

Figure 10: Caverns in the lower regions of the crucible

crucible prior to break-out confirmed the condition detected by the OCP system (Figure 9). Figure 10 is a photo of one such cavern.

8

Detection of Cracks in the Crucible

When a crucible cools down, e.g., over a weekend, numerous cooling cracks will form naturally due to volumetric contraction of the crucible material. These cracks will normally close again, due to thermal expansion of the crucible material, the next time the furnace is started up. However, an appropriate heating curve must be used to ensure this. Otherwise, the progressively melting metal may spontaneously penetrate still-open cracks and come dangerously close to the coil. The situation becomes even more critical if the furnace is filled with liquid metal before the cracks have closed. To simulate this situation, the following test was caried out: A thick-walled steel cylinder of a diameter equivalent to the inside diameter of the crucibles normally used in this application was placed in the middle of an unlined 1-tonne-furnace. This steel cylinder exhibited “artificial cracks” at a level about halfway up the furnace coil, these being in the form of 3 mm thick and 100 mm wide pieces of steel plate welded to the cylinder in the 12, 3 and 9 o’clock positions. The plate in the 3 o’clock position was welded to the surface horizontally and ended about 10 mm short of the furnace's permanent lining. An identical plate was welded on in the 12 o’clock position, but in a vertical direction. In the 9 o’clock position there was another horizontal plate which extended to the permanent lining. Finally, a 100 u 100 u 30 mm steel plate representing a cavern was welded to the cylinder in the 6 o'clock position. Figure 11 shows a sketch of this arrangement.

364

Figure 11: Thick-walled steel cylinder with welded-on steel plates simulating crucible cracks

The space between the steel cylinder and the furnace's permanent lining was filled with a quartzite dry ramming compound. The furnace was then switched on and operated at about 350 kW. This procedure was intended to simulate a cold start with existing metal-filled cracks. (It should be noted, with regard to the following illustrations, that the steel cylinder was slightly offset in a counterclockwise direction, so that the individual messages appear not exactly at 12, 3, 6 and 9 o’clock, respectively, but again, with some counterclockwise offset). Figure 12 shows the temperature situation before the furnace was heated up, with the graph on the left giving absolute temperatures and the one on the right shown in offset mode. Figure 13 illustrates the conditions measured 10 minutes after heating was started. A characteristic deformation of the temperature profile is already evident in the offset view. Figure 14 shows the temperature profiles recorded 7 minutes later. The first “uniformity” alarms are present for all “cracks”. After 30 minutes, all “crucible defects” are clearly identifiable and reported by the corresponding alarm messages (Figure 15). It should be mentioned that the small test furnace allowed us to embed only a one-layer sensor cable of limited length, which gives an inferior position resolution. The position resolution will naturally be higher on a larger furnace. However, the unusually good temperature resolution of the measuring method is impressively demonstrated.

365

Figure 12: Temperature profiles obtained before heating was commenced (9:30 a.m.)

Figure 13: Temperature profiles after 10 minutes (9:40 a.m.), showing first evidence of deformations

366

Figure 14: Temperature profiles after 17 minutes (9:47 a.m.), showing the first alarm messages

Figure 15: Temperature profiles after 30 minutes (10:00 a.m.), with all defects duly reported

367

9

Summary

The essential advantages of the Optical Coil Protection (OCP) system can thus be summarized in the following key-words: • Full protection against – Operational breakdown due to coil damage – Bodily injury and equipment damage due to molten metal breakthrough • Recording and visualisation of temperature profile over the entire crucible campaign – Indication of developments and trends of refractory wear or metal penetration – Possibility to take action in good time to extend refractory life • Direct temperature measurement, not resistance-based – Fully operational with a vast range of refractories and immediately after relining • Optical (i.e., non-electrical) measuring method • Eliminates false signals or even sensor grid damage by the magnetic field of the induction furnace • One single evaluator can monitor two furnaces, for example in tandem installations • Very high resolution, e.g. 60 spots over the circumference of an 8-tonne furnace crucible, like the second marks on a clockface • Temperature measurement with a resolution better than 1 K • This distributed optical-fibre temperature measuring method has evolved into a mature system which has been demonstrating its reliability as a central safety system for years in more than 300 installations worldwide.4), 5) In closing, it should be noted that the use of an OCP system is by now means limited to coreless induction furnaces as discussed herein. The system is also suitable for large-area temperature measurements in other industrial furnace applications, e.g., to monitor the inductors on channel-type induction furnaces.

10 [1] [2] [3]

[4] [5]

[6]

References Hopf, M.; Elektrowärme International - Edition B. Industrielle Elektrowärme 50 (1992) No. B2, B229-B23 Hopf, M.; Gießerei 80 (1993) No. 22, pp. 746-751 „Verfahren zur Auswertung optisch rückgestreuter Signale zur Bestimmung eines strekkenabhängigen Meßprofils eines Rückstreumediums“ [Process for evaluating optical backscatter signals to determine a path-dependent measuring profile from a backscatterproducing medium] EP 0692705, 1995, Dr.-Ing. U. Glombitza “Fire Protection Systems for Traffic Tunnels Under Test” Proceedings 12th International Conference on Automatic Fire Detection, Rudolf Maegerle, Siemens Building Technologies Ltd., Cerberus Division Tsakiridou, E.; “Mit Hightech gegen den Tod im Tunnel” [High-technology against tunnel fatalities], VDI-Nachrichten, No.10, p. 3, VDI-Verlag 2001

368

A New Continuous Casting Process H. Sommerhofer, P. Sommerhofer Sommerhofer Technologies, 8010 Graz, Austria

1

Abstract

Continuous casting using water as coolant is state of the art, but still has some disadvantages, one main problem is the building of vapor at low temperatures. In order to prevent these disadvantages, we use liquid metal to cool the billet. Laboratory scale experiments have been done to investigate the possibility to cast an aluminum alloy, a magnesium alloy and copper using a low melting liquid metal as cooling material. Now a pilot scale plant has been constructed by Sommerhofer Technologies. After several tests with the pilot scale plant, results on the castability of aluminum- and magnesium alloys are existing now. Advantages of the new process: Much lower crack risk; constant high heat transfer coefficients; higher casting rates; larger process window for the coolant temperatures; nearly no subsurface layer; no risk of explosions; no contaminated cooling water; exchanged heat at usable temperature level.

2

Introduction

The cooling process in continuous casting is one main influence parameter in continuous casting. Usually the cooling process consists of two or three steps. The first one is the cooling step of the melt in the mould resulting in a thin shell containing melt in it, when the billet is leaving the mould. The second cooling step is direct cooling of the billet with water, this is done in very different ways. On the one hand casting defects should be prevent, on the other hand the casting rate should be as high as possible. One problem is that water as cooling medium is boiling at very low temperatures – so the cooling effect is very strong at this low cooling water temperatures leading to a high temperature gradient between billet core and billet shell and hence to high stresses resulting in centre line cracks. Therefore an upper border for the casting rate for one casting alloy in one dimension is given by the building of centre line cracks. It is possible to weak the direct cooling step by blowing away the cooling water at a defined height under the mould exit or to withdraw the cooling water by additional units but it is not possible to change the cooling character of water, when getting in contact with a hot surface. Now, what happens, when water gets in contact with a hot surface? Looking at figure 1 we can see, the way of the casting melt trough the mould getting solid and the curve of the heat transfer number for a conventional hot top mould, how it is used for continuous casting of aluminum [1]. Looking at this graph we can see, that the heat transfer number may be diminished in the region of the hot top, immediately below the hot top the heat transfer may increase to about 10.000 W/m²K, decreasing a short distance below, where the new formed shell is strong

369

Figure 1: Heat transfer number D for heat transfer from the casting to the cooling environment during the cooling process illustrated how it is mentioned in [1] with figures for the heat transfer number

enough to shrink against the metallostatic pressure of the casting melt inside the shell. This results in an air gap reducing the heat transfer number to values of about 100 W/m²K. In the first region of direct cooling the billet surface temperatures are high, much higher than the boiling point of water. If the surface temperature is higher than the Leidenfrost temperature, a closed steam layer builds around the billet, blocking the heat transfer between coolant and casting. When the surface temperature falls below the Leidenfrost temperature the steam layer breaks partially and the heat transfer number increases to a maximum, where no steam is on the casting surface. This maximum can be at 20.000 W/m²K or higher but also lower, depending on the kind of secondary cooling (film cooling, spray water cooling, quenching or two phase cooling) and its operating parameters. Especially on the possibility to optimize the direct cooling process using water as coolant many investigations have been done and papers may be found in the literature. Figure 2 shows the possibility to change the heat transfer number for spray water cooling by controlling the cooling water impingement density. One can see that the curve of the heat transfer number may be moved upward applying higher cooling water impingement density, but the characteristically strong change of the heat transfer number when coming down from high surface temperatures (a few hundred W/m²K) to low surface temperatures (order of magnitude is about 104 W/m²K) may not be changed. So the cooling rate of the casting is changing very

370

Figure 2: The Heat transfer coefficient as a function of surface temperature and cooling water impingement density V for spray water cooling how it is shown in [2]

strong when the outer part of the casting is solidifying leading to differences in grain size and quality of the billet. So the tendency of water to boil at low temperatures is not good for the quality of the billet, furthermore too high amounts of water steam are influencing the casting process – steam may rise up into the mould. In worst case the use of water as coolant may lead to a water steam explosion or an oxyhydrogen explosion. In order to prevent the disadvantages of water as coolant one has to take another cooling material like a heat exchanger oil, gas, a molten metal or a molten salt.

3

Experiments

Weighing advantages and disadvantages these different cooling materials we decided to take a low melting metal in liquid state as coolant for the continuous casting process. A more detailed discussion about the choice of the cooling material may be found in [3]. First experiments were done with the experimental set up shown in figure 3. This is a very simple set up for investigating the possibility of casting a billet in an insulating mould, cooled in a pool of low melting liquid metal. In these tests three different classes of materials were cast each one representing a special feature of this material class. The materials cast in this first experimental set up were Magnesium (AZ91) representing the class of reactive metals, aluminum AA6063 representing the class of materials with high heat of fusion and pure copper representing the class of material with higher melting point. Experiments using the apparatus in figure 3 have shown the principal feasibility to cool a casting during the continuous casting process in a pool of low melting metal in direct contact with the billet solidifying in an insulating mould. As a consequence of the low convection in the cooling bath the cooling rates were not very high. A possibility to increase the cooling rate is to increase the convective motion in the coolant. Therefore a closed coolant cycle was applied and the coolant was brought into contact with

371

Figure 3: Experimental set up for initial tests on continuous casting with liquid metal cooling device [3]

the billet at a defined velocity and in a defined angle to the billet. This resulted in doubling the maximum casting rate of tests with the first experimental set up. The next step was to optimize the geometry of the coolant distribution unit and the whole mould design as well as testing different mould materials for casting of aluminum alloys. A disadvantage of the first test apparatus was, that the coolant temperature could not be held constant because there was no place for a heat exchanger in the small apparatus we used. One additional disadvantage was the inaccuracy of the billet withdrawal unit. The construction of a pilot scale continuous casting machine was the next step to get conditions near that of industrial work. In figure 4 a schematic depiction is given of the pilot scale plant constructed and installed by Sommerhofer Technologies. In figure 4 the coolant cycle is shown – it is a closed cooling system, which is held inert by an inert gas in order to prevent oxidation of the coolant. The coolant at its operating temperature is pumped into the cooling box, where it takes up heat from the billet. After getting out of the cooling box the coolant runs through the heat exchanger, gives up the dissipation heat and returns to the coolant storage tank at operating temperature. The dissipation heat from the heat exchanger can be used in the plant as process heat because the temperature level is higher than 200°C. Depending on the size of the casting plant this amount of heat can save much money for energy. Figure 4 shows the coolant cycle and the necessary devices for vertical continuous casting, while Figure 5 shows the set up for horizontal continuous casting. It is possible to change the casting direction at our pilot continuous casting machine by turning the roller bed with flying saw from vertical direction into horizontal direction and changing the tundish. So this continuous casting plant and this process in general is flexible concerning casting direction but also concerning different casting materials.

372

Figure 4: Schematic depiction of the pilot scale plant for continuous casting on nonferrous metals using a low melting liquid metal as cooling material

Figure 5: Arrangement of tundish, cooling box, roller bed and flying saw for horizontal continuous casting using a low melting liquid metal as coolant (depiction without cooling cycle and lubrication system)

Furthermore the addition of lubricant or gas in order to reduce the friction between mould and billet is necessary in most cases. Experiments have shown, that some aluminum alloys may be cast in an insulating mould without lubrication but not all. Anyway, casting with cooled mould allows highest casting rates, so the next step was to design a cooled mould. First experiments with a cooled mould have shown that it is impossible to cast billets in cooled moulds without lubrication. In cases where the start up succeeds, big surface cracks are the consequence of no lubrication. So initial tests with manual lubrication short before casting were done for a first determination of the best lubricant. Than the mould was adopted with a continuous lubrication system in order to enable continuous casting.

373

4

Results and Discussion

The results of the first test series may be found in Table 1, one can see that the maximum casting rates for the first experimental set up were not very high, but it could be doubled by increasing the convective motion of the coolant when getting in contact with the billet surface. It was the aim of this initial tests to get an overview about the important process parameters, the advantages and its borders. Answers to most of these questions were gained by doing the experiments with the set up in figure 3. Table 1: Conditions for the first test series and gained maximum casting speed for working with insulating mould and direct cooling in a bath of liquid metal of about 160 °C Liquidus temperature Solidus temperature Casting temperature Billet diameter Max. casting rate

[°C] [°C] [°C] [mm] [cm/min]

AA 6063 655 576 690 50 11

AZ 91 598 468 630 50 7

Pure Copper 1083 1083 1200 30 11

Table 2: Comparison of continuous casting on the pilot scale casting machine with insulating mould to casting with cooled mould

Liquidus temperature Solidus temperature Casting temperature Billet diameter Max. casting rate

[°C] [°C] [°C] [mm] [cm/min]

AlSi12Cu4NiMg Insulating Mould

AlSi12Cu4NiMg Cooled Mould

586 506 600 50 13

586 506 600 50 > 40*)

*) 40 cm/min is not the maximum casting rate for casting this alloy with cooled mould it is higher but not determined until the this paper is written.

Table 2 shows the difference in casting rates when working with insulating mould to casting with cooled mould for AlSi12Cu4NiMg. Although the maximum casting rate for a cooled mould is higher than 40 cm/min the difference is very impressive.

Figure 6: Surface of a AlSi12Cu4NiMg billet cast with cooled mould at 40 cm/min and a billet diameter of 50 mm gained by vertical continuous casting, how it can be used as forging feedstock

374

Figure 7: Macro structure of a AA6082 billet cast with insulating mould at 15 cm/min without grain refiner, billet diameter 50 mm gained by vertical continuous casting, how it can be used as forging feedstock

In Figure 6 the surface quality of a AlSi12Cu4NiMg billet of a diameter of 50 mm is shown, one can see that the region of start up is very short and the surface is smooth after a short distance. So the butt crop for this process is very low in percentage compared to the conventional VDC process. Experiments are showing that the maximum casting rate is higher than that of conventional processes for aluminum alloys and will be much higher for copper and copper alloys due to the high and constant heat transfer number from the billet to the cooling metal. In Figure 7 the uniform grain structure of a AA6082 billet cast with insulating mould without grain refiner is shown. Figure 8 shows the smooth surface of magnesium alloy AZ91, gained by casting with insulating mould without lubrication.

Figure 8: Smooth Surface of a Magnesium AZ91 billet of 50 mm diameter cast with insulating mould at a casting rate of 7 cm/min

375

Figure 9: Comparison of macro structure of pure copper cast with apparatus shown in figure 3 with different casting rates, left picture 7,5 cm/min, right picture: 11 cm/min

Figure 9 compares the macro structure of two copper billets cast in insulating mould and cooled in a pool of liquid metal with two different casting rates. It may be seen that the structure gets very fine at higher casting rate although the casting rate is small compared to that of casting with cooled mould and convective contact with the coolant. So this process is able to improve the grain structure of copper billets without grain refiner.

5

Conclusion

Much research and development has been done in order to gain a continuous casting process for industrial application showing the following advantages: • Constant high heat transfer number o higher casting rate for aluminum alloys and much higher casting rates for copper and copper alloys • Much lower crack risk (Temperature gradient between core and surface much lower) • Uniform grain structure from the core of the billet to the surface • Start up phase very short • Dissipation heat may be used as process heat (T > 200°C) • Cooling water treatment unnecessary • Explosions are impossible (Very important for reactive alloys of Mg or Al-Li) This list of advantages is the reason, why this process is a revolution for continuous casting of non ferrous metals.

376

6 [1] [2]

[3]

References E. K. Jensen, „Mathematical model calculations in level pour DC-casting of aluminium extrusion ingots” in C. J. McMinn, Light Metals, AIME, 1980, p. 631–642 R. U. Jeschar, u. Reiners et.al., „Wärmeübergang in der Sekundärkühlzone von Stranggießanlagen“ in E. Lossack, Stranggießen, DGM-Informationsgesellschaft Verlag, 1986, 91–114 H. Sommerhofer, Ph.D. Thesis, University of Leoben, 2003

Authors

Akshentsev, Yu.N. Anger, G. 87 Anglada, E. 202

265

Bach, Fr.-W. 81, 95 Bainbridge, I.F. 106 Bast, J. 112 Bender, W. 174 Bienvenu, Y. 320 Boender, W. 189 Boller, K.H. 51 Bombach, E. 112 Brandt, R. 174 Brockmann , G.J. 344 Burghardt, A. 189 Chang Hung-Ju, 213 Chao Long-sun, 213 Commandeur, B. 101 Davey, K. 70 Donsbach, F. 353 Drezet, J.-M. 151 Duck-young Hwang, 137 Dundar, M. 23, 87 Dürrschnabel, W. 51 Eberle, R. 226 Eckert, K. 296 Ellendt, N. 249 Emmerich, H. 162 Eskin, D. 271, 277, 283, 290 Frankenberg , R. Friedrich, B. 3

51

Garmendia, I. 202 Geant a, V. 42 Gottstein, G. 309 Grandfield, J.F. 106 Greenberg, B.A. 265 Gremaud, M. 151

Gruber-Pretzler, M. 194, 219 Grün, G.-U. 174 Grundmann , R. 296 Haga, T. 70, 131, 143 Hanada, K. 77 Han-shin Choi, 137 Hatsushikano, K. 77 Hepke, M. 81 Hoff, H. 353 Honsel, C. 209 Hoon Cho, 137 Hwang Weng-seng, 213 Hyung-ho Jo, 137 Jeong, H. 124 Jurgk, M. 162 Katgerman, L. 271, 277, 283, 290 Kayikci, R. 36 Kazantseva, N.V. 265 Keles, O. 23, 87 Khoury, A. 51 Kiersch, J. 290 Kim, G. 124 Kim, M. 124 Koga, N. 70 Kräutlein, C. 3 Król, J. 118 Krone, K. 3 Krug, P. 101 Kudashov, D.V. 256 Kumai, S. 131, 143 Kumar Nadella, R. 277 Lai Yi Lin, 213 Landaberea, A. 202 Lebreton, V. 320 Louhenkilpi, S. 240 Ludwig, A. 194, 219

Mäkinen, M. 234, 240 Matsuzkai, K. 77 Mayer, F. 194, 219 Moiseev, J. 194 Morishita, M. 29 Müller, H.R. 51, 256 Müller, W. 51, 329 Nagaumi, H. 182 Neuer, G. 174 Niedermair, F. 336 Nikrityuk, P.A. 296 Oelmann, H. 51 Olaru, P. 309 Pan Wensen, 213 Park, J. 124 Pedrós, P. 202 Pineau, A. 309 Porten, M. 51 Rappaz, M. 151 Ricken, H. 209 Rode, D. 51 Rossberg, A. 81, 95 Ruvalcaba, D. 290 Sadi, F. 320 Savas, O. 36 Schacht, S. 95 Schliefer, H. 51 Schmitz, W. 353 Schneider, P. 329 Schneider, St. 51 Schwarze , M. 51

Shae K.Kim, 137 Shimizu, T. 77 Siquieri, R. 162 Sommerhofer, H. 368, 368 Specht, E. 118 Stefanoiu, R. 42 Streitenberger, P. 168 Suvanchai , P. 182 Takeda, Y. 182 Tockner, J. 249 Tokuda, K. 29 Tonn, B. 194 Torisaka, Y. 77 Turchin, A.N. 283 Uhlenwinkel, V. 249 Umeda, T. 182 Uoti, M. 234 van Klaveren, E.P. 189 Vapalahti, S. 240 Väyrynen, P. 240 Virtanen, T. 314 Voiculescu, I. 42 Volkov, A.E. 265 Walter, M. 249 Watari, H. 70, 131, 143 Wolber, P. 51 Wu, M. 194, 219 Zauter, R. 256 Zeillinger, H. 336 Zöllner, D. 168

Subject index 3004, surface defects 29

automotive applications, AA6016 87

AA6016, automotive applications 87 AlBiZn alloys, numerical simulation 194 Al-4.5% Cu alloy, melt flow effects 283 alloys – 3004 29 – 6061 277 – aluminum 101, 131, 174, 271, 277, 283, 290 – binary 296 – casting speed 277 – continuous casting 240, 320 – crystallization conditions 265 – eutectic 309 – high strength 182 – hypermonotectic 194 – magnesium 70, 81 – melt flow 283 – molten 42 – numerical simulation 202 – quenching study 290 – roll casting 143 – strip casting 77 – tenary 174 – thermal conductivity 174 aluminum – 6061 277 – casting 151, 344 – conductivity 174 – DC Casting 271 – electromagnetic casting 124 – Fe alloys 309 – high strength alloy 182 – high strength alloys 101 – horizontal direct chill casting 336 – macrostructure 283 – melt treatment 3 – microporosity 36 – nickel alloys 309 – roll casting 143 – solidification 290 – strip casting 77, 131

billet – liquid-metal coding 368 – near net shape 182 – steel 124 – sump characteristics 271 binary metal alloy, solidification 296 binary tin bronze 320 bronzes, dendrite coarsening 314 bubbles, physical characteristics 42 calcium, strip casting 77 caster, universal 336 casting – aluminum 36, 131 – continuous see continuous casting – continuous strip 70 – direct chill see direct chill casting – electromagnetic 124 – high speed 143 – horizontal direct chill 336 – liquid-metal-cooling 368 – magnesium alloys 81 – non ferrous metal 112 – numerical simulation 151, 189 – preworks 3 – semi-continuous 344 – simulation 209 – strip 70, 77 – tin bronzes 314 – tin containing alloys 320 – vertical 202 casting processes, modelling 151 casting speed 277 casting table 344 casting technology, copper 329 chill casting, magnesium 95 columnar solidification, macrosegregations 219 columnar to equiaxed transition 283 conductivity, thermal 174

continuous casting 51, 112, 137, 151, 194, 209, 213, 219, 226 – alloys 194, 202 – copper 240 – dhp-copper 234 – magnesium alloy 70 – metal wires 213 – modelling 151 – new process 368 – simulation 226 – sn-bronze 219 – state of the art 344 – technology 51 – tin bronzes 314 – tin containing alloys 320 cooled mold 137 cooling process, alloys 368 copper – casting 151, 234, 240, 329 – high purity 137 – macrostructure 283 – melt treatment 3 crystallization conditions, alloys 265 DC casting, see direct chill casting defects – casting 271 – formation 271, 277 – surface 23 dendrites – coarsening, tin bronzes 314 – growth 162 dendritic length scale, changes 314 deoxidized high phosphorus copper 234 DHP-copper, see oxidized high phosphorus copper direct chill casting 189, 336 – 3004 29 – aluminum alloy 271, 277 – magnesium 95 electrical potential, binary metal alloy 296 electromagnetic casting 124 embrittlement, tin bronzes 314 equiaxed dendrites 162

eutectic alloys, properties 309 failure properties, eutectic alloys FEM simulation, alloys 182 fluid flow – effects 240 – simulation 226

309

gas, inert 42 gas bubbles, physical characteristics 42 grain growth, monte carlo simulation 168 grain refining 29, 277 – surface defects 23 grain structures, calculation 226 HDC casting, see horizontal direct chill casting HE universal caster 336 heat transfer – copper 240 – simulation 240 high purity copper rod, fabrication 137 high speed roll casting, alloys 143 high strength alloys 101 – Al-Mg-Si 182 horizontal casting, copper 329 horizontal direct chill casting 336 hot top mould design 106 hypermonotectic albizn alloys 194 induction furnaces 353 inert gas, molten metal 42 interacting dendrites 162 liquid metal cooling

368

macrosegregation 271 – Al-Cu alloy 283 – aluminum alloy 277 – modeling 219 macrostructure – Al-Cu alloy 283 – tin-bronze 256 magnesium – chill casting 95 – high strength alloy 182

– roll casting 143 – strip casting 77 magnesium alloys 70 – numerical simulation 202 – strip casting 81 melt flow, Al-Cu alloy 283 melt treatment – aluminium 3 – copper 3 melting temperature , multi-component alloys 174 metal wire 213 metals – binary alloys 296 – molten 42 – non ferrous 112 micro wrought shapes, non ferrous metal 112 microporosity, aluminum casting 36 microstructure – development in aluminum alloys 290 – eutectic alloys 309 – tin-bronze 256 – titanium alloys 265 modelling, continuous casting processes 151, 213 mold – cooled 137 – design 106 – temperature fields 234 monitoring system, optical coil protection 353 monte carlo simulation, grain growth 168 multi-component alloys, conductivity 174 non ferrous metal, continuous casting 112 numerical simulation, alloys 194, 202, 296 OCP, see optical coil protection optical coil protection 353 plastic strains, eutectic alloys 309 plasticity, titanium alloys 265 powder metallurgy, superalloys 249 process conditions, hot top mould 106

process parameters, DC casting 271 pulse electric discharging 296 quality of water 118 quenching studies, alloys 118, 290 refining, grains 277 rings, superalloy 249 rod fabrication, copper 137 rods, modelling 213 roll casting 131, – AA6016 87 – high speed 143 – horizontal 70 – twin 77 segregation, minimization 256 semi-continuous casting 344 shapes, micro wrought 112 silicone, high strength alloy 182 simulations – alloys 182, 194, 296 – continuous casting 151, 202, 209, 213, 226 – dendrite growth 162 – grain growth 168 slit mold 124 Sn-bronze, continuous casting 219 solidification – aluminum alloys 290 – simulation 226, 240 – binary metal alloys 296 – unidirectional structure 137 spray forming – aluminum alloys 101 – superalloy rings 249 – tin-bronze 256 steel, billet 124 strain distribution, eutectic alloys 309 strip casting – alloy 77 – aluminium alloy 131 – continuous 70 – magnesium 81 superalloy rings, post processing 249

surface – defects 23, 29 – roughness 118

titanium alloys, crystallization conditions 265 twin roll caster, unequal diameter 131

temperature fields, mould 234 ternary alloys, conductivity 174 ternary tin bronze 320 thermo-mechanical model, casting 189 tin bronzes – dendrite coarsening 314 – structure 256 tin containing alloys, casting 320

UDC, see upward direct chill unidirectional solidification structure upward casting – continous 202 – direct chill 95 water, quality 118 wincast-conti 209

137