Continuous Casting Edited by K. Ehrke and W. Schneider
Deutsche Gesellschaft für Materialkunde e.V.
Weinheim · New York · Chichester Brisbane · Singapore · Toronto
Dipl.-Ing. Kurt Ehrke ALUMINIUM Essen GmbH Sulterkamp 71 D-45356 Essen Germany
Prof. Dr. Wolfgang Schneider VAW Aluminium AG Forschung und Entwicklung Georg-von-Boeselager-Str. 25 D-53117 Bonn Germany
International Congress Continuous Casting held from 1315 November 2000 in Frankfurt /Main Organizer: DGM · Deutsche Gesellschaft für Materialkunde e.V. Program Committee Dipl.-Ing. Kurt Ehrke, ALUMINIUM Essen GmbH (Chairman) Dr. Hilmar R. Müller, Wieland Werke AG, Ulm Prof. Dr. Wolfgang Schneider, VAW Aluminium AG, Bonn Dipl.-Ing. Gunnar Halvorsen, Elkem Aluminium AG, Oslo (N) Dr. Dirk Rode, KM Europametal AG, Osnabrück
This book was carefully produced. Nevertheless, authors, editors and publisher do not warrant the information contained therein 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.
Cover photo: Lüko GmbH Casting Unit for Extrusion Billets VAW aluminium AG AIRSOL VEIL® technology
Library of Congress Card No. applied for A catalogue record for this book is available from the British Library Deutsche Bibliothek Cataloguing-in-Publication Data A catalogue record for this publication is available from Die Deutsche Bibliothek ISBN 3-527-30283-2 © WILEY-VCH Verlag GmbH, D-69469 Weinheim (Federal Republic of Germany), 2000 Printed on acid-free paper. 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: WGV Verlagsdienstleistungen GmbH, Weinheim Printing: betz-druck, Darmstadt Bookbinding: Schaumann, Darmstadt Printed in the Federal Republic of Germany
Preface The aim of the conference, organized by the DGM Continuous Casting Committeee, is to highlight the importance of continuous casting of aluminium, copper and magnesium to the international fabricating industry. The conference lectures, generated by the Call for Papers, cover technological advances in all sectors which are important for the manufacture of high quality continuously cast products. Besides melt treatment, casting processes and structure of continuously cast ingots, modelling of casting will be a major topic of the conference. Numerical modelling becomes more and more dominant as a research tool to improve casting processes and the resulting products. The advantages are reduced development times and development costs. The programme of the symposium reflects this with numerous papers dealing with modelling of nucleation, heat and fluid flow as well as stresses and structure. Another new approach of the conference are the supplier sessions. The organizing committee hopes that the conference programme encourages specialists of the non-ferrous industry worldwide to take part in this meeting. K. Ehrke Chairman of the Conference
Contents 25 Years of DGM Continuous Casting Research W. Schneider, VAW aluminium AG, Research and Development, Bonn (D)..............................1 Melt Treatment Hydrogen in Aluminum Containing Copper Alloy Melts – Solubility Measurement and Removal K. Neumann, B. Friedrich, K. Krone, IME Process Metallurgy and Metal Recycling, RWTH Aachen (D) J. Jestrabek, E. Nosch, Schwermetall Halbzeugwerk GmbH, Stolberg (D) .............................15 Fundamental Research About Liquid Metal Filtration B. Hübschen, J. Krüger, RWTH Aachen (D) N. Keegan, Pyrotek Engineering Materials Limited, Dudley (GB) W. Schneider, VAW Aluminium AG, Bonn (D) .........................................................................20 Impact of Grain Refiner Addition on Ceramic Foam Filter Performance N. Towsey, W. Schneider, H.-P. Krug, VAW aluminium AG, Research and Development, Bonn (D) A. Hardman , London & Scandinavian Metallurgical Co. Limited, Rotherham, South Yorkshire (GB) N. Keegan, Pyrotek Engineering Materials Ltd., Netherton, West Midlands (GB)..................26 Review of Dissolution Testing and Alloying Methods in the Casthouse G. Borge, Bostlan, S.A., Larragane, Mungia (E) P. Cooper, S. Thistlethwaite, LSM Co.Ltd., Rotherham (GB) ..................................................33 The Effect of Casting Parameters on the Metallurgical Quality of Twin Roll Cast Strip Y. Birol, Marmara Research Center, Gebze-Kocaeli (TR) G. Kara, A. Soner Akkurt, ASSAN Aluminum Works, Istanbul (TR) C. Romanowski, FATA Hunter Inc., Riverside (USA) ..............................................................40 Casting Technology and Processes – Aluminium Influence of Different Lubricants on the Friction between the Solidifying Shell and the Mould during the DC Casting of AlMgSi0.5 F. Dörnenburg, VAW aluminium AG, Research and Development, Bonn (D) S. Engler, Foundry Institute, Aachen University of Technology, Aachen (D)..........................47 Prediction of Boundary Conditions and Hot Spots during the Start-up Phase of an Extrusion Ingot Casting S. Benum, Hydro Aluminium R&D Materials Technology, Sunndalsøra (N) D. Mortensen, H. Fjær, Institute for Energy Technology, Kjeller (N) .....................................54 Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
VIII Improved Metal Distribution during DC-casting of Aluminum Alloy Sheet Ingots P. Tøndel, J. Hayes, I. Thorvaldsen, Elkem Aluminium, Mosjøen (N) G. Grealy, G. Tahitu, Corus Research Development & Technology, Ijmuiden (NL) E. Jensen, Elkem Research, Kristiansand (N) D. Brandner, Corus Aluminium Walzprodukte GmbH, Koblenz (D) .......................................61 Determination of Material Properties and Thermal Boundary Condition from Casting Trial on Alloy AA7075 J. Rabenberg, J. Storm, Corus Research Development & Technology, Ijmuiden (NL) I. Opstelten, Corus Research Development & Technology, now with TNO Bouw, Delft (NL) J.-M. Drezet, École Polytechnique Fédérale de Lausanne, Lausanne (CH)............................71 Single-roll Strip Casting of Aluminium Alloys E. Straatsma, W. Kool , L. Katgerman, Delft University of Technology, Laboratory for Materials, Delft (NL) ................................................................................................................77 Continuous Casting of Semisolid Al-Si-Mg Alloy T. Motegi, F. Tanabe, Chiba Institute of Technology...............................................................82 Casting Technology and Processes Yield and Quality Improvements for Semi-Continuously Cast Copper Alloys C.-M. Raihle, Outokumpu Process Automation, Västerås (S) ..................................................89 Continuous Casting Technology for Magnesium U. Holzkamp, H. Haferkamp, M. Niemeyer, University of Hanover, Institute of Materials Science, Hanover (D)................................................................................................................94 Local Distribution of the Heat Transfer in Water Spray Quenching F. Puschmann, E. Specht, J. Schmidt, University of Magdeburg, Institute of Fluid Dynamics and Thermodynamics, Magdeburg (D) ........................................................101 Grain Structure, Microstructure and Texture of Copper Ingots Produced during the Continuous Casting Process V. Plochikhine, V. Karkhin, H. Bergmann, Department of Metallic Materials, University of Bayreuth (D)......................................................................................................109 Technologies and Installation for Electrochemical Hardening of Wear Surfaces R. Boiciuc, V. Munteanu, G. Petrache, Uzinsider Engineering S.A., Galati (ROM).............115
IX Modelling – Heat and Fluid Flow; Nucleation Numerical Mass and Heat Flow Predictions in Aluminum DC Casting: A Comparison of Simulations with Melt Pool Measurements A. Buchholz, Corus Research Development &Technology, Ijmuiden (NL) B. Commet, Pechiney Centre de Recherches de Voreppe (F) G.-U. Grün, VAW aluminium AG, Research and Development, Bonn (D) D. Mortensen, Institute for Energy Technology, Kjeller (N) ..................................................123 Investigations of the Primary Cooling in Sheet Ingot Casting H. Fjær, D. Mortensen, Institute for Energy Technology, Kjeller (N) A. Buchholz, Corus Research Development & Technology, Ijmuiden (NL) B. Commet, Pechiney Centre de Recherches de Voreppe (F) J.-M. Drezet, Laboratoire de Métallurgie Physique, EPFL, Lausanne (CH) ........................131 Boiling Curve Approach for Thermal Boundary Conditions in DC Casting J. Zuidema jr., L. Katgerman, Netherlands Institute for Metals Research, Laboratory for Materials, Advanced Materials and Solidification Technology, Delft (NL) I. Opstelten, Corus Research Development & Technology, Ijmuiden (NL)............................138 Theoretical and Experimental Study of Vertical Continuous Casting of Copper M. Uoti, Helsinki University of Technology, Laboratory of Metallurgy, Helsinki (SF) M. Immonen, K. Härkki, Outokumpu Poricopper OY, Kuparitie, Pori (SF)..........................143 Modeling of Grain Refinement in Aluminum Alloys A. Greer, A. Tronche, University of Cambridge (GB) ............................................................149 Modeling of the Grain Refinement in Directionally Solidified Al -4.15 wt.% Mg Alloys using Cellular Automaton – Finite Element Approach M. Vandyoussefi , A. Greer, University of Cambridge (GB) ..................................................154 Modelling – Stress and Structure ContiSim™: Process and Material Modelling of Continuous Casting in Macro and Micro Scale J. Boehmer, Process Modelling and Informatics, Betzdorf/Sieg (D)......................................163 Crystal Growth Morphology during Continuous Casting C. Caesar, Munich (D)............................................................................................................169 3D-Modeling of Ingot Geometry Development of DC-Cast Aluminum Ingots during the Start-Up Phase W. Droste, G.-U. Grün, W. Schneider, VAW aluminium AG, Research and Development, Bonn (D) J.-M. Drezet, CALCOM SA, Parc Scientifique, Lausanne (CH) ............................................175
X The Influence of Casting Practice on Stresses and Strains in 6xxx Billets – A Statistical and Modelling Study B. Henriksen, S. Braathen, E. Jensen, Elkem ASA Reasearch, Kristiansand (N)...................184 Modelling of Macrosegregation in Continuous Casting of Aluminium T. Jalanti, M. Rappaz, École Polytechnique Fédérale de Lausanne, Laboratoire de Métallurgie Physique, Lausanne (CH) M. Swierkosz, M. Gremaud, Calcom SA, Lausanne (CH) ......................................................191 The Effect of the Differencing Scheme on the Numerical Diffusion in the Simulation of Macrosegregation B. Venneker, Netherlands Institute for Metals Research, Delft (NL) L. Katgerman, Delft University of Technology, Laboratory for Materials, Advanced Materials and Solidification Technology, Delft (NL) ............................................199 Application of a New Hot Tearing Analysis to Horizontal Direct Chill Cast Magnesium Alloy AZ91 J. Grandfield, Cooperative Reasearch Centre for Cast Metals Manufacturing, CSIRO Manufacturing Science & Technology, The University of Queensland (AUS) C. Davidson, J. Taylor, Department of Mining, Minerals and Materials Engineering, The University of Queensland (AUS)......................................................................................205 Micro- and Macrostructures Nucleation Studies of Grain Refiner Particles in Al-Alloys P. Schumacher, University of Oxford (GB) ............................................................................213 Effect of Solute Elements on the Grain Structures of Al-Ti-B and Al-Ti-C Grain-Refined Al Alloys A. Tronche, University of Cambridge (GB) A. Greer, University of Cambridge (GB) and Péchiney Centre de Recherches de Voreppe (F) .......................................................................................................................218 Grain Refinement Process in Aluminium Alloys Type AlZnMgZr T. Stuczyn´ski, M. Lech-Grega, Institute of Non-Ferrous Metals, Light-Metals Division, Skawina (PL) ..........................................................................................................224 Coupled Influence of Convection and Grain-refining on Macrosegregation of 1D Upwardly Solidified Al 4.5% Cu P. Jarry, Pechiney Centre de Recherches de Voreppe (F) H. Combeau, G. Lesoult, LSG2M, Ecole des Mines de Nancy (F).........................................233 Tensile Behaviour of DC-cast AA5182 in Solid and Semi-solid State W. van Haaften, W. Kool, L. Katgerman, Laboratory of Materials, Delft University of Technology, Delft (NL) ...........................................................................................................239
XI The Columnar to Equiaxed Transition in Horizontal Direct Chill Cast Magnesium Alloy AZ91 J. Grandfield, Cooperative Reasearch Centre for Cast Metals Manufacturing, CSIRO Manufacturing Science & Technologiy, The University of Queensland (AUS) C. Davidson, J. Taylor, Department of Mining, Minerals and Materials Engineering, The University of Queensland (AUS)......................................................................................245 Study of Heterogeneous Nucleation of α-Al on Grain Refiner Particles during Rapid Solidification P. Cizek, B. McKay, P. Schumacher, University of Oxford (GB)...........................................251 Effect of Instability of TiC Particles on the Grain-Refining Behavior of Al-Ti-C Inoculants in Aluminum Alloys M. Vandyoussefi, A. Greer, University of Cambridge (GB) ...................................................257 Grain Refiners for Thin Strip Twin Roll Casting R. Cook, London & Scandinavian Metallurgical Co. Limited, Rotherham (GB)...................263 Characterisation and Optimisation of Thixoforming Feedstock Material S. Engler, Gießereiinstitut, RWTH Aachen (D) D. Hartmann, EFU Gesellschaft für Ur-/Umformtechnik mbH, Simmerath (D) I. Niedick, Volkswagen AG, Braunschweig (D), (former EFU GmbH)..................................269 Experimental Study of Linear Shrinkage during Solidification of Binary and Commercial Aluminum Alloys D. Eskine , L. Katgerman, Netherlands Institute for Metals Research, Delft (NL) ................276 The Influence of the Cooling Rate on the Type of the Intermetallic Phases in the Aluminium Alloys of the 3XXX (AlMnMgSi) Group T. Stuczyn´ski, M. Lech-Grega, Institute of Non-Ferrous Metals, Light Metals Division, Skawina (PL) ..........................................................................................................................282 Suppliers Session – Aluminium Horizontal Direct Chilled (HDC) Casting Technology for Aluminium F. Niedermair, Hertwich Engineering GmbH, Braunau (A) ..................................................293 Automatic “Bleed Out” Detection and Plug Off in VDC Billet Casting M. Lück, Wagstaff Inc., Spokane WA (USA)...........................................................................300 The AIRSOL VEIL® Technology Package for Aluminium Billet Casting G. Bulian, M. Langen, VAW aluminium AG, Bonn (D) ..........................................................302 The Manufacturing, Design and Use of Combo Bag Distributors in Sheet Ingot Casting S. Tremblay, Pyrotek High-temperature Industrial Products Inc., Chicoutimi (CAN) R. Green, Pyrotek Engineering Materials Ltd., Netherton (GB) ............................................310
XII Recent Quality and Efficiency Improvements Through Advances in In-Line Refining Technology V. Dopp, S. Simmons, Pyrotek, Inc., Tarrytown, New York (USA).........................................316 Suppliers Session – Copper Horizontal Continuous Casting of Copper Alloy Billets M. Brey, SMS Meer GmbH, Demag Technica, Veitshöchheim (D)........................................325 The Outokumpu UPCAST® System L. Eklin, Outokumpu Castform Oy, Pori (SF) ........................................................................333 Author Index .........................................................................................................................341 Subject Index.........................................................................................................................343
25 Years of DGM Continuous Casting Research Wolfgang Schneider VAW aluminium AG, Research and Development, 53117 Bonn, Germany
1
Introduction
Continuous casting experts from the non-ferrous metal industry in Germany, Netherlands, Austria, Norway and Switzerland have taken part in the Continuous Casting Committee of the Deutsche Gesellschaft für Materialkunde DGM since its foundation 1972. During the formation of the committee the different technologies used in DC. casting resp. continuous casting of semis feedstock were taken into consideration. Working groups were therefore formed to deal with the vertical DC. casting of Al and Cu as well as the horizontal casting of Al, Cu and Zn. Over the years, goals were redefined and changes were made to the organisation structure when the industrial requirements made this necessary. In former times the vertical DC casting experts were associated all together in one working group, but due to the Copper DC casting experts being underrepresented, the priority of the work centred on Aluminium DC casting In 1990 an own working group specifically for the DC. casting of Cu was founded. A further important change was the foundation of the Working Group Spray Forming in 1993. It must be mentioned however that not only an expansion of the activities of the Committee took place in the past but also activities had to be stopped and working groups had to be disbanded. Among these, for example, is the working group which dealt many years with the strip casting of Al. The closing of the last strip casting unit in Germany led to this step because the working group saw little reason for a continuation of its work. The actual organisation structure of the Committee Continuous Casting can be seen in Figure 1. The current existing working groups are to be seen in this figure. The activities carried out in these working groups will however not be considered here. One working group will be described in more detail which was not mentioned previously but was an important part of the Continuous Casting Committee for many years.. This is the Working Group Research, as seen in Figure 1. In the following beside the structure and the tasks of this working group the research activities of the past will be described. By means of exemplary results, projects which have been carried out within the framework of the activities of the Working Group Research will be presented. Finally, the future outlook of the working group will be given.
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
2 DGM Committee Continuous Casting
Working Group
Working Group
Working Group
Vertical Casting D.C.
Vertical Casting D.C.
Horizontal Casting C.C.
Aluminium
Copper
Copper / Zinc
Working Group Spray Forming
Working Group Research
Figure 1: Structure of the DGM Continuous Casting Committee
2
Research Working Group: Objectives, Tasks and Structure
The foundation of the Working Group Research took place in the year 1975. Since then, the working group is in nearly unchanged structure part of the Committee Continuous Casting. A major objective of the foundation of the Working Group Research was to interest the universities in continuous casting research and to get it established there over the long term by means of the concerted execution of research projects. This has been successful for nearly two decades. The working group was, and is still today, made in such a way that, besides the members from universities, also members of the different working groups of the Committee also take part. These are the heads of the working groups whose task it is among others, to make suggestions about research projects discussed before in their working groups. The project ideas are discussed with the university representatives concerned and a research proposal is prepared for the submission to relevant funding organisations. The working group normally meets once a year. In these meetings the progress of the actual projects is reported and new project suggestions are discussed. In the beginning the working group had to ensure that adequate equipment was available at the universities for the continuous casting research and it should give support in the procurement of funding for agreed research projects. With regard to the latter it was and will be important in the future that the industry partners were and are concerned in the preparation of the research proposals. At this point, however, it has to be mentioned that the execution and the care for research projects were not alone reserved to the Working Group Research but that there have been carried out also research projects in the other working groups of the Committee. Another important aspect of the work of the working group in the past was, to awaken the interest of the university students in continuous casting. Numerous graduates from universities who carried out the research work within the scope of their doctoral thesis subsequently took up a position in the continuous casting industry on finishing their work. In the beginning the working group´s activities focused on investigations dealing with material science of DC casting. Over the years the activities were expanded with investigations into melt quality and process technology. Correspondingly the university institutes as members of the working group expanded too. In Figure 2 the structure of the
3 Working Group Research is shown and the university institutes are named which were resp. still are long-term members of the working group. Special mention has to be made here of the Gießerei-Institut of the RWTH Aachen where under the management of Professor Engler and with the assistance of the Working Group Research in the mid of the Seventies the first DC casting unit at a university was installed. Unfortunately the unit was dismantled recently and thus is no longer available for future research work. In Figure 3 the number of research projects completed at the universities since 1979 is shown. It can be seen that the number of projects has decreased since 1995, at present no research projects are being carried out within the scope of the working group. The reasons for this will be discussed later. Working Group
Research
Foundry Institute
Heads of Committee Working Groups
RWTH Aachen Prof. S. Engler
Inst. of Non Ferrous Metallurgy RWTH Aachen Prof. J. Krüger
Inst. of Metal Science TU Berlin Prof. W. Reif
Inst. of Energy Technology TU Clausthal Prof. R. Jeschar
Inst. of Metal Research
MPJ Stuttgart Prof. B. Predel, Dr. E. Fromm
Inst. of Material Sciene TU München Prof. H. M. Tensi
Number of Projects
Figure 2: Structure of the research working group
8
Working Group RESEARCH
7
DGM Continuous Casting Committee
6 5 4 3 2 1
1975
1977
1979
1981
1983
1985
1987
1989
1991
1993
1995
Year
Figure 3: Number of completed research projects per year since 1975
1997
1999
4
3
Research Working Group: Research Priorities
The list of past research projects of the Working Group Research can be divided into the following subjects: • Melt Quality • Ingot Solidification • Process Technology In the following, an overview of the research projects carried out in these subjects will be given and exemplary results from chosen works presented. Naturally a selection must be made at this point since not all projects can be dealt with. Of course this does not mean that the projects not considered here made no important contribution to the understanding of DC casting. 3.1
Melt Quality (MPI Stuttgart, RWTH Aachen)
As is well known non-ferrous metal melts show complex interactions with the surrounding atmosphere. The centre of interest of the research work carried out was therefore on diluted hydrogen in Al and Cu melts and its measuring. Moreover the oxidation of the liquid aluminium as well as the properties of oxide layers on the melt were investigated. This knowledge is of importance to the handling of melts in the cast house. In further research work the thermodynamic fundamentals of Li removal from primary aluminium as well as investigations into the removal of Na and Li from Al melts were carried out. Within the scope of the research work on the hydrogen solubility in Al and Cu melts, a measuring method was developed which allows the continuous measurement of the hydrogen content by the means of H2 partial pressure measurement. The measurement principles of this method are shown schematically in Figure 4. In this method the probe consists of a graphite disc which is connected to a pressure meter by a gas tight ceramic pipe. After immersing the probe in the melt the probe is quickly evacuated allowing the hydrogen to diffuse into the probe until the pressure in the probe and the hydrogen partial pressure in the melt are equal. The actual hydrogen content is then calculated with simultaneous temperature measurement with the help of Sievert´s equation. An essential advantage of this method in comparison to other methods is that by dosing of the probe with hydrogen the measurement time can be shortened substantially, so that a continuous hydrogen measurement is practicable. The above described method is commercially available today as measurement equipment under the trademark CHAPEL. Within the scope of the investigations on the oxidation behaviour of Al melts the strength of oxide layers as well as the initial stage of the oxidation were investigated. For this a measuring method for the measurement of the strength of the oxide layer was developed, which is shown in a schematic depiction in Figure 5. With this measuring method the maximum torque is determined that the oxide layer can put up a rotating circular stamp in a fixed ring before it tears. In extensive tests the strength of oxide layers on pure aluminium melts as well as on AlMg and AlSi melts with different Mg and Si contents was measured. It could be seen that with increasing age of the oxide layer and with increasing temperature the strength of the oxide layer increases significantly. Moreover, it could be measured, among other things that an enrichment of defined elements increases the strength of the oxide layer. To these elements belong for example Mg, Ca, Na and Li. Their effect decreases with the holding time of the melt.
5 vacuum pump data storage
computer measuring probe
valve 2
data acquisition pressure measuring instrument
temp. measuring device
valve 1 dosing valve
H2
ceramic pipe
graphite disc
thermocouple aluminium melt
Figure 4: Principles of the CHAPEL hydrogen measurement method
The investigation of the initial state of the oxidation of Al melts was carried out at pressures below 10-6 bar. In these investigations it was determined that after skimming of the melt surface, the re-covering time until the formation of a protecting oxide layer with a thickness of about 1 µm was in the time scale of 10-3 to 10-2 sec.
Figure 5: Apparatus for measuring of the strength of oxide skins
3.2
Ingot Solidification (RWTH Aachen, TU Berlin, TU München)
The as-cast structure of the DC cast ingots has as well known considerable influence on their processing and the quality of the subsequently produced semis. DC cast ingots show structural defects which are related to the casting and solidification process. The research
6 work carried out was therefore aimed at investigating the solidification processes in the DC casting of non-ferrous metals. The main focus of interest was, among others the solidification of the sub surface of the DC cast ingots. In several research projects the formation of surface segregations and the stability of solidifying shell zones were investigated. A further priority was the melt treatment for the achievement of a fine equiaxed as-cast grain structure. Here, basic investigations into the understanding of the grain refining mechanism of Al with Titanium and Boron were carried out, new effective grain refiners for Al and Cu alloys were also developed. In additional projects the starting phase of the DC casting of Al rolling ingots and the influence of fluid flow on the as-cast structure in vertical DC casting of Al rolling ingots and of CuSn alloy billets were investigated. In the following, extraordinary results from some of the above mentioned investigations will be presented. The investigations on surface segregation of DC cast ingots with specially developed test moulds showed that the essential transport mechanism for the formation of segregations is the metallostatic pressure of the melt in front of the shell zone, i.e. this is main responsible for the structure defect surface segregation. Beside this, it could also be seen that with increasing thickness and finer as cast structure of the shell zone a significant decrease in the degree of the surface segregations can be achieved. water controller
electronic controller testing procedure lifting system chill
extensometer
shell load-cell
pneumatic closing system
hydraulic cylinder
hydraulic pump
Figure 6: Test equipment for measuring the mechanical properties of solidifying shell zones
In series of interesting investigations the stability of shell zones of Al and Cu alloys solidifying in the mould were investigated. For these investigations a test equipment was developed so that by tensile test the mechanical properties of a solidifying shell zone in dependency of the solid portion in the solidification range could be measured. The developed measuring equipment is schematically shown in Figure 6. Essential items of this equipment comprise the water-cooled copper mould: from this mould the shell zone grows into the melt which is located in an insulating box. A measuring device projects through the insulating form into the hollow space and into the cast shell. This device is connected to a tensile measuring machine. The most important results of these extensive investigations can be summarised as follows. The mechanical properties of solidifying shell zones are on a low level as shown by the examples of an Al alloy and Cu alloy respectively in Figure 7. With increasing solid portion of the shell zone during its duration in the air gap region the tensile strength increases and the elongation of fracture decreases. This is seen in Figure 8.
7 CuSn30
2,5
AlCu4
2,5
547 - 551 °C 850 - 929 °C
2,0
Tensile Strength in N/mm2
2
Tensile Strength in N/mm
2,0
882 - 927 °C 845 - 937 °C 879 - 940 °C
1,5
888 - 937 °C 860 - 938 °C
1,0
Temperature in Shell Zone
854 - 942 °C 870 - 940 °C
0,5
Temperature in Shell Zone
1,5 576 - 583 °C
1,0
581 - 591 °C
597 - 603 °C
0,5
881 - 932 °C
608 - 616 °C
905 - 937 °C
0
0
0,5
1,0
608 - 614 °C
0
1,5
623 - 629 °C
0
0,2
0,4
Elongation in %
0,6
0,8
1,0
Elongation in %
Figure 7: Mechanical properties of solidifying shell zones
99
2,5
AlMgSi0,5 94 91
0,5
2,0
0,4
1,5
0,3
1,0
0,2
0,5
0,1
0
0
600
610 620 Temperature in °C
630
Elongation in %
Tensile Strength in N/mm
2
Tensile Strength Elongation
0
Figure 8: Mechanical properties of shell zones in solidification range
For defined alloys a grain refinement can lead to an increase in the elongation of the solidifying shell zone. As already mentioned the grain refinement is an important melt treatment measure for achieving a fine equiaxed as-cast structure of the DC cast ingots. Among others the deformation behaviour of the billets is improved by an equiaxed grain structure. For the DC casting people, however, it is of particular importance that with the grain refinement the crack formation can be avoided. For Al alloys with the TiB2 containing AlTiB master alloys effective grain refiners are available. Projects carried out in co-operation with the Working
8 Group Research made a substantial contribution to the understanding of the grain refining mechanism of Al with TiB2. For the first time, reports were made on the importance and differing effect of the different alloying elements. Moreover TiB2 could be proved in the grain centre as well as Ti enrichments could be analysed at the borides for the first time. In addition a new grain refiner for Al was developed on the basis of TiC. The same holds true for the grain refinement of Cu alloys. Here the working mechanism of Zirconium as grain refiner could be clarified and the achieved knowledge was used for the development of a Zirconium containing grain refiner master alloy. The effectiveness of the newly developed grain refiners is shown in exemplary manner in Figure 9. T= 700°C th = 5min
700
Average Grain Diameter in µm
Average GrainDiameter in µm
CuSn - Alloy
Al99,7
800
AlTi5B0,2
600 AlTi5B1
500 400 300 200 100
AlTi5C0,25
0
1
2
3
4
E - Cu + m%Zr GT: 1474 K CuSn4 + m%Zr GT: 1523 K CuSn8 + m%Zr GT: 1473 K
1600
1200 Holding Time 1,5min
800
400
5
Grain Refiner Addition Rate in kg/t
0
0
0,02 0,04 0,06 0,08 0,10 Zr Concentration in m%
Figure 9: Grain refinement of aluminium and copper alloys
It has been mentioned several times that within the scope of the research projects carried out to some extent new test techniques were developed to facilitate work on the set task. To these belongs a method which has been developed for the measurement of the butt curl of Al rolling ingots during the start up phase. In this method steel wires are cast into the narrow sides of the ingot butt. These are connected with at the starter block mounted inductive linear transducers so that the progress of the butt curl can be recorded during the start-up phase. With the recorded data it is then possible to evaluate the curling speed and to determine the maximum butt curl. The above described measuring method proved to be very effective and reliable and is nowadays used for butt curl measurements world wide. 3.3
Process Technology (TU Clausthal, RWTH Aachen)
Besides the metallurgical aspects of continuous casting the casting technology is also of importance for the quality of the cast products. Due to the different products different technologies in vertical and horizontal casting are used. In the centre of the research activities of the Working Group Research concerning the process technology were the direct cooling in DC casting and the processes occurring in the mould used for the horizontal DC casting of Cu alloys. Additionally a project was carried out to investigate the working mechanism of lubrication in DC casting of A alloys. Direct cooling in DC casting by impinging of water onto the hot ingot surface is a complex process. According to the metal being cast different cooling techniques are used. In the scope of long-term research the different cooling techniques and cooling conditions were simulated in laboratory-scale tests and, on the basis of the subsequent test results, calculation models
9 were elaborated. In the projects carried out the influence of the most important parameters of spray water cooling as used in the DC casting of Copper and the film cooling as is used in the DC. casting of Aluminium were investigated. The composition of the cooling water also formed part of the investigation. During the cooling of hot metal surfaces with liquids, different stages of cooling are gone through. These are obvious in Figure 10. Of these the stable film boiling stage is unwelcome because it leads to a reduction in the heat extraction and as a result to a slower overall solidification of the ingot. In the investigations on spray and film cooling stable film boiling and the Leidenfrost temperature were of particular interest. The Leidenfrost temperature is the temperature of a hot surface at which the evaporation time of the liquid in contact with the metal surface is the longest, and the extracted heat amount per time interval in dependence on the surface temperature is at a minimum. The investigations into spray cooling showed that the heat transfer in the stable film boiling stage can only be influenced by the water flow rate. The Leidenfrost temperature is moved to higher temperatures with increasing water flow rate while the heat transfer below the Leidenfrost temperature additionally depends on the thermophysical material data and the surface roughness of the ingot to be cooled. The investigations into film cooling showed similar results as those reported on above. It could be measured too that, with increasing exit speed of the cooling water at the gate of a slotted nozzle, the heat extraction is increased during the stable film boiling stage and the vapour film thickness is decreased. That means the vapour film breaks down earlier and the Leidenfrost temperature is moved to higher values. At this point it must be mentioned that stable film boiling only takes place in DC casting of steel while in DC casting of Cu the partial film boiling stage is also of importance. In DC casting of Al, however, the partial or stable film boiling can only be watched during the start-up phase, especially of rolling ingots, while in the stationary casting phase nucleate boiling is the dominant cooling mechanism because of the actual surface temperatures.
q in W/m
2
10
B
6
C
10 5
1. free convection 2. nucleate boiling 3. partial film boiling 4. stable film boiling
A 10
4
2
1 10
3
1
10
3
1
w
4 10
sa
2
10
3
in K
Figure 10: Heat flux density as a function of temperature difference between surface and water
In continuing investigations into the influence of the cooling water composition on the cooling of hot surfaces it could be detected that the influence of the salt content (e.g. NaCl, KCl) is interconnected with the gas content (O2, CO2) of the water. Due to the results of tests and theoretical considerations it could be concluded that with decreasing salt content the gas content increases. This results in a more stable vapour film and leads to a decrease of the
10 Leidenfrost temperature. Because of the aforementioned results a continuation of the investigations into the influence of lubricants in cooling water seems to be necessary. Finally, investigations into the horizontal DC casting of Cu alloys will be considered shortly. Within the scope of these investigations the solidification behaviour of different Cu alloys was examined. By means of temperature measurements in the mould the influence of different casting parameters on the heat transfer conditions was determined. Furthermore, the influence of the withdrawal parameters of the ingot on the friction mechanisms acting between shell zone and mould was determined by making force measurements. On the basis of the achieved knowledge, solidification models have been developed with which the as-cast structure quality of horizontal DC cast Cu ingots can be optimised.
4
Research Working Group: Outlook
The working Group Research is currently in a difficult situation. In Figure 3 it has been shown how the number of research projects has dropped significantly in the last years. For three years no more publicly funded projects have been carried out at the member universities under the umbrella of the working group. The few activities which are still running are bilateral projects with the industry and therefore not accessible to the public. The decrease in number of the projects is mainly related to the fact that it is more and more difficult to find project financing: in comparison to former times, significant reduced financial resources are available for the funding organisations. This has the consequence that markedly fewer research projects can be funded or longer waiting periods have to be accepted before the start of funding of the approved projects. This often has the consequence that applications are either withdrawn or not made at all because years can pass before work can begin. A further problem seems to be that the importance of continuous casting is no longer recognised. This means that suitable projects are not given the necessary priority and therefor are very often not considered for public funding. To overcome the aforementioned problems it was tried some time ago to found a private funding organisation for DC casting research with industry members. The aim was to fund research projects at the universities with the membership fees. Unfortunately the foundation of the society did not take place because the interest of the industry was not sufficient enough to get the necessary seven foundation members together. The difficult financial situation concerning project funding also decreases the interest of the universities in continuous casting research. A further problem is that a change of university members due to age considerations takes place frequently and it is hard to interest new members in continuous casting research under the circumstances described above. For a survival of the continuous casting research and of the work of the Working Group Research of the Committee Continuous Casting new efforts are required. One possible solution could be a further expansion of the Working Group Research at European level. By further internationalisation of the working group, additional funding programs for the financing of research projects could be claimed. Furthermore the importance of continuous casting would be promoted more and the funding of projects by funding organisations would be facilitated. In addition, this could mean more direct funding of research projects by the industry at the universities and the idea of the private funding organisation of continuous casting research could be taken up again.
11
5
Research Working Group: Selected Publications
Melt Quality [1] W. Kahl and E. Fromm: Examination of the Strength of Oxide Skins on Aluminium Melts, Met. Trans. 16B (1985) S. 47-51 [2] E. Fromm: Bestimmung der Wasserstoffkonzentration in Al-Schmelzen durch eine kontinuierliche Messung des H2-Gleichgewichtsdruckes, Aluminium 65 (1989), S. 12401243 [3] X.-G. Chen and S. Engler: Measuring Hydrogen Content in Molten Aluminium Alloys using the CHAPEL Technique, Cast Metals 6 (1993) 2, S. 99-108 [4] F. Patak: Untersuchungen zur Natrium- und Lithiumentfernung aus Hüttenaluminium, Dissertation. RWTH Aachen 1983 Ingot Solidification [5] R. Ellerbrock und S. Engler: Oberflächenseigerungen von Stranggußlegierungen, Metall 37 (1983), S 784-788 [6] L. Ohm und S. Engler: Festigkeitseigenschaften erstarrender Randschalen aus AlLegierungen, Gießereiforschung 42 (1990) 3, S.131-147 und 4, S. 149-162 [7] W. Droste und S. Engler: Vorgänge beim Angießen von Al-Walzbarren im Stranggießverfahren, 8. Internationale Leichtmetalltagung, Wien, 1987 [8] C.H. Dickhaus und S. Engler: Mechanische Eigenschaften erstarrender Randschalen aus Al- und Cu-Legierungen, in Stranggießen, DGM-Informationsgesellschaft GmbH 1995, S 55-66 [9] W. Schneider und W. Reif: Untersuchungen zur Deutung der Vorgänge bei der Kornfeinung von Aluminium mit AlTiB-Vorlegierungen, Gießereiforschung 32 (1980) S. 53 [10] A. Banerji and W. Reif: Development of AlTiC Grain Refiners Containing TiC, Metall. Trans. 17A (1986), S 2127 [11] R. Mannheim und W. Reif: Kornfeinung von CuSn-Legierungen mit Zirkon, Bor und Eisen sowie CuAl-Legierungen mit Titan, Bor und Zinn., in Erstarrung metallischer Schmelzen, Deutsche Gesellschaft für Materialkunde 1981, S. 109-140 Process Technology [12] H.R. Müller und R. Jeschar: Wärmeübergang bei der Spritzwasserkühlung von Nichteisenmetallen, Zeitschr. f. Metallkunde 74 (1983), S. 257-264 [13] C. Köhler, E. Specht and R. Jeschar: Heat Transfer with Film Quenching of Vapourizing Liquids, Steel Research 61 (1990) 11, S. 553-559 [14] H. Kraushaar, H. Griebel und R. Jeschar: Einfluß der Kühlwasserqualität auf den Abkühlvorgang heißer Oberflächen, in Stranggießen, DGM-Informationsgesellschaft GmbH, 1995, S 231-240 [15] D. Hartmann und S. Engler: Erstarrungsverhalten von Cu-Legierungen beim horizontalen Stranggießen, Metall 46 (1992) H.2, S. 139-144 und H.4, S. 333-340
Melt Treatment
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
Hydrogen in Aluminum Containing Copper Alloy Melts – Solubility, Measurement and Removal Karsten Neumann, Bernd Friedrich, Klaus Krone IME Process Metallurgy and Metal Recycling, RWTH Aachen, Aachen, Germany
Jürgen Jestrabek, Elmar Nosch Schwermetall Halbzeugwerk GmbH, Stolberg, Germany
1
Introduction
Aluminum containing copper alloys show a significant hydrogen pick up during melting which is caused by reduction of humidity from various sources like atmosphere, scrap and melt covering agents. Reaction products are aluminum oxide and hydrogen, which is easily taken up by the melt. In combination with the solubility change during solidification, gas porosity may arise in cast products if critical limits of hydrogen concentration in the melt are exceeded. Currently, major quantities of CuAl5Zn5Sn1 (Nordic Gold) strip are produced for the fabrication of Euro currency coins. Some production lots show material defects caused by gas pores. The current project identifies sources for increased hydrogen contents and investigates the possibility of a quality improvement by removing increased hydrogen contents from CuAl5Zn5Sn1 melts.
2
Hydrogen Solubility
In order to predict expected hydrogen contents in the melt, the hydrogen solubility in copper melts and the influence of the main alloying elements on the solubility have to be considered. In copper alloys, the alloying elements aluminum, zinc and tin all lower the hydrogen solubility in both solid and liquid state [1,2]. However, no consistent thermodynamic data set is available to describe the effects of these elements in dependence of concentration and temperature. A previously published investigation [2] contains data to estimate solubility for one temperature in each solid and liquid state. The solubility is calculated by adding up the influences of the alloying elements. (Table 1). This simplification is limited to low contents of alloying elements. Table 1. Influence of alloying element i on hydrogen solubility in mass-ppm H per mass-% i [1] Alloying Element i 700 °C 1150 °C Aluminum -0,06 ppm / [% Al] -0,52 ppm / [% Al] Tin -0,01 ppm / [% Sn] -0,22 ppm / [% Sn] Zinc -0,01 ppm / [% Zn] -0,22 ppm / [% Zn]
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
16 Using this data, the hydrogen contents in copper and CuAl5Zn5Sn1 at 700 °C and 1150 °C in equilibrium with water vapor have been calculated (Table 2). Partial pressures of 1013 hPa and of 25 hPa as present in usual atmospheres have been considered. Table 2. Estimated hydrogen solubilities for pure copper and CuAl5Zn5Sn1 in mass-ppm Material Partial pressure of water 700 °C 1150 °C 0,56 ppm 6,4 ppm p H O = 1013 hPa Copper, oxygen free 0,09 ppm 1,0 ppm p H O = 25 hPa 0,20 ppm 2,48 ppm p H O = 1013 hPa CuAl5Zn5Sn1 0,03 ppm 0,39 ppm p H O = 25 hPa 2
2
2
2
From binary alloy solution data, a solubility curve for hydrogen in CuAl5Zn5Sn1 has been estimated. The effect of the melting interval has been included (Figure 1).
Figure 1: Equilibrium solubility of hydrogen in CuAl5Zn5Sn1 in equilibrium with 25 hPa total water pressure
3
Methods for Hydrogen Determination
Hydrogen can be determined in copper alloy melts either by in-situ measurement or by sampling and subsequent analysis of the samples. The major advantage of the in-situ measurement is the elimination of hydrogen losses during casting and solidification of the samples. Two principles for in-situ-determination are currently available. One method already in use in the steel industry works by contacting a well defined volume of carrier gas with the melt; this results in an equilibrium hydrogen content in the carrier gas corresponding to the hydrogen content of the melt which can be measured by thermal conductivity. This method requires expensive disposable probes for each measurement. The second method measures the electromotive force between a reference material and the melt with a proton-conducting ceramic serving as the electrolyte. However, this method has not yet been introduced into industrial practice and it is still uncertain whether the electrolyte is resistant to the alloying elements contained in copper alloy melts. In the present work, for technical and economical reasons it was decided to use the classical two step method with sampling/solidification and subsequent analysis of samples.
17 3.1
Sampling Method
A major requirement for reproducible hydrogen determination by sampling is the rapid solidification and further cooling of samples to avoid hydrogen losses during sampling. The samples can be cast into moulds or be taken using disposable sampling probes of different shapes which are immersed into the melt. The advantages of casting samples are low sampling costs and rapid solidification caused by the heat capacity of the mould. A major disadvantage is the possible influence of different mould temperatures on the solidification speed and thus on the hydrogen contents determined. Disposable samplers may offer a higher reproducibility concerning thermal conditions and by avoiding the casting step. However, due to the insufficient surface quality of the samples and the slower solidification inside the thermally insulating probes, the cast samples typically deliver better results. Therefore, samples cast into a steel mould have been used in the current project (Figure 2).
70 mm
Sample ∅ 6 mm ca. 4 g
10 mm
Figure 2: Sample shape used
Water has been used for quenching the samples after solidification. Due to the rapid cooling (> 500 K/min), a hydrogen pick up by decomposition of water can be neglected. Liquid nitrogen has also been tested, but led to slower cooling rates due to an isolating gas film forming on the sample surface (Leydenfrost phenomenon), which caused significant hydrogen losses during cooling (Figure 3).
Figure 3: Effect of water and nitrogen quenching on hydrogen content determined in CuAl5Zn5Sn1
3.2
Analysis Procedure
The samples are cut into pieces 2-5 g of weight and are degreased using organic solvents. The hydrogen content of the samples is determined by carrier gas hot extraction, using a LECO RH-402 analyzer with an extraction time of approx. 5 minutes at 800 °C. The equipment is
18 calibrated daily by gas dosing and blank measurement. Hydrogen determination by melt extraction has also been tested, but has not been used as standard procedure because of the increased zinc evaporation during the extraction which leads to contamination of the analysis equipment and is less reproducible.
4
Influences on Hydrogen Contents
As the production batches differed substantially in hydrogen content, investigations have been made to determine possible correlations between environmental or process parameters and hydrogen contents of melts. Atmospheric water vapor and humidity of scraps are considered to be the main source for hydrogen pick up. According to the reaction gas 1 + 13 Al = H+ 61 Al2 O3 ( H,Al:species dissolved in the melt ) 2 H 2O the decomposition of water vapor leads to increased hydrogen contents in the melt. In the present study, the hydrogen contents determined in CuAl5Zn5Sn1 melts show a significant correlation with the atmospheric humidity at cast time (Figure 4).
Figure 4: Correlation between atmospheric partial pressure of water vapor and hydrogen contents in CuAl5Zn5Sn1 melts
Lubrication agents contained in chips and other scrap materials may be another possible source of hydrogen. However, a significant correlation between the amount of humid scraps and the hydrogen contents has not been determined yet. Further investigations in this respect are in progress.
5
Production Scale Testwork
Many of the material defects occurring in CuAl5Zn5Sn1 coin strip production have to be attributed to increased hydrogen contents of the melts. Because the main source for hydrogen
19 is the atmosphere, hydrogen pick up is more or less inevitable in normal melting practice. Thus, an additional melt treatment step for hydrogen removal is required. In general, hydrogen can be removed either by a vacuum treatment or by inert gas purging, the latter being the only feasible alternative for economical reasons. Gas Purging with nitrogen using an impeller injection system (Foseco) was chosen as the best suitable method. The purging operation is carried out in a 40 t short coil, crucible type induction furnace (diameter: 1,80 m; impeller immersion depth: 1,50 m). Currently, in a 30 minute treatment with 40 l/min of Nitrogen, a decrease in hydrogen content of 0,15-0,2 ppm is achieved. The hydrogen content in CuAl5Zn5Sn1 melts can be reduced below a critical value, and the quality of the rolled product can be significantly improved. Meanwhile, all production lots are treated by this gas purging operation.
6
Conclusions
Different possibilities for determination of hydrogen in zinc and aluminum containing copper alloy melts have been evaluated. A sampling procedure has been developed which allows reproducible determination of hydrogen in CuAl5Zn5Sn1 (Nordic Gold) melts by carrier gas hot extraction using a LECO RH-402 analyzer. Atmospheric water vapor has been identified as a source of increased hydrogen contents found in production melts of CuAl5Zn5Sn1. Due to this fact, a melt treatment step for hydrogen removal is advisable. Production scale tests show that impeller gas purging of CuAl5Zn5Sn1 melts in a 40 t induction furnace is an effective way to reduce hydrogen content to an acceptable level. Meanwhile, all production lots are treated with the gas purging operation developed. Further optimization is still required to enhance the reproducibility and stability of the process.
7
References
[1] R.O. Thomas, S. Harper, J.E. Bowers, Gaseous and Gas-forming Elements in Copper and Copper Alloys, International Copper Research Association, Inc., New York 1983 [2] E. Fromm, E. Gebhardt, Gase und Kohlenstoff in Metallen, Springer-Verlag, Berlin, Heidelberg, New York 1976 [3] K. Neumann, B. Friedrich, K. Krone, Wasserstoffgehalte in aluminium- und zinkhaltigen Kupferlegierungen, in NiM 2000 (Ed. D. Hirschfeld), Schriftenreihe der GDMB, Clausthal 2000
Fundamental Research About Liquid Metal Filtration Bettina Hübschen1, Joachim G. Krüger1, Neil J. Keegan2, Wolfgang Schneider3 1
RWTH Aachen, Aachen, Germany Pyrotek Engineering Materials Limited, Dudley, UK 3 VAW Aluminium AG, Bonn, Germany 2
1
Introduction
Increased quality demands for aluminum have led to the fact that in today’s casthouses filtration of molten aluminum has become a standard operation. For filtration the use of Ceramic Foam Filters (CFF) is a common method to remove inclusions. Deep bed filtration is considered as a dominant filtration mode in Ceramic Foam Filters. Inclusion capture in deep bed filters is the result of two sequential events: transport of the particle to the filter wall and attachment of the particle at the wall [1]. For both events the fluid flow in the channels of the filter is a very important parameter. As inclusions are generally smaller than the pore sizes of the filters used, they are deposited on the pore walls of these filters. So already detached particles can be carried away with the liquid and washed out of the filter during hydrodynamic perturbations, for example if flow velocity rapidly changes or if sudden vibrations occur [2]. The aim of this work was to investigate the flow behavior inside a CFF as a function of flow rate, pore size, pore shape and flow direction. For these investigations two different water models were used.
2
Fundamentals
In Ceramic Foam Filters the liquid has to find its way through a tortuous path of pores connected by channels. Since inside the filter the pore size changes, also the velocity of the melt changes and there are region of turbulent and of laminar flow. In regions of laminar flow, a deposition of particles can occur while in turbulent flow regions very small particles can agglomerate and then be attached in the laminar flow regions. However rapid increases in flow velocity due to perturbations can lead to the fact that previously laminar flow regions will become turbulent and already detached particles will be released as a consequence. Within the pores of the CFF different flow conditions and mixing behavior can occur, depending on flow velocity. The CFF can be regarded as a continuous reactor and the different volume fractions can be determined to characterize the flow behavior inside the filter. One important characterization of the fluid flow is the average residence time which is given by equation 1 [3]. V t= (1) v& where t : average residence time, V: volume of fluid in the vessel, v& : volumetric rate of fluid flow.
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
21 Usually the residence time in reactors differs from the calculated residence time and a variety of different residence times can be found. This means that some fluid elements spend a longer and others a shorter period of time in the system. This distribution of residence times is an important characteristic and describes the performance of a reactor. Tracer measurements are a tool to investigate the flow behavior. There are several methods for introducing tracer material into a system of which pulse input of tracer was used in this work. This involves the feeding of a quantity of tracer over a short time period into the system. The tracer material must not interact with the fluid in the reactor or the reactor itself and the amount of tracer has to be negligible in comparison with the amount of fluid present in the system. In addition the time period over which the tracer is introduced to the system has to be very small compared to the calculated residence time. The concentration of the tracer in the outlet stream is measured and plotted against time. The results can be plotted in dimensionless and therefore more general form by using the variables (C-Diagrams) [3]: c C= (2) Q V where Q = quantity of tracer injected, V = Volume of fluid in the vessel, C = dimensionless exit concentration. And t (3) Θ= t where: Θ = dimensionless time, t = actual time, t = average residence time. The area under each C curve must be unity since all the tracer introduced to the system must eventually leave the system. ∞
∫ CdΘ ≡ 1
(4)
0
According to the distribution of residence times it is possible to define different reactor types of which the C Diagrams have a typical design. This is to be seen in figure 1.
Figure 1: Basic reactor types: left: plug flow, middle: backmix flow, right: presence of dead volume [3]
In the case of plug flow the tracer elements introduced to the system do not mix at all while they pass the reactor and arrive at the outlet exactly at Θ = 1. So there is no spread of residence time. (figure 1,left). In a backmix flow reactor the tracer is dispersed immediately and uniformly throughout the system. This means the tracer concentration in the outlet stream is equal to the concentration inside the reactor. Thus the C diagram shows a decrease in tracer
22 concentration starting from unity during the test run. This means that a fraction of the tracer stays inside the system for a time much longer than expected while another fraction passes the system much quicker (figure 1, middle). The presence of dead volume regions is indicated by a maximum in the C diagram at a time smaller than the average residence time and C > 1 (figure 1, right). The volume of the reactor seems to be much smaller than it actually is. Real reactors usually are a mixture of plug flow, backmix and dead volume. The C diagram for such a mixed model is shown in figure 2. The volume fractions can be determined from the diagram.
Figure 2: Determination of the volume fractions in a mixed model [3]
Figure 3: Full Scale filter box model
FWhile in most cases dead volume decreases the performance of a rector, for filtration a certain amount of dead volume is essential for the deposition of particles.
3
Experimental Setup
For the testwork two different model types were used. The first one was a full scale filter box model. Tracer tests on real CFF were made to investigate the change of flow behavior with the flow rate and filter pore size. The second water model type used was a specially designed single channel model to simulate the flow in one channel of a CFF. Three different single channel models were tested to investigate the influence of pore shape and flow direction on the flow behavior. The setup of the models used for the investigations is shown in figures 3 and 4. For all tests sulphuric acid was used as tracer and also KMnO4 was added for visualization of the flow. A small amount of tracer was added at the inlet of the filter/model and an electrode at the outlet measured the electric conductivity which gives an immediate signal for any change in concentration. The signal was recorded by a computer and thus the concentration-time-curve could be plotted. A valve in the outlet stream allowed different flow rates to be adjusted.
23
Figure 4: Experimental setup and design of single channel models
4
Results
4.1
Single Channel Models
The C-diagrams for each model were evaluated and the different volume fractions determined. Thus the influence of flow rate could be presented in a flow rate-volume fraction diagram. The first single channel model tested consisted of a sequence of spherical pores (model 1). The results of the different volume fractions measured are shown in figure 5. 100
100
Vp
90
Vd
Vm
Vp
90
80
80
70
70
60
60
50
50
40
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30
30
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20
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10
Vd
Vm
0
0 0
0,2
0,4
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flow velocity, mm/s
Figure 5: Model with spherical pores
0,8
1
1,2
0
0,1
0,2
0,3
0,4
0,5
0,6
0,7
flowvelocity, mm/s
Figure 6: Model with elongated pores
The amount of dead volume was very high in this model. With rising flow rate however the amount of dead volume decreased and plug flow volume increased. Due to the fact that the pore shape of this model did not correspond to the pore shape of CFF’s, a second model with elongated pores was tested (model 2). The results are shown in figure 6. In this model the amount of dead volume was generally smaller than in the spherical pores model. At the same time the variation of the results was higher.
24 The third model tested finally considered the change in flow direction the fluid experiences during its passage through the CFF. So a specially designed tortuosity model was constructed (model 3) and the results of the measurement are to be seen in figure 7. 100
100 90
Vp
Vd
Vm
80
80
70 60
60
50 40
40
30 20
20
10 0 0
0,2
0,4
0,6
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flowvelocity, mm/s
0 10
15
20
25
Vd
Figure 7: Volume fractions as function of flow rate
30
35
40
FlowRate, gal/min Vm
Vp
Figure 8: Volume Fractions in a 80 ppi CFF
for tortuosity model
While in the model with spherical pores the amount of dead volume was much higher than the amount of mixing volume, now both parts are similar up to a certain flow rate. Then mixing volume increases rapidly and dead volume decreases at the same time. 4.2
Full Scale Model
In analogy to the single channel model tests, the volume fraction distribution in real CFF’s was measured. Figure 8 shows the results for a 80 ppi filter. In this CFF the amount of dead volume was smaller than in the single channel models. The amount of dead volume is smaller than the amount of mixing volume even for very small flow rates. From a certain flow rate on, the amount of mixing volume decreases rapidly and dead volume increases at the same time. This indicates highly turbulent conditions inside the filter. This also shows a very good correlation to the results of the single channel models. 4.3
Tests with Inclusions
Another test series with the single channel model involved the addition of Al2O3 particles to the model and the observation of their flow behavior as a function of flow rate. Results of the elongated pores model for small and high flow rates are shown in figures 9 and 10. While in figure 9 the particles flow down the model without any turbulences, in figure 10 the particles are circling in one pore for several seconds. In this case no dead volume was available to make a deposition of particles possible.
25
Figure 9: Particle flow in the elongated pores model for low flow rates
5 • • • • • • •
6
Figure 10: Particle flow in the elongated pores model for high flow rates
Conclusions The single channel models used showed an increase in mixing volume when flow velocity increased At the same time dead volume - which is essential for the deposition of particles decreased. Also the shape of the pores was found to be very important. The amount of dead volume was higher if the pores were spherical. If the liquid changes its flow direction, the amount of dead volume is smaller and decreases more rapidly with flow rate. In the real CFF the amount of dead volume was generally smaller than in the single channel models In the CFF the amount of dead volume decreased when flow velocity increased. If there are turbulent conditions inside the pores and no dead volume is available the deposition of particles becomes very unlikely.
References
[1] Eckert, C. E.; Miller, R. E., Molten Metal Filtration: Fundamentals and Models, Light Metals 1984, 1281-1304 [2] Desmoulin, J.-P., Reliability of molten metal filtration, Light Metals 1992, 1093-1099 [3] Szekely, J.; Themelis, N., Rate Phenomena in Process Metallurgy; Wiley-Interscience, 1971
Impact of Grain Refiner Addition on Ceramic Foam Filter Performance Nicholas Towsey1, Wolfgang Schneider1, Hans-Peter Krug1, Angela Hardman2, Neil J. Keegan3 1
VAW aluminium AG, Research and Development, Bonn, Germany London & Scandinavian Metallurgical Co.Limited, Rotherham, South Yorkshire, UK 3 Pyrotek Engineering Materials Ltd., Netherton, West Midlands, UK 2
1
Abstract
An extensive program of work has been carried out to evaluate the efficiency of ceramic foam filters (CFF’s) under carefully controlled conditions. The first phase of this work, reported at previous international meetings, showed that ceramic foam filters have the capacity for high filtration efficiency and consistent, reliable performance. The next phase of the program was to study their performance under conditions closer to those in production. Work was carried out to establish the influence of grain refiner (Al-3%Ti-1%B) additions made before the filter. The evaluation program was again conducted using AA 1050 alloy and metal quality was, as before, determined using LiMCA and PoDFA. Spent filters were also analyzed. In order to better understand the impact that a grain refiner addition has on filter performance, trials were also undertaken using specially produced commercial purity aluminum and Al-0.7%Ti rods.
2
Introduction
The performance of ceramic foam filters (CFF’s) under simulated production conditions has been studied extensively for an AA1050 alloy as reported previously1-3. In these studies, grain refiner was deliberately omitted in order that a baseline understanding of filter performance could be established. Under such conditions, LiMCA results showed that ceramic foam filters could give mean filtration efficiencies comparable in range to those of competitive filter systems such as bed filters and rigid media filters. The mechanism of filtration appeared to be associated with the formation of `bridges´ of inclusions across the `windows´ of the ceramic foam pore structure in 50ppi and finer filters. In the work presented here, the impact of a grain refiner addition (in rod form) on melt quality and ceramic foam filter performance has been studied for 3:1 TiBAl (Al-3%Ti-1%B) grain refiner and 50 ppi filters.
3
Experimental Procedure
Trials were carried out at the specially dedicated production scale R&D unit at VAW’s Rheinwerk plant. An AA 1050 alloy, batched using reduction line metal, was cast into ingots
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
27 by the direct chill process at a flowrate of 10 tonne/hr, whilst metal quality measurements were made in the launder. A schematic of the experimental layout employed for this work is shown in Figure 1. Grain refiner rod was fed via a guide tube into the melt as far as possible upstream of the filter box. The time available for the rod to dissolve before the metal entered the filter was approximately 1.5 minutes. LiMCA was the main technology used for assessing inclusion concentrations. PoDFA was used as a back up, with the added benefit of providing particle identification capabilities. Spent filters were also assessed for each trial. The majority of this work was carried out with 3:1 TiBAl rod as normally used in production. A number of additional trials were devised to study the impact of grain refiner on melt cleanliness in more detail. The possible `mechanical´ effects of a rod addition, such as vibration and oxide pull-in at the point of entry, were addressed by feeding a rod of commercial purity Al into the melt. A binary Al-Ti rod containing 0.7% Ti (approximately the same amount of `free´ Ti as that found in a 3:1 TiBAl grain refiner) was used to study the dissolution of TiAl3 platelets along the length of the launder in the absence of TiB2 particles. All the rod material used in this program was evaluated metallographically. In most cases the rod alloys were fed at a rate of 1kg/tonne, slightly higher than that for routine production to deliberately intensify any effects that might occur. For most trials the rod was only entered into the melt partway through the cast after about 20-30 minutes of stable LiMCA readings, to clearly highlight the effect of the rod addition on LiMCA counts. To simulate `real´ casting conditions, additional trials were conducted where the rod was fed from the start of the cast. Furnace LiMCA/PoDFA Rod feed
±4.5m
LiMCA/ PoDFA
Mold
Ceramic Foam Filter LiMCA/ PoDFA
Figure 1: Pilot plant schematic
The `loading´ on the filter, here defined as the inclusion content of the AA1050 metal flowing from the furnace, was varied by a stirring (high load) or settling (low load) practice. The impact of sustained high loading throughout casting was investigated by stirring the furnace during casting. Finally, trials were conducted without a CFF where LiMCA was used to evaluate the effects of TiAl3 dissolution rate. Here three LiMCA units along the length of the launder system were used to monitor the changes in particle concentration with increased distance from the rod insertion point. For all categories of trials, replicate casts were made to confirm reproducibility of the results.
28
4
Results and Discussion
4.1
Impact of a 3:1 TiBAl Rod on the Efficiency of a CFF
Figure 2 shows the impact of a 3:1 TiBAl addition on the performance of a CFF when the incoming inclusion loading is high. It can be seen that before the grain refiner rod was introduced (after 28mins) the high level of filter efficiency found in the previous trial series for the 50ppi CFF’s was confirmed. The CFF has once again removed the vast majority of the incoming inclusion loading. After the 3:1 TiBAl rod is introduced a steady rise in the inclusion value can clearly be seen exiting the filter. In order to get a more realistic appraisal of this effect the grain refiner was also fed from the start of some casts (Figure 3). 50ppi CFF, 3:1 T iBAl 1kg/T , stirred before cast 14
12
Before filter
Impurity Level N15 (k/kg)
10
After filter 8
6
4
ROD IN 2
0 0
10
20
30
40
50
60
70
Ca sting Time (min.)
Figure 2: 3:1TiBAl rod – fed partway through cast – “high” inclusion load 50ppi CFF, 3:1 T iBAl in at start of cast, stirred 14
12
Impurity Level N15 (k/kg)
10
8
6
Before filter After filter 4
ROD IN 2
0 0
10
20
30
40
50
60
70
80
Casting Time (min.)
Figure 3: 3:1TiBAl rod – fed from the start of casting – “high” inclusion load
Figures 3 and 4 indicate that the overall efficiency for this cast and the efficiency across the inclusion size distribution range are decreased compared to casts where grain refiner was not used1-3. These efficiency levels are now more consistent with those reported previously under production conditions4. More significantly the downstream cleanliness levels for the cast in Figure 3 were around 2 – 3k/kg as opposed to the 0.2 – 0.3k/kg level without grain refiner. The observed effect was found to occur mainly in the 15 – 35 µm size range.
29 To appreciate the magnitude and implication of this post filter effect, it is necessary to consider the pre filter effect of adding grain refiner rod. This is best seen on a well settled melt where the effect of adding the rod results in an increase in N15 of 0.5 – 1.5 k/kg – Figure 5. Comparing Figure 2 and Figure 5, it is believed reasonable to assume that the pre filter effect when the loading is `high´ is of a similar order of magnitude to when it is `low´. The post filter effect when the loading is `high´ (4-6k/kg) is thus `disproportional´ to the pre filter effect (0.5-1.5k/kg). When the loading is low as in Figure 5, the post filter effect is much milder. This very fact suggests that it is not the agglomeration of 20k/kg) no release effects at all are evident. In fact, despite there being severe metal level disturbances due to the vigor of the stirring, the 50ppi CFF displayed a very high efficiency and a stable and consistently low post filter LiMCA value (0.25k/kg). It could be concluded therefore, that inclusion loading alone at the levels investigated did not reduce the filter efficiency. This is accepting that at even higher loading levels filter saturation and diminished efficiency may occur. Figure 8 looks at the problem in another way, with the rod entered at the start, as per normal practice, but this time with stirring occurring later in the cast. Here, the same response as before can be noted. When a higher loading is introduced from the furnace in the presence of the grain refiner, a sharp decline in the filter efficiency occurs. 50 ppi CFF, NO rod, stirred during cast
18 16
Impurity Level N15 [k/kg]
14 12 10
after filter
8
STIR
6
before filter
STIR
4 2 0 0
10
20
30
40
50
60
70
80
90
Cast ing T ime (min.)
Figure 7: No grain refiner – stirred throughout cast – “high” inclusion load
In summary, it is postulated that the introduction of Ti & B containing grain refiner material alters the behavior of the ceramic foam filter in trapping and/or retaining particles thus causing them to have a diminished efficiency compared to those found in the absence of grain refiner. This was only found to be significant when the incoming inclusion loading is high. If good furnace practices are followed and the inclusion loading is low (settled melts) there appears to be only a minimal impact of the grain refiner on the filter’s performance. Metallographic assessment of spent filters suggested that the `bridges´ of inclusions across the filter cell junctions seen in the absence of grain refiner1-3, do not appear when grain refiner
32 is employed. It is believed that the interaction of the TiB2 particles and the inclusions in the metal or filter alter this mechanism of filtration and is responsible for the significant increase in inclusion counts at the filter outlet when the grain refiner was introduced later in the cast. `Bridges´ across the `window´ regions (at least for the particle types in the study) may be a form of cake filtration and appear to be associated with high filtration efficiencies. 50 ppi, LSM 3:1 TiBAl rod 1kg/T, ’pre-settled’ 90min.
14
Impurity Level N15 [k/kg]
12
after cff before cff
10 8
5min.air stir started
6 4
Rod in 1kg/T
2 0 0
10
20
30
40
50
60
70
80
90
Casting T ime (min.)
Figure 8: 3:1 TiBAl – fed from start of casting plus stirred partway through cast –“low” inclusion load
5
Conclusions
1. Ceramic foam filters have the capacity for high efficiencies in the absence of grain refiner, even under severe disturbance conditions and with sustained high loading throughout the cast. 2. At high inclusion loading the introduction of a 3:1 TiBAl grain refining rod leads to a reduced filtration efficiency. When the inclusion loading is low there appears to be a minimal impact of the grain refiner on the filter’s performance.
6
Acknowledgments
Sincere thanks are due to VAW Rheinwerk personnel and D.Gründler & N. Ozturk of VAW’s R&D division, without whose dedication the success of this program would not have been possible.
7
References
[1] N.J.Keegan, W.Schneider, H.P.Krug, Light Metals 1999, pp 1031 - 1041 [2] N.J.Keegan, W.Schneider, H.P.Krug, 6th Australasian Asian Pacific Course & Conference, Aluminium Cast House Technology : Theory & Practice (Ed.: M.Nilmani), TMS 1999, 159-174. [3] N.J.Keegan, W.Schneider, H.P.Krug, Light Metals 1997, 973-982. [4] C.Dupuis, G.Beland, J.P.Martin, Proceedings of the 32nd Annual Conference of Metallurgists, Quebec, Canada, CIM, 1993, 349-358.
Review of Dissolution Testing and Alloying Methods in the Casthouse Gregorio Borge1, Paul S. Cooper2 and Stuart R. Thistlethwaite2 1
Bostlan, S.A., Larragane, 1 E-48100 Mungia (Spain).
[email protected] LSM Co.Ltd., Fullerton Road, Rotherham, South Yorkshire S601DL (Great Britain),
[email protected];
[email protected] 2
1
Introduction
The properties of Al alloys are largely dependent on the correct addition of the alloying elements before the casting process. There are a number of ways of performing these alloying additions such as pure metal, master alloys, powder injection and a variety of compacted powders additives (tablets, mini tablets and briquettes). The choice of which alloying addition to use for each element is complex [Thistlethwaite 1992]. A number of competing factors have to be taken into account and their relative importance may vary from plant to plant and product to product. Some of the key criteria include metal temperature available, virgin:scrap ratio, furnace type and layout, addition and stirring practices, alloy change frequency and end product quality. The cost of alloying is not always easily defined. There are not only raw materials costs, but also processing, yield, quality and overhead considerations, which need to be taken into account when selecting the most appropriate alloying technique. LSM and Bostlan´s experience as worldwide suppliers of different products for the aluminium alloying industry is that in recent years consumption of compacted additives has noticeably increased. Some casthouses have stopped their injection production lines, and new facilities for this addition practice are rarely set up, mainly due to capital costs and further the strict quality control requirements of the raw materials because of safety risks when handling powders. Compared to master alloys additives, compacted powder additives are easy to handle; cold metal quantity to be added to the furnace is not very high (since the lowest concentration of the alloying metal in the compact is 75%); accurate additions for compositional adjustments can be performed if necessary; and stocking costs are reduced. In the mid 90s some studies on dissolution of compacted powders were published [Young 1993; Campbell 1994; Fisher, 1994; Perry 1994; Shafyei 1995]. Many of these works are focused on laboratory studies, so it can be said that the general mechanism and the behaviour of the compacted additives in small furnaces is known (exothermic heating of the compact, intermetallic compounds and swelling of the compact). During subsequent years, the literature concerning dissolution and recovery of alloying metals from compacted powders has significantly decreased. The most recent work covers deeper studies on the intermetallic compounds influence for the explanations of the dissolution mechanisms [Bristow 1999; Lee 2000]. All this background literature is useful for understanding the behaviour of the compacted additives, but it could be said that in general no recommendations for an industrial practice have been given. There are several issues the casthouse is interested in including dissolution Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
34 rates and final recoveries in industrial scale furnaces; production of skims/drosses and correct skimming practices for each additive; correct stirring. The challenge is to know how laboratory tests relate to industrial practice.
2
Products
Pressing mixtures of metallic powders (Mn, Fe, Ti, Cu, Cr, Ni, Pb, etc) with aluminium, a flux, or a mixture of both components produces compacted powder additives for the aluminium alloying industry. Alloying metal contents range from 75% to 85%. The most common compacted additives are tablets and mini tablets. Both are cylindrical. Tablets are nominally 90 mm diameter; height and weight depend on the alloying element, but usual figures range from 1250 to 1333 g and approximately 45 mm height. For mini tablets, 40 to 45 mm diameters are available, with 50 to 200 g weights and different heights as well. Tablets and mini tablets are produced in hydraulic presses, using special steel punches and dies. To maintain acceptable tool life requires the use of a lubricant in the formulation of the product. A comparable product is a briquette with the pressure produced between two compacting rolls with “pillow shaped” indentations in the rolls to form the tablet shape.
3
Techniques
The performance and behaviour of compacted additives in molten aluminium furnaces is usually studied by the TP-2 test as published by The Aluminum Association [Aluminum Association 1990], although many laboratories adapt the procedures or equipment to their own facilities. The TP-2 test includes the use of crucible furnaces, temperature ranges around 732±10ºC, sampling every minute during the test’s first ten minutes, and frequent stirring practices. Following the test, dissolution rates and final recoveries of the alloying metal can be studied, as well as skims produced and reactivity phenomena (bubbles, flames, fumes). Scientific literature [Perry 1994; Shafyei, 1995] mentions also microscopic techniques for the study of the intermetallic compounds, whose formation is the first step before the dissolution. The connection of these study techniques, and their corresponding application, to industrial practice can be sometimes confusing or not clearly seen, mainly due to problems arising from the different scale when working in the casthouse. New techniques or working methodologies applied to the study of the compacted additives can help to understand the dissolution process. 3.1
The Microscopic Behaviour (the Monitoring of what Happens in the Furnace)
3.1.1 Classical Monitoring Using a steel cone as in the Aluminum Association TP-1 grain refiner test 10 kg of Al is heated in a small resistance furnace. A whole or part tablet is placed on its edge in the bottom of the steel cone mould. The tablet and mould are preheated and then lowered into the bath and allowed to fill to the level of the notch. The mould is held in the bath for the time required before removing and quenching. The experiment is repeated with different hold times if necessary.
35 The cast cones are sectioned vertically, bisecting the circular face of the mini tablet. The cut faces of the samples are ground flat to reveal the structure and any undissolved remnants of the tablet. Figure 1 shows some scanned images of the cast samples for 85% Mncontaining 200 grams mini tablets.
Figure 1: Steel cone mould test for 200 g 85%Mn mini tablets. Mini tablets were extracted after 60 seconds, 120 seconds, and 240 seconds after addition
Visual information given by this technique is direct. The methodology has not been widely mentioned in the literature [Bristow 1999], but comparisons between different materials can be performed easily and at relatively low cost by this technique, especially if dissolution rates in the first few minutes are important. Thus, fast checking of the behaviour of the material is possible. As disadvantages, no micrographic studies are possible, and scaling problems could arise when applying results to an industrial furnace. As an example, there is no possibility of using complete standard tablets [only a portion]. 3.1.2 X-Ray Monitoring Recently, the Department of Materials of the University of Birmingham has developed a technique that allows a direct view through the dissolution process based on X-Ray radiations. Using this technique, a continuous monitoring of the compacted additive can be recorded on videotape. For the test a mini tablet can be added into a sand mould 125 mm wide and 250 mm deep containing 7 kg aluminium. This technique confirms directly the results obtained from others, especially those concerning the swelling and breaking down of the compacted additive. Flames in the first seconds are due to the presence of solid lubricants in the formulation of the additives for proper compaction. Further swelling phenomena are due to the formation of intermetallic compounds between the alloying element powder and the aluminium. The additive is finally broken due to the melting of the aluminium/flux within the mini tablet. This stage does not mean that the alloying metal (for example Mn) has been recovered at this time, but that the compacted structure has been broken down. The X-ray technique is costly, although direct and comparative results can be quickly obtained. However the dimensions of the sand mould are not adequate for standard tablets or for adding more than one mini tablet as the alloying level would be too high, and swelling phenomena could be uncontrolled. In addition temperature control is not possible during the experiments.
36 3.1.3 Swollen Compacted Additives The techniques described are only useful for direct monitoring of one mini tablet due to the size limitations. Real additions are never like this: many compacted additives are added together usually in the same part of the furnace, so liquid aluminium may not enter into the mini tablet so easily. Swelling phenomena and further metal recovery may thus be delayed. Swollen compacted additives (one mini tablet, one standard tablet, or some mini tablets) can be obtained from industrial scale test furnaces with a more realistic approach to the customer’s situation. As an example, Figure 2 shows some results for 75% Mn 100 grams mini tablets added at 730ºC and extracted from the furnace at the times shown. The furnace used was a rotary oxycombustion 400 kg facility and a 15 cm diameter by 8 cm height holed ladle was used for sinking and extracting the samples. Table 1 below summarizes data from the tablet samples. Table 1: Tablet samples Time (s) Temp (ºC) W (g) 0 N/A 100 30 728-725 103 45 730-728 137 60 731-729 59 70 732-730 23
h (mm) 17.55 24.95 25.00 N/A N/A
Ø (mm) 40.05 49.90 55.10 N/A N/A
ρ (gcm-3) 4.6 2.4 2.3 N/A N/A
The extracted and swollen samples can be cut and polished for examining on an optical microscope. Figure 3 shows an example for the piece extracted at 45 seconds. The intermetallic compounds can be seen. Advantages of this technique include a direct monitoring of the real process in a furnace. Samples extracted can be weighed, and directly analysed. For example, it can be seen that at 30 seconds the swelling phenomena has started but no aluminium has entered into the mini tablet yet. If related to previous techniques (X-Ray), it can be concluded that even if the mini tablet is destroyed at 70 seconds, dissolution is not complete. Conclusions from the micrographic analysis can also be obtained; in this case the samples are closer to what would be obtained from an industrial furnace. Finally, this 400 kg furnace allows also the addition and retention of swollen standard tablets or groups of mini tablets. The main disadvantage of this technique is the cost. 3.2
The Macroscopic Behaviour
In terms of the use of compacted additives, issues of importance in the casthouse include high recovery in a short time, the production of skims and/or dross due to the operation, and the reactivity (flames, fumes, bubbles) of the material added to the furnace. Knowledge of the dissolution mechanism given by the microscopic techniques helps the manufacturer to produce materials with different characteristics for achieving better results. However these results have to be accomplished in an aluminium furnace. The Aluminum Association’s TP-2 test does not specify any furnace size for the dissolution and recovery test. Crucible furnaces of any size could be used. Laboratory costs imply that small (10 to 40 kg) furnaces are generally used. The results of these tests are subject to doubt because they are on a much smaller scale:
37
Mn75% - 730ºC
0’’
30’’
45’’
60’’
70’’
Figure 2: Swollen samples extracted from the 400 kg furnace
Figure 3: Microscopy of the sample extracted at 45 seconds (x50)
•
Standard tablets cannot be directly studied, since they add too much material for the final alloying level. • Stirring practices are different to those used in the casthouse. The TP-2 test proposes a highly effective stirring system for every minute of the test for the first 10 minutes. • Dross production and reactivity have to be estimated in a very small surface. Skimming practices are thus also different from those used in the casthouse. Issues of relevance to practical dissolution of compacted additive powders cannot be explained just by experiments carried out in small crucible furnaces. With the size of the experiment being such an important factor, a methodology was developed for a 400 kg furnace in order to perform continuous and reproducible dissolution tests. The furnace is shown in Figure 4. The main aim of using an almost industrial furnace is to obtain results with no scaling problems. The content of this kind of furnace should be stirred for research work, since homogenisation of the melting bath is necessary for maintaining the temperature and for adequate and repetitive sampling. Stirring is performed before every sampling process, using a rake not hitting the tablets (or mini tablets) added. Samples are usually taken every five
38 minutes for 40 minutes (for Mn, Fe or Cr tests, for example) or even for 90 minutes (for Ti). For performing a proper comparison of the results, materials are added without any packaging, whereas industrial practice is usually to add plastic or foil wrapping, and/or cardboard boxes or paper sacks. Temperature control is performed with a thermocouple sunk in the bath; it is usually accepted to have a range of ±10ºC for each experiment, but usually a ±4ºC after material addition can be achieved.
Figure 4: A rotary 400 kg furnace for aluminium dissolution tests
Some advantages of this facility for reproducing the industrial practices are: • Realistic stirring practices. A ceramic rake is used in this case, which is longitudinally used for homogenisation. This can be taken as a very similar practice to that of many customers. • Distribution studies can be performed. Since differences arise from adding all the material in the same point or in evenly distributed points, a furnace with an adequate surface as this can be used for this kind of research. • Realistic skimming practices. Experiments can be also performed following the customer’s practices for skimming: either before or after addition, or with addition of drossing fluxes if required. • Furnace capacity ensures the researcher or the customer that any final alloying level of any aluminium series alloy required can be achieved. Addition of standard complete tablets is not a problem. The main issue with this kind of experiments is the cost, which is affected by the large aluminium quantity used, even if aluminium can be recovered after the experiment. This technique usually yields dissolution curves closer to real situations in the casthouse, but the control of the experimental factors is more difficult and costly than with a smaller (and more easily controlled) furnace.
4
Conclusions
The increase in consumption of compacted powders for alloying aluminium in the casthouse has focused recent research developments in this field. This work has presented different working methodologies and new-in-the-field applications for ascertaining and proving the behaviour of tablets and mini tablets in aluminium furnaces. It could be said that a supplier
39 controlling these diverse techniques can give a more complete answer to many problems of the customer/producer. In this sense, the most evident application in order to satisfy/answer a customer’s dissolution problem is the large capacity furnace. This furnace combines an almost industrial facility with the possibility of performing designed and highly controlled work: working conditions are closer to those of the casthouse, and macroscopic research work can be performed to obtain results concerning dissolution rates, final recoveries, dross production, and reactivity. Dissolution rate results given by this furnace are usually lower than those given by the typically used laboratory scale crucible furnaces. In most cases this should not be taken as an ineffective method or product, but as a different approach to the problem. On the other hand, the microscopic monitoring of the behaviour of the compacted additives has been improved with new technologies and different methodologies. Fast and low cost comparative analysis of the dissolution rate can be performed sinking a steel cone mould containing a mini tablet or a tablet portion into a prepared crucible furnace. A direct insight of the swelling and breaking down of a mini tablet can be performed using a sand mould and an X-Ray technique. Fast and comparative (but costly) direct results can be obtained. Finally, different types of real swollen samples obtained from a furnace without scaling problems can be classically studied with the microscope. All of these research results are useful in order to propose new developments/products by the supplier, and to understand specific customer related problems under many different situations.
5
References
[1] Aluminum Association, The, Standard Test Procedure for Measuring the Dissolution of Aluminum Hardeners, The Aluminum Association, 1990 [2] Bristow, D.J., Lockwood, S., Woodcock, T.G., and Cook, R., 128th TMS Annual Meeting and Exhibition, 1999. [3] Campbell, G.T., Bridges, R.E. and Niedzinski, M., Light Metals, 1093-1097, 1994. [4] Fisher, P., Cooper, P.S., and Thistelthwaite, S.R., Dissolution Mechanisms in Aluminium alloy additives. 123rd TMS Annual Meeting and Exhibition, 1994. [5] Lee, Y.E., and Houser, S.L., Dissolution mechanism for high melting point transition elements in aluminium melt, 129th TMS Annual Meeting and Exhibition, 2000 [6] Perry, W.H., Aluminium recovery from ‘all metallic’ hardener briquettes, Light Metals, 841-848, 1994. [7] Shafyei, A., and Guthrie, R.I.L., Dissolution mechanism of compact briquettes of high melting point additives stirred in liquid aluminium, Light Metals, 831-839, 1995. [8] Thistlethwaite, S.R., Review of alternative methods for alloying aluminium, Light Metals, 1005-1011, 1992. [9] Young, D.K., Setzer, W.C. and Boone, G.W., New concept in alloying aluminum, Light Metals, 745-751, 1993.
The Effect of Casting Parameters on the Metallurgical Quality of Twin Roll Cast Strip Yucel Birol1, Gökhan Kara2, A. Soner Akkurt2, Chris Romanowski3 1
Marmara Research Center, P.O. Box 21, 41470 Gebze-Kocaeli, Turkey ASSAN Aluminum Works, Tuzla, Istanbul 81700, Turkey 3 FATA Hunter Inc., Riverside, California 92507, USA 2
1
Abstract
31 different samples covering a range of casting parameters for the AA8006 alloy, were cast on industrial scale with 1725mm and 2184 mm wide Speed Casters. This paper describes the general trends correlating casting parameters to the metallurgical quality of the cast strip.
2
Introduction
When compared to the traditional hot mill process, the relatively low capital cost of twin roll casters, in combination with their lower energy and manning costs, have made twin roll casting an increasingly popular method of producing a wide range of aluminum flat rolled products [1]. The recent trend has been to reduce the gauge at which these casters operate to t+∆t
Calculation of microsegregation
t+∆t Tt+∆t , f lt+∆t , wt+∆t i,l , wi,s ,
End
Figure 1: Flow chart of the macrosegregation computation
3
Model Description
In the presence of shrinkage and solid transport at a uniform velocity, vs, the average conservation equations for heat, mass and solute species are given by [3]: ∂ Heat : + div - div ( grad T) = 0 (1) ∂t ∂ Mass : + div = 0 (2) ∂t ∂ Solute (i) : + div = 0 (3) ∂t where denotes the average over the liquid and solid phases of the field ξ, i.e., = ξsgs + ξl gl, where ξs and ξl are the values in the solid and liquid phases, gs and gl being the volume fractions of solid of liquid, respectively. ρ is the specific mass, h is the enthalpy per unit mass, v is the velocity, κ is the thermal conductivity, T is the temperature and wi is the mass fraction of solute species (i). For this last entity, the local averaging of the solute concentration, , is slightly different if microsegregation does not occur at equilibrium (i.e., not according to lever rule) :
193 gs
= ρlwi,lgl + ⌠ ⌡ ρswi,s(g,t)dg , i = 1, N
(4)
0
where N is the number of solute species. The momentum conservation equation is usually written only for the liquid phase. In the case of global transport of the solid, such as that encountered in DC casting, this can be done for the relative velocity of the fluid u = (v - vs) = (vsgs + vlgl) - vs = gl(vl – vs). One has : µgl ∂u ρl ρl + g (Grad u) ⋅ u - µ ∆u + K u + gl grad p = glρbg (5) ∂t l µ is the dynamic viscosity of the fluid, K the permeability of the mushy zone and ρb the Boussinesq approximation of the density of the fluid : ρb = ρl,o [1 - βT(T - To) - ∑ βi(wi,l - wi,l,o)] (6) i
ρl,o is a nominal specific mass of the liquid taken at some reference temperature, To, and concentrations in the liquid, wi,l,o. βT and βi are the thermal and solutal expansion coefficients, respectively. The schematic flowchart of the program is given in Fig. 1. Without giving the details of the formulation which can be found in [3], these equations are solved at the scale of the whole casting as follows : 1. The heat flow conservation equation is solved according to an enthalpy scheme [2], i.e., using the enthalpy as the variable and linearising the temperature-enthalpy relationship. The velocity field of the previous time step is used. 2. The mass and momentum conservation equations are solved simultaneously using a GLS (Galerkin Least Squares) formulation for the pressure-velocity fields [4]. The temperature, volume fractions, solute concentrations and variations of the average specific mass (shrinkage) are taken from the previous time step. 3. Once the relative velocity of the fluid is known, the solute conservation equation(s) are solved in order to deduce the new average solute concentrations. Knowing the average enthalpy, t+∆t, and solute concentrations, <wi>t+∆t, at all the nodes of the mesh, the new temperature, volume fractions of phases, solute concentrations in the liquid and solid phases, and average specific mass are calculated according to a local microsegregation model. In the present case, the model developed by H. Combeau and A. Mo for binary alloys has been used [5]. It allows to consider non-linear phase diagrams, eutectic reaction and back-diffusion. It should be pointed out that the ternary model developed by X. Doré within the framework of this project [6] for Al-Mg-Si alloys has also been implemented in calcoMOS. The microsegregation model is equivalent to solving at each nodal point the N equations (4) with the following ones [5,6]: Equilibrium at the interface :
*
wi,s = ki wi,l , i = 1, N
(7)
Liquidus relationship : T = TL(wi,l) (8) Enthalpy relationship : = T + Lgl (9) The ki's are the partition coefficients, TL is the equation of the liquidus, is the average volumetric specific heat and L is the latent heat of fusion. When precipitation of secondary phases occurs, the problem becomes more complex and is not detailed here [5,6]. Providing a back-diffusion model is given (i.e., evolution of wi,s(g,t) appearing in Eq. (4)), Eqs. (4,7-9)
194 provide (2N + 2) equations for the (2N + 2) unknowns : wi,s, wi,l, T and gl. This back-diffusion model can be given either by a polynomial function approximation of wi,s(g,t) [5] or by the solution of Fick’s second law in the solid phase using a 1D FDM technique [6]. In the case of DC casting of Al alloys in which one is interested mainly in the stationary solution, the time stepping is used as a means of iteration among the equations. In this case, however, the Eulerian description of the macroscopic conservation equations (i.e., in a reference frame attached to the mould) must be coupled with a Lagrangian description of microsegregation (i.e., in a reference frame attached to the dendrites). Details of this coupling can be found in [3].
4
Results
The results which are presented in this section are for an Al-Mg alloy, using the microsegregation model of [5]. Fig. 2 shows the calculated stationary 2D velocity field near the liquidus isoline for a small DC cast ingot 5-cm thick. For symmetry reason, only half the ingot has been calculated over a length of 15 cm. The metal was supposed to be injected uniformly from the top at a velocity of 1 mms-1 with a nominal concentration of 4.5% Mg. Three calculations were performed under identical conditions but considering various sources of fluid movement : buoyancy only (case (a)), shrinkage only (case (b)) and buoyancy plus shrinkage (case (c))
1 cm (a)
.01 m/s
(b)
0.0002 m/s
(c)
0.01 m/s
Figure 2: Calculated stationary field of the relative fluid flow velocity for a small DC cast Al-4.5%Mg ingot (casting speed : 1 mms-1) with isolines of fraction of solid. a) buoyancy only ; b) shrinkage only ; c) buoyancy and shrinkage. Microsegregation model of Ref. [5] with a Scheil approximation
As can be seen, the relative velocity of the fluid in the fully liquid region (i.e., above the first isoline of fraction of solid also represented in this figure) is of the order of cms-1, regardless of whether shrinkage is included or not in the calculations (compare Figs. 2a and
195 2c). It is induced by thermal buoyancy, the solutal expansion coefficient associated with magnesium being small. Despite this fairly large relative velocity of the liquid, the induced macrosegregation is almost negligible when only buoyancy is considered (see Fig. 3). This is due to the fact that the flow pattern in this case is essentially parallel to the liquidus isotherm, i.e., perpendicular to the solute gradient, and vanishes quickly in the mushy zone. As reviewed in [1] and pointed out many years ago by Flemings, it is the component of the velocity along the thermal gradient which induces macrosegregation. Although the shrinkageinduced velocity is nearly two orders of magnitude smaller than that associated with thermal buoyancy (see Fig. 2b), it has a much more pronounced influence on the final concentration profile at the exit of the ingot (Fig. 3). In the mushy zone, the relative velocity of the fluid is on the order of 0.2 mms-1 only, but it is nearly perpendicular to the isofractions of solid. (Please note that the same flow exists in the mushy zone of Fig. 2c but is not visible with the scale used to visualise the overall flow pattern). Since the streamlines of the interdendritic fluid flow deviate from the ingot centerline (Fig. 2b), this induces a negative segregation at the center of the ingot and a positive one at the surface (Fig. 3). This shrinkage-induced segregation, commonly labeled “inverse segregation” in static castings [1], increases with the depth of the liquid pool, i.e., with the casting speed [3].
(a) (c)
Figure 3: Mg concentration profiles at the exit of the small DC cast ingot for the cases involving buoyancy only (a) and buoyancy plus shrinkage (c) of Fig. 2
The same phenomenon can be observed in real-size ingot computations (Figs. 4 and 5). Unlike small DC castings for which a good resolution can be obtained with a structured mesh following the coordinate axes [3], large scale simulations have to be performed with an unstructured mesh in order to obtain a sufficient accuracy within a reasonable CPU time. A first thermal calculation allowed to determine the approximate position of the mushy zone in this 0.25 x 1 m2 domain and to refine the mesh in this region (Fig. 4a). Feeding of metal through a distribution bag was simulated by inserting a horizontal plate inside the domain. Since the Eulerian-Lagrangian algorithm implemented in calcoMOS for back-diffusion calculations requires, at present, structured meshes (i.e., mesh points aligned along verticals), the macrosegregation result shown in Fig. 5 was obtained with the lever-rule approximation.
196
10 cm (1)
(2a)
(2b)
(3a)
(3b)
Figure 4: Real size DC casting simulation for an Al-4.06%Mg alloy solidified at 1 mms-1 : Mesh size (1), velocity streamlines (2) and isolines of fraction of solid (3) for the cases with shrinkage only (a) and with buoyancy plus shrinkage (b). Microsegregation module of [5] used with the lever rule. 16417 nodes, about 24h CPU on SGI2000
In Fig. 4, the streamlines and isolines of fraction of solid are shown for the two cases of shrinkage only (case (a)) and shrinkage plus buoyancy (case (b)). In the first case (Fig. 4(2a)), the streamlines turn around the horizontal plate and then are fairly straight : they directly outline feeding of the ingot from the upper gate to the mushy zone. On the contrary, thermal buoyancy in the liquid pool is turbulent if a laminar viscosity value is used : it gives rise to a complex fluid flow pattern, which never reaches a stationary state. Since no turbulent model was implemented in calcoMOS, an artificially increased viscosity by a factor 100 was used to obtain the result shown in Fig 4(2b). Results obtained with such an increased-viscosity approximation have been compared recently with those calculated with a turbulent model [7]. One can notice the influence of the primary and secondary coolings on the isolines of fraction of solid (small cusp near the top of Fig. 4(3)). The corresponding values of the heat transfer coefficients were deduced from experimental measurements and inverse method [8]. As can be seen, thermal buoyancy slighltly modifies the depth of the molten pool. However, the two corresponding segregation profiles at the exit of the ingot calculated for this real-size casting (Fig. 5) do not differ much, thus indicating again that shrinkage-induced macrosegregation is dominant over that associated with natural convection. A negative centerline segregration is again predicted by the simulation. The amount of segregation predicted by this model compares fairly well with the concentrations measured within the framework of the EMPACT project [9].
197
(a) (b)
Figure 5: Calculated Mg concentration pofiles at the exit of the real-size DC cast Al-4.06%Mg ingot (Fig. 4) for the cases with shrinkage only (a) and with buoyancy plus shrinkage (b)
5
Conclusion
Computation of macrosegregation in real-size DC cast ingots is a real challenge, even in two dimensions. The size of the region where both the fluid flow and solute gradients are non-zero is very small compared with the overall size of the ingot. Unstructured meshes offer clearly an advantage in this respect, but they are also more complicated to implement (e.g., for a mixed Lagrangian-Eulerian description). It has been shown in the present contribution that shrinkage-induced macrosegregation can already account for a fairly large portion of the concentration inhomogeneity measured in Al-Mg ingots. Thermal buoyancy has a minor influence for this alloy. This does of course not preclude anything about the influence of other phenomena such as grain movement, deformation of the mushy zone or solutal convection induced by other alloying elements.
6
Acknowledgements
This research was undertaken as part of the European Brite-Euram program EMPACT (A European Modelling Programme on Aluminium Casting Technology). It was sponsored by the European Community under contract CEC 0112 and by the Office Fédéral de l’Education et de la Science, Bern, under contract 95.0037-2. The authors would like to thank Dr H. Combeau, Ecole des Mines de Nancy, for providing the binary microsegregation model and for his help during its implementation in calcoMOS.
198
7
References
[1] Ch. Beckermann, Modeling of Macrosegregation : Past, Present and Future. To appear in: Flemings’ Symposium (TMS, Warrendale, PA, 2001). [2] N. Ahmad, H. Combeau, J.-L. Desbiolles, T. Jalanti, G. Lesoult, J. Rappaz, M. Rappaz and C. Stomp, Met. Mater. Trans., 29A, 1998, 617-30. [3] T. Jalanti, Etude et modélisation de la macroségrégation dans la coulée semi-continue des alliages d'aluminium, EPFL PhD Thesis No 2135, Lausanne, 2000. [4] L. P. Franca and S. L. Frey, Comput. Methods Appl. Mech. Engng., 99, 1992, 209-33. [5] H.Combeau and A. Mo, Met. Mater. Trans., 28 A, 1997, 2705-14. [6] X. Doré, H. Combeau and M. Rappaz, Acta mater., 2000, to appear. [7] G.-U. Grün, A. Buchholz and D. Mortensen, in Light Metals 2000 (TMS Publ., Warrendale, PA, USA, 2000) p. 573-78. [8] J.-M. Drezet, M. Rappaz, G.-U. Grün and M. Gremaud, Met. Mater. Trans., 31A, 2000, 1627-34. [9] Joly, G.-U. Grün, D. Daloz, H. Combeau and G. Lesoult, in Materials Science Forum (Trans Tech Publ., Switzerland, 2000) p. 111.
The Effect of the Differencing Scheme on the Numerical Diffusion in the Simulation of Macrosegregation B.C.H. Venneker Netherlands Institute for Metals Research, Delft, The Netherlands
L. Katgerman Delft University of Technology, Laboratory for Materials, Advanced Materials and Solidification Technology, Delft, The Netherlands
1
Introduction
A reliable calculation of macrosegregation during the casting of alloys depends on the accurate modeling of the associated physical mechanisms. Besides that the particular microsegregation model (Scheil, lever-rule) is of importance, the relative movement of the liquid and solid phase inside the mushy zone controls the amount of macrosegregation. In solving the solute concentration equation, the accuracy of the velocity field is thus of great concern. From the literature on Computational Fluid Dynamics, we also know that in high Peclet number flows, the incorrect treatment of the convection terms causes numerical diffusion which can completely overshadow the actual physical diffusion. Throughout the history of CFD, a great number of differencing schemes for the convection term has been proposed in order to reduce the numerical diffusion. In the current research several of these schemes are examined on their ability to correctly predict macrosegregation in the DC casting of an Al-4.5wt% Cu alloy.
2
Numerical Errors in Modeling
If it is assumed that in a particular CFD simulation the right equations are solved, there is still the concern of solving the equations in the right way, see (1) for an important discussion on this topic. The (partial) differential equations of the complete model are transformed into algebraic discretisation equations. Besides that we have to be sure that we have reached convergence, a potentially source of errors lies in the transformation of differential equations to discretisation equations. Particularly the convection term needs special attention. Incorrect treatment of the convection term can result in two types of errors: numerical diffusion - the spreading out of profiles – and numerical dispersion – the appearance of wiggles (oscillations) on the profile. Roughly speaking, numerical diffusion is common to odd-order schemes when convection dominates physical diffusion (high cell-Peclet numbers). Numerical dispersion is common for even-order schemes and occurs in the vicinity of steep gradients. A special form of numerical diffusion is crosswind diffusion, which manifests itself when the flow field is not aligned with the computational mesh. Note the importance of this issue in DC-casting: the more or less parabolic profile of the mushy zone, and the circulating fluid flow in the liquid sump caused by buoyancy effects is a definite cause of the Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
200 nonalignment of flow field and mesh. An attempt to align the mesh has some positive effects, see (2), but is for 3D simulations a cumbersome task. At this point it is worth investigating how these sources of error affects predicted macrosegregation. As shown in several papers (2,3,4,5) the numerical diffusion in the solute conservation equation manifests itself as wiggles in the concentration profile. Although this may point to numerical dispersion at first sight, in (2) it was explained that the origin of wiggles is in the first cell in the upper part of the mushy zone. Moving downwards the scattering in the concentration is frozen. A better term for this phenomenon is false segregation. The location of minimum and maximum concentration can be traced back to the position of the steps in the staircase profile of the liquidus. The differencing scheme in these cases was upwind (2,5) and the power-law scheme (4). Note that the power-law scheme for Peclet numbers greater than 10, is equal to the upwind scheme. It is commonly known that upwind differencing introduces the greatest amount of artificial diffusion. In reference (4) an interesting comparison between a finite volume and finite element package is given. In the FE simulations, however, the streamline upwind Petrov Galerkin method of (6) is used. This scheme has been designed in such a way that artificial diffusion is only added in the streamwise direction, and thus eliminating crosswind diffusion. The results with that method were free of oscillations, whereas with the power-law scheme in some cases oscillations were present, see fig. 8 in that paper. In the remaining of this paper, we will look at some other schemes and will compare their behavior with that of upwind differencing.
3
Differencing Schemes
Out of the vast variety of differencing schemes available in literature, the following schemes were chosen: • upwind differencing • central differences • QUICK • QUICK-2D • higher-order upwind with Superbee flux limiter • skew upwind differencing scheme (SUDS) Each of these schemes is a representative of a particular class of schemes. Thus, upwind and central differencing are the most common used schemes of first and second order accuracy. The QUICK scheme of (7) is a third order modification of upwind differencing. QUICK-2D (8) is the local two-dimensional version of QUICK. The fifth scheme is only second order but has as special feature that the appearance of wiggles on the solution is prevented by the use of a non-linear flux limiter, in this case superbee of (9). Finally, the last scheme is like the QUICK-2D scheme also locally two-dimensional but has as special feature that the exact direction of flow is taken into account. It was especially designed by (10) to take into account the local direction of the flow at each cell face. The performance of these schemes was first tested with two benchmarks: the convection of a scalar profile over 180 degrees - the so-called Smith & Hutton problem (11) and laminar convection in a driven cavity. For both cases, it turned out that the two-dimensional QUICK scheme is the best option. Second best are SUDS for the Smith & Hutton problem, and central
201 differencing in case of the driven cavity. A known peculiarity of the Superbee flux limiter was found in case of the Smith & Hutton problem: i.e. the overcorrecting of numerical diffusion.
4
Numerical Procedures
The following set of equations was solved with the CFD-package CFX version 4.3. ∂ρ + ∇ ⋅ (ρ U ) = 0 ∂t
(1)
∂ (ρ U ) + ∇ ⋅ (ρ UU ) = − ∇p + ∇ ⋅ (µ l ∇ U ) + ρ 0 g − ρ 0 β T (T − T0 )g − µ l K (U − U s ) ∂t ∂f ∂ (ρ h ) + ∇ ⋅ (ρ Uh ) = ∇ ⋅ (λ∇ T ) − ρ L l − ρ L U s ∇ f l ∂t ∂t ∂ (ρ c ) + ∇ ⋅ (ρ Uc ) = ∇ ⋅ (ρ f l D l ∇ c ) + ∇ ⋅ [ρ f l D l ∇(c l − c )] − ∇ ⋅ [ρ (c l − c )(U − U s )] ∂t
(2) (3) (4)
The mushy zone is modeled as a porous region, with K the permeability calculated with the Kozeny-Carman relation. Liquid fraction (fl) and concentrations are calculated according to the lever rule. In order to investigate fully the effects of false segregation, no solidification shrinkage was taken into account. The calculated velocity field is therefore completely determined by the inlet and thermal buoyancy. A relative small ingot of 5 cm diameter and 10 cm length was taken as geometry. Metal entry was over the complete width of the ingot. The inlet velocity was 1.0 mm/s and as alloy Al-4.5wt% Cu. The material properties of this alloy were taken from (3). As primary cooling a heat transfer coefficient of 100 W/m2K was taken over the first 1.25 cm. For secondary cooling, we assumed a heat transfer coefficient of 500 W/m2K. The central difference scheme is unstable for the solute conservation equation, so in that case the upwind scheme was used. From the tested schemes, only QUICK-2D and SUDS are not standard in CFX 4.3. The program offers, however, the opportunity to include userdefined source terms. So, for these schemes, the upwind scheme was chosen as standard and the difference with the preferred scheme is included in the source term using the deferredcorrection method. The last term in equation (2) is included as a user-defined body force. The last term in eq.(3) and (4) are also convection terms and these were also included in the source term, with of course, the appropriate discretisation scheme. Besides the importance of the differencing schemes, the cell size of the mesh is also a crucial factor for numerical diffusion. We therefor performed simulations with each scheme on 4 different meshes: 10, 20, 40 and 80 cells in the radial direction. The cell-width in the axial direction was in all cases equal to that in the radial direction.
5
Macrosegregation Predictions
First, the effect of grid refinement is discussed. From figure 1a, we can conclude that all schemes converge to the same asymptotic concentration at the ingot center for extremely fine grids. With coarse meshes, however, the value can be either greater or smaller than the inlet composition, depending on the particular scheme.
202 All schemes predict a decrease in the concentration with decreasing cell size. It seems that predictions made by the QUICK, QUICK-2D, Superbee and SUDS scheme converge for small cell sizes towards each other, while the upwind (and central) scheme predict an ever increasing negative segregation. Simulations with an even finer mesh have to be performed in order to investigate whether all schemes finally convergence to the same asymptotic value. 5
4.6
4.4
a)
4.6
4.4
4.2
4.2
4 0.0001
Upwind Central Quick Quick-2D Suds Superbee
4.8
C_Wall [wt%]
4.8 C_Centre [wt%]
5
Upwind Central Quick Quick-2D Suds Superbee
0.001 cellsize [m]
4 0.0001
0.01
0.001 cellsize [m]
0.01
b)
Figure 1: Cu-concentration at ingot center (a) and edge (b) as function of the cell size.
Due to the easy conduction of heat by metals, cell Peclet number for the enthalpy equation are even for the coarsest grid, lower than 2.0. Numerical diffusion is than negligible compared to physical diffusion. Therefore, there is little difference in the temperature profile with the different schemes and meshes. The diffusivity in the momentum equations has an intermediate value. The result obtained with the finest mesh was more or less gridindependent, specifically for the radial momentum with the two QUICK schemes. From previous simulations of macrosegregation in literature, it is known that when only buoyancy-induced flow is considered, the amount of segregation is negligible. For the simulations in this work, this is only true with the two QUICK schemes and the Superbee scheme, see figure (1). This should come as no surprise, as both QUICK-schemes were tested as one of the best schemes for the Smith & Hutton problem. Except for near the outer edge, the predicted concentrations are close to the inlet concentration of 4.5 wt%, especially in the case of Superbee. With these two schemes, the predicted macrosegregation is quite large. Despite the obvious presence of false segregation, the profiles are surprisingly smooth. This has probably to do with the relatively small ingot in relation to the boundary conditions. These were such, that the mushy zone is fairly thick (about 5 cm), so it is described by a large number of grid cells. Only in case of SUDS, a strange long-wave oscillation is present. It is unknown what the cause of this oscillation is.
203 4.8
Upwind Central Quick Quick-2D Suds Superbee
4.7
C [wt%]
4.6
4.5
4.4
4.3
4.2
0
0.005
0.01
0.015
0.02
0.025
r [m]
Figure 2: Predicted macrosegregation profile with the finest mesh for all tested schemes.
6
Discussion and Conclusions
The performance of six differencing schemes in the calculation of macrosegregation of Cu in an Al-4.5wt%Cu alloy has been investigated. Conditions of casting were such that the true macrosegregation can be considered negligible. Any macrosegregation, predicted by the simulations, is therefore a direct consequence of numerical diffusion and is called false segregation. In terms of minimum amount of false segregation, the higher upwind scheme with the Superbee flux limiter has the best performance. It is however known that this scheme also reduces physical diffusion. The two-dimensional version of the QUICK scheme is therefore considered to be the most appropriate scheme of the six that were tested. It is emphasized here that even with the finest mesh, no grid-independent solution for the concentration field has been obtained. Given the size of the ingot and the large number of mesh cells, an even better scheme than QUICK-2D has to be found before reliable simulations of real size castings can be made. A possible candidate may be found in the finite element literature. The Streamline Upwind Petrov Galerkin is an often-used method in FE-packages, and has proved to give excellent results in terms of the absence of false segregation (11). Higher order (fourth, fifth) schemes are another possibility.
204
7
References
[1] P. J. Roache, Verification and validation in computational science and engineering, Hermosa publishers, 1998. [2] B. C. H. Venneker and L. Katgerman, Macrosegregation during DC casting of aluminium alloys: numerical issues and the effect of metal entry, MCWASP IX proceedings, Aug. 2000, 8p. [3] V. Reddy and C. Beckermann, Metall. Trans. B., v 28, 479-489, June 1997. [4] Ahmad et al., Metall. Trans. B., v 29, 617-630, Feb. 1998. [5] J. Vreeman and F. P. Incropera, Num. Heat Transf. B., v 36, 1-14, 1999. [6] N. Brooks and T. J. R. Hughes, Comp. Meth. Appl. Mech. Eng., v 32, 199-259, 1982. [7] P. Leonard, Comp. Meth. Appl. Mech. Eng., v 19, 59-89, 1979. [8] P. Leonard, Elliptic systems – finite difference methods IV, in Handbook of Numerical Heat Transfer, eds. W. J. Minkowycz et al., John Wiley & Sons, 1988. [9] L. Roe, Some contributions to the modeling of discontinuous flows, Lectures in Applied Mechanics, v 22, American Mathematical Society, 1985. [10] G. D. Raithby, Comp. Meth. Appl. Mech. Eng., v 9, 153-164, 1976. [11] M. Swierkosz et al., Numerical simulation of macrosegregation in calcoMOS, Empact report. March 2000, 64p.
Application of a New Hot Tearing Analysis to Horizontal Direct Chill Cast Magnesium Alloy AZ91
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
206
207
208
209
210
Micro- and Macrostructures
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
Nucleation Studies of Grain Refiner Particles in Al-Alloys P. Schumacher University of Oxford, Oxford, UK
1
Introduction
The observation of nucleation events in conventional solidification experiments is difficult, because of subsequent growth and impingement of crystals obscuring nucleation events. Nucleation mechanisms of α-Al on Al-Ti-B based grain refiner particles are still not very well understood and mechanisms of nucleation by either borides only or by a peritectic reaction only [1] were proposed early, but were based on indirect, post-solidification based observations. Nevertheless, in casting practice it appeared empirically that both borides and a peritectic reaction were required for sufficient grain refinement, which was empirically verified in synthetically added borides to Al melts [2]. The interpretation of the grain refiner efficacy to nucleate Al cannot be easily derived from the average grain size of as cast microstructure which is strongly affected by growth restriction effects of solute in a given alloy [1,3]. In particular, Ti can have potentially a double role as a strong growth restrictor and as a nucleation agent via a peritectic reaction with liquid and bulk Al3Ti to form Al above 0.1 at. % Ti. The difficulty of identifying nucleation mechanisms has been overcome recently by addition of conventional grain refiner particles into Al-based eutectic alloys which on rapid quenching from metallic glasses [4,5]. Nucleation of Al occurred in the undercooled melt but growth was halted when the atomic mobility rapidly decreased on cooling below the glass transition temperature. To distinguish between nucleation events on added grain refiner particles and those occurring in the glass without additions, the crystallization behaviour of the glasses (devitrification) was studied and was reported elsewhere [6,7]. The current work uses Al based metallic glasses not used previously to simulate effects of solute on the nucleation mechanism of α-Al on commercial Al-Ti-B based refiners in glass-forming model alloys. The amount of Ti and Zr is varied to study the effects of Ti as a nucleation agent and to study Zr poisoning.
2
Experimental Methods
Nominal metallic glass compositions of Al85Ni10Ce5 (in at. %) were produced by arcmelting under reduced, inert atmosphere from pure elements (99.95 at. %). The aluminium content of the arc-melted alloys was reduced to compensate for excess Al of subsequently added grain refiner, Ti and Zr additions. Grain refiner rods (supplied by London & Scandinavian Metallurgical Co. Ltd.) with a composition of Ti-B-Al 5:1 (wt. Ti %) and of stoichiometric Al-TiB2 (2.2 wt. Ti %) composition were added prior to melt-spinning. Charges were heated to 1200°C in boron nitride crucibles and held for varying times in an inert helium atmosphere before ejecting onto a rotating copper wheel with a speed of 40 ms-1. Resulting ribbons were Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
214 approximately 2 mm wide and 30 µm thick. The melt-spun ribbons were examined using transmission electron microscopy (TEM). Thin foils were prepared by electropolishing in 6% perchloric acid in ethanol at a temperature of approximately –30°C at 20 V. TEM was on a Philips 400T microscope at 120 kV.
3
Results and Discussion
3.1
Effects of Ti on the Nucleation Mechanism
In figure 1, two clustered hexagonal TiB2 particles with hypo-peritectic addition of 0.05 at. % Ti are shown embedded into a partial amorphous matrix of Al85Ni10Ce5. The hexagonal platelet shape of the particle tilted with its zone axis nearly parallel to the electron beam can be seen while the second particle is tilted with its parallel to the electron beam is viewed edge-on to the basal face. On both particles copious nucleation of Al has occurred but can only be seen on the second particle viewed with its basal face edge-on. This suggests that similar to previous observations [4,5] a distinct crystallographic orientation relationship is favored for nucleation of Al on boride particles.
Figure 1: (a) Faceted TiB2 particles (5:1) embedded in partial amorphous Al85Ni10Ce5. The hexagonal platelets are tilted close to the (1) and (2) zone axis and are separated by a thin layer (arrow). The layer nucleates copious Al on basal faces of the TiB2 only. (b) The inverted (negative) SAD pattern showing a 0-order streak indicative of the layer
Some nucleation of Al has occurred in the surrounding glassy matrix. The nucleation density of Al particles in the matrix is much lower than that on the boride particle indicating that the boride acts as a better nucleation site for Al. However, between the boride particles, marked in figure 1 with an arrow, and between the nucleated Al and the boride a layer is visible. This layer has been reported previously as a thin layer of Al3Ti in Al85Y8Ni5Co2 glasses [4,5] and is visible here in selected area diffraction (SAD) patterns as a weak streak.
215 Further identification of the layer as Al3Ti in the present case is difficult due to its limitation in thickness and in slight misorientations to the dominant orientation relationship between the boride, aluminide and Al with its close-packed planes being parallel. It is noteworthy, that the TiB2 particles do not nucleate Al directly but indirectly by a thin layer of Al3Ti, which appears to be preserved during the high melting temperatures in the metallic glass experiments. In contrast, no separate Al3Ti particles where found in the glassy matrix suggesting dissolution of these particle similar to that in casting practice. The layer does not only nucleate Al but appears to enhance cluster formation as suggested by its presence between boride particles.
Figure 2: A stoichiometric TiB2 particle embedded in glassy matrix (a) showing no nucleation of Al. and (b) a stoichiometric TiB2 particle with 0.3 at. % excess Ti showing a restored nucleation capacity for Al on basal faces
In figure 2(a) the effect of a lack of excess Ti on the nucleation mechanism of Al on boride particles is demonstrated. Stoichiometric addition of TiB2 particles did not result in copious nucleation of Al on basal faces of boride particles. This highlights that excess Ti is needed for the formation of a nucleating Al3Ti layer. In figure 2(b) stoichiometric boride particles at hyper-peritectic addition of 0.3 at. % excess Ti are embedded in an amorphous Al85Ni10Ce5 matrix. At hyper-peritectic addition levels the nucleation capacity of the boride particles is restored and nucleation of Al can be observed. The formation of a nucleating aluminide layer on hyper-peritectic addition of Ti suggests that the layer can be formed within the melt. However, for the given short melting times of 1 min the nucleation capacity of the boride particle with new layer is not restored to that of an effective boride particle of a 5:1 refiner rod seen in figure 1 suggesting that the layer may be present in the refiner addition. It is noteworthy, that the missing layer of Al3Ti on stoichiometric boride particles would only represent a few ppm of in commercial grain refinement added excess Ti with the remainder being available for growth restriction. This suggest a double role of excess Ti being responsible for the initial formation and/or stabilization of the Al3Ti nucleating layer and secondly for growth restriction of the nucleated Al crystals.
216 3.2
Effects of Zr on the Nucleation Mechanism
The effects of poisoning in Zr containing melts were simulated in Al85Ni10Ce5 glasses with Zr additions of 0.3 at. % Zr and 0.05 at. % excess Ti coming from 5:1 grain refining rod. Grain refiner addition and Zr additions where held at 1200°C for 1, 10 and 30 min before meltspinning. The processing temperature is somewhat higher than in conventional casting but is believed to be compensated by a lower atomic mobility in glass forming melts and shorter processing times than in conventional casting permitting a qualitative comparison.
Figure 3: (a) TiB2 particle embedded in a nominal Al85Ni10Ce5 glass with 0.3 at. % Zr addition held for 1 min at 1200°C showing copious nucleation of Al (arrow) and (b) TiB2 particle held for 30 min at 1200°C showing no nucleation of Al on basal faces
In figure 3(a) no significant difference in nucleation efficacy to that in figure 1 without Zr additions can be detected after holding the melt for 1 min before melt-spinning. Nucleation of Al on the TiB2 particle is copious on basal faces of the boride only. Faint streaks of the 0order spot in the SAD pattern, caused by the shape of the layer in reciprocal space along the TiB 2 direction, indicate the presence of a layer on these basal faces. In contrast, in figure 3(b) after 30 min of holding at 1200°C before melt-spinning no nucleation of Al can be detected on boride particles and no obvious streaks are apparent in SAD indicating the absence of a layer on boride particles. At intermediate holding times at 10 min, a mixture of boride particles with and without copious nucleation of Al can be found. This indicates that the surface properties of boride particles are affected by a time dependent reaction not affecting all boride particles simultaneously. Interestingly, Al3Ti and Al3Zr have very similar crystal structures DO22 and DO23 respectively and have limited solid solubility in each other [8]. In contrast to experiments where Ta was present [4], the absence of a layer at longer holding times suggests that Al3Ti has not been replaced by a layer of Al3Zr. Work on Al87Ni10Zr3 glasses simulating high hyper-peritectic Zr concentrations have shown that TiB2 particles are readily transformed into ZrB2 particles which do not provide sufficient lattice matching to support an aluminide layer [9]. It appears that at lower concentrations this
217 reaction occurs slower and not simultaneously, suggesting in conventional grain refinement of Zr containing alloys a sufficiently long time gap at lower casting temperature for sufficient grain refinement. However, further work is required to elucidate transformation kinetics of the temperature and time dependent Zr-poisoning by boride transformation.
4
Conclusions
Similar to previous studies using an Al-Y-Ni-Co glass [4,5], separate bulk Al3Ti particles dissolved as in commercial grain refining practice at hypo-peritectic addition levels and do not play a role as a nucleant. Instead, only TiB2 particles, coated with a previously identified thin layer of Al3Ti and visible as a streak in SAD pattern, nucleate copious amounts of Al. The nucleated Al crystals are somewhat smaller due to the higher undercoolings achieved in metallic glasses while in commercial casting a boride particle of 1 µm in diameter would be expected to be occupied by one Al crystal alone at approximately 1°C of undercooling. Some excess Ti is needed to form a layer of Al3Ti on boride particles, amounts beyond that will be available for growth restriction. Addition of Zr results in a temperature and time dependent reaction to transform TiB2 into ZrB2 which is preferred over the substitution of Zr in Al3Ti or the formation of Al3(Zr,Ti). Further work is required to understand the detailed mechanism of Zr- poisoning. The nucleation mechanism observed on coated boride particles appears not to be dependent on the glass-forming model alloy chosen and is consistent with layers of Al3Ti observed in grain refiner rods analysised in detail by TEM [10,11].
5
Acknowledgements
The authors gratefully acknowledge financial support from the EPSRC in conjunction with London & Scandinavian Metallurgical Co. Ltd. and Alcan International Ltd.
6 [1] [2] [3] [4]
References
D. G. McCartney, Int. Mater. Rev. 1989, 34, 247 - 260. P.S. Mohanty and J.E. Gruzleski, Acta Metall. Mater. 1995, 43, 2001 - 2012. J.A. Spittle and S. Sadli, Mater. Sci Technol. 1995, 11, 533 - 537. P. Schumacher and A.L. Greer, Light Metals 1995 (Ed.: J. Evans), TMS, Warrendale, PA, 869 - 877. [5] P. Schumacher et al., Mater. Sci. Technol. 1998, 14, 394 - 404. [6] R.F. Cochrane, P. Schumacher et al., Mater. Sci. Eng. A 1991, 133, 367 - 370. [7] P. Schumacher and A.L. Greer, Mater. Sci. Eng. A 1997, 226, 794 - 797. [8] S. Tsurekawa and M.E. Fine, Scr. Metall. 1982, 16, 391 - 395. [9] P. Schumacher, P. Cizek and A.M. Bunn, Light Metals 2000 (Ed.: R.D. Peterson), TMS, Warrendale, PA, 839 - 844. [10] B. Mckay, P. Cizek and P. Schumacher, Light metals 2000 (Ed.: R.D. Peterson), TMS, Warrendale, PA, 839 - 844. [11] P. Cizek and P. Schumacher, this conference.
Effect of Solute Elements on the Grain Structures of Al-Ti-B and Al-Ti-C Grain-Refined Al Alloys A. Tronche1,2 and A. L. Greer1 1
University of Cambridge, Cambridge, UK Péchiney Centre de Recherches de Voreppe, Voreppe, France
2
1
Abstract
The influence of solute elements on the structure of inoculated Al alloys, and the best parameter to represent solute effects, are assessed in this work. Experimental grain sizes were measured on Al-Ti-C grain-refined samples under TP-1 conditions [1], and results obtained in other studies [2] of Al-Ti-B refiners were analyzed also. The evolution of the structure with the solute content is found to be independent of the nature of the nucleation centers. The parameter Q=m(k−1)C0 [3,4] is found to represent better the effect of the solutes than the parameter P=Q/k also used in this type of study [2]. Nevertheless, when the solute alters the nucleation process, like Zr in the case of Al-Ti-B refiners [5], the growth-restriction parameter does not allow to predict the structure of the final product. It was shown in the case of Al-TiC refiners that, for example, the nucleation centers, i.e. the TiC particles [6], are poisoned by Si. The measured data were compared to results computed using the “free-growth” model [7].
2
Introduction
Grain refinement of Al and its alloys is a widespread industrial technique, which is currently mainly achieved by addition of Al-Ti-B master alloys. The final structure of the cast product is known to depend on the nature and number of the nucleation sites, and on the growth restriction imposed by the solutes present in the melt. The latter effect was extensively studied in the case of Al-Ti-B refiners [2,3,4]. Several parameters have been used to represent the effect of the solute elements; the constitutional-supercooling parameter P=m(k1)C0/k [2] (m, k and C0 being respectively the liquidus slope, the partition coefficient and the solute content of the alloy considered), the parameter Q=kP and the parameter U=DelementQ/DAl which takes account of the diffusion coefficient of the solute, Delement [8]. Even if solute diffusion is an important process during the solid growth, coefficients of diffusion are difficult to assess and thus the parameter U is not considered in this work. The usefulness of the parameters in representing the measured evolution of the grain size is assessed for Al-Ti-B inoculated alloys, and for the newly developed Al-Ti-C grain refiners. The experimental evolutions of the grain size as a function of the growth restriction are compared to computed data. Recently, a model for grain refinement, based on that of Maxwell and Hellawell [3], was developed and tested for Al-Ti-B grain refiners [7]. This model, applied in this work, is valid
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
219 only to predict the grain size of equiaxed structures. Other models can be used to predict whether the structure is equiaxed or not [9].
3
Experimental Procedure
The influence of various solutes on the grain structures of Al alloys inoculated with Al3wt%Ti-0.15wt%C-1wt%Fe (produced by LSM Co Ltd) has been studied. The solute elements considered correspond to those considered by Spittle and Sadli [2] and represent the main elements found in commercial wrought alloys. The alloys studied are listed in Table 1. The alloys are prepared from commercial-purity Al. The solutes are added either using commercial master alloys or as the elements. Table 1: Chemical compositions of the inoculated alloys studied (2 kg tonne−1 Al-3wt%Ti0.15wt%C-1wt%Fe) and corresponding P and Q growth-restriction parameters, using the data reported in earlier articles [7]. The chemical compositions are measured by optical emission spectroscopy. Solute content: P Q (in wt% except Ti in ppm) (K) (K) Cr Cu Fe Mg Si Ti Zn Zr Mn 1 0 0 0.17 0.98 0.03 64 0 0 0 24.0 5.2 2 0 0 0.16 0 0.03 73 0 0 0.52 20.3 2.4 3 0.032 0 0.19 0 0.03 69 0 0.178 0 20 2.0 4 0 0 0.16 0 0.24 76 0 0 0 26.95 3.0 5 0 0 0.15 5.02 0.03 63 0 0 0 47.4 15.8 6 0 0 0.17 0 0.02 75 3 0 0 22.7 3.4 7 0 0 0.21 0 0.02 71 4.84 0 0 27.6 5.1 8 0.08 0 0.21 0 0.04 68 0 0 0 21.9 0.98 9 0 0 0.17 0 0.5 69 0 0 0 40.6 4.94 10 0 0 0.17 0 0.77 63 0 0 0 53.7 6.51 11 0 0 0.17 0 1.05 63 0 0 0 67.3 8.14 12 0 0 0.17 0 1.7 65 0 0 0 98.8 11,9 13 0 0 0.17 0 2.76 71 0 0 0 150.3 18.1 14 0 0 0.17 0 3.95 69 0 0 0 208.1 25.0 15 0 0 0.17 0 5.11 69 0 0 0 264.4 31.8 The grain size is assessed in TP-1 tests under standard conditions. The grain refiner is held in the melt at 730°C for 5 minutes. The mean linear intercept is measured on the TP-1 examination planes [1]. Calculations are carried out using the “free-growth” model. Any particle of size d gives birth to a grain when the solid cap on the particle can freely grow. At this stage the cap and the particles have the same diameter. In such a model the largest particles are active first since the nucleation undercooling related to these particles is lower. The efficiency of the grain refiner is fixed by the onset of recalescence. The size distribution of the nucleation centers is entered in the program and at each undercooling the number of activated particles is calculated. The temperature evolution depends on the heat released by the nucleation and the growth processes. The model and the calculation algorithm are described elsewhere [7]. The calculations are carried out assuming that the cooling rate at the TP-1 examination plane equals 3.5 K s−1 [1]. The main input of the program is the size distribution of the nucleation
220 sites. The size of the particles is measured on polished sections and transformed into a 3D distribution using standard stereological techniques. This had been done previously for the TiB2 particles [7], and was done in this work for the TiC clusters found to be acting as nucleation centers in the Al-Ti-C refined Al alloys [6]. The model predicts the grain size only in equiaxed structures. The type of structure (equiaxed or columnar) can be predicted prior to computations [9]. In the test used in this work, the temperature gradient roughly equals 1 K mm−1, a low value favoring equiaxed growth.
4
Experimental Results
The evolutions of the measured and computed grain sizes of the Al-Ti-B refined samples are reported in Figure 1 as a function of P=m(k−1)C0/k and Q=kP. 1000
1000
600 C om puted G rain Size
400
800 Grain Size ( µ m )
Grain Size ( µ m )
Measured Grain Size
Measured Grain Size
800
600
200
200 0
C om puted Grain Size
400
0
0
20 40 60 80 100 G row th-R estriction Param eter, P (K)
0
5
10
15
20
Growth-R estriction Param eter, Q (K)
Figure 1: Evolution of the grain size of Al-5wt%Ti-1wt%B inoculated alloys (2 kg tonne−1) as a function of the growth-restriction parameters P and Q
The grain size of Al-Ti-B refined alloys decreases when the growth restriction imposed by the solutes increases, and reaches a plateau at high P or Q values. The data are less scattered around the trend line when the parameter Q is used. The study was repeated under Al-Ti-C inoculation and the evolution of the grain size is plotted as a function of Q in Figure 2. Measured G rain Size
Grain Size ( µ m )
1000 800
8
15 14 10
600 400 200 0
9
2 3 1 4 7 6
13
11 12 5
Predicted
0 5 10 15 20 25 30 35 Grow th-Restric tion Param eter, Q (K)
Com puted G rain Size (µ m )
1000
1200
800 600 400 1
200 0
5 0
6 2
10
11 12 4 3 14 6 9 15 13 7 200 400 600 800 M easure d Grain Size (µ m )
1000
Figure 2: Evolution of the grain size of Al-3wt%Ti-0.15wt%C-1wt%Fe inoculated alloys (2 kg tonne−1) as a function of the growth-restriction parameter Q. Predicted and measured sizes are also compared directly. The numbers reported on the graphs correspond to the numbers in Table 1 and represent the alloys
221 As observed with Al-Ti-B master alloys, the grain size first decreases sharply with the growth-restriction parameter and then reaches a plateau. The computed match fairly the measured data for any alloy except the Si-containing ones. Unexpectedly, the grain size of these Al-Si alloys increases when the growth restriction increases (alloys 9 to 15). The evolution of the grain size of the Al-Si alloys is reported as a function of the Si content in Figure 3. 900
Grain Size ( µ m )
800 700 600 500 400 300 0
1
2
3
4
5
6
7
8
Si Content (wt% )
Figure 3: Evolution of the grain size of Al-3wt%Ti-0.15wt%C-1wt%Fe inoculated Al-Si alloys (2 kg tonne−1) as a function of the Si content
The evolution can be divided into two parts: for Si content < 3 wt%, the grain size increases slowly, while it increases quickly above this value.
5
Discussion
The evolution of the grain size with the growth restriction imposed by the solutes is independent of the representing parameter considered, P or Q, and is the same for the Al-Ti-B and Al-Ti-C grain refiners. The grain size first decreases sharply with the solute content, and then reaches a plateau. Under the experimental conditions (2 kg tonne−1 of Al-5wt%Ti1wt%B or Al-3wt%Ti-0.15wt%-C-1wt%Fe) the value of the grain size on the plateau is smaller for the B-containing master alloy. In this refiner, the higher particle volume fraction compared to the Al-Ti-C master alloy leads to a higher number of nucleation sites, despite the TiB2 particles being larger, and thus leads to the finer grain size. The trends of the measured and the computed grain sizes are clearer when the parameter Q is used to represent the effect of the solutes. In the phase diagram of the relevant solute at C0 solute content, P is the equilibrium freezing range. It represents the operating condition at the solid-liquid interface when the solid composition is the far-field composition. This case occurs under steady state planar-growth conditions, and is not related to the equiaxed growth problem. Q on the other hand represents the variation of the fraction solid fs with temperature at the liquidus. From the the Scheil-Gulliver approximation it follows that [10]: df s dT
= Tliq
1 1 = mC 0 (k − 1) Q
(1)
In the melt undergoing solidification, heat is released by the nucleating and growing solid. The rate of solid formation dfs/dt, and thus the rate of heat released, is according to Eq. (1)
222 related to the growth restriction factor Q: a higher Q reduces the absolute value of dfs/dT and in consequence reduces the rate of latent heat released. The recalescence is thus delayed allowing more grains to be nucleated. Thus under equiaxed growth, the number of grains which can nucleate, i.e. the final grain size, seems to be directly related to the parameter Q. As shown in Figure 1 and 2, it is possible from the value of Q to predict the grain size of the cast product, which is well modeled by the “free-growth” model as shown in Figure 2 for Al-Ti-C. The model fails to predict the grain size when the poisoning of the nucleation centers occurs, as it happens in Al-Ti-C refinement of Al-Si alloys. The evolution of the grain size of Al-Si alloys with the Si content reveals two stages. Microstructural observations, linked to thermodynamic considerations, reveal that for Si < 3 wt% the TiC clusters transform into particles rich in Al, Si & C. This dissolution/precipitation reaction is, like the same transformation of TiC into Al4C3 particles, sluggish [11]. Above 3 wt% Si, the particles transform into SiC. This reaction occurs quickly, leading to a complete disappearance of the nucleation sites, i.e. complete loss of the grain-refining effect.
6
Conclusions
Solute elements are of great importance in grain refinement. They restrict the solid growth since they need to partition between the solid and the liquid. The final grain size of a cast product largely depends on the growth restriction imposed by the solutes, which is well represented by the parameter Q under equiaxed growth. In addition, solute elements facilitate the columnar-to-equiaxed transition. Nevertheless, solute elements can also act during the nucleation stage either by destroying the nucleation centers (poisoning of TiB2 by Zr, of TiC by Si as shown in this study), or by enhancing the nucleation potency of the particles (formation of an Al3Ti layer on TiB2 particles) [12]. It was shown in this work that the effect of the solute elements can be well predicted by the “free-growth” model for both Al-Ti-B and Al-Ti-C refiners unless poisoning of the nucleation centers occurs.
7
Acknowledgements
AT acknowledges financial support from the LSM Co Ltd and Péchiney (CIFRE studentship), and useful discussions with A. Hardman and P. S. Cooper at LSM, with P. Jarry at Péchiney Centre de Recherches de Voreppe, and with M. Vandyoussefi at the University of Cambridge.
8
References
[1] Aluminum Association Standard Test Procedure for Aluminum Grain Refiners, The Aluminum Association, Washington D.C. 20006, 1987. [2] J. A. Spittle, S. Sadli, Mater. Sci. Technol., 1995, 11, 533 - 537. [3] Maxwell, A. Hellawell, Acta Metall., 1975, 23, 229 - 237.
223 [4] L. Backerud, M. Johnsson, Light Metals 1996, (Ed.: W. Hale), TMS, Warrendale, PA, 1996, 679 - 685. [5] J. A. Spittle and S. Sadli, Cast Metals, 1994, 247 - 253. [6] Tronche, A. L. Greer, submitted to Phil. Mag. Letters. [7] L. Greer, A. M. Bunn, A. Tronche, P. V. Evans, D. J. Bristow, Acta Mater., 2000, 48, 2823 - 2835. [8] P. Desnain, Y. Fautrelle, J. L. Meyer, J. P. Riquet, F. Durand, Acta Metall., 1990, 38, 1513 - 1523. [9] M. Vandyoussefi , A. L. Greer, these proceedings. [10] W. Kurz, D. J. Fisher, The Fundamentals of Solidification, 3rd edition, Trans Tech Publications, 1992, p. 285. [11] Tronche , A. L. Greer, Work in progress. [12] P. Schumacher, A. L. Greer, J. Worth, P. V. Evans, M. A. Kearns, P. Fisher, A. H. Green, Mater. Sci. Technol., 14, 1998, 394 - 404.
Grain Refinement Process in Aluminium Alloys Type AlZnMgZr 7RPDV]6WXF]\ VNLDQG0DU]HQD/HFK*UHJD Institute of Non-Ferrous Metals, /LJKW0HWDOV'LYLVLRQXO3LáVXGVNLHJR6NDZLQD3RODQG
1
Abstract
The results of grain refining process in the aluminium AlZnMgZr alloys containing up to 0.20% wgt. of Zr are presented. In this experiences grain refiners type AlTiB and AlTiC have been used. Investigation have shown that grain refiners type AlTiC in more resistant to poissoning by Zr and under defined casting condition will be have efficient grain refiner than AlTiB. The metallographic studies have shown that intermetallic phases type Al3(ZrxTi1-x) have occurred into microstructure. This fact would be useful to explain of Zr poisson effect during grain refineme4nt process. The laboratory scale results have been confirmed on the industry plant.
2
Introduction
It is commonly known that the grain refining process has been used in preparation of liquid aluminium and its alloys for almost fifty years, i.e. since the first paper by A. Cibula, entitled: „The Mechanism of Grain Refinement of Sand Castings in Aluminium Alloys”, published in the Journal Institutes of Metals in 1949 [1]. Despite such a long practice and almost continuous studies of the process carried out by all research centres related to the aluminium casting, no uniform and watertight theory of the grain refinement in aluminum alloys has been presented yet. Many phenomena occurring in the industrial practice prove that a comprehensive control of this process is still a distant objective. A good example of this situation may be the process of grain refining in the alloys made on the basis of aluminum and containing boron in the form of AlB2, as the so called „excess boron” [2] and in the alloys, containing Zr or Cr as the alloy additive. This situation have also refered to common aluminium alloys such as siluminum The literature data [3-16] indicate the „poisoning effect” of Zr in the grain refining processes, carried out by adding to the liquid metal the controlled amounts of Ti and B, in the form of the inter-metallic compounds, Al3Ti and TiB2, using the AlTiB master alloys. The term „poisoning effect”, describing the phenomenon of perturbation of the grain refinement process by some elements present in the alloy is commonly used in the papers devoted to grain refinement in aluminium alloys. This term is not precisely defined however. Jones and Pearson [17] use this term when the size of grains obtained after the addition of grain refiner is considerably larger in the presence of some elements, e.g. Zr, than in their absence. This means that the „poisoning element” reduces the efficiency of the grain refinement process. The presence of a „poisoning element” may cause an earlier inhibition of
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
225 the grain refinement process, which may be observed as an increase of the grain size above the acceptable level after a short time of maturing of the metal bath. Another symptom of poisoning may be an incomplete change of the columnar structure to the equi-axial one, following the addition of the grain refiner. Nowadays, there are 4 hypotheses attempting to explain the poisoning of the grain refinement by Zr, in the presence of AlTiB as the grain refiner. The hypothesis put forward by Jones and Pearson [17] assumes that the active particles of TiB2 become covered with a layer of the inter-metallic compound ZrB2, which reduces their activity during the grain refinement. This hypothesis has not been confirmed, because ZrB2 has not been found yet. On the other hand, Abdel-Hamid [18] assumes that some portion of Ti in the inter-metallic compounds, TiAl3 and TiB2, which plays an active role in grain refinement, is replaced by such elements as Zr or Cr, forming the inactive tri-component compounds, passive in the grain refinement. According to Johnsson [19], the poisoning effect of Zr is related to Ti dissolution in the inter-metallic ZrAl3 compounds. A complex (Ti1-xZrx)Al3 is formed, thus reducing Ti activity in grain refinement. A. Arjuna [20] suggests that the poisoning effect on the grain refinement is caused by the formation of the complex compounds of ZrAl3, Ti and Fe, Al3Ti, containing Zr and Fe and Al3(Fe Ti Zr), which inactivates the particles potentially active in the grain refinement processes. Between 1997 and 1999, research was conducted at the Light Metal Department of the Non-Ferrous Metals Institute concerning the process of grain refining in high-strength aluminium alloys of the AlZnMg and AlCuMg type with a Zirconium addition up to 2% of the weight. [21] The research aimed on the one hand at analysing the mechanisms of the "poisoning effect" brought about by the introduction of Zirconium into aluminium alloys and its influence on the reduction or elimination of the refining processes arrived at by the use of AlTiB master alloys, and on the other at researching ways of neutralising the negative influences of the Zirconium addition on the grain refining process. The conducted analysis brought about positive results, which allow the following theses to be posited: • a Zirconium addition up to 2% of the weight of the molten metal does not bring about a grain refining effect in the solidification process of a given type of aluminium alloy; • the presence of Zirconium in aluminium alloys does interfere with grain refining processes with the use of classical AlTiB grain refiners; • the use of AlTiC master alloys brings about the stabilisation of grain refining processes in aluminium alloys with the addition of Zirconium; • the "poisoning effect" brought about by the presence of Zirconium consists in the creation of complex compounds Al3(ZrxTl1-x), whose presence was detected in the analysed alloys, and which lead to the distortion of the optimum Ti-B proportion which is indispensable for the proper conduction of the grain refining process. The results corroborate the Johnson hypothesis in this case. It means that the formation of complex compounds Al3 (Zr, Ti, Fe) reduces the number of active Al3Ti particles, which together with TiB2 compounds are necessary for the effective grain refinement. At the current state of our knowledge, in conformity with the nucleus theory of Cibula [1] it is assumed that due to the close lattice parameters of TiCx compounds to the α- Al phase, AlTiC is a direct
226 nucleus, as opposed to the TiAl3 phase, which is activated by the presence of TiB2. As the research data show, the necessary condition for a proper grain refinement using AlTiB is the presence of both latter phases in the appropriate proportions. These results of tests conducted in laboratory conditions and in semi-technical conditions provided a basis for the successive stage of the research, i.e. applied tests conducted in an industrial environment. The present paper will discuss the results of those and the conclusions reached in the process of applying the AlTiC master alloy as the grain refiner in the process of semi-continuous production of circular φ220 mm ingots of the AlZn4.5Mg1.3ZrCr alloy.
3
Test Conduction
In the existing industrial environment, the preparation and casting of aluminium alloy ingots with diameters exceeding φ200 mm is conducted in the following manner. The melting and alloying processes are conducted in an induction channel furnace, into which a grain refiner in the form of nuglets is introduced before the pouring-in is started, as required by the existing technologic procedures. After the metal is melted and the metal bath is confirmed to have reached the chemical composition required by the factory norm, the molten metal is poured into a resistance holding furnace.
induction furnace
holding furnace
filtr
semi-cont cast unit
Figure 1: Scheme of casting unit
Then the metal is transported from the holding furnace through a foam ceramic filter of the 40ppi type to a number of DC type mould. Ingots are cast according to the following parameters: • ingot diameter - φ220 mm • number of ingots cast – 8 • casting temperature – 707°C • casting speed – 60 mm/ min • water amount 48 cu. m./ h Typical macrostructure of AlZn4.5Mg1.3ZrCr ingots produced according to the above procedure is presented in Fig. 2.
227
Figure 2: Macrostructure of AlZn4.5Mg1.3ZrCr alloy ingot
Figure 3: A schema of sample selection for microstructure analysis
Figure 4 present the microstructures detected by Barker solution etching of samples selected from the particular macros according to the schema presented in Fig. 3.
surface
½ radius
Center Figure 4: Grain in ingot
In order to improve the quality of ingots cast from AlZn4.5Mg1.3ZrCr alloy, i.e. to reduce the grain size and eliminate twinned grain, it was decided to apply AlTiC master alloy as the grain refiner in the process of producing this alloy.
228 In order to optimise the grain refining process with the use of the AlTiC master alloy, it was decided to conduct the experiment in three different variations. Variation I assumed the application of the AlTi3C0.15 master alloy in the quantity sufficient for the introduction of 100 ppm Ti, which was melted into the metal bath in an induction melting furnace. In Variation II, the same amount of the grain refiner was introduced into the molten metal stream during the pouring of the metal bath from the melting furnace to the maturation furnace. In Variation III, the AlTi3C0.15 grain refiner was introduced into the molten metal stream, as in Variation II, but its amount was increased to 150 ppm Ti. The casting process parameters remained intact and reflected the existing technological procedure regarding etching of controlled melts. Examples of macrostructures of ingots cast during controlled melts in the three variations are presented in Fig. 5. Figures 6-8 present the micostructures of the samples separated from the macros (Fig. 5), and selected according to the schema in Fig. 3 The samples were etched by Barker solution and observed in polarised light in order to detect grain distribution and size in the analysed microstructure.
variant I
variant II
variant III Figure 5: Mactrostructure of AlZn4.4Mg1.3Zr0.17Cr0.15 alloy ingot with AlTiC as grain refiner
229
850 µm Surface
850 µm ½ radius
850 µm Center Figure 6: Grain in ingot - variant I
850 µm
Surface
850 µm
½ radius
850 µm center Figure 7: Grain in ingot – variant II
230
850 µm Surface
850 µm ½ radius
850 µm
center Figure 8: Grain in ingot – variant III
In the course of the holding of the molten metal and its casting, molten metal samples were taken according to the TP-1 test procedure adapted for the purpose. Examples of sample macrostructures from the TP-1 test and the macrostructure typical for the entire experiment are presented in Fig. 9.
Figure 9: Mactrostructure of Test TP1 samples
231 This stage of the analysis aimed at analysing the possible effects of grain refining reduction during the lengthy process of preparing and casting the molten metal.
4
Results
The results obtained during the analysis of AlZn4.5Mg1.3Zr0.17Cr0.15 ingot production, conducted in an industrial environment, reveals that the application the AlTiC type grain refiner in the process of producing alloys with a Zr addition assures the obtaining of properly refined grain throughout the length of the cast ingot. These results fully confirm the results obtained in laboratory conditions and in semi-technical conditions. The analysis also demonstrated that the existing procedure of introducing the grain refiner to the metal bath in the induction channel furnace is improper. This finding is demonstrated by the comparison of the macrostructure of the ingots analysed in Variations I and II of the experiment described here. The same amount of grain refiner is observed to be much more efficient when introduced into the molten metal stream during its pouring from the melting furnace to the holding furnace. This is demonstrated by the total elimination of the twin crystals growth zone in the ingots (the comparison of the microstructures presented in Fig. Fig. 6 and 7 above), which were produced from the molten metal produced according to the Variation II procedure. This finding was also confirmed for the preparation and casting processes of other alloys. It has to be emphasised, however, that the change concerns the place where the refiner is introduced, rather than its composition. It was also found that there is no rationale for the use of higher amounts of AlTiC master alloy to be introduced. The comparison of microstructures (Fig. 7 and Fig. 8) does not reveal considerable differences in grain distribution and size when 100 and 150 ppm Ti are revealed by the AlTi3C0.15 master alloy. The analysis of the samples taken with the use of the TP-1 Test procedure (Fig. 9) manifests that the grain refining effect obtained thanks to the use of the AlTiC master alloy is not reduced even when the time span between the introduction of the master alloy to the casting of the last part of the molten metal exceeds 200 minutes. In summary, we can conclude that the analysis of AlZn4.5Mg1.3Zr0.17Cr0.15 molten metal production in an industrial environment fully corroborate the efficiency of the application of the AlTiC type grain refiner for this type of alloys. The research also brought about a modification of the existing procedures of molten metal preparation and the casting of aluminium alloy ingots with the addition of Zirconium.
5
Conclusion
1. Grain refining in aluminium alloys with the addition of Zirconium should be conducted with the application of the AlTiC master alloy, which was confirmed by analyses in semitechnical and industrial conditions. 2. If the grain refining process is conducted properly as far as the place of application and the amount of the AlTiC master alloy is concerned, homogeneous fine-grain structure can
232 be reproduced regularly throughout the cross section of the aluminium alloy ingot with the addition of Zirconium, and the twin crystal zone can be eliminated.
6
References
[1] Cibula, The Mechanism of Grain Refinement of Sand Castings in Aluminium Alloys Journal Institutes of Metals, 1949 [2] - :R QLFND 7 6WXF]\ VNL 5R]GUDEQLDQLH ]LDUQD Z VWRSDFK Z\WZDU]DQ\FK QD ED]LH aluminium oSRG\Z V]RQHM ]DZDUWR FL ERUX 0DWHULDá\ .RQIHUHQF\MQH Ä$/80,1,80 98”, Zakopane 1998 [3] S. R. Thistlethwaite, Recent Developments in Grain Refiner Technology, 4th Australasian Asian Pacific and Conference Aluminium Cast House Technology, The Minerals, Metals & Materials Society, 1995 [4] M. Johnsson, Influence of Zr on the Grain Refinement of Aluminium Carl Hanser Verlag, München 1994 [5] )DUULRU '& %ULOOKDUW $ 3UDFWLFDO 0HWKRG IRU (YDOXDWLQJ *UDLQ 5HILQHPHQW 0DWHULDá\ firmy Kawecki Berylco [6] P.S. Mohanty, J.E. Gruzleski, Mechanism of Grain Refinement in Aluminium Acta metall. mater. Vol. 43, no 5 pp. 2001-2012, 1995 [7] P. Hoefs, W. Reif, Recent Developments in Grain Refining of Aluminium and Aluminium Alloys Solidification of Metals and Alloys, No 28, 1996 [8] R. Cook, P.S. Cooper ,Benefits of Master Alloy Melt Treatments in the Aluminium Foundry Industry Light Metals 1996 [9] R. By, A.P. Fielding, Recent Developments in Grain Refining Technology & Standardization Light Metals Age, June 1997 [10] J.A. Spittle, S.B. Sadli, The Influence of Zirconium and Chromium an the Grain Refining Efficiency of Al-Ti-B Inoculants Cast Metals, V. 7, 1994 pp 247-253 [11] M.A. Kearns, P.S. Cooper, Effects of Solute Interactions on Grain Refinement of Commercial Aluminium Alloys Light Metals Age, June 1997 [12] P.C. van Wiggen Al-Ti-B Grain, Refner - the Consistent Ingredient Light Metals Age, June 1997 [13] P. Hoefs, W. Reif, Development of an Improved AlTiC Master Alloy for the Grain Refinement of Aluminium Light Metals Age, June 1997 [14] M.A. Hadia, A.A. Ghaneya, Development and Evaluation of Al-Ti-C master Alloys as Grain Refiner for Aluminium Light Metals 1996 [15] M.A. Kearns S.R. Thistlethwaite,Recent Advances in Understanding the Mechanism of Aluminium Grain Refinement by TiBAl Master Alloys Light Metals 1996 [16] 36&RRSHU0DWHULDá\.RQIHUHQF\MQHPDU]HF706&RQIHUHQFH6DQ$QWRQLR [17] G.P. Jones, J. Pearson, Metall Trans B 1976, 7B, 223-234 [18] A.A. Abdel-Hamid, Z. Metallkd 80, 1989, 566-569 [19] M. Johnsson, Z. Metallkd 85, 1994, 786-789 [20] Arjuna Rao, Materials Science and Technology, Sept. 1997, vol. 13 769-777 [21] 76WXF]\ VNL0/HFK*UHJD7KH*UDLQ5HILQLQJ3URFHVVLQ+LJK6WUHQJWK$OXPLQLXP Alloys Containing Addition of Zr, Proceedings of the International Conference "Light Alloys and Composites", 13-16 May 1999, Zakopane, Poland
Coupled Influence of Convection and Grain-refining on Macrosegregation of 1D Upwardly Solidified Al 4.5%Cu Ph. Jarry1, H. Combeau2 and G. Lesoult2 1
Pechiney Centre de Recherches de Voreppe, France LSG2M, Ecole des Mines de Nancy, France
2
1
Abstract
It has been shown in previous studies that even in the simple configuration of a 1D upward solidification set-up, inverse macro-segregation is influenced by convective phenomena. In order to better identify the respective contributions of dendrite settling, dendrite transport and thermo-solutal convection to inverse segregation in the vicinity of the chill, a comparison has been made between Al4,5%Cu ingots obtained with or without grain refinement, and solidified in a 1D upward solidification set-up either located in the laboratory at room temperature, or within a heated oven in order to inhibit convection. The macrosegregation results evidence the combined effects on the macro-segregation pattern, of dendrite mobility, as governed by the grain refinement and Ti concentration, and of convective behavior, as induced by the temperature around the mold.
2
Introduction
In order to get a better understanding of the parameters involved in inverse segregation buildup of 1-D solidified aluminum ingots, a few original experiments have been designed. Since it has been demonstrated [1] that the interaction of convection with equiaxed grains does bring a contribution to the concentration profile, it was natural to try to strongly vary the parameters governing grain mobility on the one hand, and convection on the other hand.
3
Experimental Procedure
In order to do so, a 1-D solidification set-up was used which allows to bottom-cool the liquid metal contained in an insulated cylindrical mold (50mm in diameter, 150mm in height), itself placed inside an oven maintained at a chosen temperature. The chill is a water-cooled block of copper and the cooling water is injected from outside the oven into the block. In order to act upon the formation and morphology of equiaxed grains, the grain refinement was varied between 0 and 3 kg/T of Al5%Ti1%B, and in some cases a hyper-peritectic quantity of Ti (0.2%) was introduced to favor the growth of Al3Ti at the liquidus and hence the settling of grains of solid solution of aluminum grown over those primary intermetallic particles. Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
234 Classical optical metallography was performed. The Cu concentration profiles were measured along the central vertical axis of the ingots on longitudinal polished sections by analytical scanning electron microscopy (ASEM) between 15 and 20kV accelerating voltage; for each position along the axis of the ingots, a 200µmx200µm square was scanned. The complete procedure is described in [1].
4
Results
4.1
Terminology
In the following, ingots1D-solidified in molds located in the laboratory at room temperature will be referred to as “RT-solidified”, whereas the ingots solidified in the 1D set-up placed within a heated oven will be referred to as “oven-solidified”. 4.2
Grain Structure and Sizes
The grain structures obtained are summed up in Table 1. Table 1. List of ingots studied with solidification conditions and type of as-cast structure Sample Grain refining T around Grain structure reference (with Al 5Ti 1B) mold during freezing # 18 3 kg/ton + 0.20%Ti 750° Equiaxed (400µm close to the chill) # 19 not inoculated 680°C Coarse-grained equiaxed (2500µm) # 20 3 kg/ton + 0.20%Ti 680°C Equiaxed (600µm close to the chill) #1 not inoculated 20°C Columnar then coarse-grained equiaxed #3 0,1kg/ton 20°C Coarse-grained equiaxed #5 0,3kg/ton 20°C Fine grained equiaxed #6 0,3kg/ton 20°C Fine grained equiaxed #7 3kg/ton 20°C Fine grained equiaxed #9 3kg/ton + 0.2% Ti 20°C Fine grained equiaxed 4.3
Copper Concentration Profiles
As previously described [1,2] a continuously decreasing Cu concentration over the length of the ingot (except sometimes at the upper end) is obtained in the case of the grain refined Al4,5Cu ingots solidified under usual conditions (environment at room temperature), see Figure 1a. A prolonged positive segregation is observed up to 50 to 75mm from the chill. A concentration plateau at (or near) the nominal Cu concentration is obtained in the well known case of columnar growth (ingot 1), or in the case of the very little grain-refined (0,1kg/ton) ingot #3 with coarse equiaxed grains. A striking feature is that a plateau is also obtained in ingots solidified within a heated oven regardless of the grain refinement, see Figure 1b. In this case the extension of the zone with positive segregation on the chill side is typically 25mm. The samples with hyperperitectic Ti follow the same trend except when solidified in a mold at room temperature: the Al4,5%0.2%Ti sample exhibits a more complex
235 concentration profile, see Figure 1a: after a very similar behavior up to 25mm, a positive plateau is obtained followed by a rather sharp drop in concentration. 0.2
0.1 0.05 0 0
25
50
75
100
125
ingot 19 : 0kg/T 680°C ingot 20 : 0.2Ti 3kg/T 680°C
0.15
150
-0.05 -0.1
relative segregation
0.15 relative segregation
0.2
ingot 6 : 0.3kg/T 20°C ingot 5 : 0.3kg/T 20°C ingot 7 : 3kg/T 20°C ingot 9 : 0.2Ti 3kg/T 20°C
ingot 18 : 0.2Ti 3kg/T 750°C ingot 1 : 0kg/T 20°C
0.1
ingot 3 : 0.1 kg/T 20°C 0.05 distance from the chill (mm)
0 0
25
50
75
100
125
150
-0.05 -0.1
-0.15
-0.15
distance from chill (mm)
Figure 1: Copper concentration profiles along the vertical axis of the ingots; a/ profiles obtained for equiaxed Al4.5%Cu samples in molds at room temperature; b/ profiles obtained for samples with no or little inoculation at room temperature or for “oven solidified” samples, whatever the inoculation.
4.4
Comparison of Macrosegregation Intensities
The compared intensities of macrosegregation in the different ingots are summed up in Figure 2, showing the average of absolute values of relative segregation in copper measured between 12mm and 137mm from the chill. These limits have been chosen in order to eliminate both the zone affected by chill inverse macrosegregation, where the concentrations measured are very scattered, due to the discrete detection of the presence of constituent phases, and a symmetrical zone at the top of the ingots affected by shrinkage porosity. This bar graph clearly shows that equiaxed samples solidified under usual conditions are twice as segregated as either non-inoculated samples or grain refined samples solidified in molds held at high temperature. ingot 6 : 0.3kg/T 20°C ingot 7 : 3kg/T 20°C ingot 5 : 0.3kg/T 20°C ingot 9 : 0.2Ti 3kg/T 20°C ingot 3 : 0.1 kg/T 20°C ingot 20 : 0.2Ti 3kg/T 680°C ingot 19 : 0kg/T 680°C ingot 18 : 0.2Ti 3kg/T 750°C ingot 1 : 0kg/T 20°C 0.01
0.015
0.02
0.025
0.03
0.035
0.04
0.045
Figure 2: Average of the absolute values of relative segregation measured between 12 and 137mm from the chill for the different ingots. Striped bars illustrate the influence of convection on the macrosegregation behavior of the same highly grain refined alloy. The alloy with 0.1kg/T grain refiner exhibits an intermediate behavior, because it is equiaxed but with a coarse grained structure that reduced the mobility of dendrites.
236
5
Discussion
5.1
Driving Forces at Work
Once the well-known chill inverse segregation transient is finished, a steady state regime could ideally take place in which liquid migration towards the cold isotherms would exactly compensate for volumetric contraction, in such a way that the resulting mean Cu concentration would equal the nominal concentration. Now, any mechanism that can either bring more grains to the front, take more grains from the front, or modify the flowing regime of the eutectic-enriched liquid near the mushy zone will certainly alter the Cu concentration profile of the solidified ingots. A very likely possibility stemming from [1,2] is that thermosolutal convection movements suspend crystals in the undercooled liquid, removing them from the front or preventing them from settling onto the front, a mechanism likely to prolong solidification at a Cu concentration higher than the nominal one. This mechanism is very likely to occur in the case of grain refined ingots solidified under usual conditions and exhibiting a steadily decreasing Cu concentration starting from the chill. A good clue that this mechanism actually brings such a contribution, is that grain refined ingots solidified within an oven, where convection movements are inhibited, exhibit the concentration plateau usually known to form in columnar solidification. Potentially mobile equiaxed dendrites are no longer prevented from sitting on the solidification front. In the case of Al4.5Cu0.2Ti, dendrite transport is probably enhanced by the precipitation of Al37LDQGDVVRFLDWHGSULPDU\ $OXPLQXP 5.2
Dendrite Mobility
Dendrites respond to convective driving forces through their mobility, itself linked to grain refinement. Grain refinabiliy [4] can be understood through the notion of gradient of solid fraction fs near the liquidus which is equal to the product of dfs/dT at the liquidus temperature by the thermal gradient. The increase in solid fraction when temperature decreases under the liquidus can be derived from the Gulliver-Scheil law for linear approximations of the liquidus lines of the different binary alloys corresponding to each alloying element: df − s dT
= Tliquidus
∑
1 m i (k i − 1)C oi
where mi is the liquidus slope at the liquidus temperature in each binary alloy Al-i, ki the corresponding partition coefficient, and Coi the nominal concentration of element i. Dendrite mobility can thus be expressed by: df µ d = G th. s dT
liquidus
−1
=
∑ m (k i
i
− 1)C oi
G th
It is well known [3] that among the different alloying elements of aluminium, Titanium plays a very special role, linked with the high value of the product m(k-1) for Ti in Al: m(k1)§. The weaker the increase in solid fraction near the liquidus, the easier the nucleation and growth of equiaxed grains. Now, if both the grain refinability is high and the temperature gradient is low close to the front, nucleation and growth have room to occur in the undercooled region and convection movements can affect the suspended crystals.
237 Table 2: Compendium of mechanisms accounting for the observed profiles. Grain refined Non or little grain refined Mold at RT Mold in oven Mold at RT Mold in oven Convective driving force Yes Little or none Yes Little or none Dendrite mobility Suspension of dendrites
High
Settling of dendrites Segregation index*
Mediocre (high gradient)
Low (poor G R)
Low
None or little
No
No
Yes, at the end; Al3Ti may play a role
None or little
May play a role after C.E.T.
No
§
§
§
§
Yes, at the beginning
* Average of absolute values of relative segregation between 12 and 137mm
On the other hand, under too high temperature gradient conditions, no room is available for nucleation on a large number of sites, or for a large number of equiaxed grains to grow: fewer grains, thus larger ones, are formed, of impaired mobility. As a consequence, by using a mould held at high temperature, not only the driving force for natural convection is decreased or suppressed, but also the dendrite mobility is decreased because the temperature gradient is strongly increased at the front; it is true even for the highly refinable alloys with high Titanium concentration and strong inoculation. This framework allows qualitative interpretation of macrosegregation results, see Table 2.
6
Conclusion
The influence of dendrite movement on macrosegregation profiles in upwardly solidified ingots of Al4.5%Cu is confirmed. Dendrite mobility is boosted by grain refinement and Ti content but is inhibited by too high temperature gradients. When dendrite mobility and convection can have a coupled action, a prolonged positive segregation starting from the chill is observed, steadily decreasing over the length of the ingots, then turning into a negative segregation at around mid-height of the ingots. Suppression of dendrite mobility or of the convective driving force results into a macrosegregation plateau type profile typical of what is observed or calculated for columnar growth, even with grain refined samples. As similar phenomena operate within the far more complex context of vertical continuous casting, the presented results are believed to help provide a reference of mechanisms for qualitatively interpreting the macrosegregation patterns obtained in DC Casting.
238
7
Acknowledgments
The technical contribution of J.-M. Burtin and V. Chastagnier, Pechiney Centre de Recherches de Voreppe, is gratefully acknowledged. Part of this work was conducted within the framework of Brite Euram III Project BE-1112 (European Modeling Program on Aluminium Casting Technology).
8
References
[1] V. Albert, Ph.Jarry, H. Combeau, and G. Lesoult, “Influence of the grain structure on macrosegregation in Al-Cu ingots obtained by upward 1-D solidification,” presented at SP'97, Sheffield, 1997. [2] V. Albert, “PhD work: Macroségrégation et mouvement des cristaux équiaxes lors de la solidification d'alliages,” Ecole des Mines de Nancy, Institut National Polytechnique de Lorraine, 1998. [3] J. Moriceau, “Discussion des mécanismes d'affinage de l'aluminium par le titane et le bore,” Revue de l'aluminium, pp. 977-988, 1972. [4] P. Desnain, Y. Fautrelle, J. L. Meyer, J. P. Riquet, and F. Durand, “Prediction of equiaxed grain density in multicomponent alloys stirred electromagnetically,” Acta Metallurgica, vol. 38, pp. 1513-1523, 1990.
Tensile Behaviour of DC-cast AA5182 in Solid and Semi-solid State W.M. van Haaften, W.H. Kool, L. Katgerman Laboratory of Materials, Delft University of Technology, Rotterdamseweg 137, 2628 AL, Delft, The Netherlands
1
Abstract
Hot tearing is still one of the major problems during DC casting of aluminium alloys. It may result in cracks running through the entire ingot, which require costly remelting and cause loss of productivity. Numerical models are developed to simulate and improve the casting process but there is a lack of mechanical data of the alloys, especially at elevated temperatures. Therefore tensile tests were carried out with AA5182 at high temperature including the semisolid range. A mathematical description was found for the subsolidus temperatures and this description could be extended to above-solidus temperatures by assuming that the constitutive behaviour is determined by the solid network. This was done by considering not only the fraction liquid but also the percentage of grain boundary area covered by this liquid phase. Taking this approach, the mechanical behaviour of the alloys at semi-solid temperature can be described, at least for the low liquid fractions during which hot tearing occurs.
2
Introduction
Nowadays, a wide variety of aluminium alloys is being cast on a routine base by means of the DC casting process. However, in search for a better corrosion resistance or a higher strength for example, new compositions are being developed which may impose strict constraints on the process window. Not only the process parameters such as casting speed, casting temperature, cooling rate should be controlled accurately, also a profound knowledge of thermomechanical behaviour of the alloy during the casting process is required to produce a defect free slab or billet. This study focuses on this thermomechanical behaviour at high and even semi-solid temperatures. The main defect which occurs at these high temperatures during casting is hot tearing. Hot tears are cracks which initiate during the last stage of solidification, i.e. in the mushy zone, where solid and liquid co-exist. The hot tearing problem has since long been investigated [1-5] and many cracking criteria have been defined. From these studies it has become clear that the two key parameters are 1) inadequate interdendritic feeding and 2) stress and strain development in the dendritic network. The aim of this research is to determine the stresses that can be supported by the solid network at different liquid fractions and to find a mathematical description for the alloy behaviour in this semisolid state. As a starting point we will use the equation often applied to describe solid state creep:
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
240 Q (1) ε& = Aσ n exp − RT where ε& is the plastic strain rate, A is a constant, σ is the true stress, n is the stress exponent,
Q is the apparent creep activation energy, R is the universal gas constant and T is the temperature.
3
Experimental
The material investigated is an AA5182 alloy with the following composition: Mg: 3.6, Mn: 0.16, Si: 0.21, Fe: 0.26, Al: bal. (wt%). It is sampled from a rolling slab which was DC cast at an industrial research facility [6]. The slab did not receive any additional heat treatment after casting. The tensile tests are carried out with a Gleeble 3500 thermomechanical simulator. The specimens with the same shape as used earlier [7] are taken from the slices with their tensile direction parallel to the casting direction. They are secured between two water-cooled jaws of which one is moving. The specimen is heated at 50°C/s via Joule heating during which it is kept at zero force. The test temperatures are from 370°C to 580°C. The solidus temperature of the material is 510°C and the maximum liquid fraction, calculated with the Alstruc model is 0.07 [8,9]. True strain εp is measured with a dilatometer which monitors the diameter D of the specimen: εp = -2 ln(D/D0). Strain rates are from 3·10-5 to 3·10-3 s-1, which are comparable to values in the DC casting process. Most experiments were carried out in force control. Only at 560°C and 580°C the tests were carried out in stroke control as at these temperatures the measured force becomes too low for accurate control. The solid state creep results were fitted to Eq. 1. The specimens tested at semi-solid temperatures often fractured and their fracture surface was studied by SEM.
4
Results
The stress vs. strain rate data measured by the tensile tests are shown in Fig. 1. The strain rate increases with increasing stress and increasing temperature. At the lower temperatures steady state was reached quickly and the corresponding stress and strain rate were plotted in Fig. 1. In the semi-solid temperature range, steady state was not always achieved and in that case the peak stress with the corresponding strain rate was used. The data obtained at 370°C - 470°C were used to determine the parameters of the creep equation (Eq. 1, Table 1). The fracture surface of a specimen broken at 560°C, i.e. with a fL = 0.04, shows dendrite arms covered with a thin film, which was liquid at the time of fracture (Fig. 2). It also shows some rough regions due to solid state rupture. The fracture mechanism proposed [10] is that of fluid film separation and rupture of solid bridges. To describe the material behaviour in this semi-solid range the creep law is modified by taking into account the solid network only (see next section).
241
420
-1
Strain rate (s )
370
10
470
-3
500 510 520 540
10
-4
550 556 560 580
10
-5
1
10
Stress (MPa) Figure 1: Strain rate vs. stress at the temperatures investigated
Table 1: Parameters of the creep law (Eq. 1) for AA5182 in solid state. Parameter Value A 30 (Mpan s)-1 Q 120 kJ/mol n 3.3
Figure 2: SEM micrograph of fracture surface. Fracture took place at 560°C. fL = 0.04. SB1,2 = solid bridge
242 1.0 0.8
fLGB
0.6 0.4 0.2 0.0 0.00
0.05
0.10
fL Figure 3: Fraction of grain boundary area covered with a liquid film vs. total fraction liquid. Wetting angle θ = 0°
5
Discussion
To describe the thermomechanical behaviour of the alloy in semi-solid state it was assumed that the liquid cannot carry any load and that the load is transferred to the existing solid dendrite network. The same assumption was made previously by Drezet and Eggeler [11] and they modified the creep law by dividing the stress by the fraction solid (fS). In this way they explained the high apparent activation energies observed in mushy zone behaviour. However, this modification of the creep law cannot explain the sudden drop in strength as observed in our experiments when passing the solidus temperature. Therefore we propose a further modification of the creep law by taking into account the fraction of the grain boundary area which is covered by liquid. The resulting equation is: σ ε& = A 1 − f LGB
n
Q exp − RT
(2)
where fLGB is the fraction of grain boundary area covered with liquid. Assuming that the liquid is well wetting and that the liquid pockets have a tetrahedral symmetry the fLGB can be calculated from the fraction liquid fL if the liquid fraction is small [12]. The result is shown in Fig. 3 and indicates that the fraction of the grain boundary area covered by a liquid film rises steeply with increasing liquid fraction. In Fig. 4 the normalised stress is plotted vs. temperature for all experimental data. The normalised stress is the ratio of the measured stress and the stress predicted by Eq. 1 for the same temperature with the parameters listed in Table 1. Clearly, a steep decrease in stress level can be observed when entering the semi-solid range. Also shown are the predictions by the modified creep law. The thin line indicates the modification by [11] i.e. dividing the stress in the creep law by the fraction solid fS. The thick line is the prediction by Eq. 2 and follows the experimental data quite well. This indicates that the creep behaviour of this alloy can be extended to the semi-solid range by taking into account the fraction of grain boundary area covered with liquid. Deviations between experimental data and the model are due to the following reasons. Firstly, experimental inaccuracies may arise from the low forces and low strain rates involved.
243 Secondly, small variations in temperature lead to variations in liquid content and to large variations in the fraction of liquid grain boundaries, which have a strong effect on the strength of the material. Thirdly, creep data are extrapolated over a large temperature interval.
Normalised stress
Tsol 1
0.5
error bar 0 350
400
450
500
550
600
Temperature (°C)
Figure 4: Normalised stress vs. temperature. ∆ strain rate 10-3, ◊ strain rate 10-4, × strain rate 10-5. Thin line: creep law modified with fS, thick line: creep law modified with 1- fLGB (Eq. 2). Error bar holds for semi-solid temperatures where the uncertainty is large due to reasons explained in the text.
From Fig. 4 there are indications that in the semi-solid range the normalised stress is somewhat lower for lower strain rates. It is noted that at high strain rate the liquid will mainly be present as drops at triple junctions which leaves a mainly solid grain boundary area, whereas at low strain rates, there is more time for the liquid to spread over the grain boundary area. This is further promoted by the tensile stress state, which will enlarge the effect. Such an effect is not accounted for in the simple model presented and will lead to a stress dependent term in the calculation of the fraction of liquid grain boundaries. The Alstruc model was used to calculate the fraction liquid curve of the material after solidification. The cooling rate was taken as 3°C/s, representative for the specimen location in the ingot. Back diffusion of Mg into the grains will take place during cooling of the ingot to room temperature, thereby removing low melting eutectic phases. This was also taken into account. Reheating before the tensile test was so fast that further diffusion was neglected. However, the cooling rate during solidification strongly effects the liquid fraction curve and the solidus temperature. The prediction by Eq. 2 is thus also strongly dependent on the cooling rate. Taking a cooling rate of 1°C/s for example, the solidus temperature shifts to 550°C.
6
Conclusion
By taking into account the fraction of grain boundary area covered with liquid, the creep law can be used to describe the tensile behaviour of AA5182 during the last stage of solidification. The underlying assumption (i.e. a load bearing solid network with liquid in between) is supported by microstructural observations of the fracture surface which shows the presence of a thin liquid film and the remains of solid bridges.
244
7
Acknowledgements
This research was carried out as part of the EMPACT Brite-Euram project (BRPR-CT950112). Funding by the European Committee is gratefully acknowledged. The co-operation with Mr. T. Iveland (Hydro Aluminium, Norway) on the Gleeble 3500 was very fruitful. We thank Ms. A.L. Dons (Sintef, Norway) very much for the Alstruc calculations and VAW (Germany) for providing the material.
8
References
[1] W.S. Pellini, Foundry, 1952, 80, 125-199. [2] U. Feurer, Giessereiforschung, 1976, 28, 75-80. [3] T.W. Clyne, G.J. Davies in Proc. Conf. Solidification and Casting of Metals, Sheffield, UK, 1979, 275-278. [4] L. Katgerman, Journal of Metals, 1982, 34, 46-49. [5] M. Rappaz, J.-M. Drezet, M. Gremaud, Metall. Trans. A 1999, 30A, 449-455. [6] VAW, Germany. [7] W.M. van Haaften, B. Magnin, W.H. Kool, L. Katgerman in Light Metals 1999 (Ed.: C.E. Eckert), TMS, Warrendale, PA, USA, 1999, 829-833. [8] Private comm. with A.L. Dons (Sintef, Norway), E.K. Jensen (Elkem Research, Norway) and A. Hakonsen (Hydro Aluminium, Norway). [9] A.L. Dons, E.K. Jensen, Y. Langsrud, E. Trømborg, S. Brusethaug, Met. Trans. A, 1999, 30A, p. 2135-2146. [10] W.M. van Haaften, W.H. Kool, L. Katgerman, Materials Science Forum, 2000, 331-337, 265-270. [11] J.-M. Drezet, G. Eggeler, Scr. Metall. et Mater. 1994, 31, 757-762. [12] P.J. Wray, Acta Met. 1976, 24, 125-135.
The Columnar to Equiaxed Transition in Horizontal Direct Chill Cast Magnesium Alloy AZ91
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
246
247
248
249
250
Study of Heterogeneous Nucleation of α-Al on Grain Refiner Particles during Rapid Solidification P. Cizek, B.J. McKay and P. Schumacher University of Oxford, Oxford, UK
1
Introduction
In the Al casting industry, it is common practice to add grain refiner particles to the melt in order to promote nucleation of α-Al and, thus, to ensure grain refinement of the as-cast microstructure. Commercial Ti-B-Al grain refiner master alloys contain TiB2 and Al3Ti particles, which serve as potential sites for the heterogeneous nucleation of α-Al during casting. Attempts to relate the microstructure of these master alloys to their grain refining performance have proven inconclusive, as conventional grain refinement tests have a difficulty to identify a nucleation mechanism for α-Al conclusively [1]. Recent investigations involving Al-based metallic glasses [2], as slow-motion analogues to the undercooled Al melts, have revealed thin Al3Ti DO22 layers adsorbed on the surfaces of TiB2 particles. Nucleation of α-Al was found to occur epitaxially on these layers on the basal faces of boride particles, with a low-index orientation relationship (OR) [2]. The aim of the present work was to elucidate whether the Al3Ti layers originate directly from grain refiner master alloys and whether they undergo changes during re-melting of these alloys. Rapid quenching, using melt spinning, was used to study the nucleation behaviour of α-Al and Al3Ti on boride particles and the thermal stability of the aluminide layers. The resulting large undercoolings facilitated multiple nucleation events on individual boride particles, as a result of small critical nucleus sizes, in contrast to small undercoolings obtained during conventional solidification.
2
Experimental Methods
A commercial Ti-B-Al 5:1 wt.% grain refiner rod (supplied by London & Scandinavian Metallurgical Co. Ltd.) was used as an experimental material. Rapid solidification of the rod was achieved by melt spinning in an inert helium atmosphere. The charge was heated to 1050°C in a boron nitride crucible and held for 5 minutes before ejecting onto a copper wheel rotating at a speed of 20 ms-1. Resulting ribbon was approximately 3 mm wide and 30 µm thick. Both the as-received grain refiner rod and the melt-spun ribbon were examined using TEM. Thin foils were made by electropolishing using a solution of 25% nitric acid and 75% methanol at a temperature of approximately –30°C at 20 V. TEM investigation was performed using a Philips CM20 microscope operated at 200 kV.
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
252
3
Results and Discussion
3.1
As-Received Ti-B-Al Grain Refiner Rod
The microstructure of the grain refiner rod contained TiB2 and Al3Ti particles randomly distributed within the Al matrix. The boride particles exhibited a facetted, hexagonal platelet morphology with the {001} and {100} planes as their external faces (Fig. 1).
Figure 1: TiB2 particle, found in the grain refiner rod, embedded in an Al matrix: (a) TEM bright-field micrograph; (b) dark-field micrograph obtained from the streak located between the center and (001) spots which shows a thin layer (arrowed) coating the particle; (c) SAD pattern showing the streaks (arrowed) and indices of the corresponding diffraction spots (zone axis [100]TiB2, subscripts A and B indicate the Al and TiB2 crystal lattices respectively); (d) same as in (a) after re-tilting by several degrees to obtain the neighboring Al grains in contrast.
Their width-to-thickness ratio was approximately 4:1 with the width of the majority of the particles being less than one micron. Corresponding diffraction patterns (Fig. 1c) confirmed a TiB2 hexagonal crystal structure with parameters a = 0.303 nm and c = 0.323 nm [3]. Boride particles were embedded in a matrix composed of fine Al grains (Fig. 1d), which were largely separated by high-angle boundaries, with no apparent distinct OR observed between the borides and the neighboring Al grains (Fig. 1c). The above fine Al grains appeared to be a result of recrystallisation of the as-solidified α-Al matrix, which occurred during the
253 manufacturing process of the refiner rod. Randomly orientated fine recrystallised grains may have replaced the original as-solidified large grains, thus removing any possible pre-existing well-defined OR between the borides and the neighboring aluminium matrix. Selected-area diffraction (SAD) patterns obtained from the boride particles showed two sets of streaks perpendicular to the (001) and (010) planes (Fig. 1c). These streaks are indicative of a thin layer covering both their basal and non-basal faces. The presence of the layer, about 1 nm in thickness, was confirmed by dark-field imaging, in which the objective aperture was placed on the streak between the center and the (001) spot (Fig. 1b). This suggests that there is a strong similarity between the layers found in the grain refiner rod and those observed on TiB2 particles within the Al-based glasses [2] and identified as Al3Ti DO22 phase. This finding, indicating that the Al3Ti layer might be already present in the rod, is significant as it shows that the layer is not an artifact of the metallic glass technique. It also indicates that, in most instances, the aluminide layer does not have to be formed after adding the refiner rod to the melt, although it might be enhanced or depleted [4]. The Al3Ti particles in the rod were rather small, largely less than 2 µm in diameter, and displayed roughly equiaxed shapes (Fig. 2a). Corresponding diffraction patterns (Fig. 2b) confirmed the equilibrium Al3Ti DO22 crystal structure with the parameters a = 0.385 nm and c = 0.861 nm [5].
Figure 2: Al3Ti particle found in the grain refiner rod and related to one of the neighboring α-Al grains (labeled) by a near cube-to-cube OR: (a) TEM bright-field micrograph; (b) corresponding SAD pattern showing the equilibrium Al3Ti DO22 phase (zone axis [100]).
These particles were frequently related to Al grains by a well-defined, close to the cube-tocube, OR (Fig. 2b). Arnberg et al. [3] have identified three distinct morphologies for aluminide particles (block, flake and petal) in the as-solidified grain refiner master alloys. During casting, these morphologies were found to have a marked influence on the grainrefinement contact-time as a result of differing dissolution times. The roughly equiaxed Al3Ti morphology, observed in the present study, differs from the as-solidified shapes [3] and could be a product of recrystallisation which occurred during the rod manufacturing process. It is possible that the observed well-defined OR between the aluminides and the aluminium matrix was already present in the as-solidified material and was maintained during the recrystallisation process.
254 3.2
Melt-Spun Ribbon
The microstructure of the melt-spun ribbon was composed of fine equiaxed Al grains, largely less than 2 µm in diameter. A majority of these grains contained small petal-shaped particles, with diameters less than 0.5 µm, in their interiors (Fig. 3a). Diffraction analysis (Fig. 3b) indicated that the petal-shaped particles corresponded to metastable Al3Ti, having an ordered Ll2 crystal structure with the lattice parameter a = 0.400 nm [6], as expected for rapid quenching of the present hyper-peritectic Al-Ti alloy. The particles systematically displayed an epitaxial, near cube-to-cube OR with Al (Fig. 3b), which has also been reported in [6]. This suggests that the metastable Al3Ti Ll2 particles, after their formation from the melt, served as potent substrates for subsequent heterogeneous nucleation of α-Al, presumably due to the negligible lattice mismatch between the corresponding crystal lattices.
Figure 3: Petal-shaped metastable Al3Ti Ll2 particle, found in the melt-spun ribbon, located in an α-Al grain center: (a) TEM bright-field micrograph; (b) SAD pattern obtained from the entire grain area, showing a near cube-to-cube OR between the respective crystal lattices (subscripts A and P with the diffraction spot indices indicate the Al and Al3Ti Ll2 lattices respectively).
TEM analysis showed that the thin, presumably Al3Ti DO22 [2], layer covering the TiB2 particles was inherited from the refiner rod and remained essentially unaffected by the high temperature during melt spinning. As shown in Fig. 4, the boride particles were systematically observed to nucleate α-Al, as well as the metastable Al3Ti Ll2 phase, each with the same lowindex OR, found previously using the metallic glass technique [2] In this OR, the {001}, {112}, and {111} planes and the , or , and directions corresponding to the TiB2, Al3Ti DO22 layer and Al (or metastable Al3Ti Ll2 structure), respectively, are approximately parallel to each other [2]. Several α-Al grains were frequently found to nucleate on individual boride particles and they displayed either one or the other of the two possible orientation variants of the above OR (Fig, 4). Thus, borides covered with a thin aluminide layer appeared to serve as potent substrates for heterogeneous nucleation of α-Al (as well as the structurally similar, metastable Al3Ti Ll2 phase), which is in agreement with the findings of the metallic glass technique [2]. This appears to be consistent with observations in conventional casting practice showing that excess Ti (beyond TiB2 stoichiometry), which facilitates the formation of the aluminide layers, is necessary for an effective grain refinement [2,4].
255
Figure 4: TiB2 particle, found in the melt-spun ribbon, nucleating several α-Al grains and metastable Al3Ti Ll2 particles with the low-index OR found in [2]: (a) TEM bright-field micrograph (A1 and A2 indicate α-Al grains corresponding to the two possible variants of the observed OR); (b) large-area SAD pattern; (c) schematic showing indices of the corresponding diffraction spots (zone axes [100]TiB2 and [011]Al or [0-1-1]Al, subscripts A1 and A2 are the same as in (a), subscript B indicates the TiB2 crystal lattice).
4
Conclusions
A detailed study of a commercial Ti-B-Al grain refiner rod, both in an as-received state and after rapid solidification using melt spinning, was undertaken in the present work. The microstructure of the as-received rod appeared to be dominated by recrystallisation rather than solidification processes. TiB2 particles were covered with thin Al3Ti DO22 layers that were already present in the as-received rod and appeared to remain unaffected by the high melt spinning temperature. These particles proved to be very potent sites for heterogeneous nucleation of α-Al during quenching, competing successfully with the metastable Al3Ti Ll2 particles having a negligible lattice mismatch with α-Al.
5
Acknowledgements
The authors gratefully acknowledge financial support from the EPSRC in conjunction with London & Scandinavian Metallurgical Co. Ltd. and Alcan International Ltd.
256
6 [1] [2] [3] [4]
References
D. G. McCartney, Int. Mater. Rev. 1989, 34, 247 – 260. P. Schumacher et al., Mater. Sci. Technol. 1998, 14, 394 – 404. L. Arnberg, L. Bäckerud, H. Klang, Met. Technol. 1982, 9, 7 – 13. 4. A. M. Bunn et al. in Solidification Processing 97 (Ed.: J. Beach and H. Jones), University of Sheffield, Sheffield, UK, 1997, 264 – 267. [5] S. Hashimoto, K. F. Kobayashi, S. Miura, Z. Metallkd. 1983, 74, 787 – 792. [6] W. T. Kim et al., Int. J. Rapid Solidif. 1992, 7, 245 - 254.
Effect of Instability of TiC Particles on the Grain-Refining Behavior of Al-Ti-C Inoculants in Aluminum Alloys M. Vandyoussefi and A. L. Greer University of Cambridge, Cambridge, UK
1
Abstract
According to the Al-Ti-C phase diagram, TiC is only a stable phase in Al-melts rich in Ti and C, and should convert to Al4C3 at the normal addition levels of Al-Ti-C refiners to melts. The effects of melt composition on TiC stability and on grain-refiner performance are examined using TP-1 refining tests on super-purity Al and using thermodynamic modeling. Effective grain refinement with little fading of performance on holding the refiner in the melt is observed only for melt compositions within the liquid + TiC phase field. In contrast when Al4C3 is stable, the refining is poor even at short holding times, and fades considerably at longer holding times. This change in grain-refining behavior is related to the decomposition of TiC.
2
Introduction
In spite of wide use of the conventional Al-Ti-B grain refiners [1-2], their disadvantages i.e. poisoning [3] and agglomeration of the boride particles [4], have led to interest in the production of Al-Ti-C based refiners [5-8]. Although the Al-Ti-C refiners have overcome these limitations, their effectiveness is lost on holding in the melt, especially at higher temperatures. This could be related to the intrinsic thermodynamic instability of the TiC phase in melts poor in titanium [9-11]. TiC particles could convert at their surfaces to, e.g., Ti2AlC or Al4C3, or dissolve completely and cause the precipitation of Al4C3 [9-10]. Consequently, nucleation potency would decrease through loss of the lattice matching of TiC and ccp-Al [10-11]. The formation and decomposition of Al4C3 have been experimentally confirmed [12-13], but under conditions different from normal grain refinement practice. Both nucleation and solutal restriction of crystal growth control grain refinement [14-17]. In aluminum melts, titanium plays a more important role in the restriction of crystal growth than other solutes. Usually, it enters into melts through addition of grain refiners which have more Ti than required to form TiB2 or TiC. The stability of the TiC phase depends on the dissolved titanium content in the melt. Thus the performance of Al-Ti-C refiners should be analyzed based on the titanium content. The aim of the present study is to show the effect of the stability of TiC particles on the grain-refining performance, through thermodynamic modeling and experiments.
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
258
3
Experimental Procedure
The TP-1 test [18] was applied to evaluate the performance of the Al-Ti-C refiners. Superpurity (SP) aluminum (99.999%) was chosen to highlight the role of dissolved titanium from the refiner itself, avoiding the interference by other solutes. To investigate the effects of phase stability on grain refinement, different phase fields in the Al-Ti-C phase diagram were sampled by selecting two refiners added to various levels. The Ti contents of refiner 1 and refiner 2 are respectively 4.91 and 1.82 wt.% and C contents are 0.14 and 0.17 wt.%. The melts were held at 700±5Û& EHIRUH DQG DIWHU DGGLWLRQ RI UHILQHU 7DEOH VXPPDUL]HV WKH experimental conditions and the chemical compositions of the alloys. Concentrated Tucker’s reagent (25% H2O, 45% HCl, 15% HNO3, 15% HF) was used to reveal the grain structure. Areas of 2 x 2 cm were cut from the centers of TP-1 sections and after polishing were anodized in Barker’s reagent (98% H2O, 2% HBF4). The grain size was measured in polarized-light microscopy on these samples. Table 1: Experimental conditions for the TP-1 tests on super-purity Al. Refiner 1 is used throughout, except for the S45 series (45 ppt addition of refiner 2). The shortest holding time in the melt is 5 min and the longest time is 180 or 360 min. The chemical compositions are measured by optical emission spectroscopy. The compositions of non-refined alloys are all very similar and are given only for the S1 series. The carbon contents are calculated from the known refiner addition levels. Q is the growth-restriction factor which is used to describe the effect of different solutes on grain size.† Sample series
Addition Holding time level (ppt) [min] Ti S1-0 0 — 0 S1-1 to S1-5 1 5, 30, 60, 180, 360 0.005 S5-1 to S5-4 5 5, 30, 60, 180 0.027 S17-1 to S17-4 17 5, 30, 60, 180 0.094 S30-1 to S30-4 30 5, 30, 60, 180 0.17 S45-1 to S45-4 45 5, 30, 60, 180, 360 0.08
Solute content (wt.%) C 0 0.0001 0.0007 0.0023 0.0041 0.0073
Si 0.0006 0.0004 0.003 0.085 0.0031 0.0049
Fe 0.0004 0.0009 0.0015 180 MPa, elongation > 15 %. Standard production is high pressure die casting with heat treatment of the components which results in distortion and therefore expensive dressing. The self-hardening alloy AlMg5Si2Mn (Magsimal 59) was chosen as a promising candidate and two modifications (AlMg3,5Si1,4Mn; AlMg2Si0,8Mn) were included to study the influence of silicon and magnesium.
3
Feedstock Production
Feedstock material for the investigations was produced by EFU’s vertical continuous DCcaster with electromagnetic stirrer (Fig. 1). The main components of the DC-caster are a tilting induction furnace (capacity: approx. 250 kg aluminium) and a vertical hydraulically driven casting machine. The furnace is electrically powered by medium frequency current Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
270 (500 Hz). The stirring method is circumferential. All experiments were performed with a degassing step that was carried out in the furnace bath before casting. The used impeller works with Argon as refining gas. The treatment was performed in two subsequent periods (2 x 15 min, each 12l Ar/min). After degassing the furnace is also used as pouring furnace and the melt is continuously poured into the launder from where it flows into the water cooled DC-casting mould. The mould design used is a Hot-Top system which allows the additional use of electromagnetic mould stirring. From one batch (250 kg) of molten metal up to four billets (length 3,5 m) could be cast subsequently.
Figure 1: Vertical continuous DC-caster with electromagnetic stirrer
Generally the requirements for thixoforming feedstock material and the influence of casting speed and stirring intensity may be summarized as follows: 3.1
Microstructure
The reheating process will cause the rosette-like structures still present in the as-cast billet to become fully globular, thereby creating the material structure typical for the process. The best flow behaviour for forming is achieved when the microstructure consists of very fine and very round globular grains. In addition the amount of liquid phases entrapped within the grains should be very low, because it could not participate to the deformation of the material and therefore is to be considered as solid. The grain size and the entrapped liquid decreases with increasing stirring intensity. The casting speed has no influence on the entrapped liquid and a slight one on the grain size; decreasing grain size with increasing casting speed. The form factor (quotient of circumference of a circle with the same area than the grain and circumference of the grain) increases with increasing stirring intensity and is not influenced by the casting speed.
271 3.2
Macrosegregation
Macrosegregation in the feedstock material could lead to parts with uneven mechanical properties. Near the surface the highest concentration of the alloying element silicon could be measured while the casting speed is high or the stirring intensity is high. The marked degree of the so called reversed block-segregation in the centre of the material increases with increasing casting speed and increasing stirring intensity. 3.3
Melt Loss
It is generally noted that various feedstock types lose different amounts of liquid phase during the inductive heating process [4]. Due to economically reasons it is important to ensure that the melt loss is 180 MPa; elongation > 15 %) which should be achieved without heat-treatment the feedstock material must have excellent flow and die filling behaviour, because this type of components involves long flow distances with very thin walls. Therefore the self-hardening alloy AlMg5Si2Mn was examined (Table 1). Table 1: chemical composition of alloy AlMg5Si2Mn (Magsimal 59) Si Fe Cu Mn Mg Zn Ti others Min. 1,8 0,5 5,0 0,1 Max. 2,5 0,13 0,05 0,8 5,5 0,08 0,2 0,06 The reheated microstructure (Fig. 4) consists of fine globular primary phase surrounded by eutectic phases. The α-Phase is less globular then known from alloy A356.
100 µm
100 µm
as cast reheated Figure 4: Microstructure of alloy AlMg5Si2Mn before and after reheating
The thixoformability was tested with various test parts, representing thin and thick walled cross sections. The alloy AlMg5Si2Mn showed excellent flow and die filling behaviour and is even better than of alloy A356 [10]. The mechanical properties measured in a thin walled plate (thickness = 3 mm) and in a door handle (wall thickness 5 – 20 mm) are a little bit to high concerning the yield stress, sufficient concerning the ultimate tensile strength and slightly too low concerning the elongation (Fig. 5).
274
YS; UTS (MPa) 500 required properties 450 properties of the plate 400 properties of the door handle 350 297 311 300 250 120 - 177 160 > 180 200 150 150 100 50 0 YS UTS
elongation (%) 20 15 18 13 > 15 16 14 12 10 8 6 4 2 0 elongation
Figure 5: Mechanical properties of alloy AlMg5Si2Mn
The content of eutectic phases after reheating of the alloy modifications AlMg3,5Si1,4Mn and AlMg2Si0,8Mn is less than for alloy AlMg5Si2Mn, due to the decreased content of silicon and magnesium. The thixoformability and the mechanical properties were tested with the connecting rod (Fig. 2). The flow and die filling behaviour is unsatisfactory, while the material showed non-rewelded flow fronts at the cores. The yield stress of the alloy modifications was reduced as required, but the elongation in the as cast condition could not be improved.
6
Conclusion
Alloy AlMgxSix has a sufficient thixoformability concerning its flow behaviour and its ability of rewelding flow fronts. The alloy has no hot cracking tendency. The mechanical properties measured in tensile test specimens dissected from the demonstrator part after T6-heat treatment were significantly improved compared to alloy A356). The best agreement with the required properties for thin walled space-frame-knots is alloy AlMg5Si2Mn. Also this alloy showed excellent die filling capacity (judged even better than alloy A356 [10]).
275
7
References
[1] G. Chiarmetta, 4th Int. Conference on Semi-Solid Processing of Alloys and Components, Sheffield (England), (1996), p. 204 – 207 [2] H. E. Pitts, H. V. Atkinson, 5th Int. Conference on Semi-Solid Processing of Alloys and Components, Golden (Colorado), (1998), p. 97 - 104 [3] G. Hirt, R. Cremer, T. Witulski, VDI-Berichte Nr. 1324, (1997), p. 55 – 73 [4] F. Niedermaier, J. Langgartner, G. Hirt, I. Niedick, 5th Int. Conference on Semi-Solid Processing of Alloys and Components, Golden (Colorado), (1998), p. 407 – 414 [5] S. Nikitin, R. Ellerbrok, S. Engler, Gießerei 76, Nr. 9, (1989), p. 297 – 299 [6] M. P. Kenney, J. A. Courtis, R. D. Evans, G. M. Farrior, C. P. Kyonka, A. A. Koch, Metals Handbook, 9th edition, ASM International, Volume 15: Casting, (1988), p. 327 – 338 [7] B. Wendinger, 4th Int. Conference on Semi-Solid Processing of Alloys and Components, Sheffield (England), (1996), p. 239 – 241 [8] N.N., Thixocasting, Aluminium Pechiney, (1994) [9] J.-P. Gabathuler, R. Jaccard, R. Röllin, Ch. Ditzler, , VDI-Berichte Nr. 1235, (1995), p. 81 – 105 [10] Niedick (ed.) ”Thixotec”-Abschlußbericht, ISBN 3-89653-651-6, p. 12-27
Experimental Study of Linear Shrinkage during Solidification of Binary and Commercial Aluminum Alloys Dmitri Eskine and Laurens Katgerman Netherlands Institute for Metals Research, Delft, The Netherlands
1
Introduction
The process of direct-chill (DC) casting is a very common way to produce ingots and billets for further deformation. Although this technology is used in aluminum industry since 1950s, there are still many problems related to the process. Among these problems, there are the formation of structure and properties of the mushy zone and, the most important for industry, the causes of defects. Hot tearing, porosity and macrosegregation are the major defects occurring during casting. Obviously, the flow of liquid, the diffusion of solute elements, the structure formation, and the development of strength in the mushy zone, being in complex interaction, may result in the formation of defects. Hot tearing or hot cracking is one of the most common problems encountered in DC casting of aluminum alloys. The main cause of this defect is that the stresses and strains built up during solidification are too high as compared to the actual strength of the semi-solid material. This type of defects occurs in the lower part of the solidification range, close to the solidus, when the solid fraction is more than 0.9 [1]. At this moment, the mushy zone is definitely coherent, but the liquid film still exists between most of grains. The term coherency (or coherency temperature) should be used with caution. Usually, the coherency marks the formation of a continuous dendritic network, when the material starts to develop the strength [2, 3]. At temperatures above the coherency point, the grains are free to move with respect to each other and do not transfer any forces. Moreover, before the coherency point is reached the liquid phase can easily flow between grains and, therefore, the melt feeding and the redistribution of solute elements occur without much difficulties. Another meaning of the coherency point designates it as a point at which the measured temperature dependence of the viscosity of the liquid-solid mixture changes the slope. Depending on the alloying system and, what is very important, on the technique of measurement, the coherency point may occur in the broad range of temperatures (or solid fractions). Different authors assign this point to the solid fraction in the range from 0.1 to 0.5%. The question is when hot tearing really occurs, and what are the driving forces for hot tearing. To answer the first question, the terms of the effective solidification range [2, 4] or the vulnerable part of the solidification interval [5] were introduced in 1940s–1950s. The upper boundary of this range is the point where the solidification shrinkage starts [2]; the lower boundary is the solidus (equilibrium or nonequilibrium, depending on the solidification conditions). As for the driving forces, it is well adopted that the collapse of a dendritic network is caused by the developing stresses, which result from the shrinkage and driving pressure for liquid flow. The necessary condition for hot tearing is the existence of thin, continuous, interdendritic liquid film alongside the low permeability of the mushy zone. This Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
277 condition is usually fulfilled at large volume fractions of solid, 0.9 to 0.99 [6]. Therefore, there is a clear gap between the coherency point and the temperature below which hot tearing occurs. Novikov [2] suggested to determine the upper boundary of the effective solidification range by measuring the linear shrinkage. Hence, the upper temperature of this solidification range is the temperature, at which the linear shrinkage starts. To avoid the confusion, we need to specify the term “linear shrinkage”. There are different terms in which the word “shrinkage” is used. The casting shrinkage is a technological parameter describing the difference in dimensions or volume between the original cavity of the mold and the final casting cooled to the room temperature. This shrinkage, by definition, includes the volume differences between liquid and solid phases, the shrinkage occurring in the solidification range, and the contraction of the completely solid casting. The solidification shrinkage is the shrinkage (usually volume shrinkage) occurring in the solidification range, from 100% liquid to 100% solid. The linear solidification shrinkage is the horizontal change in linear dimensions of a casting during solidification [2]. Above the temperature of the linear shrinkage onset, the alloy is fluid as there is not continuous (between the walls of the mold) network of dendrites. In this stage of solidification, the thermal shrinkage cannot manifest itself as the horizontal contraction of the casting. All volumetric changes appear as the decreasing level of the melt in the permanent mold or does not appear at all during DC casting, due to the continuous supply of the melt to the mold. However, the linear shrinkage appears, and can be measured, when the fluidity of the alloy drastically drops, and the rigid skeleton of the solid phase forms. Starting from that moment, the alloy acquires the capability to retain its shape, and the thermal shrinkage of the solid phase displays itself as the linear contraction in the horizontal direction. One can also call this parameter “a one-dimensional solidification shrinkage”, keeping in mind that it is measured in the horizontal dimension of the casting. A special technique was developed to measure the linear shrinkage (and pre-shrinkage expansion) upon solidification [2, 7, 8]. Several designs of an experimental setup were suggested, all sharing the following features: graphite mold (providing low friction and high thermal conductivity) with one moving wall; water-cooled base (for high cooling rates comparable with those upon chill casting); and simultaneous temperature and displacement measurements [2]. Despite numerous published studies by Novikov et al., e.g. [2, 7–9], there are only few data available on the contraction behavior of solidifying wrought alloys, which are mainly produced by DC casting. The aim of this paper is to describe the experimental technique used for measuring the linear shrinkage upon solidification of binary and commercial wrought alloys; to discuss the factors influencing the measured property; and to present the experimental results obtained using the developed technique.
2
Experimental
The experimental setup used in measurements of the linear shrinkage upon solidification comprised the following parts: a graphite mold (Fig. 1) with one moving wall, a water-cooled bronze base; a K-type thermocouple (0.25 mm wires); and a linear displacement sensor (a Schaevitz DC-DC LVDT). To attach the solidifying metal to the moving wall on the one
278 side and to the permanent wall on the other side, we used a metallic rod with a thread embedded into the moving head and a T-shaped cavity on the other side of the mold. The cross-section of the T-shaped cavity was thinner than that of the main cavity, which allowed the melt to solidify faster. The metallic rod fixed in the moving head was frozen in the sample during experiments. The cross-section of the main cavity was 25 × 25 mm with a gauge length of 100 mm. The data was acquired using a Keithley KPCI-1801HC interface card and a LabVIEW software.
Figure 1: Experimental mold for measuring linear contraction behavior in the solidification range
The dimensions of the mold were chosen in agreement with Novikov et al. [8], who showed that these dimensions made the measured property not scale-dependent. The linear shrinkage was determined as follows: ε = {(ls + ∆lexp – lf)/ls} × 100%, where ls is the initial length of the cavity (100 mm); lf is the length of the sample at a temperature of solidus, and ∆lexp is the pre-shrinkage expansion. The linear displacement was measured accurate to 6 µm or 0.006% with 3 to 5 samples measured for each point. The distance between the thermocouple tip and the bottom of the mold was about 1.5 mm. Otherwise, there were problems with filling the gap. Evidently, the measured property should depend on the structure parameters; chemical composition; and the force applied to the moving wall of the mold. In order to examine the effects of these factors we (i) varied the melt temperature from 700 to 800 oC (grain size) and the cooling rate from 7 to 15 K/s (dendritic parameter); (ii) used alloys of different alloying systems (Table 1); and (iii) changed the friction force applied to the moving part from 0.11 to 0.83 N. In most cases, the liquidus and solidus temperatures could be derived from the cooling curve. Note, however, that we determined the linear shrinkage in the entire solidification range which, at the used cooling rates, extends down to the lowest possible eutectic temperature. The liquidus and solidus temperatures (both equilibrium and nonequilibrium) are given in Table 1. After acquiring the primary data, temperature and displacement against time (Fig. 2a), we processed the cooling curve in order to obtain information about critical temperatures. After that, we reconstructed the data to receive the direct dependence of displacement on
279
o
700 600 500 400 300 200 100 0 -100 -200 -300 -400
0.00
1
Displacement, %
Temperature, C Dispacement, µm
temperature (Fig. 2b). From this dependence the linear pre-shrinkage expansion, the linear solidification shrinkage and the temperature of its onset can be extracted.
2
-0.40
-0.80 0
10
20
Time, s
30
400
40
a)
500
600 o
Temperature, C
700
b)
Figure 2: Examples of data obtained during experiments: a, primary data, 1 for temperature and 2 for displacement (an Al-4% Cu alloy; melt temperature 750 oC; cooling rate 12 K/s; friction force 0.164 N) and b, temperature dependence of displacement (an Al–4% Cu alloy; melt temperature 750 oC; cooling rate 11 K/s; friction force 0.12 N)
Table 1: Nominal chemical composition and critical temperatures (tl – liquidus; tes and tnes – equilibrium and nonequilibrium solidus, respectively) for examined alloys. Alloy Chemical composition, wt% tl, oC tes, tnes, o o C C Cu Mg Zn Mn Fe Si Al Al–Cu 1050 2024(a) 5182(a) 6082(a) 7075(a) 1.
3
4.0 4.4 0.1 0.1 1.6
1.5 4.5 0.8 2.5
5.6
0.6 0.35 0.7 0.2
0.15 0.3 0.3 0.2 0.3 0.3
0.1 0.2 0.3 0.2 1.0 0.3
bal. bal. bal. bal. bal. bal.
652 657 638 638 649 635
582 646 502 577 566 532
548 502 ~530 557 477
Contained also Ti and Cr.
Results and Discussion
An Al–4% Cu alloy was chosen for examination because of the well-known phase diagram, structural features and some reference data on its contraction behavior. The results are summarized in Fig. 3, where the combined effects of structure and applied force are given as a 3-D plot along with contours of the linear solidification shrinkage. The structure factor used in these plots is a combined function of the grain size (D) and the dendritic parameter (d): d/D. The examination of structure showed that the actual grain size varied from 120 to 220 µm, the structure being mostly equiaxed across the sample. For this type of macrostructure, the finer is the grain and the coarser is the dendrite, the larger is linear solidification shrinkage.
280 In further experiments with commercial alloys, the friction force was 145 mN. The linear solidification friction measured at considerably low friction forces is about 0.2%. This value agrees well with that previously reported (0.25%) [2]. Korol’kov [9] reported that the measured linear solidification shrinkage was always higher than the calculated values. This is in favor of our results. The important parameter, that can be derived from our experiments, is the temperature of the linear shrinkage onset (tso). This temperature is about 560 oC, being independent of the friction force and structure. This temperature, being between the temperatures of the equilibrium solidus and nonequilibrium solidus, shows that the continuous dendritic network is formed at very high solid fractions. Moreover, the amount of nonequilibrium eutectic liquid should be taken into account when considering the solidification at high cooling rates. The results on the linear solidification shrinkage for several commercial alloys are given in Table 2. Structure factor (x 1000)
5.0 4.5
0.17
4.0
0.15 0.13
3.5
0.11
3.0
0.09
2.5
0.07
2.0
0.05
1.5
0.03
1.0 0.0
0.01
0.2
0.4
0.6
0.8
1.0
Friction force, N Figure 3: The effect of structure and friction force on the linear solidification shrinkage of an Al–4% Cu alloy
Obviously, commercial alloys behave differently upon solidification, with respect to an alloying system. Commercial aluminum (1050) shows almost no shrinkage in the solidification range, which is no surprise. However, Al–Mg and Al–Mg–Si alloys (5182 and 6082, respectively) also demonstrate negligible linear solidification shrinkage and start contraction only at the solidus temperature or even below. The largest linear solidification shrinkage is observed in a 7075 (Al–Zn–Mg–Cu) alloy, the temperature of shrinkage onset being between the equilibrium and nonequilibrium solidus. Table 2: Temperature of shrinkage onset (tso), pre-shrinkage expansion (+∆l) linear solidification shrinkage (εs) and the total casting shrinkage (εc) of commercial alloys (melt temperature 720–730 oC; cooling rates 12–18 K/s). Alloy tso, oC +∆l, % εs, % εc, % 1050 657 0 0.03 2.3 2024 515 0.035 0.2 1.5 5182 528 0.08 0 1.5 6082 563 0.025 0.02 1.5 7075 525 0.03 0.27 1.5
281
4
Conclusions
1. A technique to adequately measure the linear solidification shrinkage is developed and used for binary and commercial aluminum alloys. 2. The linear solidification shrinkage depends on the chemical composition, structure, and applied force. 3. The finer the grain structure and the coarser is the internal dendrite structure, the larger the linear solidification shrinkage. 4. The temperature of the shrinkage onset depends on the chemical composition.
5
References
[1] J. Campbell, Castings, Butterworth-Heinemann, Oxford, 1991. [2] I.I. Novikov, Goryachelomkost tsvetnykh metallov i splavov (Hot Tearing of Nonferrous Metals and Alloys), Nauka, Moscow, 1966 (in Russian). [3] A.K. Dahle, L. Arnberg, Acta Mater. 1997, 45, 547–559. [4] A.A. Bochvar, Izvestiya Akad Nauk SSSR, OTN 1942 (9) 31 (in Russian). [5] W.S. Pellini, Foundry 1952, 125–133; 192–199. [6] T.W. Clyne, G.J. Davies, Br. Foundrymen 1981, 74, 65–73. [7] A.N. Yakubovich, I.I. Novikov, G.A. Korol’kov, Russian Castings Production 1969, (10), 472–474. [8] I.I. Novikov, G.A. Korol’kov, A.N. Yakubovich, Russian Castings Production, 1971, (8), 333–334. [9] G.A. Korol’kov, Liteinoe Proizvod. 1986, (1), 6–7 (on Russian).
The Influence of the Cooling Rate on the Type of the Intermetallic Phases in the Aluminium Alloys of the 3XXX (AlMnMgSi) Group Tomasz Stuczyn´ski and Marzena Lech-Grega Institute of Non-Ferrous Metals, /LJKW0HWDOV'LYLVLRQXO3LáVXGVNLHJR6NDZLQD3RODQG
1
Abstract
The results of investigations of the role of cooling rate, within the range of 0.6 to 28°C/s, and of chemical composition on the type of intermetallic phases in the aluminium alloys of the 3xxx (AlMnMgSi) type are presented. The procedure introduced and published by the Alcoa Research Centre was used in the studies. It was shown that the fraction of phases of the Al6(FeMn) and a-AlFeMnSi type varies in the investigated alloy depending on the cooling rate and the Fe/Si ratio. It was found that the AlFeMnSi phase, required from the point of view of the application of the studied alloy (beverage cans) predominates at the Si content twice the Fe content. On the other hand the a-AlFeMnSi phase is the only one ferricmanganese phase at the cooling rate of 2.6°C/s. The results obtained on the laboratory scale have been confirmed on the pilot plant scale, during casting the ingot from the studied alloy, using a semi-continuous method.
2
Introduction
A characteristic feature of the process of ingot casting of aluminium and its alloys by a semicontinuous method is that each elementary volume of the cast metal is solidified with different rate, the magnitude of which depends on the employed parameters of the casting process, i.e., the temperature of the melt, casting speed, the amount of cooling water and on the position in relation to the mould walls and the axis of the ingot being casted. It means that this phenomenon is critical in the case of casting of large ingots of rectangular cross-section used as the charge for making rolled products. The studies on the thermal processes occurring upon the casting of the melt into an ingot have shown that the cooling rate expressed indirectly through the intensity of the temperature drop of the given elementary volume of the cast metal by the semi-continuous method into classical crystallizers of the DC type ranges between 0.5 to 20°C/sec. A typical distribution of the cooling rates in an ingot is presented in Fig. 1. In the plot depicting the distribution of cooling rates in the ingot, the four zones can be distinguished: Zone I: the area of the highest cooling rate placed at the surface of the ingot. Responsible for the kinetics of the cooling process in this area is the intensity of carrying away of heat through the walls of the mould forming the cast ingot. Zone II: the area of reduced cooling rate placed between the end of the area of the primary cooling impact - through the mould walls - and the beginning of the area of the secondary cooling caused by the direct cooling water spray flowing out of the mould. Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
283
Figure 1: Typical distribution of cooling rates in an ingot
Zone III: the area of increased cooling rate placed in the area of the impact of the secondary cooling. Zone IV: includes the middle part of the ingot. The size of this area depends on the intensity of carrying away heat along the axis of the cast ingot and for the most part it is the function of the casting speed. The occurrence of these four zones in the distribution of cooling rates is reflected in the pattern of the solidification front in the casted ingot (Fig. 2).
Figure 2: Pattern of solidification front
There exists a known relationship between the cooling rate Vk and the casting speed VL expressed by the equation: r r vK = vL cos α where: α is the angle between the normal to the solidification front line and a line parallel to the direction of casting. In the foundry practice it was found that in many cases, even the changes in the casting speed of 50% do not affect the distribution of intermetallic compounds, the example of which
284 can be the result obtained on the 5182 alloy. Nevertheless there exists a group of alloys in which the occurrence of differentiated conditions in the solidification process in the ingot, as manifested by a different cooling rate of different elements of volume dx⋅dy⋅dz, gives rise to significant differentiation in the structure of the alloy at the level of its microstructure. This group of alloys includes mainly the alloys of the series 1xxx and 3xxx according to ASTM designation. In ingots made of these alloys, a significant differentiation occurs in the form and type of intermetallic compounds based on Al, Fe, Mn and Si depending on the position of the studied area in relation to the side walls of the ingot, i.e., on local parameters of the solidification process. Since the studies on the temperature field in ingots cast by the semi-continuous method are difficult because of the apparatus problems and a lack of a mathematical 3D model of the process of ingot solidification, in the investigations on the determination of the distribution of cooling rates, model studies are used. In the presented studies, the research procedure for the determination of the impact of the cooling rate and the composition of intermetallic compounds was employed as introduced and published by the Alcoa Research Centre. [1]
3
Experimental
Three variations of 3104 alloy (AlMnMg) were selected for the studies. The chemical composition of the studied alloys is presented in Table 1. Table 1: Chemical composition of studied alloys. Alloy No. Si Fe Cu Mn Mg Zn 1 0.27 0.29 0.24 1.23 1.18 0.22 2 0.28 0.69 0.24 1.23 1.21 0.24 3 0.57 0.28 0.24 1.25 1.20 0.23
Cr 0.08 0.11 0.10
Ti 0.04 0.05 0.04
The alloys were prepared in a resistance crucible furnace of crucible volume of ca. 8 kg of the melt. The weight of the melt used in the tests was equal 5 kg. The charge was A1E grade aluminium produced by Aluminium Plant "Konin" whereas alloy additions were introduced in the form of master alloys produced AlTab type of by LSM. On this stage of investigation the grain refiner was not. The temperature of the melt prior to the casting was constant and equal to 700°C. The melt was cast into a permanent mould the schematic diagram of which is presented in Fig. 3. Prior to proper tests, preliminary studies were carried out in order to determine the intensity of the cooling of the cast metal in fixed places positioned in relation to the copper plate constituting the main chill of the permanent mould. A computer-aided thermal analysis stand was employed in the studies. Selected values of cooling intensities are presented in Fig. 3 depicting the schematic diagram of the construction of the experimental permanent mould. It was found that the employed research procedure allows to determine the impact of the cooling rate in the range of 0.5 to 28°C/sec on the form and shape of the components of the structure of the studied alloy. It means the entire range of the cooling rates occurring in the rolling ingots cast by the semi-continuous method into the DC type mould is included.
285
Figure 3: Schematic diagram of test set-up [1]
The research methods were employed which allowed to determine the morphology of the phases, the grain size of α-solution and the identification of phase components. The observation of the microstructure was carried out by a Steresscan 420 Scanning Microscope and microanalysis by a LINK ISIS 300 Microanalyser.
4
Results
The samples cast into the permanent mould shown schematically in Fig. 3. were cut and from the areas in which the cooling rate was determined, samples for metallographic studies were prepared. After metallographic microsections had been prepared, the samples were etched in Barker’s reagent in order to disclose the grain size and thus to determine the effect of the cooling rate on this structure describing parameter. In Fig. 4. a typical relationship between the grain size and the cooling rate is depicted (without grain refiner). Next, the same samples were polished once again and new metallographic microsections were made for the studies allowing to identify the single phases occurring at the boundaries of the solid solution α. As a result of the microscopic observations and the studies of the chemical composition of intermetallic phases it was found that in the studied alloys occur the following metallic phases: Al.(FeMn), α-Al.(FeMn)Si and Mg2Si. Comparing the number of the identified phases, their relative amount (%) in the studied microareas was estimated. The results are summarised in Table 2.
286
28oC/sek
1.6oC/sek
5.0oC/sek
2.6oC/sek
0.9oC/sek
Figure 4: Relationship between the grain size and the cooling rate (50x)
In Figs. 5 to 7, typical structures of the studied alloy 3104 as a function of the Fe and Si content for given cooling rates are illustrated.
Phase α−Al(FeMn)Si a, e, f, g, h
Al 6 (FeMn) b, c, d, i, k, l
Figure 5: Alloy No. 1, cooling rate 0.9°C/sec, Fe/Si = 1
Phase α−Al(FeMn)Si b
Al 6 (FeMn) a, b, c, d
Figure 6: Alloy No. 2, cooling rate 0.9°C/sec, Fe/Si = 2.46
287
Phase α−Al(FeMn)Si a, c, d, e, f, g,h
Al 6 (FeMn) k
Figure 7: Alloy No. 3, cooling rate 0.9°C/sec, Fe/Si = 0.49
Table 2: The identified phases relative amount in the microarea of studied alloys Alloy Phase Cooling rate [°C/sec] 28 5.0 2.6 1.6 0.9 1 α-Al(FeMn)Si 20 30 40 30 50 Fe/Si =1 2 Fe/Si = 2.5 3 Fe/Si=0.49
5
Al6(FeMn) α-Al(FeMn)Si
80 5
70 20
60 20
70 20
50 5
Al6(FeMn) α-Al(FeMn)Si
95 60
80 80
80 100
80 80
95 90
Al6(FeMn)
40
20
-
20
10
Discussion
Analysing the obtained results, it should be remembered above all that the model tests carried out, based on the Alcoa procedure, include the entire range of the changes of the cooling rates occurring in practical conditions during the casting of the studied aluminium alloy of the AlMnMg type into rolling ingots by the semi-continuous method. The above statement makes good reason to put a thesis that the observed correlations between the cooling rate, the chemical composition, particularly the ratio of Fe to Si content, and the type of intermetallic compounds in the structure of the studied alloy and the grain size are reflected in real conditions in industrial practice. The disclosing of the dependence of the grain size on the cooling rate indicates in practice the importance of employing the grain refinement treatment, the correct carrying out of which can result in the decrease of this dependence. As a result of the application of the grain
288 refinement, a homogeneous distribution of the grain size in the entire cross-section of the ingot should be obtained, irrespective of the position, thus of the cooling rate. The results indicate that the percent fraction of the given intermetallic phases mainly depends on the ratio of iron to silicon content within the chemical composition according to the standard in force. This finding offers the possibility of the conscious control of the structure formation of a given alloy taking into account its application and further processes occurring during the working and heat treatment. It was proved by the experimental cast with the objective to form a rolling ingot of the 3104 alloy, with the predominating fraction of the α-Al.(FeMn)Si phase in its structure. According to the research studies, the chemical composition presented in Table 3 was selected. Table 3: Chemical composition of slab ingot 300 × 125 mm of 3104 alloy. No. Chemical composition [wt%] Fe Si Mg Mn Cu Ti 2 Casting 18 0.25 0.51 1.13 1.20 0.28 0.016
Zn 0.02
After the casting in the Institute of Non-Ferrous Metals - Light Metals Division Skawina at the production-experimental stand of semi-continuous casting, the ingot was subjected to metallographic studies. A typical picture of the obtained structure together with the disclosed intermetallic compounds is shown in Fig. 8.
a b Al(FeMn)Si Al 6(MnFe)
c Al(FeMn)Si
d Al(FeMn)Si /Al 6(MnFe)
e Al(FeMn)Si/ Al 6(MnFe
Figure 8: Identification of compounds in ingot cast in the Institute of Non-Ferrous Metals - Light Metals Division
As can be seen, according to the expectations, a structure was obtained in which αAl.(FeMn)Si constitutes the predominating intermetallic phase which was the aim of the experiment carried out. Summarising, it should be noted that the presented studies created the basis for the formulation of definite technological requirements allowing to consciously control the formation of the structure of the studied alloys of 3104 type cast into rolling ingots in industrial conditions.
289
6
Conclusions
Summarising the obtained results it can be noted that: 1. To a large extent, the type of intermetallic compounds is determined by the ratio of iron to silicon content. 2. In the case of the ratio of Fe/Si = 2, the predominating intermetallic compound in the structure of the 3104 alloy is Al6(MnFe) phase, whereas if this ratio is equal 0.5, the predominating compound becomes the α-Al.(FeMn)Si phase. 3. In the studied range of cooling rates (0.5 to 28°C/sec), the cooling rate affects mainly the grain size in the structure of the 3104 alloys and its impact on the type of intermetallic compounds is insignificant, but higher cooling rate is conductives to growth of intermetallic compound type Al6(MnFe). 4. The obtained results on model studies carried out according to the procedure developed by Alcoa were confirmed in industrial conditions upon the studies of the structure of rolling ingots cast by a semi-continuous system.
7
References
[1] P.N.Anyalebechi, Light Metals 1991, p. 821-850
Suppliers Session – Aluminium
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
Horizontal Direct Chilled (HDC) Casting Technology for Aluminium F. Niedermair Hertwich Engineering GmbH, Braunau/Inn, Austria
1
The Universal Caster
HDC casting has well earned its place in modern Al-casthouses, and is still gaining momentum. Hertwich Engineering has successfully commissioned 37 Horizontal Continuous Casting Plants 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: • Foundry ingot • T-bar • Busbar • Extrusion billet • SSM-feedstock etc.
2
Foundry Ingot and T-bar
Over the past few years especially the mass producers of remelt product have discovered the Hertwich Horizontal Casters 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 casting on Vertical Direct Chilled Casting Machines (VDCs) is loosing ground to the over the years developed Horizontal Direct Chilled casting process from Hertwich Engineering. Figure 1 shows T-bars produced on the Hertwich HDC casting machine.
Figure 1: T-bars produced on Hertwich HDC Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
294 Traditionally T-bars were mainly produced on Vertical Direct Chilled casting machines (VDCs). The VDC process has the following drawbacks compared to the HDC process: • Higher costs of the VDC, especially due to higher cast house necessity of a 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 demands more labour. Whereas with the Hertwich HDC, continuous production runs of 3 to 20 days are common. For T-bar production only one operator per shift is needed. • On VDC plants sawing is not integrated in the process, so that an additional sawing station plus operator is required. Horizontal casting employs a automatic flying saw, which cuts the T-bars to length without disturbing the casting process. • The fully continuous HDC process is ideal for automation. This advantage has been fully exploited by Hertwich Engineering. The downstream equipment is fully in line with the casting process and no additional personnel is required. Sawing, weighing, hard stamping, ink marking, labelling, stacking and strapping is carried out fully automatic. Figure 2 shows casting of foundry ingots, Figure 3 shows automatic marking, stacking, strapping and weighing of foundry ingots.
Figure 2: Casting of foundry ingots
Sows, pigs and ingots were traditionally produced employing the open mould technology. Although this technology was improved over the past years, dross formation and inclusions are unavoidable. The HDC process, however, is absolutely free of any 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 closed mould. Thus leaving no chance for oxides and other impurities getting into the product. In contrast to that stands the open mould technology. 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,
295 pouring height and pouring rate. Values achieved during production of pure aluminium sows are shown in table 1.
Figure 3: Foundry ingot automatic marking, stacking, strapping and weighing
Table 1: Values achieved during production of pure aluminium sows. Pouring height [m] Temperature [°C] Dross formation [kg per ton of poured metal] Approx. 0,2 to 0,3 700 –770 0,2 - 0,4 > 800 0,3 - 0,6 Approx. 0,6 to 1,0 approx. 750 2,5 – 4 approx. 850 to 900 5 - 7 The horizontally 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 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. Both items can be delivered by Hertwich Engineering to obtain foundry ingot and T-bars free from porosity and inclusions.
296
3
Extrusion Billet and SSM-feedstock
Especially for extrusion mills and “thixoforming die casters”, a Compact Type Remelt Plant, which includes a Horizontal Direct Chilled casting machine, offers commercial and technological advantages for in-house recycling. The capacity of these plants covers a range of 2.000 – 20.000 tpy. Figure 3 shows the schematic of a Compact Type Remelt Plant.
Figure 4: Schematic of compact type remelt plant
In-house generated extrusion or thixoforming scrap can be charged by means of a 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 permits to remelt profile scrap at a melting loss of lower than 0,5%. Primary metal and clean scrap from the market may be remelt as well. Contaminated scrap, like painted profiles, can be processed in a recently developed 3 chamber furnace. This furnace evaporates and combusts the hydrocarbons from the paint prior to melting. Thus avoiding additional metal losses, 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 and CFF to the Horizontal Direct Chilled Casting machine. Extrusion billets up to 10“ and SSM feedstock up to 6“ are cast as single or multiple strands. For extrusion billets, the logs are directly fed into Hertwich Continuous Homogenizing Furnace for heat treatment. SSM feedstock is produced on the HDC, by fitting a electromagnetic stirrer around the mould, as shown in figure 4. The solidifying metal is stirred in a helicoidal manner. This avoids growth of fringe crystals and leads to a fine grained and uniform globular structure throughout each slug. The HDC Rheocaster produces high quality SSM feedstock, with reproducible thixotropic properties to an attractive low price.
297
Figure 5: HDC Rheocaster casting 4” SSM feedstock
4
Casting of Forging Stock
Forging stock of 25mm to 125mm diameter can be directly produced in 5 to 30 strands casting process. For many forged products forging stock being free from extrusion grain texture is of great advantage. For many applications the cast bars will require a scalping process. However, HDC casting of forging stock plus scalping is in any case more cost effective than billet casting and extruding.
5
Plant Description and Capacities of the Universal Caster
The mechanical structure of the plant embraces the following major parts: Casting conveyor, flying saw and saw run out system (figure 6).
Figure 6: Major assembly groups of a HDC
298 The Horizontal Caster is the key machine in the effective HERTWICH Compact Type Remelt Plant (figure 7)
Figure 7: Layout of a compact type remelt plant for production of extrusion billets from clean and contaminated scrap
In the primary aluminium field, initially a HDC plant is often bought for producing busbar for potline construction. 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. The caster can produce up to 13 tonnes per hour T-bar or 8 t/h foundry ingot. A wide range of alloys can be produced, for instance from pure aluminium to 11% silicon and 5% magnesium. Each product type follows its own exit downstream the flying saw (Figure 8).
Figure 8: Typical layout of the Universal Caster
Once all these exit systems have been installed, a product change can be undertaken within one shift, by changing to a different tundish and mould as well as loading the new applicable cast recipe on the PC. Even a changeover to SSM feedstock would be possible by adding a stirrer.
299 The HERTWICH HDC plants are highly automated. They require only one to two operators per shift. Over the past years the plants were improved consistently and feature now automatic cast starts and stops as well as automatic tundish adjustment. The plant is controlled by the Hertwich PCPLC system, which offers an error manager system as well as menu-type casting recipes. Besides, all important plant parameters are monitored and controlled.
6
Conclusion
The Universal Caster from Hertwich Engineering has become a familiar sight in cast houses and extrusion mills. Its versatility, the low investment costs involved, the high quality products produced and requiring few personnel makes this plant unique.
Automatic “Bleed Out” Detection and Plug Off in VDC Billet Casting Manfred Lück Wagstaff Inc., Spokane WA, USA
1
Casting of Extrusion Billets
The Vertical Direct Chill (VDC) semi continuous casting of extrusion billets employs a water-cooled mold through which molten aluminium is solidified and afterwards cooled by the direct impingement of circular water jets or a water curtain. The solidification of the billet shell, providing physical stability for the billet exiting the mold, takes place within the mold bore (indirect cooling). In this area the shell created has to withstand mechanical forces caused be friction, contraction and shrinkage, including the danger of billet surface damage. The advantages of minimal billet shell zones to the extrusion process led to the development of the hot top casting technology, including the use of shorter mold bores resulting in less time for indirect billet cooling. Consequently the achievable shell thickness decreases, and mold bore friction was minimized by installing continuous oil lubrication systems, using graphite as a mold bore material or a combination of both principals. Modern billet casting systems, which today are ‘state of the art’ technology, in addition to this implant an air cushion between the mold bore and the solidifying billet shell, avoiding any direct metal surface/mold bore contact and so again reduce the heat extraction within this area. Consequently the billet’s shell thickness is reduced further, making it more vulnerable to any mechanical damage which can cause an opening of the billet surface and a run out of molten metal. Such defect, known as a “bleed out” , can be a hazardous situation with the high potential of molten metal/water reactions. The start of a cast includes a higher risk for “bleed outs”, which can be related to variations from established casting practice parameters or to the casting equipment not operating correctly. The practice of preventive maintenance and thorough training of operators have combined to significantly reduce, but not fully eliminate , the possibility of “bleed outs”. Localizing the bleeding position can be challenging and the common practice to manually plug of a bleeding billet by hand with a stop off cone brings an operator close to molten metal. To enhance operator safety and to take a further step into the direction of fully automated billet casting Wagstaff developed a technology called Wagstaff StopCastTM Automatic Bleed Out Detection System (further referred to as StopCast)which includes several new patents. Three major goals were identified at the start of the development project in 1997: • A “bleed out” within a casting table needs to be positively identified as quick as possible. • The “bleed out” identification needs to be reliable. • The “bleed out” identified must be plugged off promptly. The new system was dedicated to become an option to Wagstaff’s MaxiCastTM Billet Casting Systems of which more than 550 – mostly equipped with the Wagstaff AirSlipTM Air Casting Process – are operated worldwide today. During the development within Wagstaff’s research department it turned out to be necessary introducing a new trough design, providing Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
301 metal to each mold individually, to make the system work reliable. The StopCast system itself consists of individual dams, located at the entry of each mold, which are raised and lowered by a guided air cylinder. At the start of a cast, the dams are in the up position. With an identified bleed out the air cylinder is activated, rotates the dam to a position above the mold entry, lowers the dam and so shuts off the metal flow to the affected mold. Each dam position is connected to the “bleed out” detection system which is integrated into conventional AirSlip billet moulds.(Although not demanded yet MaxiCast moulds can be equipped with this feature too). The mold has a groove machined into its alignment lugs around the mold’s circumference where a sacrificial nylon tube is placed to act as a detection circuit providing 360° of bleed out detection. The pressurized tube ruptures when contacted by molten metal and the releasing air pressure activates the related StopCast dam via a pilot valve. A control logic assures a lowered dam cannot be opened again until the cast is finished. StopCast was first installed at Alcoa Lafayette, Indiana, USA in June 1998, and has been developed there in a production environment . A second system was commissioned at Aluminium Ranshofen (AMAG) in April 2000 and is in constant operation since, a third system for this plant will be commissioned in December 2000.
2
Summary
A „bleed-out“ is a condition where molten metal runs out of an opening of the billet surface during VDC casting. A „bleed-out“ can be a hazardous situation with the potential of water/metal reactions. When a „bleed-out“ occurs, the bleeding position must be positively identified and plugged off promptly. The use of modern billet casting systems, the practice of preventive maintenance, and thorough training of operators have combined to significantly reduce, but not fully eliminate, the possibility of bleed-outs. Localizing the bleeding position can be challenging and the common practice to manually plug off a bleeding billet by hand with a stop off cone brings an operator close to molten metal. As an advanced step in VDC billet casting a system for automatic bleed-out detection and plug off is presented. Experimental results and field results from the first European application in a production environment are reported.
3
Acknowledgements
Wagstaff Inc. wishes to acknowledge and thank the Aluminium Ranshofen Hüttengießereigesellschaft m.b.H., Austria, especially Mr. Carl van Gils and Mr. Helmut Suppan for strongly supporting this work and contributing to this presentation.
The AIRSOL VEIL® Technology Package for Aluminium Billet Casting Gerd W. Bulian VAW Aluminium Technologie GmbH, Bonn, Germany
Manfred Langen VAW aluminium AG, Bonn, Germany
1
Summary
Mould systems working with air are the dominant technology for casting aluminium billets. To achieve high billet quality, it is important to have a reliable air feed to the mould and a uniform distribution of the melt temperature within a multiple casting unit. The AIRSOL VEIL® Technology Package, consisting of moulds, designed with the aid of pressure-loss calculations and a computerised air-feeding system, guarantees uniform air distribution over the mould circumference as well as optimum feeding of the necessary air-flow rate during casting. Also incorporated in the mould system unit is an optimised metal distribution system (launder) and an advanced design of starter blocks. The launder system, optimised by means of computer modelling, reduces the temperature variation across the casting unit to less than 10° Celsius. The Technology Package is rounded off by a starter block design which avoids the formation of starting cracks in the billets. This and the uniform melt temperature distribution lead to a significant reduction in the cut-off length of the billet butt.
2
The AIRSOL VEIL® Technology Package
The first production unit for extrusion billets and forging ingots came on stream in 1985 [1]. It soon became clear, however, that an air-regulating system operated by conventional flow adjustment alone was insufficient. On the one hand, the correct backpressure in the system must be at the exactly same level as the metallostatic pressure formed by the melt height in the mould. On the other hand, the required airflow rates should remain within a defined range depending on the mould diameter. [2] Conventional multiple casting systems operate with constant airflow rates and the same amount of air in every mould. As noted previously, this should be controlled individually to provide optimum billet quality. The AIRSOL VEIL® Technology Package consists of a specially designed mould, a computerised air feeding system, an optimised launder system and a specially designed starter block. The different parts of the technology will be described in more detail in the following sections of this paper.
Continuous Casting. Edited by K. Ehrke and W. Schneider. © WILEY-VCH Verlag GmbH, Weinheim. ISBN: 3-527-30283-2
303
3
The AIRSOL VEIL® Mould
The AIRSOL VEIL® mould is an advanced VAW Hot Top Mould with an optimised indirect and direct cooling system . This type of mould was introduced into VAW casthouses between 1971 and 1974. An advantage of the mould is that there is no dependence on costly special spare parts, such as graphite rings. This greatly simplifies the maintenance procedures. Figure 1 shows two current generation AIRSOL VEIL® moulds, one round and one rectangular. These moulds display a major modification compared with the Hot Top System, i.e. the addition of an air and oil distribution ring. This is shown in Figures 2 and 3.
Figure 1: AIRSOL VEIL® mould
Figure 2: Configuration of rings
304 The channels in the oil and air distribution ring were designed using pressure-loss calculations [3,4]. To ensure uniform air and oil supply, all channel sizes have been carefully optimised. Another advantage is the fact, that no mineral oil is used as a lubricant but only vegetable oils such as rape seed oil, which does not generate any cracked, cancerogenic vapour during casting. Oil residues in the water can be effectively skimmed off by a hydrocyclone system in the water basin. Top of Plate
Bottom of Plate
Air distribution
Lubricant distribution
Figure 3: Air and oil distribution plate
4
Computerised Air-Feeding System
Besides the mould system, the main reason for the high billet quality is the sophisticated aircontrol system, which allows individual pressure control of the air supply to each mould [5]. Using this system, the respective air pressure is always equal to the metallostatic pressure in every mould. In conventional casting systems with an air cushion, the airflow can become too high, which creates overpressure. The resulting surplus air escapes from the mould by bubbling through the melt surface and creating oxides in the melt. If the air flow is lower than the metallostatic pressure, the heat flow through the mould wall is higher; this results in a thicker surface segregation layer. In AIRSOL VEIL® moulds, these quality-impairing effects are eliminated by means of the computerised air-control system, which regulates each individual mould. Figure 4 shows an air-control cabinet for up to 48 moulds. This system is based on a process control computer, which allows the required pressure of the air-volume flow rate to all moulds to be set simultaneously and independently during the continuous casting process. Depending on the selected casting phase, the relevant control-system variable is either the required air-volume rate to the mould (start-up phase with a constant, high air-volume rate, end phase with a constant low air-volume rate) or the pending air pressure in the mould (constant air pressure during the casting phase). The desired values for the maximum or minimum air-volume flow rates, as well as for the air pressure, can be set by means of the relevant input menu on the operator control panel.
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Figure 4: Air-control cabinet
A special feature of this system lies in the fact that the control-loop parameters are adjustable via the operator interface. This facilitates the individual adjustment of the control loop to the specific requirements, which can vary according to the composition of the alloy and the layout of the compressed-air system. In this way, the pressure loss in the tubing between the control cabinet and the moulds is adjusted to zero by the system.
Figure 5: Visualisation of process parameters
306 Figure 5 shows the main control menu during casting. The screen displays 4 moulds with their parameters simultaneously. This is the optimum number since more would make the screen unintelligible.
Figure 6: Mould menu during the stationary pressure-control phase. Channel #21 is disturbed
The display, which normally appears on the screen during casting, is shown in Figure 6. This figure shows the mould parameters during the stationary pressure-control phase of the cast. In this case, the condition indicator reports a disturbance in Channel #21 (red alarm light). Detailed information about the operating parameters of the affected channel can be obtained through the Control menu.
5
Launder System
The launder system was designed using numerical modelling of the starting or filling time as well as for the casting period [6,7,8]. The optimisation resulted in very even filling of all moulds and also gives a very small variance in the temperature of the metal entering each individual mould. Figure 7 shows the simulation of the metal velocity in the launder. Figure 8 shows the filling of a 48-strand casting unit before and after optimisation. The filling behaviour of the new launder design is considerably improved. Figure 9 shows the photo of a launder system of a 72-strand 7" casting unit. The temperature difference between the coldest and warmest region of this unit is less than 10 °C.
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Figure 7: Metal velocity of a 48-strand launder
Original
Improved
Figure 8: Filling of a 48-strand launder system before and after optimisation
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Figure 9: Launder system of a 72-strand casting unit (K 7