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© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
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ASM Handbook Volume 15 Casting Prepared under the direction of the ASM International Handbook Committee
Editorial Committee Srinath Viswanathan, University of Alabama, Chair Diran Apelian, Worcester Polytechnic Institute Raymond J. Donahue, Mercury Marine Babu DasGupta, National Science Foundation Michael Gywn, ATI Inc John L. Jorstad, J.L.J. Technologies, Inc Raymond W. Monroe, Steel Founders’ Society of America Mahi Sahoo, CANMET Materials Technology Laboratory Thomas E. Prucha, American Foundry Society Daniel Twarog, North American Die Casting Association Steve Lampman, Project Editor Charles Moosbrugger, Editor Eileen DeGuire, Editor Madrid Tramble, Senior Production Coordinator Ann Britton, Editorial Assistant Diane Whitelaw, Production Coordinator Kathryn Muldoon, Production Assistant Scott D. Henry, Senior Product Manager Bonnie R. Sanders, Manager of Production
Editorial Assistance Elizabeth Marquard Heather Lampman WPR Indexing Service
Materials Park, Ohio 44073-0002 www.asminternational.org
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
Copyright # 2008 by ASM InternationalW All rights reserved No part of this book may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording, or otherwise, without the written permission of the copyright owner. First printing, December 2008
This book is a collective effort involving hundreds of technical specialists. It brings together a wealth of information from worldwide sources to help scientists, engineers, and technicians solve current and long-range problems. Great care is taken in the compilation and production of this Volume, but it should be made clear that NO WARRANTIES, EXPRESS OR IMPLIED, INCLUDING, WITHOUT LIMITATION, WARRANTIES OF MERCHANTABILITY OR FITNESS FOR A PARTICULAR PURPOSE, ARE GIVEN IN CONNECTION WITH THIS PUBLICATION. Although this information is believed to be accurate by ASM, ASM cannot guarantee that favorable results will be obtained from the use of this publication alone. This publication is intended for use by persons having technical skill, at their sole discretion and risk. Since the conditions of product or material use are outside of ASM’s control, ASM assumes no liability or obligation in connection with any use of this information. No claim of any kind, whether as to products or information in this publication, and whether or not based on negligence, shall be greater in amount than the purchase price of this product or publication in respect of which damages are claimed. THE REMEDY HEREBY PROVIDED SHALL BE THE EXCLUSIVE AND SOLE REMEDY OF BUYER, AND IN NO EVENT SHALL EITHER PARTY BE LIABLE FOR SPECIAL, INDIRECT OR CONSEQUENTIAL DAMAGES WHETHER OR NOT CAUSED BY OR RESULTING FROM THE NEGLIGENCE OF SUCH PARTY. As with any material, evaluation of the material under end-use conditions prior to specification is essential. Therefore, specific testing under actual conditions is recommended. Nothing contained in this book shall be construed as a grant of any right of manufacture, sale, use, or reproduction, in connection with any method, process, apparatus, product, composition, or system, whether or not covered by letters patent, copyright, or trademark, and nothing contained in this book shall be construed as a defense against any alleged infringement of letters patent, copyright, or trademark, or as a defense against liability for such infringement. Comments, criticisms, and suggestions are invited, and should be forwarded to ASM International. Library of Congress Cataloging-in-Publication Data ASM International ASM Handbook Includes bibliographical references and indexes Contents: v.1. Properties and selection—irons, steels, and high-performance alloys—v.2. Properties and selection—nonferrous alloys and special-purpose materials—[etc.]—v.21. Composites 1. Metals—Handbooks, manuals, etc. 2. Metal-work—Handbooks, manuals, etc. I. ASM International. Handbook Committee. II. Metals Handbook. TA459.M43 1990 620.1’6 90-115 SAN: 204-7586 ISBN-13: 978-0-87170-711-6 ISBN-10: 0-87170-711-X
ASM InternationalW Materials Park, OH 44073-0002 www.asminternational.org Printed in the United States of America Multiple copy reprints of individual articles are available from Technical Department, ASM International.
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
Dedicated to the memory of JOSEPH R. DAVIS (1954–2007) Editor, Metals Handbook, 1982–1991
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
Foreword
In this revision of ASM HandbookW Volume 15 on casting, ASM International is indebted to the volunteer efforts of the Volume 15 Editorial Committee and over 150 participants who helped as authors or reviewers. Their professional commitment and efforts represent a continuing devotion to the practice of metalcasting and the publication of peer consensus information on it. Special thanks are extended to Srinath Viswanathan, University of Alabama, for recruiting an outstanding Editorial Committee with Diran Apelian, Worcester Polytechnic Institute; Babu DasGupta, National Science Foundation, Raymond J. Donahue, Mercury Marine; Michael Gywn, ATI Inc.; John L. Jorstad, J.L.J. Technologies, Inc.; Raymond W. Monroe, Steel Founders’ Society of America; Thomas E. Prucha, American Foundry Society; Kumar Sadayappan and Mahi Sahoo, CANMET Materials Technology Laboratory; Edward S. Szekeres, Casting Consultants Incorporated; and Daniel Twarog, North American Die Casting Association. We thank them and the other contributors for this publication.
Dianne Chong President ASM International Stanley C. Theobald Managing Director ASM International
iv
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Policy on Units of Measure
By a resolution of its Board of Trustees, ASM International has adopted the practice of publishing data in both metric and customary U.S. units of measure. In preparing this Handbook, the editors have attempted to present data in metric units based primarily on Syste`me International d’Unite´s (SI), with secondary mention of the corresponding values in customary U.S. units. The decision to use SI as the primary system of units was based on the aforementioned resolution of the Board of Trustees and the widespread use of metric units throughout the world. For the most part, numerical engineering data in the text and in tables are presented in SI-based units with the customary U.S. equivalents in parentheses (text) or adjoining columns (tables). For example, pressure, stress, and strength are shown both in SI units, which are pascals (Pa) with a suitable prefix, and in customary U.S. units, which are pounds per square inch (psi). To save space, large values of psi have been converted to kips per square inch (ksi), where 1 ksi = 1000 psi. The metric tonne (kg 103) has sometimes been shown in megagrams (Mg). Some strictly scientific data are presented in SI units only. To clarify some illustrations, only one set of units is presented on artwork. References in the accompanying text to data in the illustrations are presented in both SI-based and customary U.S. units. On graphs and charts, grids corresponding to SI-based units usually appear along the left and bottom edges. Where appropriate, corresponding customary U.S. units appear along the top and right edges. Data pertaining to a specification published by a specification-writing group may be given in only the units used in that specification or in dual units, depending on the nature of the data. For example, the typical yield strength of steel sheet made to a specification written in customary U.S.
units would be presented in dual units, but the sheet thickness specified that specification might be presented only in inches. Data obtained according to standardized test methods for which the standard recommends a particular system of units are presented in the units of that system. Wherever feasible, equivalent units are also presented. Some statistical data may also be presented in only the original units used in the analysis. Conversions and rounding have been done in accordance with IEE/ ASTM SI-10, with attention given to the number of significant digits in the original data. For example, an annealing temperature of 1570 F contains three significant digits. In this case, the equivalent temperature would be given as 855 C; the exact conversion to 854.44 C would not be appropriate. For an invariant physical phenomenon that occurs at a precise temperature (such as the melting of pure silver), it would be appropriate to report the temperature as 961.93 C or 1763.5 F. In some instances (especially in tables and data compilations), temperature values in C and F are alternatives rather than conversions. The policy of units of measure in this Handbook contains several exceptions to strict conformance to IEEE/ASTM SI-10; in each instance, the exception has been made in an effort to improve the clarity of the Handbook. The most notable exception is the use of g/cm3 rather than kg/m3 as the unit of measure for density (mass per unit volume). SI practice requires that only one virgule (diagonal) appear in units formed by combination of several basic units. Therefore, all of the units preceding the virgule are in the numerator and all units following the virgule are in the denominator of the expression; no parentheses are required to prevent ambiguity.
v
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Preface
From the preceding Volume 5 (1970) of the 8th Edition Metals Handbook and the 9th Edition Metals Handbook Volume 15 (1988) on casting, this Handbook provides an update on the continuing advances in casting technologies and applications. Casting, as both a science and practical tool in art and technology, is enormously varied in scope. It is impossible to capture the full scope of casting technology in one volume.
on sand casting is expanded relative to the previous edition with the intention of providing a reference that may be helpful as a communication tool between product designers and metalcasters in developing successful and economical products. This Volume consists of 18 sections. The first section introduces the historical development of metal casting, as well as to the advantages of castings over parts produced by other manufacturing processes, their applications, and the current market size of the industry. This includes an article on “Metalcasting Technology and the Purchasing Process” written by Al Spada and the technical staff of the American Foundry Society. Then, the principles and practice of melt processing are described in the next three sections followed by a section on principles of solidification including nucleation kinetics, fundamentals of growth, transformation behavior, and microstructure development. Solid-state processing of casting, such as heat treat treatment and hot isostatic pressing, are also introduced. This is followed by a section on the “Modeling and Analysis of Casting Processes.”
The main focus of this Volume is on the products and processes of foundry (shape) casting, although primary (ingot or continuous casting) of steel and aluminum are also covered. In addition, continuous casting of copper is described in an article, as copper continuous casting was a precursor to steel and aluminum continuous casting in some respects. Some of the articles on melt processing, such as the articles on “Electric Arc Furnace Melting” and “Steel Melt Processing,” also briefly describe primary production of cast metal. Shape casting of metal is dominated by cast iron, which constitutes just over 70% of the worldwide production of castings on a tonnage basis (See Table 2 in the first article “History and Trends of Metal Casting”). This is followed by steel, copper-alloy, and aluminum-alloy castings, which make up about 25% on the worldwide tonnage of casting production. Magnesium and zinc are on the order of 1% or less. The dominance of just a few alloys in shape casting is due to the fact that successful and economic shape casting typically involves alloy compositions near a eutectic. The lower melting points and narrower freezing range of near-eutectic compositions promote better castability.
Like the previous edition, traditional subjects such as patterns, molding and casting processes, foundry equipment, and processing considerations are extensively covered in the next sections. As noted, coverage on sand casting has been consolidated and expanded. For example, the major method of shell molding is described in an article—based on an update of a still largely valid 1970 handbook (Volume 5) article. New updates are also provided on processes growing in use, such as squeeze casting, lost-foam casting, semisolid metal forming, and low-pressure casting. The latter is particularly important in producing quality products, as described by John Campbell in the article “Filling and Feeding Concepts.”
Since the 1988 edition of Volume15, several developments have occurred (see Table 5 in the first article “History and Trends of Metal Casting”). Of these developments, computer technology continues to shorten development time and help simulate the casting process. Automation and robotic technology also has improved the productivity and process control of casting. In terms of processes, semisolid processing, squeeze casting, lost-foam, vacuum molding, and various dies casting technologies continue to improve and finds new applications. These important topics are updated in this Volume. For nonferrous alloys, high-pressure die casting of aluminum is a major area of expansion and update in this volume.
Finally, the last five sections describe the major types of cast alloys in term of processing and the properties and characteristics of cast ferrous and nonferrous alloys. Emphasis is placed on cast iron, cast steel, aluminum, copper, and zinc. The last section covers the quality aspects of cast products and the processing of castings. It is hoped that this Handbook is a useful work of peer-consensus reference information for the producers, designers, and buyers of castings. Many thanks are extended to all the contributors and the editors who worked on this Volume. This publication would not have been possible without their commitment and effort.
In addition, coverage on sand casting is expanded and consolidated in this Volume with major articles on “Green Sand Molding,” No-Bake Sand Molding,” and “Shell Molding and Shell Coremaking.” Bonded sand mold casting, although well-established for many years, is the most widely used method of casting on a tonnage basis. Improvement in methods and materials continue to provide better yields, productivity, and product quality. The sand system is also a major factor in the economics of large-volume, production casting. Coverage
Srinath Viswanathan University of Alabama Chair, Volume 15 Editorial Committee
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© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Officers and Trustees of ASM International (2007–2008)
Dianne Chong President and Trustee The Boeing Company Roger J. Fabian Vice President and Trustee Bodycote Thermal Processing Lawrence C. Wagner Immediate Past President and Trustee Texas Instruments Inc (retired) Paul L. Huber Treasurer and Trustee Seco/Warwick Corporation Stanley C. Theobald Secretary and Managing Director ASM International
Trustees
Sue S. Baik-Kromalic Honda of America Christopher C. Berndt Swinburne University of Technology Brady G. Butler University of Utah Danniel P. Dennies The Boeing Company Leigh C. Duren Worcester Polytechnic Institute Pradeep Goyal Pradeep Metals Ltd.
Digby D. Macdonald Penn State University Subhash Mahajan Arizona State University Charles A. Parker Honeywell Aerospace Megan M. Reynolds Washington State University Mark F. Smith Sandia National Laboratories Jon D. Tirpak ATI
Members of the ASM Handbook Committee (2007–2008) Larry D. Hanke (Chair 2006–; Member 1994–) Materials Evaluation and Engineering Inc. Kent L. Johnson (Vice Chair 2006–; Member 1999–) Engineering Systems Inc. Viola L. Acoff (2005–) University of Alabama David E. Alman (2002–2008) National Energy Technology Laboratory Tim Cheek (2004–) DELTA (v) Forensic Engineering Lichun Leigh Chen (2002–) Technical Materials Incorporated Sarup K. Chopra (2007–) Nook Industries Craig Clauser (2005–) Craig Clauser Engineering Consulting Incorporated Craig V. Darragh (1989–)
The Timken Company Jon L. Dossett (2006–) Consultant David U. Furrer (2006–) Rolls-Royce Corporation Lee Gearhart (2005–2008) Moog Inc. Ernest W. Klechka (2006–) Lloyd Register Capstone Inc. Alan T. Male (2003–) University of Kentucky William L. Mankins (1989–) Metallurgical Services Inc. Dana J. Medlin (2005–2008) South Dakota School of Mines and Technology Joseph W. Newkirk (2005–) University of Missouri-Rolla Cory J. Padfield (2006–) Hyundai America Technical Center Inc.
Toby V. Padfield (2004–) ZF Sachs Automotive of America Charles A. Parker (2007–) Honeywell Aerospace Elwin L. Rooy (2007–) Elwin Rooy & Associates Karl P. Staudhammer (1997–) Los Alamos National Laboratory Kenneth B. Tator (1991–) KTA-Tator Inc. George F. Vander Voort (1997–) Buehler Ltd.
Chairs of the ASM Handbook Committee J.F. Harper (1923–1926) (Member 1923–1926) W.J. Merten (1927–1930) (Member 1923–1933) L.B. Case (1931–1933) (Member 1927–1933) C.H. Herty, Jr. (1934–1936) (Member 1930–1936) J.P. Gill (1937) (Member 1934–1937) R.L. Dowdell (1938–1939) (Member 1935–1939) G.V. Luerssen (1943–1947) (Member 1942–1947) J.B. Johnson (1948–1951) (Member 1944–1951) E.O. Dixon (1952–1954) (Member 1947–1955) N.E. Promisel (1955–1961) (Member 1954–1963)
R.W.E. Leiter (1962–1963) (Member 1955–1958, 1960–1964) D.J. Wright (1964–1965) (Member 1959–1967) J.D. Graham (1966–1968) (Member 1961–1970) W.A. Stadtler (1969–1972) (Member 1962–1972) G.J. Shubat (1973–1975) (Member 1966–1975) R. Ward (1976–1978) (Member 1972–1978) G.N. Maniar (1979–1980) (Member 1974–1980) M.G.H. Wells (1981) (Member 1976–1981) J.L. McCall (1982) (Member 1977–1982) L.J. Korb (1983) (Member 1978–1983)
vii
T.D. Cooper (1984–1986) (Member D.D. Huffman (1986–1990) (Member D.L. Olson (1990–1992) (Member 1989–1992) R.J. Austin (1992–1994) (Member W.L. Mankins (1994–1997) (Member M.M. Gauthier (1997–1998) (Member C.V. Darragh (1999–2002) (Member Henry E. Fairman (2002–2004) (Member Jeffrey A. Hawk (2004–2006) (Member Larry D. Hanke (2006–2008) (Member
1981–1986) 1982–) 1982–1988, 1984–1985) 1989–) 1990–2000) 1989–) 1993–) 1997–) 1994–)
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Authors and Contributors
University of Birmingham, UK Kevin R. Anderson Mercury Marine
Paul G Campbell Metaullics Systems Division Pyrotek, Inc.
Princewill N. Anyalebechi Padnos College of Engineering & Computing, Grand Valley State University
Robert D Carnahan Thixomat Inc.
Diran Apelian Metal Processing Institute, Worcester Polytechnic Institute
Sujoy Chaudhury Honeywell International, Inc. K. K. Chawla University of Alabama At Birmingham
Alan P. Druschitz University of Alabama at Birmingham Cor van Ettinger Gieterij Doesburg Henry “Ed” Fairman Cincinnati Metallurgical Consultants Edward W. Flynn General Motors
Leigh Chen Technical Materials Incorporated
Jim Frost ACIPCO (American Cast Iron Pipe Company)
Henry Bakemeyer Die Casting Design and Consulting
Yeou-Li Chu Ryobi Die Casting
Richard Fruehan Carnegie Mellon University
Bryan Baker Vulcan Engineering Co.
Craig D. Clauser Craig Clauser Engineering Consulting Incorporated
Rafael Gallo Foseco Metallurgical
Lars Arnberg Norwegian University Of Science & Technology
Steve L. Cockroft The University of British Columbia
Charles-Andre´ Gandin Centre de Mise en Forme des Mate´riaux (CEMEF)
Chris Cooper Vulcan Engineering Co.
George Giles Morganite Crucible Inc
James D. Cotton The Boeing Co.
Peter C. Glaws The Timken Company
Randall Counselman Oldenburg Group Incorporated
Martin E. Glicksman University of Florida
Joe Bigelow CONTECH U. S. LLC.
Jonathan A. Dantzig University of Illinois-Urbana
George Goodrich Stork Climax Research Services
Ron Bird Stainless Foundry & Engineering
Atef Daoud Central Metallurgical Research and Development Institute, Cairo, Egypt
Frank E. Goodwin International Lead Zinc
Stewart A Ballantyne Allvac Allegheny Technologies Company Scott A. Balliett Latrobe Specialty Steel Company Christoph Beckermann University Of Iowa, The Michel Bellet Ecole des Mines de Paris, Centre de Mise en Forme des Mate´riaux (CEMEF)
Malcolm Blair Steel Founders Society of America William Boettinger NIST James A. Brock Sr. Alcoa Harold D. Brody University Of Connecticut Craig C. Brown Stork Technimet, Inc. Zach Brown CONTECH U.S. LLC Andreas Buhrig-Polackzek Gieberei-Institut, Germany Willam A. Butler General Motors (Retired) John Campbell
Craig V. Darragh The Timken Company Babu DasGupta National Science Foundation Bob Dawson Honeywell Aerospace MCOE Raymond Decker Thixomat Incorporated Raymond J Donahue Mercury Marine Daniel Dos Santos Gießerei Institut der RWTH Aachen Jon Dossett Metallurgical Consultant Jean Marie Drezet Ecole Polytechnique Fe´de´rale de Lausanne
viii
Lee Gouwens CMI Novacast Inc. Ronald Graham ATI Wah Chang Patrick Grant Oxford University Daniel E.Groteke Q.C. Designs Inc. Mustafa Guclu Stanley Associates Inc. Richard Gundlach Stork Climax Research Services Nikhil Gupta Polytechnic University Mike Gwyn ATI, Inc. John Hall
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
Hall Permanent Mold Machines Qingyou Han Purdue University Larry D. Hanke Materials Evaluation and Engineering Inc Justin Heimsch Madison Kipp Corporation
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Makhlouf M. Makhlouf Metal Processing Institute, Worcester Polytechnic Institute William L. Mankins Metallurgical Services Incorporated Stephen J. Mashl Bodycote HIP Andover
Intermet-PCPC Robert Pischel FOSECO Metallurgical Inc. David R. Poirier University of Arizona David Poweleit Steel Founders’ Society of America
Juan Carlos Heinrich The University of New Mexico
Ragnvald H. Mathiesen Norwegian University of Science and Technology
Gregg Holman Wah Chang
Robert J. McInerney BuhlerPrince
Thomas E. Prucha American Foundry Society
Robert A. Horton Precision Metalsmiths Incorporated
Scott McIntyre Grede Foundrie
Donald J. Hurtuk H.C. Starck Inc.
Brad McKee Deloro Stellite Inc.
Peter Quested Industry and Innovation Division, National Physical Laboratory, UK
Alain Jacot Laboratoire de Simulation des Mate´riaux
Harold T. Michels Copper Development Association
John L. Jorstad J.L.J. Technologies, Inc.
Tony Midea Foseco Metallurgical Inc.
Ursula R. Kattner N I S T/Metallurgy Division
Paul H Mikkola Metal Casting Technology Incorporated
Seymour Katz S. Katz Associates
Asbj rn Mo SINTEF Materials and Chemistry
Jay S. Keist Worcester Polytechnic Institute
Raymond W. Monroe Steel Founders’ Society of America
Graham Keough Consarc Corporation
Y.V. Murty Cellular Materials International, Inc.
Matthew Krane Purdue University
Laurentiu Nastac Concurrent Technologies Corporation
Mary Beth Krusiak Sand Technology Co. LLC
David V. Neff Metaullics Systems Division Pyrotek, Inc.
T.A. Kuhn Alcoa Primary Metals
Charles Nelson Morris Bean and Company
Wilfried Kurz Ecole Polytechnique Federale De Lausanne (EPFL) Switzerland
Itsuo Ohnaka Osaka Sangyo University, Japan
Selcuk Kuyucak CANMET Materials Technology Laboratory Vic La Fay Hill & Griffith Company Robin A. Lampson Retech Systems LLC
Cory Padfield American Axle & Manufacturing Qingyue Pan SPX Corporation
William L. Powell ThyssenKrupp-Waupaca (Retired)
Wayne Rasmussen Consultant Martin Reeves Striko-Dynarad Allen Richardson SPX Corporation Pradeep Rohatgi University of Wisconsin-Milwaukee Elwin L. Rooy Elwin Rooy and Associates Gary Ruff SMW Automotive Corporation Kumar (Muthukumarasamy) Sadayappan CANMET Materials Technology Laboratory Mahi Sahoo CANMET Materials Technology Laboratory Kenneth Savage Magnesium Elektron Gary Scholl Metal Casting Technology Incorporated Ted Schorn Enkei America, Inc. Benjamin F. Schultz University of Wisconsin-Milwakee David Schwam Case Werstern Reserve University
Charles A. Parker Honeywell Aerospace
William D. Scott AAA Alchemy
David Lee Howmet Corp.
Kent Peaslee Missouri University of Science and Technology (Missouri S&T)
Ben Q. Li University of Michigan Dearborn
Pehlke, Robert D. Univ of Michigan
Sumanth Shankar McMaster University
Alan Luo General Motors R&D Center
John H. Perepezko University of Wisconsin, Madison
Geoffrey K. Sigworth Alcoa Primary Metals
Dan Maas Ex-One
Ralph Perkul Foundry Solutions and Design LLC
Steve Sikkenga Cannon-Muskegon Corporation
James F. Major Alcan International Ltd
Gordon H. Peters
Brian V. Smith
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John Serra Carpenter Brothers Inc.
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
General Motors Corporation, Powertrain Division Jerry Sokolowski University of Windsor Al Spada American Foundry Society Karl Staudhammer Los Alamos National Laboratory Doru M. Stefanescu The Ohio State University Ingo Steinbach RWTH-Aachen, ACCESS e. V. Karthik Subbiah SPX Product Solutions Group Erin H. Sunseri Iowa State University Tony Suschil Foseco Inc Richard Sutherlin ATI Wah Chang Edward S.Szekeres Casting Consultants Incorporated Brian G. Thomas
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University of Illinois Michael L Tims Concurrent Technologies Corporation David W. Tripp TIMET
Srinath Viswanathan University of Alabama Birmingham Robert C. Voigt Penn State University Vaughan R. Voller University of Minnesota
Rohit Trivedi AMES Laboratory Paul Trojan University of Michigan-Dearborn John Troxler Rolls Royce
Tom Waring ME Elecmetal Sufei Wei Centrifugal Casting Machine Company Incorporated David White Schaefer Furnace
Nao Tsumagari American Showa, Inc. Daniel Twarog North American Die Casting Association Derek E. Tyler Olin Metals Thin Strip (retired) Don Tyler General Aluminum Manufacturing Company Stephen Udvardy North American Die Casting Association Juan J. Valencia Concurrent Technologies Corporation
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Lawrence Whiting CANMET Materials Technology Laboratory Richard Williams FOSECO Metallurgical Inc. Greg G.Woycik Woycik Consulting Oscar Yu RTI International Metals, Inc. Keith Zhang Teck Cominco Metals Ltd.
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Contents
Applications and Specification Guidelines. . . . . . . . . . . . . . . . . . . 1 Raymond W. Monroe, Steel Founders’ Society of America; Thomas E. Prucha, American Foundry Society; Daniel Twarog, North American Die Casting Association History and Trends of Metal Casting. . . . . . . . . . . . . . History of Metal Casting . . . . . . . . . . . . . . . . . . . . Metal Casting Methods. . . . . . . . . . . . . . . . . . . . . . Production Trends . . . . . . . . . . . . . . . . . . . . . . . . . Market Trends . . . . . . . . . . . . . . . . . . . . . . . . . . . . Metalcasting Technology and the Purchasing Process . . . . The Purchasing Process . . . . . . . . . . . . . . . . . . . . . Define Requirements and Develop a Purchasing Plan. Requesting and Evaluating Quotations/Bids . . . . . . . Select a Supplier and Negotiate Contract Terms . . . . Contract Fulfillment and Continuous Improvement . . Considerations in Purchasing Metal Castings . . . . . .
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Principles of Liquid Metal Processing . . . . . . . . . . . . Chemical Thermodynamics and Kinetics . . . . . . . . . . . . Chemical Thermodynamics . . . . . . . . . . . . . . . . . . Chemical Kinetics . . . . . . . . . . . . . . . . . . . . . . . . Kinetics of Alloys Additions . . . . . . . . . . . . . . . . . Thermodynamic Properties of Iron-Base Alloys . . . . . . . Purification of Ferrous Melts. . . . . . . . . . . . . . . . . Thermodynamics of Ferrous Systems . . . . . . . . . . . Multicomponent Iron-Carbon Systems . . . . . . . . . . Structural Diagrams . . . . . . . . . . . . . . . . . . . . . . . Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys . . . . . . . . . . . . . . . . . . . . Activities and Thermal Properties . . . . . . . . . . . . . Ineraction Coefficients . . . . . . . . . . . . . . . . . . . . . Thermal Properties for Hypothetical Standard State . Phase Diagrams . . . . . . . . . . . . . . . . . . . . . . . . . . Gases in Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cast Iron. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aluminum and Its Alloys . . . . . . . . . . . . . . . . . . . Copper and Its Alloys . . . . . . . . . . . . . . . . . . . . . Magnesium and Its Alloys . . . . . . . . . . . . . . . . . . Overcoming Gas Porosity . . . . . . . . . . . . . . . . . . . Inclusion-Forming Reactions . . . . . . . . . . . . . . . . . . . . Inclusion Types . . . . . . . . . . . . . . . . . . . . . . . . . . Physical Chemistry . . . . . . . . . . . . . . . . . . . . . . . Control of Inclusions . . . . . . . . . . . . . . . . . . . . . . Inclusions in Ferrous Alloys . . . . . . . . . . . . . . . . . Inclusions in Nonferrous Alloys . . . . . . . . . . . . . .
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56 56 56 58 58 64 64 67 70 71 73 74 74 75 76 77 81
Melting and Remelting . . . . . . . . . . Electric Arc Furnace Melting . . . . . . . Power Supply . . . . . . . . . . . . . . Arc Furnace Components . . . . . . Weighing and Chemical Analysis Scrap and Alloy Storage . . . . . . .
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Acid Melting Practice . . . . . . . . . . . . . . . . . Basic Melting Practice . . . . . . . . . . . . . . . . . Newer Technologies. . . . . . . . . . . . . . . . . . . Cupola Furnaces . . . . . . . . . . . . . . . . . . . . . . . . . Progressive Refinements in Cupola Equipment Cupola Construction and Operation . . . . . . . . Charge Materials . . . . . . . . . . . . . . . . . . . . . Cupola Control Principles . . . . . . . . . . . . . . . Control Tests and Analyses. . . . . . . . . . . . . . Specialized Cupolas . . . . . . . . . . . . . . . . . . . Induction Furnaces . . . . . . . . . . . . . . . . . . . . . . . Types of Furnaces . . . . . . . . . . . . . . . . . . . . Electromagnetic Stirring . . . . . . . . . . . . . . . . Power Supplies . . . . . . . . . . . . . . . . . . . . . . Water Cooling Systems . . . . . . . . . . . . . . . . Lining Material . . . . . . . . . . . . . . . . . . . . . . Melting Operations . . . . . . . . . . . . . . . . . . . Vacuum Induction Melting . . . . . . . . . . . . . . . . . Process Description . . . . . . . . . . . . . . . . . . . VIM Metallurgy . . . . . . . . . . . . . . . . . . . . . Production of Nonferrous Materials . . . . . . . . Electroslag Remelting . . . . . . . . . . . . . . . . . . . . . ESR Furnaces . . . . . . . . . . . . . . . . . . . . . . . Ingot Solidification. . . . . . . . . . . . . . . . . . . . Process Variations . . . . . . . . . . . . . . . . . . . . Steel ESR . . . . . . . . . . . . . . . . . . . . . . . . . . Superalloy ESR . . . . . . . . . . . . . . . . . . . . . . Vacuum Arc Remelting . . . . . . . . . . . . . . . . . . . . Process Description . . . . . . . . . . . . . . . . . . . Solidification . . . . . . . . . . . . . . . . . . . . . . . . Process Variables. . . . . . . . . . . . . . . . . . . . . Superalloy VAR . . . . . . . . . . . . . . . . . . . . . Titanium VAR . . . . . . . . . . . . . . . . . . . . . . Skull Melting . . . . . . . . . . . . . . . . . . . . . . . . . . . Vacuum Arc Skull Melting . . . . . . . . . . . . . . Furnaces . . . . . . . . . . . . . . . . . . . . . . . . . . . Induction Skull Melting . . . . . . . . . . . . . . . . Electron Beam Melting . . . . . . . . . . . . . . . . . . . . Process Description . . . . . . . . . . . . . . . . . . . Drip Melting . . . . . . . . . . . . . . . . . . . . . . . . Cold Hearth Melting . . . . . . . . . . . . . . . . . . Plasma Melting and Heating . . . . . . . . . . . . . . . . Plasma Torches . . . . . . . . . . . . . . . . . . . . . . Furnace Equipment . . . . . . . . . . . . . . . . . . . Atmosphere Control . . . . . . . . . . . . . . . . . . . Melting Processes . . . . . . . . . . . . . . . . . . . . Crucible Furnaces . . . . . . . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . . Furnace Types . . . . . . . . . . . . . . . . . . . . . . . Design and Operation. . . . . . . . . . . . . . . . . . Reverberatory and Stack Furnaces . . . . . . . . . . . . Furnace Types . . . . . . . . . . . . . . . . . . . . . . . Reverberatory Furnace Practice . . . . . . . . . . . Molten Metal Circulation . . . . . . . . . . . . . . .
3 4 8 9 14 16 16 17 21 22 22 22
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91 95 98 99 99 100 103 106 106 107 108 109 109 111 111 113 116 116 119 122 124 125 128 129 129 131 132 132 134 135 137 137 139 139 139 140 142 142 143 145 149 149 149 151 151 151 155 155 156 158 160 161 163 165
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
Molten Metal Processing and Handling . . . . . . . . . . . . . Transfer and Treatment of Molten Metal—An Introduction . Launders. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Tundishes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Molten-Metal Pumps . . . . . . . . . . . . . . . . . . . . . . . . Pouring and Dosing . . . . . . . . . . . . . . . . . . . . . . . . . Degassing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aluminum Foundry Degassing . . . . . . . . . . . . . . . . . Degassing of Magnesium . . . . . . . . . . . . . . . . . . . . . Degassing of Copper Alloys . . . . . . . . . . . . . . . . . . . Molten-Metal Filtration . . . . . . . . . . . . . . . . . . . . . . . . . . Sources of Inclusion. . . . . . . . . . . . . . . . . . . . . . . . . Removal of Inclusions . . . . . . . . . . . . . . . . . . . . . . . Filter Characteristics . . . . . . . . . . . . . . . . . . . . . . . . Types of Molten-Metal Filters . . . . . . . . . . . . . . . . . . In-Mold Filtration . . . . . . . . . . . . . . . . . . . . . . . . . . In-Furnace Filtration of Molten Aluminum . . . . . . . . . Filtered Metal Quality . . . . . . . . . . . . . . . . . . . . . . . Steel Melt Processing . . . . . . . . . . . . . . . . . . . . . . . . . . . Furnaces and Refractories . . . . . . . . . . . . . . . . . . . . . Converter Metallurgy . . . . . . . . . . . . . . . . . . . . . . . . Ladle Metallurgy . . . . . . . . . . . . . . . . . . . . . . . . . . . Aluminum Fluxes and Fluxing Practice . . . . . . . . . . . . . . . Aluminum Fluxing. . . . . . . . . . . . . . . . . . . . . . . . . . Categories of Fluxes . . . . . . . . . . . . . . . . . . . . . . . . Foundry Practices . . . . . . . . . . . . . . . . . . . . . . . . . . Modification of Aluminum-Silicon Alloys . . . . . . . . . . . . . Heat Treated Castings . . . . . . . . . . . . . . . . . . . . . . . Porosity and Shrinkage Formation . . . . . . . . . . . . . . . Modification Mechanisms . . . . . . . . . . . . . . . . . . . . . Eutectic Nucleation . . . . . . . . . . . . . . . . . . . . . . . . . Eutectic Grain Structure . . . . . . . . . . . . . . . . . . . . . . Recommended Foundry Practices . . . . . . . . . . . . . . . Growth of Silicon Eutectic . . . . . . . . . . . . . . . . . . . . Effect of Other Elements . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Grain Refinement of Aluminum Casting Alloys . . . . . . . . . Measurement of Grain Size. . . . . . . . . . . . . . . . . . . . Mechanisms of Grain Refinement . . . . . . . . . . . . . . . Boron as a Grain Refiner . . . . . . . . . . . . . . . . . . . . . Best Practices for Grain Refinement. . . . . . . . . . . . . . Benefits of Grain Refinement . . . . . . . . . . . . . . . . . . Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . Refinement of the Primary Silicon Phase in Hypereutectic Aluminum-Silicon Alloys . . . . . . . . . . . . . . . . . . . . . Importance of Primary Silicon Refinement . . . . . . . . . Accomplishing Primary Silicon Refinement . . . . . . . . Principles of Solidification . . . . . . . . . . . . . . . . Diran Apelian, Worcester Polytechnic Institute Thermodynamics and Phase Diagrams . . . . . . . . Thermodynamics . . . . . . . . . . . . . . . . . . . . Phase Diagrams . . . . . . . . . . . . . . . . . . . . . Solidification . . . . . . . . . . . . . . . . . . . . . . . Nucleation Kinetics and Grain Refinement . . . . . Thermodynamics of Solidification . . . . . . . . Nucleation Phenomena . . . . . . . . . . . . . . . . Inoculation Practice . . . . . . . . . . . . . . . . . . Grain Refinement Models . . . . . . . . . . . . . . Master Alloy Processing . . . . . . . . . . . . . . . Transport Phenomena During Solidification . . . . . Fundamentals . . . . . . . . . . . . . . . . . . . . . . Transport and Microstructure . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . Plane Front Solidification . . . . . . . . . . . . . . . . . Transients and Steady State . . . . . . . . . . . . Morphological Stability . . . . . . . . . . . . . . .
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171 173 174 174 175 178 185 185 188 189 194 194 196 197 199 199 201 204 206 206 208 212 230 231 233 235 240 241 242 243 244 245 247 249 249 250 255 255 255 257 258 260 260
Rapid Solidification Effects. . . . . . . . . . . . . . . . . . . . . . Solidification Microstructure Selection Maps . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Non-Plane Front Solidification . . . . . . . . . . . . . . . . . . . . . . . Cellular Microstructures . . . . . . . . . . . . . . . . . . . . . . . . Dendrite Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nonplanar Eutectic Growth. . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Eutectic Solidification—An Introduction . . . . . . . . . . . . . . . . Binary Eutectic Phase Diagram . . . . . . . . . . . . . . . . . . . Eutectic Structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solidification of Eutectic Alloys—Aluminum-Silicon . . . . . . . The Aluminum-Silicon Equilibrium Phase Diagram . . . . . Classification of the Aluminum-Silicon Eutectic Structure Microstructure of the Aluminum-Silicon Eutectic . . . . . . Theories of Solidification of the Aluminum-Silicon Eutectic . . . . . . . . . . . . . . . . . . Theories of Chemical Modification of the Aluminum-Silicon Eutectic . . . . . . . . . . . . . . . . . . Solidification of Eutectic Alloys—Cast Iron. . . . . . . . . . . . . . Structure of Liquid Iron-Carbon Alloys . . . . . . . . . . . . . Nucleation of Eutectic in Cast Iron . . . . . . . . . . . . . . . . Nucleation of Primary Phases in Cast Iron . . . . . . . . . . . Growth of Eutectic in Cast Iron. . . . . . . . . . . . . . . . . . . Cooling Curve Analysis . . . . . . . . . . . . . . . . . . . . . . . . Peritectic Solidification . . . . . . . . . . . . . . . . . . . . . . . . . . . . Peritectic Solidification. . . . . . . . . . . . . . . . . . . . . . . . . Peritectic Reaction . . . . . . . . . . . . . . . . . . . . . . . . . . . . Peritectic Transformation and Direct Precipitation . . . . . . Primary Metastable Precipitation of Beta . . . . . . . . . . . . Peritectic Cascades . . . . . . . . . . . . . . . . . . . . . . . . . . . Peritectic Transformations in Multicomponent Systems . . Microsegregation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solidification Models . . . . . . . . . . . . . . . . . . . . . . . . . . Binary Isomorphous Systems, k > 1: Titanium-Molybdenum . . . . . . . . . . . . . . . . . . . . . Binary Eutectic Systems, k > 1: Hypoeutectic Aluminum-Copper Alloys . . . . . . . . . . . . . . . . . . . Binary Eutectic Systems: Hypoeutectic Aluminum-Silicon Alloys . . . . . . . . . . . . . . . . . . . Binary Peritectic Systems: Copper-Zinc a-Brass . . . . . . . Multicomponent, Multiphase Eutectic Systems: Hypoeutectic Al-Si-Cu-Mg Alloys . . . . . . . . . . . . . . . Multicomponent Multiphase Steel: Fe-C-Cr . . . . . . . . . . Multicomponent, Multiphase Nickel-Base Superalloy. . . . Macrosegregation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Macrosegregation Induced by Flow of the Liquid . . . . . . Macrosegregation Induced by Movement of the Solid . . . Interpretation and use of Coolng Curves—Thermal Analysis . . Quantitative Thermal Analysis . . . . . . . . . . . . . . . . . . . Deviation from Equilibrium . . . . . . . . . . . . . . . . . . . . . Differential Thermal Analysis . . . . . . . . . . . . . . . . . . . . Determining Liquidus and Solidus Temperature . . . . . . . X-Ray Imaging of Solidification Processes and Microstructure Evolution . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . X-Ray Imaging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . In-House Laboratory X-Ray Imaging of Solidification Processes . . . . . . . . . . . . . . . . . . . . . X-Ray Imaging Studies of Solidification Processes with Synchrotron Radiation . . . . . . . . . . . . . . . . . . . . Further Reading. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . The Oxide Film Defect—The Bifilm. . . . . . . . . . . . . . . . . . . Bifilms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Entrainment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Furling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Unfurling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Reliability and Reproducibility of Mechanical Properties .
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© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
www.asminternational.org
Central Role of Bifilm Defects in Fracture . . . . . . . . . . . Shrinkage Porosity and Gas Porosity. . . . . . . . . . . . . . . . . . . Shrinkage Porosity . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas Porosity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Factors Affecting Porosity Formation . . . . . . . . . . . . . . . Castability—Fluidity and Hot Tearing. . . . . . . . . . . . . . . . . . Fluidity. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hot Tearing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Semisolid Metal Processing . . . . . . . . . . . . . . . . . . . . . . . . . Flow Behavior of SSM. . . . . . . . . . . . . . . . . . . . . . . . . Semisolid Microstructural Formation . . . . . . . . . . . . . . . Semisolid Processes: Thixocasting versus Rheocasting . . . Spray Casting. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Principles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Atomization and the Droplet Spray . . . . . . . . . . . . . . . . Deposition and Grain Mulitplication . . . . . . . . . . . . . . . Microsegregation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Macrosegregation and Coarsening . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rapid Solidification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microstructural Modification . . . . . . . . . . . . . . . . . . . . . Solidification During Casting of Metal-Matrix Composites . . . Incorporation of Reinforcements . . . . . . . . . . . . . . . . . . Reinforcement-Metal Wettability . . . . . . . . . . . . . . . . . . Solidification Processing of MMC . . . . . . . . . . . . . . . . . Solidification Fundamentals . . . . . . . . . . . . . . . . . . . . . Modeling of Particle-Pushing Phenomenon . . . . . . . . . . . Nanocomposites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solidification Research in Microgravity. . . . . . . . . . . . . . . . . Availability of Reduced Gavity: Past, Present, and Future Solidification Research in Microgravity . . . . . . . . . . . . . Key Solidification Experiments . . . . . . . . . . . . . . . . . . . The Future of Microgravity Solidification . . . . . . . . . . . . Homogenization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sources of Inhomogeneity. . . . . . . . . . . . . . . . . . . . . . . Benefits of Homogenization Treatment. . . . . . . . . . . . . . Characterization Methods . . . . . . . . . . . . . . . . . . . . . . . Computational Modeling . . . . . . . . . . . . . . . . . . . . . . . Heat Treatment. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Designations of Heat Treatment . . . . . . . . . . . . . . . . . . Solution Heat Treatment. . . . . . . . . . . . . . . . . . . . . . . . Quenching . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aging . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Emerginhg Heat Treating Technologies . . . . . . . . . . . . . Microstructural Changes due to Heat Treatment of Cast Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hot Isostatic Pressing of Castings. . . . . . . . . . . . . . . . . . . . . History of HIP. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Reasons for Using HIP . . . . . . . . . . . . . . . . . . . . . . . . . Overview of a HIP System . . . . . . . . . . . . . . . . . . . . . . Effect of HIP on Mechanical Properties . . . . . . . . . . . . . Effect of Inclusions on As-HIPed Properties . . . . . . . . . . Effect of HIP on the Shape and Structure of Castings . . . Problems Encountered in HIP . . . . . . . . . . . . . . . . . . . . Economics of HIP . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Modeling and Analysis of Casting Processes. . . . . . . . Srinath Viswanathan, University of Alabama Numerical Methods for Casting Applications . . . . . . . . . Fundamentals of CFD . . . . . . . . . . . . . . . . . . . . . Numerical Solution of the Fluid-Flow Equations . . . Grid Generation for Complex Geometries. . . . . . . . Casting Applications . . . . . . . . . . . . . . . . . . . . . . Modeling of Transport Phenomena and Electromagnetics Conservation of Thermal Energy . . . . . . . . . . . . . .
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406 408 408 409 409 411 412 413 414 415 415
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419 419 420 420 422 425 425
Conservation of Solute . . . . . . . . . . . . . . . . . . . . . . . . . Equations Cast in Advection-Diffusion Form . . . . . . . . . Basic Equations for Mass and Momentum Conservation . Modeling the Momentum in the Solid + Liquid Mushy Region . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Modeling Turbulence . . . . . . . . . . . . . . . . . . . . . . . . . . Phenomena Restraining or Driving Flow . . . . . . . . . . . . Dealing with Fronts and Free Surfaces . . . . . . . . . . . . . . Simulation of Electromagnetic Effects in Castings. . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Direct Modeling of Structure Formation . . . . . . . . . . . . . . . . Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Direct Microstructure Simulation Using the Phase Field Method. . . . . . . . . . . . . . . . . . . . . . . . Direct Grain Structure Simulation Using the Cellular Automation Method . . . . . . . . . . . . . . . . . Coupling of Direct Structure Simulation at Macroscopic Scale . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Modeling of Microsegregation and Macrosegregation . . . . . . . Microsegregation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microsegregation with Diffusion in the Solid . . . . . . . . . Continuum Model of Macrosegregation . . . . . . . . . . . . . Modeling of Stress, Distortion, and Hot Tearing . . . . . . . . . . Governing Equations . . . . . . . . . . . . . . . . . . . . . . . . . . Thermomechanical Coupling . . . . . . . . . . . . . . . . . . . . . Numerical Solution . . . . . . . . . . . . . . . . . . . . . . . . . . . Model Validation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . Example Applications . . . . . . . . . . . . . . . . . . . . . . . . . Hot-Tearing Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Practical Issues In Computer Simulation of Casting Processes . Setting Simulation Objectives, Modeling, and Selection of Simulation Code . . . . . . . . . . . . . . . . . . . . . . . . Modeling of Shape and Phenomena . . . . . . . . . . . . . . . . Initial and Boundary Conditions and Physical Properties . Enmeshing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Evaluation of Simulation Results . . . . . . . . . . . . . . . . . . Estimation of Structure and Properties . . . . . . . . . . . . . . Optimization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermophysical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . Sources and Availability of Reliable Data . . . . . . . . . . . Limitations and Warning on the Use of Data . . . . . . . . . Methods to Determine Thermophysical Properties . . . . . . Specific Heat Capacity and Enthalpy of Transformation . . Enthalpy of Melting, Solidus, and Liquidus Temperatures Coefficient of Thermal Expansion . . . . . . . . . . . . . . . . . Density . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Surface Tension. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Viscosity. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Electrical and Thermal Conductivity . . . . . . . . . . . . . . . Emissivity. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Typical Themophysical Properties Ranges of Some Cast Alloys . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Principles and Practices of Shape Casting . . . . . . . Shape Casting Processes—An Introduction . . . . . . . . Patterns and Patternmaking . . . . . . . . . . . . . . . . . . . Types of Patterns . . . . . . . . . . . . . . . . . . . . . . . Additional Features of Patterns . . . . . . . . . . . . . Pattern Allowances . . . . . . . . . . . . . . . . . . . . . Pattern Materials . . . . . . . . . . . . . . . . . . . . . . . Selection of Pattern Type . . . . . . . . . . . . . . . . . Trends in Pattern Design and Manufacture . . . . . Casting Practice—Guidelines for Effective Production of Reliable Castings . . . . . . . . . . . . . . . . . . . . .
xiii
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428 428 429 429 431 432 435 435
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441 443 445 445 445 446 449 449 451 451 453 454 456 458 462
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462 462 463 464 465 466 466 468 468 468 468 470 470 471 471 471 474 474 474
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483 485 488 488 490 491 492 494 495
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© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
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Rule 1: Good-Quality Melt . . . . . . . . . . . . . . . . Rule 2: Liquid Front Damage . . . . . . . . . . . . . . Rule 3: Liquid Front Stop. . . . . . . . . . . . . . . . . Rule 4: Bubble Damage . . . . . . . . . . . . . . . . . . Rule 5: Core Blows . . . . . . . . . . . . . . . . . . . . . Rule 6: Shrinkage Damage . . . . . . . . . . . . . . . . Rule 7: Convection Damage . . . . . . . . . . . . . . . Rule 8: Segregation . . . . . . . . . . . . . . . . . . . . . Rule 9: Residual Stress. . . . . . . . . . . . . . . . . . . Rule 10: Location Points . . . . . . . . . . . . . . . . . Filling and Feeding System Concepts . . . . . . . . . . . . Entrainment Defects in Gravity-Poured Castings . Traditional Filling Designs . . . . . . . . . . . . . . . . Prepriming Techniques . . . . . . . . . . . . . . . . . . . Naturally Pressurized Filling System . . . . . . . . . Filters. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Performance of Alternative Filling Systems . . . . Recommended Foundry Practice . . . . . . . . . . . . Processing and Finishing of Castings . . . . . . . . . . . . Shakeout and Core Knockout . . . . . . . . . . . . . . Cleaning Operations. . . . . . . . . . . . . . . . . . . . . Automating Gate Removal and Grinding . . . . . . Flame Cutting . . . . . . . . . . . . . . . . . . . . . . . . . Blast Cleaning of Castings . . . . . . . . . . . . . . . . Inspection . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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Expandable-Mold Casting Processes with Permanent Patterns. . . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc Introduction—Expendable Mold Processes with Permanent Patterns . . . . . . . . . . . . . Aggregates and Binders for Expendable Molds Mold Media . . . . . . . . . . . . . . . . . . . . . Foundry Sands. . . . . . . . . . . . . . . . . . . . Clays for Green Sand Molding . . . . . . . . Inorganic Binders . . . . . . . . . . . . . . . . . Organic Binders. . . . . . . . . . . . . . . . . . . Other Expendable Mold Media . . . . . . . . Green Sand Molding . . . . . . . . . . . . . . . . . . . Sand Molding . . . . . . . . . . . . . . . . . . . . System Control and Formulation . . . . . . . Influence of Molding Equipment . . . . . . . Maintaining Sand System Quality . . . . . . Molding Methods. . . . . . . . . . . . . . . . . . Green Sand Media Preparation . . . . . . . . Molding Problems . . . . . . . . . . . . . . . . . Mold Finishing . . . . . . . . . . . . . . . . . . . Shakeout. . . . . . . . . . . . . . . . . . . . . . . . Sand/Casting Recovery. . . . . . . . . . . . . . Sand Reclamation . . . . . . . . . . . . . . . . . No-Bake Sand Molding . . . . . . . . . . . . . . . . . No-Bake Sand Processing . . . . . . . . . . . . CO2 -Cured Sodium Silicate . . . . . . . . . . Air-Setting Sodium Silicates . . . . . . . . . . Air-Setting Organic Binders . . . . . . . . . . Gas-Cured Organic Binders. . . . . . . . . . . Sand Reclamation and Reuse . . . . . . . . . Coremaking . . . . . . . . . . . . . . . . . . . . . . . . . Binders for Sand Coremaking . . . . . . . . . Basic Principles of Coremaking. . . . . . . . Core Sands and Additives. . . . . . . . . . . . Compaction. . . . . . . . . . . . . . . . . . . . . . Curing of Compacted Cores . . . . . . . . . . Core Coatings (Core Washes) . . . . . . . . . Assembly and Core Setting . . . . . . . . . . . Design Considerations . . . . . . . . . . . . . . Coring of Tortuous Passages . . . . . . . . . .
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497 497 499 500 501 501 502 503 504 505 506 506 506 507 508 511 511 512 513 513 516 517 519 519 520
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525 528 528 529 535 538 539 546 549 549 551 552 553 554 558 560 560 561 562 564 567 567 569 572 573 576 577 581 581 582 583 584 585 587 588 589 592
Green Sand versus Dry Sand Cores . . . . . Cores in Permanent Mold Castings . . . . . Cores in Investment Castings . . . . . . . . . Designing to Eliminate Cores . . . . . . . . . Coring versus Drilling . . . . . . . . . . . . . . Shell Molding and Shell Coremaking . . . . . . . Applicability . . . . . . . . . . . . . . . . . . . . . Properties and Selection of Shell Sands . . Resins . . . . . . . . . . . . . . . . . . . . . . . . . Preparation of Resin-Sand Mixtures . . . . . Sand Reclamation . . . . . . . . . . . . . . . . . Control Testing of Resin-Sand Properties . Patterns and Core Boxes. . . . . . . . . . . . . Production of Molds and Cores . . . . . . . . Mold Assembly . . . . . . . . . . . . . . . . . . . Casting Process . . . . . . . . . . . . . . . . . . . Dimensional Accuracy . . . . . . . . . . . . . . Causes and Prevention of Casting Defects Comparisons with Green Sand Molding . . Slurry Molding . . . . . . . . . . . . . . . . . . . . . . . Plaster Molding . . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . Plaster Mold Compositions . . . . . . . . . Patterns and Core Boxes . . . . . . . . . . . Mold-Drying Equipment . . . . . . . . . . . Metals Cast in Plaster Molds . . . . . . . . Conventional Plaster Mold Casting. . . . Match Plate Pattern Molding . . . . . . . . The Antioch Process . . . . . . . . . . . . . . Foamed Plaster Molding Process . . . . . Ceramic Molding . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . Shaw Process . . . . . . . . . . . . . . . . . . . Unicast Process . . . . . . . . . . . . . . . . . No-Bond Sand Molding. . . . . . . . . . . . . . . . . Magnetic Molding . . . . . . . . . . . . . . . . . Vacuum (V-Process) Molding . . . . . . . . .
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Expandable-Mold Casting Processes with Expendable Patterns . . . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc Introduction—Expendable Mold Processes with Expendable Patterns . . . . . . . . . . . . . . . . . Lost Foam Casting . . . . . . . . . . . . . . . . . . . . . Process Technique . . . . . . . . . . . . . . . . . . Processing Parameters . . . . . . . . . . . . . . . Advantages of Lost Foam Casting . . . . . . . Casting Quality . . . . . . . . . . . . . . . . . . . . Investment Casting . . . . . . . . . . . . . . . . . . . . . Pattern Materials . . . . . . . . . . . . . . . . . . . Patternmaking . . . . . . . . . . . . . . . . . . . . . Pattern and Cluster Assembly . . . . . . . . . . Ceramic Shell Mold Manufacture . . . . . . . Manufacture of Ceramic Cores . . . . . . . . . Pattern Removal . . . . . . . . . . . . . . . . . . . Mold Firing and Burnout . . . . . . . . . . . . . Melting and Casting. . . . . . . . . . . . . . . . . Postcasting Operations . . . . . . . . . . . . . . . Inspection and Testing . . . . . . . . . . . . . . . Design Advantages of Investment Castings . Design Recommendations . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . Special Investment Casting Processes. . . . . Replicast Molding. . . . . . . . . . . . . . . . . . . . . . Process Details . . . . . . . . . . . . . . . . . . . . Process Capabilities . . . . . . . . . . . . . . . . .
xiv
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593 593 594 595 596 598 599 599 600 601 603 603 604 604 607 610 611 614 615 617 617 617 617 618 618 618 619 620 620 621 622 622 622 624 628 628 628
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637 640 640 642 645 645 646 646 648 650 651 654 654 654 655 656 656 657 657 657 657 662 662 663
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
Centrifugal Casting . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc Centrifugal Casting . . . . . . . . . . . . . . . . . . Centrifugal Casting Methods . . . . . . . . Equipment . . . . . . . . . . . . . . . . . . . . Molds. . . . . . . . . . . . . . . . . . . . . . . . Defects in Centrifugal Castings . . . . . . Applications . . . . . . . . . . . . . . . . . . . Horizontal Centrifugal Casting . . . . . . . . . . Equipment . . . . . . . . . . . . . . . . . . . . Casting Process . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . Vertical Centrifugal Casting. . . . . . . . . . . . Mold Design . . . . . . . . . . . . . . . . . . . Process Details . . . . . . . . . . . . . . . . .
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Permanent Mold and Semipermanent Mold Processes . . . . . . . . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc Permanent Mold Casting . . . . . . . . . . . . . . . . . . Gravity Casting Methods . . . . . . . . . . . . . . Advantages and Disadvantages . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . Casting Design . . . . . . . . . . . . . . . . . . . . . Mold Design . . . . . . . . . . . . . . . . . . . . . . . Mold Materials . . . . . . . . . . . . . . . . . . . . . Mold Coatings. . . . . . . . . . . . . . . . . . . . . . Mold Life . . . . . . . . . . . . . . . . . . . . . . . . . Mold Temperature . . . . . . . . . . . . . . . . . . . Control of Mold Temperature . . . . . . . . . . . Dimensional Accuracy . . . . . . . . . . . . . . . . Surface Finish . . . . . . . . . . . . . . . . . . . . . . Low-Pressure Die Casting . . . . . . . . . . . . . . . . . Conventional Low-Pressure Casting . . . . . . . Counterpressure Casting . . . . . . . . . . . . . . . Vacuum Riserless/Pressure Riserless Casting Low Pressure Counter Gravity Casting . . . . . . . . Mold Filling . . . . . . . . . . . . . . . . . . . . . . .
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Die Casting Tooling . . . . . . . . . . . . . . . . . . . . . . . Die Casting Machines and Dies. . . . . . . . . . . . Product Design and the Process. . . . . . . . . . . . Tooling and Metal Flow . . . . . . . . . . . . . . . . . Heat Removal . . . . . . . . . . . . . . . . . . . . . . . . Materials Selection for Tooling . . . . . . . . . . . . Automation in High-Pressure Die Casting . . . . . . . . Die Casting Cycle . . . . . . . . . . . . . . . . . . . . . Automatic Pouring . . . . . . . . . . . . . . . . . . . . . Injection Process . . . . . . . . . . . . . . . . . . . . . . Solidification and Heat Removal . . . . . . . . . . . Casting Extraction and Insert Loading . . . . . . . Automatic Spray . . . . . . . . . . . . . . . . . . . . . . Finishing Automation . . . . . . . . . . . . . . . . . . . Die Casting Internal Quality Check Automation
667 667 669 670 670 671 674 674 675 678 680 680 684
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High-Pressure Die Casting. . . . . . . . . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc High-Pressure Die Casting . . . . . . . . . . . . . . . . . . . . . . . . . . Die Casting Alloys and Processes . . . . . . . . . . . . . . . . . Advantages of High-Pressure Die Casting. . . . . . . . . . . . Disadvantages of High-Pressure Die Casting. . . . . . . . . . Product Design for the Process . . . . . . . . . . . . . . . . . . . Metal Injection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hot Chamber Die Casting . . . . . . . . . . . . . . . . . . . . . . . . . . Melting Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Injection Components. . . . . . . . . . . . . . . . . . . . . . . . . . Distinctions Between Hot and Cold Chamber Processes . . Gate and Runner Design. . . . . . . . . . . . . . . . . . . . . . . . Temperature Control . . . . . . . . . . . . . . . . . . . . . . . . . . Ejection and Postprocessing . . . . . . . . . . . . . . . . . . . . . Cold Chamber Die Casting . . . . . . . . . . . . . . . . . . . . . . . . . Machine Components . . . . . . . . . . . . . . . . . . . . . . . . . . Process Parameters. . . . . . . . . . . . . . . . . . . . . . . . . . . . Squeeze Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Squeeze Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Casting and Tooling Design for Squeeze Casting . . . . . . Factors Affecting Squeeze-Cast Products . . . . . . . . . . . . Squeeze-Cast Applications . . . . . . . . . . . . . . . . . . . . . . Commonly Used Alloys, Properties, and Microstructures . Vacuum High-Pressure Die Casting . . . . . . . . . . . . . . . . . . . Conventional Die Casting . . . . . . . . . . . . . . . . . . . . . . . Vacuum Die Casting . . . . . . . . . . . . . . . . . . . . . . . . . . High-Vacuum Die Casting . . . . . . . . . . . . . . . . . . . . . .
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689 690 690 691 691 692 693 695 697 697 698 698 699 700 700 704 706 709 709
. . . 713 . . . . . . . . . . . . . . . . . . . . . . . . . .
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715 715 716 717 717 717 719 719 719 720 721 721 722 724 724 725 727 727 727 728 729 730 732 732 732 732
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Semisolid Casting . . . . . . . . . . . . . . . . . . . . . . . . . John L. Jorstad, J.L.J. Technologies, Inc Semisolid Casting — Introduction and Fundamentals . Advantages of SSM Processing . . . . . . . . . . . . . SSM Processing Routes . . . . . . . . . . . . . . . . . . Thixocasting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bar-Making Processes . . . . . . . . . . . . . . . . . . . Billet Sawing and Volume Control . . . . . . . . . . Reheating and Billet Handling. . . . . . . . . . . . . . Rheological Tests . . . . . . . . . . . . . . . . . . . . . . Basics of Semisolid Tooling Design. . . . . . . . . . Casting Process and Parameters. . . . . . . . . . . . . Rheocasting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thixomolding. . . . . . . . . . . . . . . . . . . . . . . . . . . . . Process Description . . . . . . . . . . . . . . . . . . . . . Emerging Developments . . . . . . . . . . . . . . . . . . Mechanical Properties and Microstructure . . . . .
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Cast Irons . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thomas E. Prucha, American Foundry Society Introduction to Cast Irons . . . . . . . . . . . . . . . . . . . Classification of Cast Irons . . . . . . . . . . . . . . . Composition of Cast Irons . . . . . . . . . . . . . . . Solidification of Cast Irons . . . . . . . . . . . . . . . Graphite Morphologies . . . . . . . . . . . . . . . . . . Matrix Structure Development. . . . . . . . . . . . . Physical Properties of Cast Irons . . . . . . . . . . . Thermal Properties. . . . . . . . . . . . . . . . . . . . . Conductive Properties. . . . . . . . . . . . . . . . . . . Magnetic Properties . . . . . . . . . . . . . . . . . . . . Acoustic Properties . . . . . . . . . . . . . . . . . . . . Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . Welding of Cast Irons . . . . . . . . . . . . . . . . . . Machining and Grinding . . . . . . . . . . . . . . . . . Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . Materials Selection . . . . . . . . . . . . . . . . . . . . Cast Iron Foundry Practices . . . . . . . . . . . . . . . . . . Cast Iron Melting Practice . . . . . . . . . . . . . . . Gray Iron Foundry Practice. . . . . . . . . . . . . . . Ductile Iron Foundry Practice . . . . . . . . . . . . . Compacted Graphite Iron Foundry Practice. . . . Malleable Iron Foundry Practice . . . . . . . . . . . Foundry Practice for High-Nickel Ductile Irons Foundry Practice High-Silicon Ductile Irons . . . Foundry Practice for High-Silicon Gray Irons . . Foundry Practice for High-Alloy White Irons . . Gray Iron Castings . . . . . . . . . . . . . . . . . . . . . . . . Classes of Gray Iron . . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . . . Castability. . . . . . . . . . . . . . . . . . . . . . . . . . .
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734 734 735 737 740 742 747 747 748 749 751 753 754 755 757
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761 761 762 764 764 766 767 768 769 770 773 773 777 777 778 780
. . . . . . . . . . 783 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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785 785 788 791 792 794 794 795 797 800 801 802 805 806 809 809 812 812 815 820 826 829 830 830 830 831 835 835 835 835
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . Section Sensitivity . . . . . . . . . . . . . . . . . . . . . . . Prevailing Sections. . . . . . . . . . . . . . . . . . . . . . . Test Bar Properties . . . . . . . . . . . . . . . . . . . . . . Specifications . . . . . . . . . . . . . . . . . . . . . . . . . . Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Fatigue Limit in Reversed Bending . . . . . . . . . . . Pressure Tightness . . . . . . . . . . . . . . . . . . . . . . . Impact Resistance . . . . . . . . . . . . . . . . . . . . . . . Machinability . . . . . . . . . . . . . . . . . . . . . . . . . . Wear . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Resistance to Wear . . . . . . . . . . . . . . . . . . . . . . Elevated-Temperature Properties . . . . . . . . . . . . . Dimensional Stability . . . . . . . . . . . . . . . . . . . . . Effect of Shakeout Practice. . . . . . . . . . . . . . . . . Alloying to Modify As-Cast Properties. . . . . . . . . Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . . Physical Properties. . . . . . . . . . . . . . . . . . . . . . . Ductile Iron Castings . . . . . . . . . . . . . . . . . . . . . . . . Classes and Grades of Ductile Iron . . . . . . . . . . . Factors Affecting Mechanical Properties. . . . . . . . Hardness Properties . . . . . . . . . . . . . . . . . . . . . . Tensile Properties . . . . . . . . . . . . . . . . . . . . . . . Shear and Torsional Properties . . . . . . . . . . . . . . Compressive Properties . . . . . . . . . . . . . . . . . . . Fatigue Properties . . . . . . . . . . . . . . . . . . . . . . . Fracture Toughness . . . . . . . . . . . . . . . . . . . . . . Physical Properties. . . . . . . . . . . . . . . . . . . . . . . Austempered Ducfile Iron (ADI) . . . . . . . . . . . . . Compacted Graphite Iron Castings . . . . . . . . . . . . . . . Graphite Morphology . . . . . . . . . . . . . . . . . . . . . Chemical Composition . . . . . . . . . . . . . . . . . . . . Castability. . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mechanical Properties and Corrosion Resistance . . Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . Malleable Iron Castings. . . . . . . . . . . . . . . . . . . . . . . Melting Practices . . . . . . . . . . . . . . . . . . . . . . . . Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . Current Production Technologies . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . Ferritic Malleable Iron . . . . . . . . . . . . . . . . . . . . Pearlitic and Martensitic Malleable Iron . . . . . . . . White Iron and High-Alloyed Iron Castings. . . . . . . . . Nickel-Chromium White Irons . . . . . . . . . . . . . . High-Chromium White Irons. . . . . . . . . . . . . . . . High-Alloy Graphitic Irons . . . . . . . . . . . . . . . . . . . . High-Silicon Irons for High-Temperature Service . Austenitic Nickel-Alloyed Gray and Ductile Irons . High-Silicon Irons for Corrosion Resistance . . . . .
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Steel Castings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Raymond W. Monroe, Steel Founders’ Society of America; Thomas E. Prucha, American Foundry Society Steel Ingot Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Pouring Method. . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ingot Design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hot Tops . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ingot Molds and Stools . . . . . . . . . . . . . . . . . . . . . . . Ingot Solidification. . . . . . . . . . . . . . . . . . . . . . . . . . . Ingot Stripping . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solidification Simulation. . . . . . . . . . . . . . . . . . . . . . . Steel Continuous Casting. . . . . . . . . . . . . . . . . . . . . . . . . . Historical Aspects of Continuous Casting of Steel . . . . . General Description of the Process . . . . . . . . . . . . . . . Plant Layout . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Near-Net Shape Casting . . . . . . . . . . . . . . . . . . . . . . . Tundish Metallurgy . . . . . . . . . . . . . . . . . . . . . . . . . .
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836 837 838 839 840 841 843 844 846 846 848 848 849 849 850 851 851 853 ... 856 857 859 859 861 861 862 865 867 869 872 872 872 873 874 879 884 884 885 887 888 889 890 896 896 899 904 904 906 907
. . . . 909 . . . . . . . . . . . . . .
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911 911 912 913 913 914 916 916 918 918 918 919 919 920
Pouring Stream Protection . . . . . . . . . . . . . . . . . . Molds for Continuous Casting of Steel. . . . . . . . . . Productivity Improvements . . . . . . . . . . . . . . . . . . Quality Improvements . . . . . . . . . . . . . . . . . . . . . Horizontal Continuous Casting . . . . . . . . . . . . . . . Future Developments . . . . . . . . . . . . . . . . . . . . . . Shape Casting of Steel . . . . . . . . . . . . . . . . . . . . . . . . Castability of Steel. . . . . . . . . . . . . . . . . . . . . . . . Melting Furnaces . . . . . . . . . . . . . . . . . . . . . . . . . Metlting Practice . . . . . . . . . . . . . . . . . . . . . . . . . Pouring Practice . . . . . . . . . . . . . . . . . . . . . . . . . Foundry Factors in Design . . . . . . . . . . . . . . . . . . Design Factors. . . . . . . . . . . . . . . . . . . . . . . . . . . Avoidance of Hot Spots . . . . . . . . . . . . . . . . . . . . Thin-Wall Steel Castings . . . . . . . . . . . . . . . . . . . Tolerances . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Size of Cores . . . . . . . . . . . . . . . . . . . . . . . . . . . Casting in Graphite Molds . . . . . . . . . . . . . . . . . . Steel Castings Properties . . . . . . . . . . . . . . . . . . . . . . . Carbon Steels . . . . . . . . . . . . . . . . . . . . . . . . . . . Low-Alloy Steels . . . . . . . . . . . . . . . . . . . . . . . . . Wear-Resistant Steels . . . . . . . . . . . . . . . . . . . . . . Corrosion-Resistant Steels. . . . . . . . . . . . . . . . . . . Heat-Resistant Steels . . . . . . . . . . . . . . . . . . . . . . Common Alloys . . . . . . . . . . . . . . . . . . . . . . . . . Carbon and Low-Alloy Mechanical Properties . . . . Mechanical Properties of Wear-Resisting Steels . . . Mechanical Properties of Corrosion-Resisting Steels Mechanical Properties of Heat-Resisting Steels . . . . Ferrite in Cast Stainless Steels . . . . . . . . . . . . . . . Carbon and Low-Alloy Heat Treatments . . . . . . . . High-Alloy Heat Treatment. . . . . . . . . . . . . . . . . . Corrosion-Resistant Applications . . . . . . . . . . . . . . Heat-Resistant Applications . . . . . . . . . . . . . . . . . Selection and Evaluation of Steel Castings . . . . . . . . . . Design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Manufacturing . . . . . . . . . . . . . . . . . . . . . . . . . . . Nondestructive Examination . . . . . . . . . . . . . . . . . Purchasing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Research and Development . . . . . . . . . . . . . . . . . .
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921 922 923 924 924 924 926 926 928 930 933 934 936 938 943 946 947 948 949 949 950 950 951 952 953 953 960 961 966 968 970 971 972 973 975 975 980 982 986 986
Casting of Nonferrous Alloys . . . . . . . . . . . . . . . . . . . . . Nonferrous Casting — An Introduction . . . . . . . . . . . . . . . Nickel Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Titanium Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lead. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Dross, Melt Loss, and Fluxing of Light Alloy Melts. . . . . . Dross Formation . . . . . . . . . . . . . . . . . . . . . . . . . . . Influence of Melter Type on Dross Generation . . . . . . Influence of Charge Materials on Melt Loss . . . . . . . . Influence of Operating Practices on Melt Loss . . . . . . Economic Implications of Dross . . . . . . . . . . . . . . . . In-Plant Enhancement or Recovery of Dross Metallics . Ways to Reduce Dross Formation . . . . . . . . . . . . . . . Aluminum Alloy Ingot Casting and Continuous Processes. . Ingot Forms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Molten-Metal Processing . . . . . . . . . . . . . . . . . . . . . Ingot Casting Processes . . . . . . . . . . . . . . . . . . . . . . Continuous Processes . . . . . . . . . . . . . . . . . . . . . . . . Solidification in the DC Process . . . . . . . . . . . . . . . . Postsolidification Processes . . . . . . . . . . . . . . . . . . . . Safety. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aluminum Shape Casting . . . . . . . . . . . . . . . . . . . . . . . . Melting and Melt Handling . . . . . . . . . . . . . . . . . . . . Casting Process Selection . . . . . . . . . . . . . . . . . . . . . Pressure Die Casting of Aluminum . . . . . . . . . . . . . . Premium Engineered Castings . . . . . . . . . . . . . . . . . .
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987 989 989 989 990 992 992 993 994 994 998 998 999 1001 1001 1001 1002 1004 1006 1007 1007 1009 1009 1010 1013 1016
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© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
Copper Continuous Casting . . . . . . . . . . . . . . History . . . . . . . . . . . . . . . . . . . . . . . . . Vertical Continuous Casting . . . . . . . . . . Horizontal Continuous Casting . . . . . . . . Strip Casting . . . . . . . . . . . . . . . . . . . . . Wheel Casting . . . . . . . . . . . . . . . . . . . . Ohno Continuous Casting Process . . . . . . Casting of Copper and Copper Alloys . . . . . . . Castability. . . . . . . . . . . . . . . . . . . . . . . Melting and Melt Control . . . . . . . . . . . . Fluxing of Copper Alloys . . . . . . . . . . . . Degassing of Copper Alloys . . . . . . . . . . Deoxidation of Copper Alloys . . . . . . . . . Grain Refining of Copper Alloys . . . . . . . Filtration of Copper Alloys . . . . . . . . . . . Melt Treatments for Group I to III Alloys Production of Copper Alloy Castings . . . . Casting Process Selection . . . . . . . . . . . . Gating . . . . . . . . . . . . . . . . . . . . . . . . . Feeding . . . . . . . . . . . . . . . . . . . . . . . . Casting of Zinc Alloys . . . . . . . . . . . . . . . . . Control of Alloy Composition . . . . . . . . . Melt Processing . . . . . . . . . . . . . . . . . . . Die Casting of Zinc Alloys . . . . . . . . . . . Other Casting Processes for Zinc Alloys. .
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Nonferrous Alloy Castings . . . . . . . . . . . . . . . . . . . . . . . Mahi Sahoo, CANMET Materials Technology Laboratory Aluminum and Aluminum Alloy Castings . . . . . . . . . . . . . Chemical Compositions . . . . . . . . . . . . . . . . . . . . . . Effects of Alloying and Impurity Elements . . . . . . . . . Structure Control . . . . . . . . . . . . . . . . . . . . . . . . . . . Modification and Refinement of Aluminum-Silicon Alloys . . . . . . . . . . . . . . . . . Foundry Alloys for Specific Casting Applications . . . . Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Properties of Aluminum Casting Alloys . . . . . . . . . . . Copper and Copper Alloy Castings . . . . . . . . . . . . . . . . . . Alloying Additions. . . . . . . . . . . . . . . . . . . . . . . . . . Alloying Systems and Specifications . . . . . . . . . . . . . Physical Metallurgy . . . . . . . . . . . . . . . . . . . . . . . . . Types of Copper Castings. . . . . . . . . . . . . . . . . . . . . Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cupronickels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Zinc and Zinc Alloy Castings . . . . . . . . . . . . . . . . . . . . . Alloying Additions. . . . . . . . . . . . . . . . . . . . . . . . . . Physical Metallurgy . . . . . . . . . . . . . . . . . . . . . . . . . Die Casting and Aging . . . . . . . . . . . . . . . . . . . . . . . Specialty Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . Finishing of Die Castings . . . . . . . . . . . . . . . . . . . . . Metalworking and Machining . . . . . . . . . . . . . . . . . . Joining . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications for Zinc Die Castings . . . . . . . . . . . . . . Magnesium and Magnesium Alloys . . . . . . . . . . . . . . . . . Magnesium Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . General Applications . . . . . . . . . . . . . . . . . . . . . . . . Melting Furnaces and Auxiliary Pouring Equipment . . Melting Procedures and Process Parameters . . . . . . . . Sand Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Investment Casting. . . . . . . . . . . . . . . . . . . . . . . . . . Permanent Mold Casting . . . . . . . . . . . . . . . . . . . . . Die Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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1019 1019 1020 1022 1023 1024 1024 1026 1026 1027 1029 1032 1035 1037 1037 1038 1039 1040 1041 1044 1049 1049 1049 1051 1054
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1059 1059 1059 1063
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1066 1069 1071 1077 1077 1085 1085 1087 1088 1091 1091 1091 1091 1091 1095 1095 1096 1097 1097 1098 1098 1099 1097 1099 1100 1102 1102 1103 1106 1109 1109 1109
Direct Chill Casting . . . . . . . . . . . . . . . . . . . . . High-Temperature Alloys . . . . . . . . . . . . . . . . . Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cobalt and Cobalt Alloy Castings. . . . . . . . . . . . . . . Physical Metallurgy . . . . . . . . . . . . . . . . . . . . . Foundry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Postcasting Treatment . . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . . . . Nickel and Nickel Alloy Castings. . . . . . . . . . . . . . . Compositions . . . . . . . . . . . . . . . . . . . . . . . . . Structure and Property Correlations . . . . . . . . . . Melting Practice and Metal Treatment . . . . . . . . Foundry Practice . . . . . . . . . . . . . . . . . . . . . . . Pouring Practice . . . . . . . . . . . . . . . . . . . . . . . Gating Systems . . . . . . . . . . . . . . . . . . . . . . . . Risers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Heat Treatment . . . . . . . . . . . . . . . . . . . . . . . . Specific Applications . . . . . . . . . . . . . . . . . . . . Titanium and Titanium Alloy Castings . . . . . . . . . . . Historical Perspective of Casting Technology . . . Physical Metallurgy . . . . . . . . . . . . . . . . . . . . . Alloy Ti-6Al-4V . . . . . . . . . . . . . . . . . . . . . . . Casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Postcasting Practice . . . . . . . . . . . . . . . . . . . . . Product Applications . . . . . . . . . . . . . . . . . . . . Zirconium and Zirconium Alloy Castings . . . . . . . . . Zirconium Reactivity Considerations . . . . . . . . . Melting Processes . . . . . . . . . . . . . . . . . . . . . . Casting Processes. . . . . . . . . . . . . . . . . . . . . . . Processing. . . . . . . . . . . . . . . . . . . . . . . . . . . . Physical Properties of Zirconium Castings . . . . . Impact Properties of Zirconium Castings . . . . . . Synthesis and Processing of Cast Metal-Matrix Composites and Their Applications . . . . . . . . . . Classification of MMCs . . . . . . . . . . . . . . . . . . Solidification Processing of MMCs and Possible Effects of Reinforcement on Solidification. . Property Motivation for Using MMCs . . . . . . . . Components of MMCs Currently in Use. . . . . . . Development of Select Cast MMCs . . . . . . . . . . Research Imperatives . . . . . . . . . . . . . . . . . . . .
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1112 1112 1112 1114 1114 1116 1117 1117 1119 1119 1121 1123 1124 1124 1124 1125 1125 1125 1126 1128 1129 1129 1132 1134 1137 1138 1143 1143 1144 1144 1145 1147 1147
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Quality Assurance, Nondestructive Evaluation, and Failure Analysis. . . . . . . . . . . . . . . . . . . . . . . Babu DasGupta, National Science Foundation Approaches to Measurement of Metal Quality. . . . . . Metal Cleanliness Assessment Techniques . . . . . . . . Nondestructive Testing of Components . . . . . . . . . . . . . . Liquid Penetrant Inspection. . . . . . . . . . . . . . . . . . . Radiographic Inspection . . . . . . . . . . . . . . . . . . . . . Fluoroscopic Inspection and Automated Defect Recognition . . . . . . . . . . . . . . . . . . . . . Ultrasonic Inspection . . . . . . . . . . . . . . . . . . . . . . . Eddy Current Inspection . . . . . . . . . . . . . . . . . . . . . Process-Controlled Resonant Testing (Contributed by Quasar International, Inc.) . . . . . . . . . . . . . . . . Leak Test . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Electrical Conductivity Measurement . . . . . . . . . . . . Casting Fracture Characteristics . . . . . . . . . . . . . . . . . . . General Fracture Examination . . . . . . . . . . . . . . . . . Fracture Features . . . . . . . . . . . . . . . . . . . . . . . . . . Casting Failure Analysis Techniques and Case Studies . . . Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Background Information . . . . . . . . . . . . . . . . . . . . .
xvii
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1150 1155 1156 1160 1162
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1167 1167 1174 1174 1174
. . . . . 1174 . . . . . 1174 . . . . . 1175 . . . . . . . . .
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1175 1177 1177 1178 1178 1179 1183 1183 1183
© 2008 ASM International. All Rights Reserved. ASM Handbook Volume 15: Casting (#05115G)
Planning and Sample Selection . . . . . . . . . . . . . . . . Preliminary Examination . . . . . . . . . . . . . . . . . . . . Material Evaluation . . . . . . . . . . . . . . . . . . . . . . . . Data Analysis and Report Preparation . . . . . . . . . . . Case Studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Common Defects in Various Casting Processes . . . . . . . . Case Studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Repair Welding of Castings . . . . . . . . . . . . . . . . . . . . . . Repair Welding of Ferrous Castings. . . . . . . . . . . . . Repair Welding of Nonferrous Castings . . . . . . . . . . Quality Planning Tools and Procedures . . . . . . . . . . . . . . Advanced Product Quality Planning (APQP). . . . . . . Maintenance, Repair, Alterations, and Storage of Patterns and Tooling. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
www.asminternational.org
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1184 1184 1184 1187 1187 1192 1199 1203 1203 1205 1207 1207
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Ownership . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Maintenance Responsibility. . . . . . . . . . . . . . . . . . . . . Common Failure Mechanisms . . . . . . . . . . . . . . . . . . . Maintenance Interval . . . . . . . . . . . . . . . . . . . . . . . . . Maintenance Records . . . . . . . . . . . . . . . . . . . . . . . . . Preventive Maintenance Program. . . . . . . . . . . . . . . . . Repair . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Alterations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Storage—Physical Protection, Cataloging, and Tracking. Reference Information. . . Metric Conversion Guide. . Abbrieviations & Symbols . Index. . . . . . . . . . . . . . . .
xviii
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1213 1213 1213 1213 1214 1214 1214 1214 1215
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1217 1219 1222 1225
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 3-15 DOI: 10.1361/asmhba0005186
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
History and Trends of Metal Casting Thomas E. Prucha, American Foundry Society Daniel Twarog, North American Die Casting Association Raymond W. Monroe, Steel Founders’ Society of America
CASTING is one of the oldest manufacturing methods known to humankind and a very direct method of producing metal parts. It involves pouring molten metal into a cavity that is close to the final dimensions of the finished form, and it is used to manufacture many types of complex components of any metal, ranging in weight from less than an ounce to single parts weighing several hundred tons. Typical uses of cast components include infrastructure and structural components, water distribution systems (pipes, pumps, and valves), automotive Table 1
components (engine blocks, brakes, steering and suspension components, etc.), prosthetics, jewelry, and gas turbine engine hardware. In one study (Ref 1), castings were estimated to be used in 90% or more of all manufactured goods and in all capital goods machinery used in manufacturing. Some of the major market areas of castings are listed in Table 1 (Ref 2). Economically, casting processes are capable of producing highly reliable, cost-effective components ranging from low-volume, singlepart prototype production runs to economies
of scale for millions of parts. In terms of component design, casting also permits the formation of streamlined, intricate, integral parts with varied shape and size capabilities for designers. Metal casting is very flexible in terms of configuration design, and if a pattern can be made for a part, it can be cast. For example, the flexibility of metal casting, particularly sand molding, may permit the use of difficult design techniques, such as undercuts and curved or reflex contours that are not possible with other high-production processes. Tapered
Other major markets include machine tools, construction
Major markets for metal castings
Ferrous castings Gray iron Ingot molds Construction castings Motor vehicles Farm equipment Engines Refrigeration and heating Construction machinery Valves Soil pipe Pumps and compressors Pressure pipe Other major markets include machine tools, mechanical power transmission equipment, hardware, home appliances, mining machinery, and oil and natural gas pumping and processing equipment(a) Malleable iron Motor vehicles Valves and fittings Construction machinery Railroad equipment Engines Mining equipment Hardware Other major markets include heating and refrigeration, motors and generators, fasteners, ordnance, chains, machine tools, and general industrial machinery(a) Ductile iron Pressure pipe Motor vehicles
Farm machinery Engines Pumps and compressors Valves and fittings Metalworking machinery Construction machinery Other major markets include textile machinery, woodworking and paper machinery, mechanical power and transmission equipment, motors and generators, refrigeration and heating equipment, and air conditioning(a)
Steel Railroad equipment Construction equipment Mining machinery Valves and fittings General and special industrial machinery Motor vehicles Metalworking machinery Other major markets include steel manufacturing, spring goods, heating and air conditioning, recreation equipment, industrial materials handling equipment, ships and boats, and aircraft and aerospace(a) Nonferrous castings Aluminum Auto and light truck Aircraft and aerospace Other transportation Engines Household appliances Office machinery Power tools Refrigeration, heating, and air conditioning
equipment, mining equipment, farm machinery, electronic and communication equipment, power systems, and motors and generators(a) Magnesium Power tools Sporting goods Anodes Automotive Other major markets include office machinery, health care, and aircraft and aerospace(a)
Copper base Valves and fittings Plumbing brass goods Electrical equipment Pumps and compressors Power transmission equipment General machinery Transportation equipment Other major markets include chemical processing, utilities, desalination, and petroleum refining(a)
Zinc Automotive Building hardware Electrical components Machinery Household appliances Other major markets include scientific instruments, medical equipment, radio and television equipment, audio components, toys, and sporting goods(a)
(a)Ranked in approximate order of tonnage shipped, but in some cases, the total of other major markets is larger as a whole than the individual markets listed. Source: Ref 2
4 / Applications and Specification Guidelines sections with thickened areas for bosses and generous fillets are routine. The inherent design freedom of metal casting also allows the designer to combine what would otherwise be several parts of a fabrication into a single piece. This is significant when exact alignment must be held, as in high-speed machinery, machine tool parts, or engine end plates and housings that carry shafts. Various surface finishes on a part are also possible, ranging from a rougher as-cast surface of sand-molded casting to the smoother surfaces obtained through shell molding, investment casting, or other casting methods. In general, casting can provide various functional advantages for component design that include: Design both internal and external contours
independently to almost any requirement
Place metal in exact locations where it is
needed for rigidity, wear, corrosion, or maximum endurance under dynamic stress Produce a complex part as a single, dependable unit Readily achieve an attractive appearance This article briefly introduces the history of metal casting, general capabilities and methods of casting, and industry trends. In terms of value and volume, metal casting ranks second
only to steel rolling in the metal-producing industry. Metal casting is one of the ten largest industries when rated on a value-added basis, and worldwide casting shipments have continued to increase in recent years. Production statistics for 2005 are summarized in Table 2 (Ref 3), with comparison to 2004 production statistics from Modern Casting’s “Annual Census of World Casting Production.”
History of Metal Casting The casting of metal is a prehistoric technology. Exactly when the casting of metals began is not known, but it appears relatively late in the archaeological record. Archaeologists give the name Chalcolithic to the period in which metals were first being mastered, and they date this period, which immediately preceded the Bronze Age, very approximately to between 5000 and 3000 B.C. (Table 3, Ref 4). There were many earlier fire-using technologies, collectively called pyrotechnology, that provided a basis for the development of metal casting. Among these were the heat treatment of stone to make it more workable, the burning of lime to make plaster, and the firing of clay to produce ceramics. At first, it did not include
smelting, because the metal of the earliest castings appears to have been native. Cast iron appeared in China in approximately 600 B.C. Its use was not limited to strictly practical applications, and there are many examples of Chinese cast iron statuary. Most Chinese cast irons were unusually high in phosphorus and, because coal was often used in smelting, high in sulfur as well. These irons therefore have melting points that are similar to those of bronze, and when molten are unusually fluid. The iron castings, like the Chinese cast bronzes, are often remarked upon for the thinness of their wall sections. From these early beginnings, development and application of metal casting continued (Table 4, Ref 5). During the wars of medieval times, metalcasters produced cannons. In peace, they recast the metal into bells. In the Middle Ages, church bells were cast by priests, abbots, or bishops who also were trained metal founders. As the metal was melted, the brethren stood around the furnace intoning psalms and prayers. The molten metal was then blessed, and divine protection was asked for the bell, which usually bore the name of a saint (Ref 5). During the 1800s, there were significant advances in technology. Metallurgy also began
Table 2 World casting production for 2004–2005 40th census of world casting production for 2005, metric tons Country
Austria Belgium Brazil Canada China Croatia Czech Republic Denmark Finland France Germany Great Britain Hungary India Italy Japan Korea Lithuania Mexico Netherlands(a) Norway Poland Portugal Romania Russia Slovenia(a) Slovakia South Africa Spain Sweden Switzerland Taiwan Thailand Turkey Ukraine (i) United States Totals
Gray iron
47,500 70,700(a) 2,460,000(b)(d) 483,000(a)(b)(d) 12,303,963 25,306 269,493 46,932 47,817 896,500 2,469,058 531,000 35,307 4,116,000 923,700 2,782,509 960,100 14,300 600,000 63,100 16,348 420,700 27,285 81,578 3,480,000 75,900 41,520 160,000 491,600 171,700(b) 26,400 827,932 170,000 567,000 626,610 4,457,905 40,788,763
Ductile iron
131,000 15,600(a)(b) ... ... 5,838,753 16,252 49,524 40,333 60,123 957,700 1,487,234 362,000 32,200 618,000 512,200 1,919,435 565,200 200 270,000 78,241 49,207 112,500 69,792 10,500 720,000 27,800(b) 8,760 65,000 551,600 83,300 40,400 220,413 70,000 327,000 40,000 4,241,088 19,591,355
Malleable iron
17,500 ... ... ... 514,129 50 5,435 0 0 1,400 52,713 12,000 10 56,000 8,400 57,851 46,500 ... 2,000 6,209 0 21,740 0 926 280,000 ... ... 10,000 17,400 ... ... ... 30,000 8,000 10,000 75,296 1,233,559
Steel
18,000 42,100 293,891 117,600(a) 3,224,374 2,219 109,680 0 18,532 113,700 200,409 112,000 5,385 805,000 73,324 276,589 149,600 30 75,000 438 3,723 59,560 12,096 37,900 1,200,000 22,700 4,300 136,000 83,200 23,400 ... 73,619 30,000 125,000 266,060 1,287,295 9,002,724
Copper base
5,459 ... ... 18,585(a) 416,097 1,120 1,724 1,126 3,999 26,264 84,386 15,000 2,110 0 16,800 97,794 23,200 5 180,000 ... 4,667 6,300 8,700 2,700 160,000 6,727(f) 2,160 14,900 7,886 11,200 2,689 41,959 28,600 16,000 11,000 292,113 1,511,270
Aluminum
103,347 19,428(c) 240,457(e) 324,973 1,886,000(c) 15,038 80,252(c) 4,000 6,565 318,445(c) 727,139 206,000(c) 75,720 516,000 857,500 1,477,349(c) 144,300 56 660,000 ... 22,926 174,300 20,250 20,000 920,000 30,183 26,260 56,000 139,312(c) 44,400 17,785 250,366 100,000 66,500 20,500 2,080,174 11,718,025
Magnesium
6,580 ... ... ... ... 0 ... 0 0 ... 27,283 ... 3,050 0 12,000 ... 10,300(f) ... 500 ... 0 30 150 10 70,000 ... ... 0 ... 1,700 ... 6,427 ... 500 ... 100,697 239,227
Zinc
12,997 ... ... ... 237,889 502 2,656 1 434 25,066 56,439 21,000 3,050 0 70,000 36,216 ... ... 350 ... 0 8,600 900 700 20,000 ... 1,800 2,600 16,051 5,900 1,616 69,316 16,900 12,700 ... 312,978 936,661
Other nonferrous
... ... ... ... ... 1,300 279 5.022 0 3,461 3,346 2,000 0 0 67,600 8,102 0 ... ... ... ... 770 120 5 50,000 ... 5 0 433 ... ... 4,418 ... ... ... 48,987 195,848
Total
342,383 124,507(a) 2,968,600 863,373(a) 24,421,205 61,787 519,043 97,414 137,470 2,342,536 5,108,007 1,261,000 156,832 6,111,000 2,541,524 6,655,845 1,899,200 14,591 1,787,850 147,988(g) 96,871 804,500 139,293 154,319 7,620,000(h) 163,310 84,805 444,500 1,307,482 341,600 88,890 1,494,450 445,500 1,122,700 974,170 12,896,533 85,741,078
Totals for 2004
325,205 124,507 2,829,916 863,373 22,420,452 56,449 522,400 95,949 133,325 2,465,617 4,984,473 1,273,000 164,892 4,623,000 2,441,282 6,386,449 1,857,300 18,715 2,185,200 147,988 86,757 804,500 122,750 207,023 6,300,000 163,310 4,800 444,500 1,309,249 321,100 104,077 1,451,768 235,850 982,000 974,170 12,314,121 79,745,467
Note: “. . .” indicates unreported data. (a) 2004 tonnage. (b) Includes malleable iron. (c) Includes magnesium. (d) Includes ductile iron. (e) Includes all nonferrous. (f) Includes zinc. (g) Includes only ferrous. (h) Includes 720,000 tons of special iron. (i) 2002 tonnage. Source: Modern Casting
History and Trends of Metal Casting / 5 Table 3
Chronological list of developments in the use of materials
Date
Development
9000 B.C. 6500 B.C. 5000–3000 B.C.
3000–1500 B.C.
3000–2500 B.C. 2500 B.C. 2400–2200 B.C. 2000 B.C. 1500 B.C.
Earliest metal objects of wrought native copper Earliest life-size statues, of plaster Chalcolithic period: melting of copper; experimentation with smelting Bronze Age: arsenical copper and tin bronze alloys Lost-wax casting of small objects Granulation of gold and silver and their alloys Copper statue of Pharoah Pepi I Bronze Age Iron Age (wrought iron)
Location
Near East Jordan Near East
Near East
Near East Near East Egypt Far East Near East
Date
700–600 B.C. 600 B.C. 224 B.C.
Development
Etruscan dust granulation Cast iron Colossus of Rhodes destroyed 200–300 A.D. Use of mercury in gilding (amalgam gilding) 1200–1450 A.D. Introduction of cast iron (exact date and place unknown) Circa 1122 A.D. Theophilus’s On Divers Arts, the first monograph on metalworking written by a craftsman 1252 A.D. Diabutsu (Great Buddha) cast at Kamakura Circa 1400 A.D. Great Bell of Beijing cast 16th century Sand introduced as mold material
Location
Italy China Greece
Date
Development
1709
Cast iron produced with coke as fuel, Coalbrookdale Boring mill or cannon developed Great Bell of the Kremlin cast Cast steel developed by Benjamin Huntsman Cast iron used as architectural material, Ironbridge Gorge Zinc statuary Electrodeposition of copper
1715 Roman world Europe
1735 1740
Germany
1779
Japan
1826 1838
China France
1884
Electrolytic refining of aluminum
Location
England
Switzerland Russia England England
France Russia, England United States, France
Source: Ref 4
Table 4
Timeline of casting developments to 1899
B.C. 9000 B.C.—Earliest metal objects of wrought native copper are produced in the Near East. 3200 B.C.—The oldest casting in existence, a copper frog, is cast in Mesopotamia. 3000 B.C.—Early metalcasters cast bronze tools and weapons in permanent stone molds. 3000–2500 B.C.—Small objects are cast via lost-wax (investment casting) process in the Near East. 1500 B.C.—Wrought iron is discovered in the Near East. 600 B.C.—The first cast iron object, a 270 kg (600 lb) tripod, is cast by the Chinese. 233 B.C—Iron plowshares are cast. 200 B.C.—Oldest iron castings still in existence are produced during the Han Dynasty. A.D. 500—Cast crucible steel is produced in India. 1200s—Loam or sweep molding is used by European metalcasters to cast bells for cathedrals. 1252—The colossal statue, the Great Buddha at Kamakura, Japan, is cast in high-lead tin bronze. The project began in the 700s, and its head alone weighed 140 tons. 1313—The first cannon is cast in bronze by a monk in the city of Ghent. 1400s—During the siege of Constantinople, heavy guns are cast from bronze “on the spot,” virtually under the walls of the besieged city. Movable, cast lead type for printing presses revolutionizes the world’s methods of communication. 1480—Vannoccio Biringuccio (1480–1539), the first true metalcaster and the “father of the metal casting industry,” is born. The founder of the Vatican, his De La Pirotechnia is the first written account of proper metal casting practice. 1500s—Sand is introduced as a molding material in France. 1600s 1612—Mined from under the sea, seacoal is mentioned for the first time by German metalcaster and inventor Simon Sturtevant. 1619—North America’s first iron furnace is built at Falling Creek, Va., on a branch of the James River, 100 km (60 miles) from the Jamestown colony. Three years later, Native Americans destroy it during a raid. 1645—Earliest recorded use of the term foundry appears in the Oxford English Dictionary in its variant “founderie.” 1646—America’s first iron metal casting facility (and second industrial plant), Saugus Iron Works, near Boston, pours the first American metal casting, the Saugus pot. The Saugus River site was selected by Richard Leader and was built to produce iron products for Massachusetts and England. 1661—First U.S. copper deposits are discovered by Gov. Winthrop in Middletown, Conn. Source: Ref 5; see also Selected References
1700s 1709—Two important developments by Abraham Darby from Coalbrookdale, England, improve casting methods. He developed the first true metal casting facility flask to modernize molding practices (which had been carried out in pits on the floor by use of pattern boards tied together or in crude box frames). He would later initiate the use of coke as a furnace fuel for iron production. 1722—A.F. de Reamur, recognized as the world’s first metallurgical chemist, develops malleable iron, known today as European whiteheart malleable. 1750—Benjamin Huntsman reinvents the cast crucible steel process in England, a process that disappeared after first being developed in India. The English parliament prohibits the refining of pig iron or the casting of iron in the American colonies, contributing to the start of the American Revolution. 1756—Coalbrookdale’s Richard Reynolds oversees the invention of cast iron tram-road rails, replacing wooden rails. 1775—Revolutionary patriot Paul Revere, who operated a bell-and-fittings foundry in Boston, rides from Boston to Lexington warning colonists of the British invasion. 1776—Metalcasters Charles Carroll, James Smith, George Taylor, James Wilson, George Ross, Philip Livingston, and Stephen Hopkins sign the American Declaration of Independence. 1779—First iron bridge ever erected (above England’s Severn River) was cast and constructed at Coalbrookdale Works. 1794—John Wilkinson of England invents the first metalclad cupola furnace, using a steam engine to provide the air blast. 1797—First cast plow in the United States is invented by Charles Newbold of Sauk, N.J. 1800s 1809—Centrifugal casting is developed by A.G. Eckhardt of Soho, England. 1815—The cupola furnace is introduced to the United States (Baltimore). 1817—The first iron water line in the United States, 120 m (400 ft) long, is laid in Philadelphia. 1818—The first U.S. cast steel is produced by the crucible process at historic Valley Forge Foundry. 1825—Aluminum, the most abundant metal in the Earth’s crust, is isolated from aluminum chloride by Denmark’s Hans Oerstad. 1830s—Seth Boyden of Newark, N.J., produces the first blackheart malleable iron in the United States. 1832—Nickel-bronze is produced commercially in England. 1837—The first reliable molding machine on the market is made and used by the S. Jarvis Adams Co., Pittsburgh. 1847—Cast steel guns are made by Krupp Works in Germany. Asa Whitney, Philadelphia, obtains a patent on
a process for annealing chilled-iron car wheels cast with chilled tread and flange. 1847—John Deere commissions Jones and Quiggs Steel Works, Pittsburgh, to cast and roll a steel plow, which it accomplishes at one-half the previous cost of the product. 1849—A manually operated die casting machine is patented to supply rapidly cast lead type for newspapers. 1850—Drop-bottom cupola is developed. 1863—Metallography is developed by Henry C. Sorby of Sheffield, England, enabling metalcasters to polish, etch, and microscopically examine metal surfaces for physical analysis. 1867—James Nasmythe, inventor of the steam hammer, develops a gear-tilted safety ladle to prevent pouring accidents. 1870—Sandblasting is developed for large castings by R.E. Tilghman of Philadelphia. 1874—The Colliau cupola, the first commercially made cupola in America, is introduced. 1876—The first authenticated aluminum castings were produced by Col. William Frishmuth at his Philadelphia foundry. Assembled to produce an engineer’s transit, these castings were made from $1/oz chemically produced aluminum. Manganese-bronze patent is granted to Parsons in England. Tobin Bronze begins developing manganesebronze in the United States. 1880–1887—W.W. Sly, Cleveland, develops the first casting cleaning mill, greatly reducing hand-chipping and grinding to allow a custom-finished product. 1884—The first architectural application of aluminum, a cast aluminum pyramid that was produced by Frishmuth, is mounted on the tip of the Washington Monument. 1886—Charles M. Hall, a 22 year old student at Oberlin College, discovers a process of aluminum reduction through electrolysis. The invention replaced chemical reduction and lowered the metal cost (from $15/lb in 1884 to $0.50/lb in 1890), spurring a new industry of aluminum applications. 1887—Eli Millett invents a core oven for drying small cores in individual drawers. 1890—The first motor-driven mold conveyor is installed, integrating molding, pouring, and cooling operations. 1897—Iowa dentist B.F. Philbrook adapts the lost-wax investment casting process for producing dental inlays, the first nonart application of the process in the modern metal casting age. 1898—Poulson and Hargraves (U.K.) produce the first sand molds bonded with sodium silicate. Germany’s Imperial Navy recommends copper-nickel alloys containing 4–45% Ni for saltwater piping system. 1899—Electric arc furnace, developed by France’s Paul Heroult, begins commercial production.
6 / Applications and Specification Guidelines to achieve prominence in 1889 when nickel was alloyed to make steel stronger. In terms of casting, centrifugal casting was developed by A.G. Eckhardt of Soho, England, in 1809. This method involves the pouring of molten metal into a rapidly rotating metallic mold. The method was soon adopted by the pipe foundries and was first used in Baltimore, Maryland, in 1848. Sir Henry Bessemer, famed for his converter, used centrifugal casting to remove gases and was the first to pour two or more metals into a single rotating mold. The centrifugal casting of steel was first attempted in 1898 at the plant of the American Steel Foundries in St. Louis, Missouri. Railroad car wheels were spun cast in 1901 at a rotation speed of 620 rpm. Following the early development of the centrifugal method, a permanent mold method known as slush casting was introduced. Slush casting is a process in which molten metal is poured into a split metal mold (generally made of bronze) until the mold is filled; then, the mold is immediately inverted, and the metal that is still liquid is allowed to run out. The time required for this casting operation is sufficient to freeze a metal shell in the mold, corresponding to the shape of the cavity wall. The thickness of the wall of the casting depends on the time interval between the filling and the inverting of the mold, as well as on the chemical and physical properties of the alloy and the temperature and composition of the mold. Usually, lead and zinc alloy castings are
Table 5
produced by slush casting. The process is limited to the production of hollow castings and was used to produce lamp bases (Ref 6). Manually operated die casting machines were patented as early as 1849 (Sturgiss) and 1852 (Barr) in an effort to satisfy the insatiable demands of a growing reading public by way of rapidly cast lead type. These early inventions led to Ottmar Mergenthaler’s Linotype, an automatic casting machine in which molten lead is forced by piston stroke into a metal mold. The first die casting machine bearing the Linotype name was patented in 1905 by H.H. Doehler. Two years later came E.B. Wagner’s casting machine, a prototype of the now familiar hot chamber die casting machine. It was first used on a large scale during World War I for binocular and gas mask parts. Zinc alloys were used for die casting as early as 1907 but were not competitive until Price & Anderson developed the Zamak die casting alloy in 1929. In 1851, James Bogardus’s factory in Chicago, which was constructed with cast iron supports, opened the way for what many art historians considered to be America’s only original contribution to the arts of the world— the skyscraper. One of the oldest casting techniques, investment (lost-wax) casting, was also rediscovered in 1897 by B.F. Philbrook of Iowa, who used it to cast dental inlays. Industry paid little attention to this sophisticated process until the urgent military demands of World War I overtaxed the machine tool industry. Shortcuts were then needed to provide finished tools and
precision parts, avoiding time-consuming machining, welding, and assembly. Other efforts at this time included the first fully automated plant in the United States (one of the first in the world), which was a Rockford, Illinois, foundry that cast hand grenades for the U.S. Army in 1918. The first separation of aluminum occurred in 1825, but important engineering applications did not occur until the cost of aluminum steadily declined as a result of smelting process improvements and scale of operations through the end of the 19th century. Aside from additions for aluminum-bronze, zinc-aluminum alloys, and the deoxidation of steel, castings were the first important commercial market for aluminum. The process was simple, low cost, and addressed the limited range of product interests. At first, applications were limited to curiosities such as house numbers, hand mirrors, combs, brushes, tie clasps, cuff links, hat pins, and decorative lamp housings that emphasized the low density, silvery finish, and novelty of the new metal. Lightweight, corrosion-resistant and heat-conductive aluminum cookware could be produced using the same patterns used to cast iron and brass pots, pans, and kettles (Ref 7).
The 20th Century Casting technology during the 20th century advanced in many ways (Table 5), with “more casting progress since World War II than in the previous 3000 years” (Ref 8). For more than
Timeline of casting developments in the 20th century
1900s 1900—Brinnell hardness test machines introduced. Aluminum-bronze in regular production in the United States. Early 1900s—First patent for low-pressure permanent mold casting process issued to England’s E.H. Lake. 1901—American Steel Foundries (St. Louis) produces the first centrifugal cast rail wheels. 1903—The Wright Brothers’ first successful machinepowered aircraft contains a cast aluminum block and crankcase (together weighing 69 kg, or 152 lb), produced either at Miami Brass Foundry or the Buckeye Iron and Brass Works. 1905—H.H. Doehler patents the die casting machine. 1906—The first electric arc furnace is installed in the United States at Halcomb Steel Co., Syracuse, N.Y. First lowfrequency induction furnace is installed at Henry Diston & Sons, Tacony, Pa. 1907—Alfred Wilm discovers that the properties of cast aluminum alloys can be enhanced through heat treating and artificial aging. 1908—Stockham Homogenous Sand Mixer Co., Piqua and Newark, Ohio, releases the sand cutter. 1910s 1910—Matchplates are developed, fostering the viability of jolt-squeeze machines. 1911—Metallurgical microscope is obtainable. First electric arc furnace for metal casting is installed at Treadwell Engineering Co., Easton, Pa. 1912—The first muller with individually mounted revolving mullers of varying weights is marketed by Peter L. Simpson. Sand slinger is invented by E.O. Beardsley & W.F. Piper, Oregon Works.
1915—Experimentation begins with bentonite, a colloidal clay of unusually high green and dry strength. Ajax Metal Co., Philadelphia, installs first low-frequency induction furnace for nonferrous melting. 1916—Dr. Edwin Northrup, Princeton University, invents the coreless induction furnace. 1917—Alcoa completes a great deal of early development work in aluminum as World War I generates a great demand for high-integrity castings for aircraft engines. 1918—The first fully automated metal casting facility in Rockford, Ill., casts hand grenade hulks for the U.S. Army. 1920s 1921—Modification of the silicon structure in aluminum begins as Pacz discovers that adding metallic sodium to molten aluminum just prior to pouring greatly improves ductility. Copper-silicon alloys are prepared in Germany as a substitute for tin-bronzes. 1924—Henry Ford sets a production record of 1 million autos in 132 working days. Automotive manufacturing will grow to consume one-third of casting demand in the United States. 1925—X-ray radiography is established as a tool for checking casting quality. By 1940, all military aircraft castings require x-ray inspection prior to acceptance. American Brass, Waterbury, Conn., installs first mediumfrequency induction furnace in the United States. 1928—Alcoa develops the first aluminum vehicle wheel, a sand-cast 355 alloy designed for truck trailers. 1930s 1930—First high-frequency coreless electric induction furnace is installed at Lebanon Steel Foundry, Lebanon, Pa. Spectrography is pioneered by University of Michigan (continued)
Source: Ref 5; see also Selected References
professors for metal analysis. Davenport and Bain develop the austempering process for iron castings. 1937—Applied Research Laboratories’ founder Maurice Hasler produces the first grating spectrograph for the Geological Survey of California. Spectrometers begin finding their way into foundries by the late 1940s, replacing the previous practice of metallurgists estimating chemical compositions with a spectroscope and welder’s arc. The austempered microstructure in cast iron is recognized. 1940s 1940—Chvorinov develops the relationship between solidification time and casting geometry. Early 1940s—Statistical process control is first employed as a quality control tool in U.S. machine shops, principally to control dimensional tolerances. Inoculation of gray iron becomes common as high-quality cast irons replace scarce steel. 1941—U.S. Lt. Col. W.C. Bliss tells the American Foundrymen’s Association St. Louis Chapter that “the side which maintains the larger production of war goods is going to win the war.” The War Production Board reports later that each U.S. soldier requires 2200 kg (4900 lb) of steel compared to 40 kg (90 lb) in World War I. 1942—The use of synthetic sands increases as a replacement for many war materials. 1943—Keith Millis, a 28 year old metallurgist working at the International Nickel Company searching for a replacement to chrome due to interrupted supply, discovers that magnesium alloy in molten iron produces a spheroidal graphite structure. In 1949, he, Albert Gagnebin, and Norman Pilling would receive a U.S. patent on ductile iron production via magnesium treatment.
History and Trends of Metal Casting / 7 Table 5 (continued) 1940s 1944—The first heat-reactive, chemically cured binder is developed by Germany’s Johannes Croning to rapidly produce mortar and artillery shells for Axis troops during World War II. Two years after the war, his shell process is discovered among other inventions in the German patent office and made public. Croning is recognized with an American Foundrymen’s Society (AFS) Gold Medal in 1957 for his invention. 1946—Allied investigators uncover German foundry research on high-temperature alloys. Having “heard” of the Croning discovery (prior to the release of the report), Ford’s Ed Ensign and E.I. Valyi, Navy Bureau of Ships, attempted to replicate the process and produced shellmolded castings at the Midwest Foundry, Coldwater, Mich. Late 1940s—Thermal sand reclamation is applied to core sands and, to a limited degree, clay-bonded sands. 1948—The first nonlaboratory ductile iron casting is produced at Jamestown Malleable Iron Co., Jamestown, N.Y., as a 168 cm (66 in.) test bar is poured. Industry’s first ductile iron pipe is cast at Lynchburg Foundry, Lynchburg, Va. 1949—Keel blocks, diesel engine parts, a pressure cylinder, an 20 cm (8 in.) cube, and two cylinder liners become the first commercial castings of ductile iron at CooperBessemer, Grove City, Pa. Development of magnesiumferrosilicon makes ductile iron treatment far easier. Buffalo Pipe & Foundry Co., Tonawanda, N.Y., is the first U.S. operation to pour castings using Croning’s shell process. 1950s Early 1950s—Experimentation in high-pressure molding begins as metalcasters begin to increase the air pressure in air squeeze-molding machines to increase mold hardness (density). Fast-drying core oils are introduced. The pneumatic scrubber is developed to reclaim clay-bonded sands. Several wet reclamation systems also are in operation. 1951—Ford Motor Co., Dearborn, Mich., converts 100% of its crankshaft production to ductile iron. 1952—D-process is developed for making shell molds with fine sand and fast-dry oil by Harry Dietert. Sodiumsilicate/CO2 system is introduced. 1953—Hotbox system of making and curing cores in one operation is developed, eliminating the need for dielectric drying ovens. 1954—The CO2 process, a novel mold and coremaking process, is introduced from Germany. Working closely with General Motors, B&P develops a method for coating individual sand particles with resin binder. It also introduces a coreblower capable of producing cores with resin-coated shell sand—a modification to the Croning process. 1955—Ductile iron pipe is introduced to the marketplace by People’s Gas of Chicago, the first to install ductile iron gas mains. Mid-1950s—The squeeze-casting process originates in Russia. 1956—The first Betatron is installed in a U.S. foundry at ESCO Corp., Portland, Ore., for radiography of heavy steel castings. 1957—The vertically parted flaskless green sand molding production machine is invented by Vagn Aage Jeppesen, a 40 year old professor at the Technical University of Denmark. He was granted a patent in 1959, which was purchased by Dansk Industri Syndikat in 1961. 1958—Harold F. Shroyer obtains a patent for the full mold process, a process developed by artists in which simple patterns and gating systems are carved from expanded polystyrene and placed into a green sand mold. The process, known today as lost-foam casting (using loose, unbonded sand), is patented a short time later. Phenolic and furan acid-catalyzed no-bake binder systems are introduced. Ductile iron desulfurization via shaking ladles is developed in Sweden. 1959—General Electric uses the transient heat-transfer digital computer program and successfully applies the finite difference method to heavy steel casting production.
1960s 1960—Furan hotbox binders are developed for core production. Deep bed filters are used commercially for aluminum casting at Alcoa and British Aluminum in the United Kingdom. Compactibility and methylene blue clay tests are developed for green sand control. 1961—Alcohol-borne shell coating process is introduced (warm-coated). 1962—New CO2 sand testing method is introduced for sands bonded with sodium silicate and cured with CO2. Beardsley & Piper’s Al Hunter, Bob Lund, and Angello Bisinello develop the first automated green sand molding machine. In their design, the cope and drag are side-blown simultaneously and then hydraulically squeezed. The birth of automated matchplate molding reportedly improved metal casting productivity by levels as much as 60% in a short amount of time. Phenolic hotbox binders are introduced. 1963—Shell flake resin is introduced, eliminating the need for solvents. 1964—Dell & Christ’s paper on mold inoculation spurs the development of many of today’s forms of mold and late stream inoculation. The first vertically parted green sand machine (max. 240 molds/h) is delivered to United Danish Iron Foundries in Frederiksvaerk, Denmark. Early adopters report man-hour per ton improvements on the order of 50%. 1965—Oil urethane no-bake binder systems are used for cores and molds. General Electric’s Jim Henzel and Jack Keverian predict freezing patterns in large steel castings via computer. Cast metal-matrix composites are first poured at International Nickel Co., Sterling Forest, N.Y. 1968—The coldbox process is introduced by Larry Toriello and Janis Robins and introduced to the metal casting industry by Ashland Chemical Co. for high-production coremaking. Germany’s Daimler-Benz foundry in Mannheim is the first to run the process for automotive parts. John Deere Silvas Foundry, Moline, Ill., is the first to use the process for mass production in North America. 1969—The Chevrolet Vega is introduced by General Motors, featuring the first all-aluminum block with no cast iron cylinder liners. A total of 2.5 million blocks were produced during the vehicle life cycle. Late-1960s—Scanning electron microscope is invented in England. Thermal analysis begins to be used in iron foundries for the rapid determination of carbon equivalent and phosphorus contents, making it possible to study the transformation of an alloy during cooling. Manganese Bronze & Brass Co. and J. Stone & Co. join to promote nickel-aluminum bronze propellers. 1970s 1970—The sodium-silicate/ester catalyzed no-bake binder system is introduced for cores and molds. Safety-critical ductile iron steering knuckles are introduced on Chevrolet’s Cadillac. A new phenolic urethane no-bake process is introduced by Ashland Chemical, replacing oil sand dump box operations and significantly reducing energy requirements for core/mold production throughout the 1970s. Diran Apelian’s doctoral work at the Massachusetts Institute of Technology (MIT) leads to the development of foam filters for metal casting by Olin Metals. Commercial ceramic foam filters will be in use in metal casting facilities by 1974. 1971—The vacuum-forming or V-process molding method of using unbonded sand with a vacuum is developed in Japan. MIT doctoral candidate David Spencer performs experiments leading to semisolid molding, a process in which a partially frozen metal (fluidity similar to machine oil) can be poured into a die cavity. Following advancement in metal slurry consistencies with Professor Mert Flemings, Newton Diecasting Co., New Haven, Conn., produces the first semisolid castings. 1972—A 0.5 kg (1 lb) crankshaft for a refrigerator compressor produced at Wagner Castings (designed and engineered by Tecumseh) becomes the first productionvolume austempered ductile iron (ADI) component. Wagner also produces the first as-cast ductile iron ductile rods for passenger cars. The Canada Center for Mineral and Energy Technology (CANMET) uses real-time (continued)
Source: Ref 5; see also Selected References
1970s radiography to study the flow of steel in molds. Hitchiner Manufacturing, Milford, N.H., patents countergravity (vacuum) casting process. 1973—The first U.S. foundry argon oxygen decarburization unit is installed at ESCO Corp. 1974—Fiat introduces the in-mold process for ductile iron treatment. The phenolic urethane no-bake binder system is introduced for mold production. Mid-1970s—Alcoa and Union Carbide commence rotary degassing for wrought aluminum. Reading Foundry Products would apply this technology to aluminum foundries in the mid-1980s. Digital codes are developed to simulate solidification and fluid flow analysis. Ultrasonic verification of ductile iron nodularity is developed. Metal casting facilities examine new beneficial reuse routes for spent foundry sand, leading to applications such as cement and paving products, bricks, and flowable fill. 1976—Foote Mineral Co. and the British Cast Iron Research Association (U.K.) develop compacted graphite iron. Acidslag cupola practices plus external desulfurization with CaC2 begin to replace basic slag cupolas. 1977—General Motors installs ADI rear differential sets in passenger cars. The alumina phosphate no-bake binder system, an inorganic nonsilicate binder, is introduced for mold production. 1978—The furan/SO2 binder system is developed for core and mold making. Polyurethane no-bake binder system is introduced for aluminum applications. 1980s Early 1980s—Tundish ladle is embraced by industry as favored practice of nodularizing ductile iron. 3-D relational parameters are developed for Computer-aided design solid models. 1981—High-production lost-foam casting begins at General Motors’ Massena, N.Y., plant for aluminum cylinder heads. 1982—Warmbox binder system is introduced. 1983—Air impulse molding process is developed. Free radical cure/SO2 binder system is introduced. 1984—Charles Hull applies for a patent on stereolithography process. Other rapid prototyping techniques emerge shortly after. Phenolic ester no-bake binder is introduced. Thermal analysis makes breakthrough in molten aluminum processing for determination of grain refinement and silicon modification. 1985—Phenolic ester coldbox binder is developed. New automaker Saturn makes a strategic decision to select lostfoam process for its aluminum cylinder blocks and heads and ductile iron crankshafts and differential cases. Mid-1980s—Computer solidification software is commercialized. Amine recycling is introduced to enhance the environmental benefits of the amine-cured coldbox process. Ube Machinery introduces first squeeze-casting equipment. Aikoh’s (Japan) flux injection technology is initiated into U.S. aluminum foundry market. Late 1980s—3-D visualization techniques are developed. CaO/CaF2 desulfurization of cupola-melted ductile base iron begins to replace CaC2 method. Lanxide, Newark, Del., develops pressureless metal infiltration process for particulate-reinforced metal bodies. Magnesium wire injection method for ductile iron treatment is first tested. Casting solidification modeling software gains acceptance, allowing metal casting facilities to optimize quality, production, and cost prior to actually pouring a casting. 1988—Rapid prototyping and Computer-aided design/ computer-aided manufacturing technologies combine in a breakthrough to shorten tooling development time. Ford adapts Cosworth process precision sand casting process for high production. Metaullics Systems combines flux injection/rotary degassing technologies for aluminum processing. 1989—IMI (Yorkshire, U.K.) begins experimenting with bismuth as a lead substitute in copper alloys. 1990s 1990—Equipment for semisolid casting is introduced by Alumax Engineered Materials and Buhler, Inc. Foseco patents a direct-pour system that permits casting production without conventional gating/risering. Major automotive application comes in 1995 with CMI International’s upper intake
8 / Applications and Specification Guidelines Table 5
(continued)
manifold. Precision sand casting and casting quality for engine blocks are improved in mass production by major automotive companies with the Cosworth and Zeus processes for aluminum and the Loramendi Key Core process for precision sand-cast iron applications. 1991—“Dry ice” CO2 process is developed for cleaning coreboxes and foundry tooling. A noncontact gage for accurate dimensional analysis of lost-foam patterns and sand cores is developed through the AFS Lost-Foam Consortium. Eight years later, the consortium develops an instrument to measure the gas permeability of lost-foam pattern coatings (which controls flow of metal and has a dominant effect on casting quality). 1993—First foundry application of a plasma ladle refiner (melting and refining in one vessel) occurs at Maynard Steel Casting Co., Milwaukee. 1994—Use of low-expansion synthetic mullite sand for lost foam is patented by Brunswick Corp., Lake Forest, Ill., to enable precision casting of large components. Delphi Chassis Systems and Casting Technology Co. begin a program on squeeze-cast aluminum front knuckle, the
first high-volume production (1.5 million cars) two-cavity squeeze-cast aluminum conversion of its kind. Mid-1990s—Microstructure simulation is developed, contributing to better understanding of metallurgy effects and the prediction and control of mechanical properties in castings. Semisolid casting makes commercial inroads and market penetration. 1996—Cast metal-matrix composites (brake rotors) are used for the first time in production model automobiles, the Lotus Elise. Environmentally friendly (fluorine-free) fluxes are developed at the Worcester Polytechnic Institute and are commercialized. General Motors Corp. introduces GMBond, a water-soluble biopolymer-based core sand binder that is nontoxic and recyclable. 1997—AFS Consortium research at CANMET, Ottawa, Canada, results in the development and commercialization of lead-free copper alloys using bismuth and selenium. Following tests at Germany’s Aachen University, American Cast Iron Pipe Co., Birmingham, Ala., builds a continuously operated electric arc furnace for cast iron production.
Late 1990s—Stress and distortion simulation introduces the benefits of controlling casting distortion, reducing residual stresses, eliminating hot tears and cracking, minimizing mold distortion, and increasing mold life. 2000s 2001—The National Aeronautics and Space Administration (NASA) and the Dept. of Energy/Office of Industrial Technologies released a physics-based software tool to accurately predict the filling of expanded polystyrene patterns and sand cores as numerous variables are changed. Mercury Marine installs North America’s first pressurized lost-foam casting line at its new facility in Fond du Lac, Wis. 2002—Math-based modeling is used to develop a new sand core blowing simulation capability for process modeling software. 2003—The AFS Magnesium Division completes a research project proving that magnesium could be cast via the lost-foam casting process. Similar trials confirmed the castability of magnesium via the V-process.
phenolics led the way to urea and the dielectric process and then to furans and urea-free resins. The continued development of binders for the production of chemically bonded cores and molds is being directed toward increasing productivity as well as achieving the dimensional repeatability necessary to meet the new challenges of net shape and near-net shape casting requirements (Table 6). Many patterns were made of epoxy resins and polyurethane and other expendables such as polystyrene. Current understanding of sand casting now addresses the fundamentals of clay mineralogy, sand preparation, sand compaction, and the physical properties of molding and core sands. A better understanding of molding sand technology has resulted in sands of a higher degree of uniformity being prepared for the repetitive green sand (clay-bonded) molding sand. This high degree of achievement could only be possible with the great advances in sand testing produced by the foremost researchers and developers of sand testing instrumentation. Permanent-mold methods, which preceded the loam mold and the sand mold by centuries, also progressed for die casting (high pressure, low pressure, and gravity), centrifugal casting (vertical and horizontal), and hybrid processes such as squeeze casting and semisolid metal casting. Other contributions to the continued advancement of metal casting include:
Cast metal-matrix composites Argon oxygen decarburization for steel refining
Source: Ref 5; see also Selected References
Table 6 Development of core and mold processes Process
Core oil Shell: liquid and flake Silicate/carbon dioxide Airset oils Phenolic acid-catalyzed no-bake Furan acid-catalyzed no-bake Furan hot box Phenolic hot box Oil urethane no-bake Phenolic/urethane/amine cold box Silicate ester-catalyzed no-bake Phenolic urethane no-bake Alumina phosphate no-bake Furan sulfur dioxide Polyol urethane no-bake Warm box Free radical cure sulfur dioxide Phenolic ester no-bake Phenolic ester cold box
Approximate time of introduction
1950 1950 1952 1953 1958 1958 1960 1962 1965 1968 1970 1974 1977 1978 1978 1982 1983 1984 1985
Source: Ref 9
400 years, foundry processes and materials often relied on the methods developed by Vannoccio Biringuccio, the 16th century “father of the foundry industry,” who recommended using the dregs of beer vats and human urine as binders for molding sand, both of which were in use well into the 20th century. Until the 1920s, sand testing consisted of squeezing a handful of sand to judge its ability to compact and stick together. Early in that period, a sand research committee of the American Foundrymen’s Society began to develop sand test methods. By 1924, standards were established that covered the various properties of molding sands. Since World War II, experimentation also accelerated in organic and chemical sand binders for the thermosetting of molds and cores. Beginning with the Croning process (shell process),
Foundry automation A scientific approach to the gating and riser
ing of castings using computer simulation of solidification Evaporative (lost)-foam casting and semisolid casting Computer-aided design and manufacture of castings Integrated foundry systems The melting of metals using the plasma arc cupola
Metal Casting Methods Metal casting is unique among metal forming processes, with a variety of casting processes (Table 7) (with Fig. 1, Ref 10). Very broadly, casting processes fall into two groups (Fig. 2): expendable-mold and permanent-mold processes. Each manufacturing process has the capability of producing a part with a certain range of tolerance, surface finish, and complexity. Manufacturing costs depend on the required dimensional tolerance, surface detail, and the quantity of parts required. For each process, there is a minimum batch size below which it is not economical to go because of costs of tooling, fixtures, and equipment. Also related to part cost is the production rate or the cycle time, the time required to produce one part. The most commonly used casting processes are evaluated with respect to these characteristics in Table 8 (Ref 11). Virtually any metal that can be melted can and is being cast. Choice of process and mold material (Table 9) is greatly influenced by the melting point of the alloy. A vital issue is avoidance of voids (due to insufficient feeding of melt) and porosity (due to dendritic solidification and the presence of gases). Success depends, to a large extent, on the skill of the mold designer, but there are very strong influences of the metal (alloy) and mold material. Even with the best design, solidification on the mold walls closes passage of melt at a certain distance, and this leads to a minimum allowable section thickness. There are large differences in the fluidity of alloys, and therefore, minimum allowable section thickness depends on the alloy being cast. Lower minimum thickness can be allowed with zinc, aluminum, and cast
History and Trends of Metal Casting / 9 Table 7
General characteristics of casting processes
See Fig. 1 for shape descriptions Casting process Characteristic
Part Material (casting) Porosity and voids(a) Shape(b) See Fig. 1 for shapes
Green sand
Resin-bonded sand
Plaster
Lost foam
All C–E All
All D–E All
Zn to Cu D–E All
Al to cast iron C–E All
All E All
Minimum section, mm (in.) Minimum core diameter, mm (in.) Surface detail(a)
0.01–300,000 (0.02–660,000) 3–6 (0.12–2.4) 4–6 (0.16–0.24) C
0.01–100 (0.02–220) 2–4 (0.08–0.16) 3–6 (0.12–0.24) B
0.01–1000 (0.02–2200) 1 (0.04) 10 (0.4) A
0.01–100 (0.02–220) 2–4 (0.08–0.16) 4–6 (0.16–0.24) C
Cost Equipment(a) Die (or pattern)(a) Labor(a) Finishing(a)
C–E C–E A–C A–C
C B–C C B–D
C–E C–E A–B C–D
Production Operator skill(a) Lead time Rates (piece/h mold) Minimum quantity
A–C Days 1–20 1–100
C Weeks 5–50 100
A–B Days 1–10 10
Size, kg (lb)
Permanent mold
Investment
Die
Zn to Cu A–C Not T3, 5, F5
0.01–100 (0.02–220) 1 (0.04) 0.5–1 (0.02–0.04) A
Zn to cast iron B–C Not T3, 5, F5 with solid core 0.1–100 (0.2–220) 2–4 (0.08–0.16) 4–6 (0.16–0.24) B–C
B–C B–C C C–D
C–E B–C A–B C–D
B B C B–D
A A E C–E
C Weeks–months 1–20 500
A–B Hours–weeks 1–1000 10–1000
C Weeks 5–50 1000
C–D Weeks–months 20–200 100,000
<0.01–50 (0.02–110) 0.5–1 (0.02–0.04) 3 (Zn: 0.8) (0.12; Zn: 0.03) A–B
(a) Comparative ratings, with A indicating the highest value of the variable, E the lowest (e.g., investment casting gives very low porosity, produces excellent surface detail, involves moderate to low equipment cost, medium to high pattern cost, high labor cost, medium to low finishing cost, and high operator skill. It can be used for low or high production rates and requires a minimum quantity of 10 to 1000 to justify the cost of the pattern mold). (b) From Fig. 1. Source: Adapted from Ref 10
iron than with steel. Streamlined flow also is helpful; thus, sharp corners are avoided. Thicker sections solidify last and must be fed adequately. Metals and alloys have different castability characteristics based on their combination of liquid-metal properties and solidification characteristics that promotes accurate and sound castings. Castability factors include fluidity, shrinkage, and resistance to hot cracking: Shrinkage is whether the alloy has a high
tendency to form shrinkage porosity.
Resistance to hot cracking is the ability of
the alloy to withstand stresses developed by the contraction while cooling from the hotshort temperature range. Fluidity is the ability of the liquid alloy to flow readily in the mold and to fill thin sections. Fluidity is usually evaluated by pouring the metal into a narrow spiral mold cavity and measuring the length of penetration of the metal before it solidifies. Fluidity is inversely related to the temperature range over which solidification occurs. Thus, pure metals and eutectic alloys, with short freezing range, have greater fluidity Because the viscosity of molten metals is very low, fluidity is determined more by the solidification dynamics of the metal and the mold. A high surface tension of the liquid metal reduces fluidity, and oxide films forming on the surface of the molten metal thus have a significantly adverse effect on fluidity. Insoluble inclusions also reduce fluidity. Castability ratings from the three factors are often described on a scale of 1 to 5 for each of the important factors. Table 10 shows typical
applications for metal alloys and relative ratings for castability, weldability, and machinability. As an example of a castability comparison in terms of fluidity, Fig. 3 shows three fluidity spiral samples of gray, compacted graphite, and ductile (nodular) iron. They were poured from the same base metal at the same temperature and have comparable chemical compositions. They differ only in their melt treatment. Gray (flake graphite) iron shows the best fluidity, and ductile iron shows the worst. Compacted graphite iron occupies an intermediate position. However, because compacted irons are stronger than gray iron with the same carbon equivalent (CE), a higher CE can be used to obtain the same strength, allowing greater fluidity and easier pouring of thin sections.
Production Trends The use of metal castings is pervasive throughout the economies of all developed countries, both as components in finished manufactured goods and as finished durable goods (Ref 1). As indicated earlier, castings are used in 90% of all manufactured goods and in all capital goods machinery used in manufacturing. They are also extensively used in transportation, building construction, municipal water and sewer systems, oil and gas pipelines, and a wide variety of other applications. In the broadest sense, foundries are categorized into two general groups: ferrous foundries and nonferrous foundries. For marketing purposes, foundries are also categorized according to the nature of their operations as being either captive (producing castings for their own use)
or jobbing (producing castings for sale). The market is sometimes further broken down by major casting processes when they can be readily identified or are particularly significant. This is done more often in the case of nonferrous castings; in this case, the product is usually categorized as being produced in sand, permanent mold, or die cast. Because 90% or more of all iron and steel castings are produced in some form of sand medium, distinguishing the process used is not quite as significant. Ferrous castings shipments are usually classified by market category. For example, iron castings are generally categorized as engineered (designed for specific, differentiated customers) and nonengineered (produced in large volumes of interchangeable units, usually consisting of ingot molds and pressure and soil pipe). The diversity among the various foundries makes it difficult to determine the exact structure of the industry. For example, it is not unusual for a single operating foundry to produce a variety of metals and alloys, both ferrous and nonferrous, in the same plant. Some also use a variety of processes in their operations. Many aluminum foundries, for instance, use both sand and permanent-mold processes, and some even produce die castings in the same facility. In addition, the foundry industry consists of a variety of large and small facilities. Table 11 gives recent census results on the number of casting facilities in major industrialized countries. Ferrous castings easily make up the largest production sector of the foundry industry in terms of volume, with gray iron outdistancing the other major metals constituting this group (Table 2). In comparison, iron and steel
10 / Applications and Specification Guidelines Increasing spatial complexity
Type (abbreviation)
0 Uniform cross section
1 Change at end
2 Change at center
3 Spatial curvature
4 Closed one end
5 Closed both ends
6 Transverse element
7 Irregular (complex)
Round (R)
Bar (B)
Section, open (S) Semiclosed (SS)
Tube (T)
Flat (F)
Spherical (Sp)
Fig. 1
The choice of possible manufacturing methods is aided by classifying shapes according to their geometric features. Source: Ref 10
Shape-casting processes
Expendable mold
Permanent pattern
Sand molding Clay Green Skin-dry Dry Carbon dioxide Cement Oil Resin Cold box Hot box Shell Flaskless
Fig. 2
Permanent mold
Expendable pattern
Slurry molding Plaster Ceramic
Semipermanent (sand cores)
Gravity Slush
Permanent (metal cores)
Low pressure
Hot chamber Sand (full-mold)
Slurry Lost wax Centrifuged
True centrifugal (outer mold)
Semicentrifugal (complete mold)
General classification of shape-casting processes. Source Ref 10
Centrifuged (complete mold)
Squeeze
Die
Cold chamber
foundries in North America shipped an average of 15.5 million tons of castings annually in the years from 1947 to 1986. The following decade (1967 to 1976) saw a significant upturn as the average annual shipments of iron and steel castings rose to slightly more than 17.5 million tons. Iron casting shipments have averaged approximately 14 million tons per year during the last four decades. From 1977 to 1986, ferrous casting shipments from American foundries fell to approximately 9.5 million tons by 1982 (Fig. 4). In this same time period, casting production around the world has shown a significant decline, generally reflecting the state of manufacturing. World production of ferrous castings fell some 27%, from 64.6 million tons in 1979 to 47.3 million tons in 1982 (Table 12). Nonferrous shipments followed a similar path in this time frame (Fig. 5). Recent Production Trends (1995 to 2005). Currently, metal casting production is dominated by China, along with large growth rates
History and Trends of Metal Casting / 11 Table 8
Casting processes and their attributes Surface roughness
Process
Dimensional accuracy
Complexity
Production rate
Production run
Relative cost
Size (projected area)
Pressure die casting Centrifugal casting Sand casting Shell mold casting Investment casting Powder metallurgy
L M H L L L
H M M H H H
H M M H H H
H/M L L H/M L H/M
H M/L H/M/L H/M H/M/L H
H H/M H/M/L H/M H/M H/M
M/L H/M/L H/M/L M/L M/L L
Key: H M
>250 >63 and <250
High Medium
>100 >10 and <100
>5000 >100 and <5000
High Medium
>0.5 >0.02 and <0.5
L Units
<63 min.
<0.005 >0.005 and <0.05 >0.05 in.
Low ...
<10 Parts/h
<100 Parts
Low ...
<0.02 m2
Source: Adapted from Ref 11
Table 9
Process and mold material options for casting
Process
Sand casting Investment casting Die casting
Cast iron
Carbon steel
Alloy steel
Stainless steel
Aluminum and aluminum alloys
Copper and copper alloys
Zinc and zinc alloys
Magnesium and magnesium alloys
Titanium and titanium alloys
Nickel and nickel alloys
Refractory metals
A B
A A
A A
A A
A A
A A
B B
A B
B B
A A
B B
C
C
C
C
A
B
A
A
C
C
C
A, normal practice; B, less-common practice; C, not applicable. Source: Adapted from Ref 12
Table 10 Typical applications for castings and general rating of castability, machinability, and weldability Note that these are general overall ratings for a class of alloys. Within each class, different alloys may vary considerably with respect to specific castability characteristics of resistance to hot cracking, fluidity, and shrinkage tendency.
Table 11 Number of operating metal casting facilities by nation in 2005
Casting alloy type
Country
Applications
Aluminum Copper Gray iron Magnesium Malleable iron Nickel Nodular iron Steel (carbon and low alloy) Steel (high alloy) White cast iron Zinc
Castability Machinability Weldability
Pistons, clutch housings, exhaust manifolds Pumps, valves, marine propellors Engine blocks, gears, brake disks and drums, machine bases Crankcases, transmission housings Farm and construction machinery, heavy-duty bearings, railroad rolling stock Gas-turbine blades, pump and valve components for chemical plants Crankshafts, heavy-duty gears Die blocks, heavy-duty gear blanks, railroad wheels
1–2 1–3 1 1–2 2
1–2 1–2 2 1 2
3 3 5 2 5
3
3
3
2 3
2 2–3
5 1
Gas-turbine housings, pump and valve components, rockcrusher jaws Ball mill liners, shotblasting nozzles, railroad brake shoes, crushers Door handles, radiator grilles, carburetor bodies
3
3
1
2
4
4
1
1
5
1, excellent; 2, good; 3, fair; 4, very poor; 5, difficult. Source: Ref 13
FG iron
CG iron
SG iron
10 cm
Fig. 3
Spiral fluidity samples poured from gray (flake graphite), compacted graphite, and ductile (spheroidal graphite) iron. Source: Ref 14
Austria Belgium Brazil Canada(a) China Croatia Czech Republic Denmark Finland France Germany Great Britain Hungary India Italy Japan Korea Lithuania Mexico Netherlands(a)(b) Norway Poland Portugal Romania Russia(e) Slovenia(e) Slovakia South Africa Spain Sweden Switzerland Taiwan(a) Thailand Turkey Ukraine(e) United States
Iron
Steel
Nonferrous
Total
12 18 467 60 17,000 16 106 10 16 115 212 218 41 ... 180 457 499 8 380 16 7 173(c) 50 58 ... 16 12 95 67 36 19 520 230 857 400 619
4 9 139 24 4,700 5 32 0 6 38 53 46 27 ... 27 78 140 3 10 0 3 36(c) 10 41 ... 4 7(f) 42 35 15 0 45 26 73 233 262
29 9 658 66 4,300 23 56 11 15 352 365 216 143 ... 870 1,173 208 5 250 5 11 245(d) 38 65 ... 33 32 119 57 84 32 350 220 361 437 1,499
45 36 1,264 150 26,000 44 194 21 37 505 630 480 211 4,500 1,077 1,708 847 16 640 21 21 454 98 164 1,900 53 51 256 159 135 51 915 476 1,291 960 2,380
Note: “. . .”indicates unreported data.(a) 2004 data. (b) Includes only members of AVNeG (General Assoc. of Dutea Foundries). (c) Includes approximately 24 metal casting facilities making both iron and steel. (d) Includes approximately 40 metal casting facilities working with both ferrous and nonferrous metals. (e) 2002 data. (f) Includes investment casting facilities. Source: Modern Casting
12 / Applications and Specification Guidelines 25 000
20 000 Thousands of tons
Table 12 Ferrous casting production from 1982 to 1986
Gray iron Ductile iron Malleable iron Steel Total
15 000
Figures given in thousands of metric tons Year Country
United States Japan West Germany France United Kingdom Brazil Canada(a) Korea Taiwan
10 000
5 000
0 78
Fig. 4
79
80
81
82 83 Year
84
85
86
87
1982
1983
1984
1985
1986
8,651,455 5,735,001 3,501,299 2,230,264 1,621,700 1,207,700 671,913 639,000 401,990
9,313,866 5,515,999 3,311,600 2,014,387 1,549,100 986,142 781,722 762,000 492,150
10,856,200 5,128,896 3,387,300 1,855,634 1,490,500 ... 981,730 805,000 599,122
9,991,393 5,240,261 3,499,900 1,902,608 ... 1,407,324 856,077 848,000 637,660
8,564,725 4,860,591 3,452,000 1,797,960 1,222,000 1,543,147 865,777 934,000 778,930
(a) Canada’s tonnage is estimated to be 75% of actual. Source: Ref 15
U.S. ferrous casting shipments from 1978 to 1987. Source: Ref 2
24 2500
21 Aluminum base Copper base
18
Zinc base Total
Castings, millions of tons
Thousands of tons
2000
China
1500
1000
500
U.S.
15 12
Russia
9
Japan
6 0 77
78
79
80
81 82 Year
83
84
85
3
86
Germany
India
0 1998
U.S. nonferrous casting shipments from 1977 to 1986. Source: Ref 2
in India (Fig. 6). The market for castings is likely to shift more to Asia and, to a lesser extent, to Central and Eastern Europe. However, despite the contraction of industry production bases in the West, high-quality manufacturing should continue to meet the needs of home markets and the demands of more productivity, energy efficiency, and environmental sustainability. German metal casting sites are reported to have the highest productivity in 2005, followed by the United States and France (Fig. 7). From 1995 to 2005, global output in the foundry industry grew by some 25%, topping 85 million tons. This is due in part to strong growth in its traditional markets but also to increasing use of these materials in a wide variety of other applications. These include the use of cast iron products in wind farms and large motors for power generation and ships, as well as the increasing popularity of light materials in automotive engineering, especially aluminum. During this period, output of cast steel continued to grow at an average pace, although there was something of a downward trend until 2004. Cast steel tends to be used less often in automotive and mechanical engineering than in
1999
2000
2001
2002
2003
2004
2005
6.0 5.5 5.0
Castings, millions of tons
Fig. 5
4.6 4.0 India
3.5
Brazil
3.0
France
2.5 Italy
2.0 1.5
Korea
1.0 1998
Fig. 6
1999
2000
2001
2002
2003
2004
Tonnage of major casting-producing nations from 1998 to 2005. Source: Ref 3
2005
History and Trends of Metal Casting / 13
8,108
Productivity per metal casting site
5,419
Germany (630 sites)
4,639 France (505 sites) Russia (1900 sites)
4,010 3,097 2,360 2,348 2,242 1,358 939 0
Fig. 7
U.S. (2380 sites)
Japan (1708 sites)
Italy (1077 sites) Brazil (1264 sites) Korea (847 sites)
India (4500 sites)
China (24,000 sites) 1
2
3
4
5
6
7
8
9
Productivity per metal casting site in 2005 for major industrial nations. See Table 11 for breakdown of site count. Source: Ref 3
1995 20%
United States
17%
China
16%
Commonwealth of Independent State
10%
Japan
5%
India
6%
Germany
26%
Other countries
30%
China
14%
United States
8%
India
8%
Japan
7%
Russia
7%
Germany
26%
Other countries
2005
Fig. 8
National shares in cast iron production. Source: Ref 3
power plant and factory construction. It benefited from robust growth in its main markets in 2005.
Cast Iron. China dominates cast iron production (Fig. 8). In terms of volume, cast iron and cast steel account for the bulk of foundry
industry production, comprising approximately 80% of total output. Global production has increased since 1995 by some 15%, with Asia recording the fastest rate of growth: by 2005 the region was already responsible for over half of world output (Fig. 8). Conversely, production in North and South America and Europe declined slightly. However, the situation varies considerably from country to country. In North America, production of cast iron in the United States and Canada fell by approximately 15%, likewise reducing their share of the world market (Fig. 8). Production volume has increased again since 2002, although it still remains below the 1995 figure. Even so, the United States continues to be the world’s second-largest producer. By contrast, Latin American countries such as Brazil and Mexico boosted their output by over 80%. From 1995 to 2005, China doubled its share in the global market. In Japan, on the other hand, production volume dropped by nearly 10%. Growing at a fast pace, output in India rose by approximately 80%, surpassing Japanese production for the first time in 2005. Together, these three countries account for approximately 90% of Asian production of cast iron. Nonferrous Castings. Nonferrous metal casting have sustained an upward trend. Global output of nonferrous metal casting products grew by some 75% from 1995 to 2005 (Fig. 9). Although production in regions around the globe gained ground, growth in North and South America (50%) lagged behind the rest of the world, where the volume of production more than doubled. In the Western hemisphere, the United States accounted for a good two-thirds of production, its 17% rate of growth pulling down the regional average. Copper and copper alloys as well as zinc products contracted by 6 and 20%, respectively, whereas output of all other materials expanded. Particularly notable is the robust increase in aluminum production in Canada, where Alcan Inc., one of the world’s top aluminum producers, boosted capacity. In Mexico, copper output increased tenfold, evidently due to the transfer of production from its northern neighbor. The volume of production in Asia increased by over 130%, driven first and foremost by rapid expansion in China, where output nearly tripled, and India, which likewise succeeded in significantly widening its share of the world market. Although production in Japan also rose, the country’s relative importance in the world market diminished. Aluminum continues to dominate nonferrous metal casting. Worldwide, aluminum accounts for approximately 80% of nonferrous metal casting, trailed by copper and copper alloys with 10%; the share of zinc is slightly over 6%. Production of aluminum cast parts doubled during the period under review, while other metals posted total gains of “only” approximately 10%. This is due primarily to greater use of aluminum in the automotive sector. Owing to its low specific weight, aluminum is an increasingly common
14 / Applications and Specification Guidelines In million tons 14 12 10 8 6 4 2 0 1995
2000
2001
2002
2003
2004
2005
Asia Europe North and South America
Fig. 9
Trends in production of nonferrous metal castings. Source: Ref 3
1995 30%
United States
19%
Japan
9%
China
8%
Germany
8%
Italy Commonwealth of Independent State
3% 1%
India
22%
Other countries
18%
United States
16%
China
13%
Japan
8%
Russia
7%
Italy
6%
Germany
4%
India
28%
Other countries
2005
Fig. 10
National shares in production of aluminum castings. Source: Ref 3
material in lightweight engineering. In the aviation industry, too, it continues its rapid ascent. The relative share of the traditional industrial nations in global aluminum production declined markedly from 1995 to 2005 (Fig. 10). While the United States and Japan increased their production volumes by approximately 30%, China’s output practically tripled. India, starting from a very slender base, also expanded its share massively. By 2005, China’s
production volume had nearly reached that of the United States, which it is likely to over-take during the next two years.
Market Trends The two largest sectors of metal castings are the automotive industry and machinery manufacturers. In the automotive industry, the
world’s major industrial economies account for approximately 85% of global car production. However, vehicle output in China and the nations of Central and Eastern Europe has been growing at a much faster pace. The next ten years are likely to see a continuation of this trend. In 2015, the world’s traditional car-making nations will probably be responsible for just 70% of global production; China’s share is likely to increase from approximately 10% to a more than 15% (Ref 16). Another important market for castings is the machinery/instrumentation sector, with products ranging from massive hubs, housings, and shafts for power generators to miniscule components for semiconductors. The world’s two leading makers of machines are the United States and Japan, which both lost ground from 2000 to 2005. In context, German manufacturers widened their share of the world market (Fig. 11). The Chinese sector is also expected to grow from its fourth place ranking. Although China remains one of the world’s biggest importers of plant and machinery, it is increasing its presence as an exporter as well. In 2005, China’s machine-tools sector (with annual sales of approximately £5 billion) exceeded that of Italy. Even with a distinct slowdown in the growth of imports, foreign suppliers meet 60% of China’s total demand for machine tools.
REFERENCES 1. “Competitive Assessment of the U.S. Foundry Industry,” USITC Publication 1582, U.S. Department of Commerce, Sept 1984, p xiii 2. D.P. Kanicki, Casting Advantages, Applications, and Market Size, Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 37–46 3. Annual Census of World Casting Production, Mod. Cast., Dec 2005 and 2006 4. M. Goodway, History of Casting, Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 15–23 5. Timeline of Casting Technology, Mod. Cast. Cast Expo Issue, May 2005 6. G.L. Werley, Slush Casting, Forging and Casting, Vol 5, Metals Handbook, 8th ed., American Society for Metals, 1970, p 334 7. E. Rooy, A History of Aluminum Foundry Alloy Development, AFS International Conference on Structural Aluminum Castings: Recent Advancements for the Design and Manufacturing of Reliable Structural Aluminum Castings, Nov 2–4, 2003 (Orlando, FI), p 87–98 8. B.L. Simpson, in History of the Metalcasting Industry, American Foundrymen’s Society, 1948 9. E.L. Kotzin, Development of Foundry Technology in the United States, Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 15–23
History and Trends of Metal Casting / 15 SELECTED REFERENCES (1980–1990) AFS/CMI Conference on Green Sand TechUnited States
Japan
Germany
China Italy
France Great Britain
South Korea billion
0
50
100
150
200
250
300
350
2000 2005
Fig. 11
Global engineering sales of machinery/machine-tool sector. Source. Ref 17
10. J.A. Schey, Introduction to Manufacturing Processes, 2nd ed., McGraw-Hill, 1987 11. E.B. Magrab, Integrated Product and Process Design and Development, CRC Press, Inc., 1997 12. G. Boothroyd, P. Dewhurst, and W. Knight, Product Design for Manufacture and Assembly, Marcel Dekker, Inc., 1994 13. S. Kalpakjian, Manufacturing Processes for Engineering Materials, 2nd ed., Addison-Wesley, 1991, p 262 14. D.M. Stefanescu, R. Hummer, and E. Nechtelberger, Compacted Graphite Irons, Casting, Vol 15, ASM Handbook, ASM International, 1988 15. Census of World Casting Production, Mod. Cast., American Foundrymen’s Society, 1987 16. Foundry Industry, IKB Deutsche Industriebank, June 2007 17. Verband Deutscher Maschinen und Anlagenbau, German Engineering Federation SELECTED REFERENCES (published after 1990) A Tribute to Keith Millis and the Unveiling of
Ductile Iron 50 Years Ago, Mod. Cast., Oct 1998
J.R. Bodine, 1999 Charles Edgar Hoyt Memo-
rial Lecture: From a Monument to the Vega: The Journey of the Aluminum Casting Industry, Mod. Cast., May 1999, p 60–63 R.M. Cooper, The History of the Soho Foundry, GEC Rev., Vol 11 (No. 2), 1996, p 114–119 J. Gerin Sylvia, Cast Metals Technology, AFS, 1990 F. Habashi, Milestones in the History of Metallurgy, BUMA-V: Fifth International Conference on the Beginnings of the Use of Metals and Alloys, April 21–24, 2002 (Gyeongju, Korea), p 53–60 E.L. Kotzin, 1997 Hoyt Memorial Lecture: Lost Wax to Lost Foam: Reflections on the Past, Present and Future, Mod. Cast., July 1997, p. 42–45 Modern Casting’s A Look Back at the 20th Century Series, 2000–2001 Proceedings of 1992 Coreless Induction Melting Conference, AFS, 1992 N. Wukovich, A Short History of the Steel Foundry, Trans. Am. Foundrymen’s Soc., Vol 104, April 20–23, 1996 (Philadelphia, PA), p 643–649 1991 AFS International Sand Reclamation Conference, AFS, 1991 1995 Lost Foam Technology and Applications Conference Proceedings, AFS, 1995
nology—Productivity for the ’80s, AFS, 1983 N.N. Bubtsov, History of Foundry Practice in USSR, Moscow, 1962; trans., Washington, D.C., 1975 Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988 J. Foster, “The Iron Age Moulds from Gussage All Saints,” Occasional Paper 12, British Museum, London, 1980 Foundry Trade J., Sept 11, 1986 High-Pressure Molding, 2nd ed., AFS, 1987 E.L. Kotzin, Metalcasting & Molding Processes, AFS, 1981 J.P. LaRue, Basic Metalcasting, AFS, 1989 C.S. Smith, The Early History of Casting, Molds, and the Science of Solidification, A Search for Structure: Selected Essays on Science, Art, and History, MIT Press, Cambridge, MA, and London, 1981, p 127–173 R.F. Tylecote, History of Metallurgy, The Metals Society, London, 1976 R.F. Tylecote, The Early History of Metallurgy in Europe, Longman, London, 1987 C.F. Walton and T.J. Opar, Iron Castings Handbook, Iron Castings Society, Inc., 1981 T.A. Wertime and J.D. Muhly, The Coming of the Age of Iron, Yale University Press, New Haven, 1980 T.A. Wertime and S.F. Wertime, Early Pyrotechnology: The Evolution of the First FireUsing Industries, Washington, D.C., 1982
SELECTED REFERENCES (published prior to 1980) L. Aitchison, A History of Metals, New York,
1960
Aluminum: Design and Application, Vol II,
American Society for Metals, 1967
J.A. Gitzen, Progress in the Development
and Utilization of Additives in the Foundry Industry, AFS, 1957 P. Knauth, The Metalsmiths, New York, 1974 C.A. Sanders and D.C. Gould, History Cast in Metal, AFS, 1976 B.L. Simpson, Development of the Metal Castings Industry, American Foundrymen’s Association, Chicago, 1948 B.L. Simpson, History of the Metalcasting Industry, AFS, 1948
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 16-28 DOI: 10.1361/asmhba0005188
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Metalcasting Technology and the Purchasing Process American Foundry Society Technical Department
CAST METAL COMPONENTS are not items that are commonly bought out of catalogs; most often they are tailor-made objects, designed and produced to the specific needs of a particular customer. As such, cast components cannot and should not be treated as low-value, high-volume secondary commodity items, where price is the sole purchase criterion. Cast components should be considered highvalue manufacturing products, where the metalcasting facility is an extension of the buyer’s manufacturing capability, involving long-term relationships and open, on-going communication at all levels. As high-value items, the purchasing criteria for cast components include a balance of quality, cost, service, and delivery. Many original equipment manufacturing (OEM) component buyers are recognizing that a carefully considered purchasing policy is the key to bestvalue procurement of metal castings. The policy should define the buyer-seller relationship, based on the following principles:
Long-term relationships based on shared
benefits and best value for all parties
Qualified cast component suppliers who
meet production quality standards, deliver on schedule, provide expertise and knowledge to support component design and improvement, and provide best-value pricing Open and continuous communication at executive, purchasing, and technical levels to support concurrent engineering/design, design for manufacturability, and continuous improvement Over the last 20 years, significant change has occurred in the industrial purchasing environment—the old purchasing rules and situations no longer apply. Today’s (2008) manufactured product market is marked by six major changes that impact both the buyers and the sellers: Globalization with new competition and new
markets: Today (2008), information, goods, and services move around the globe at an unprecedented speed. The OEM purchasing
managers can find sources for manufactured products in Malaysia, Mexico, or Poland as easily as they find a source in Ohio, Georgia, or California. The flip side is that a producer in Wisconsin can market components in Europe, Japan, and Brazil. Products can be delivered the next day from around the world. Best-value purchasing: The competitive market is driving everyone to focus on providing the best value to the final customer. Best value means that buying decisions are made not just on cost but also on a balance of quality, service, price, and availability/delivery. Recognizing how subcontractors impact final cost and quality, smart OEM manufacturers purchase their components and supplies from vendors who consistently can meet the best-value requirements. Value-added services or “one-stop shopping” for castings is common. Outsourcing: In the old days, many OEM manufacturers used internal production divisions to supply components. Today (2008), more OEM companies are recognizing that it is smarter to outsource their components through strategic partnerships with sources that have the cutting-edge expertise and advanced technology to provide a best-value product. Supplier management and quality assurance: When the quality and cost of a manufacturer’s end product depend critically on components from an outside supplier, the manufacturer will take a supply-management approach to working with the source. In supply management, buyers work with all their purchasing sources, developing long-term relationships, providing oversight on quality assurance, and working cooperatively with the sources in product-improvement teams. Continuous improvement: Nothing stands still for long anymore. In a competitive market, if a company does not improve its product and services, its competitors will improve theirs and win the contract. The supplier must work with the customer to find
and exploit opportunities to improve quality, reduce cost, and cut cycle/production time. That kind of innovative effort applies to working on current production components as well as on new product designs. Information availability and velocity: The internet and electronic commerce have changed the way that information moves and circulates. Both buyers and sellers are expanding their electronic commerce reach in locating and assessing sources and opportunities for purchases/sales. Electronic marketplaces serve as locations for posting requests for quotes, advertising production capabilities, and finding matches between requirements and services. In a similar manner, suppliers are posting information about their capabilities and expertise on the internet, providing instant information to potential customers. Electronic commerce extends beyond the marketing function. All types of information (orders, contracts, technical, and financial) move electronically rather than by paper and mail. The speed of information movement is measured in minutes rather than days, cutting cycle time, reducing data errors, and overcoming geographic separation between buyers and suppliers. The metalcasting market is not isolated from these changes. Cast component buyers and sellers recognize that they must respond to the new market imperatives if they are to stay in business and remain competitive.
The Purchasing Process Purchasing is a critical process in any business. A company purchases raw materials, subcomponents, and support services that are essential to the production and delivery of the finished product to their customer. There is tremendous value in developing a purchasing process that is tailored to a the, specific needs of a business and that obtains components and services from outside sources in an effective, efficient, and timely manner.
Metalcasting Technology and the Purchasing Process / 17
Fig. 1
Four basic steps of purchasing process
Whether it is a product or service, the purchasing process consists of four basic steps (Fig. 1): 1. Define requirements and develop a purchasing plan. 2. Request and evaluate bids from potential sources. 3. Select a source and negotiate contract terms. 4. Carry out the contract and pursue continuous improvement. These basic steps of the purchasing process are described in the next four sections of this article. It is also important to remember that purchasing does not end with the selection of a source and contract agreement. The process is not complete until the required product/service is received and meets the requirements for quality, cost delivery, and service. Purchasing continues through product delivery (Fig. 1). The following sections provide detailed guidance on purchasing cast components, going through each step and describing specific issues and approaches that have proved useful in purchasing castings.
Define Requirements and Develop a Purchasing Plan The first step in any purchasing process is to define the product requirements and develop a purchasing plan. For cast components, product requirements are based on technical performance and targets for quality, quantity, schedule and cost. With this information, a complete and accurate request for quotation can be drafted and distributed to metalcasting facilities for consideration. To be a smart buyer, one needs more than just a requirements list. A purchasing plan that outlines and defines the following is also required: Objectives and issues in the purchase Purchase requirements How sources will be identified, screened,
and selected How the purchase will be contracted and managed for quality, cost, and schedule This plan may be simple and straightforward for small, low-value purchases; it may be very detailed and extensive for a high-value, critical purchase.
Request a Quotation Purchasing cast components is generally not a simple matter of transferring pattern numbers and quantity desired from the engineering bill of materials to a purchase order. Careful consideration of technical quality, quantity, and schedule factors is necessary to ensure that all pertinent and important information is supplied to the metalcasting facility. When necessary information is not provided, this slows the procurement and results in an inaccurate cost estimate or in a cast component that does not fully meet requirements. When possible, the metalcasting facility should be supplied with Computer-aided design files or blueprints for both the as-cast and the machined part. If the metalcasting facility is to provide the most suitable and useful cast components, it is important that the buyer indicate the areas to be machined and the locating points for machining. In addition, the buyer must know and understand the critical factors in specifying cast components and how those factors impact production and costs in the metalcasting facility. A request-for-quotation (RFQ) form is used to good advantage by many metalcasting facilities and cast component buyers. It provides a convenient way to develop the component requirements in a comprehensive and detailed approach. A model RFQ form is provided in Fig. 2. The RFQ form supplies the metalcasting facility with complete information on component requirements and enables the metalcasting facility to make an accurate casting price quotation for the most suitable casting for a particular application.
The Purchasing Plan With the detailed requirements for the casting defined in a draft RFQ, the buyer can develop a purchasing plan (Fig. 3.) that specifies: Major purchasing objectives Purchasing strategy Deliverables expected from the metalcast-
ing facility at each stage of the purchasing process Desired purchaser-metalcasting facility working relationship to ensure that quality, schedule, and cost objectives are met Specific selection criteria Purchasing tasks, responsibilities, and schedule
Purchasing Objectives. The buyer must clearly define the overall requirements and priorities in purchasing the cast component. Those requirements must be described in hard numbers that are clearly understandable and easy to communicate. The overall requirements package will include the engineering requirements as well as cost targets, trade-off considerations, and the company supply-chain management strategy. Each metalcasting purchase will have different answers to the following questions: What are the quality and performance require-
ments for the cast component?
How many cast components will be purchas-
ed over a certain period?
What is the cost target for the cast compo-
nents?
What is the delivery schedule for the cast com-
ponents?
How critical is the casting to the project/busi-
ness in terms of performance, risk, schedule, and cost? How will the different objectives (quality, performance, cost, and schedule) be balanced and prioritized? This requires a trade-off analysis (either informal or formal). How does the cast metal component purchase fit into the overall supply-chain strategy for the company? The first priority is to meet the quality, quantity, service, schedule, and cost requirements for a given cast part. A cast component order is seldom a one-time purchase for a company. More often, cast components are integral to a company product line, and the purchase must be considered in the context of a company overall supply-chain strategy. A corporate supplychain strategy considers how to maximize the value and efficiency of all purchases through coordination and integration, rather than simply maximizing the value of individual purchases. Purchasing Strategy. Purchasing is nothing less than asset management, where the assets are suppliers and the goal is to maximize performance and value while minimizing risk across the casting purchase process. As with any investment, the buyer should develop a strategy that will meet objectives with a balance between risk and return. For a purchasing strategy, the buyer must ask the following questions in light of the particular requirements for a cast component: How much benefit is there in finding the best
possible source?
What is the level of risk in purchasing from
a nonperforming source?
What are the cost drivers for the component,
and where can the metalcasting facility have a positive or negative impact on those drivers? What investment of time and resources is necessary to identify a suitable source versus a best-possible source? How does the purchase of a given cast metal component fit into the overall purchasing
18 / Applications and Specification Guidelines
Fig. 2
Casting request-for-quotation form
Metalcasting Technology and the Purchasing Process / 19
Fig. 3
Basic elements of a purchasing plan and strategy
strategy for castings within a division or company? How will the company work with the supplier to maximize value during all stages of the purchase cycle?
cast components will include more than just the receipt of qualified castings at the shipping dock. It also should include details for technical partnership services during the entire purchasing process that include:
Answers to these questions will provide a framework for determining the value of a well-qualified metalcasting facility source to the success of a given project/program as well as the overall supply-chain management effort. In that framework, the purchasing agent then must consider the strategic questions for a particular purchase:
Initial design services by the metalcasting
Will the cast part be purchased from a single
Working Relationship Requirements. The working relationship between the customer and the metalcasting facility is a critical factor in defining a purchasing strategy that will provide best value for the casting purchase. The buyer must recognize from the first step that a cooperative team effort between the production design engineer and the metalcasting facility engineer is the key to a cast component design and production effort that meets the desired quality, cost, and schedule requiements. The buyer should work with the production design engineer during the initial purchase plan to define the goals, tasks, level of effort, and schedule required for a successful buyer/metalcasting facility working plan. The plan must consider consultation, technical services, and information exchange at each stage of the purchasing process. This plan should include:
source or multiple sources? What levels of metalcasting facility capability and production experience are necessary to successfully produce and deliver the casting on time and at cost? Do current casting suppliers have the capability and capacity to meet the new purchasing requirement? How will the performance of the supplier be evaluated in terms of quality, schedule, and cost? How will the supplier be integrated into the initial design process and future productimprovement efforts? What is the plan for developing and maintaining an open exchange of information between the purchaser and the supplier? How will the benefits, risks, and costs of the purchase be shared between the purchaser and the supplier? What kind of contract (firm-fixed price, cost plus fee, contract with performance/penalty fees) will best meet the purchase objectives and the overall purchasing strategy?
Deliverables Expected. If cast components were simple commodities that were bought off the shelf, then a contract would require only the delivery of a finished qualified component at a certain location, on a specific schedule, in a known quantity, and for a given price. However, each casting is unique, and each metalcasting facility has different capabilities, skills, and experience to bring to the table. Engineering service and supplier partnerships are critical factors in the design, production, and delivery of reliable, quality cast parts. For that reason, a deliverables list in a contract for
Fig. 4
Selection criteria of a purchasing plan
facility
Scheduled information exchange on quality Schedule and cost achievements and issues Quality planning and assurance during
production
Quality and cost-improvement initiatives
with shared benefits during production
Initial casting design (for performance and
manufacturability)
Tooling design and construction Casting development and qualification Component production with value enhan-
cement Selection Criteria. A purchasing plan always will include the criteria for evaluating, prioritizing, and selecting the bids from the responding metalcasting facilities. The criteria for qualifying and/or selecting a metalcasting facility and its bid for a specific casting fall into five categories (Fig. 4): Production facilities (capabilities, capacities,
and expertise)
Direct experience and past performance with
the required alloy, similarly sized and
complex parts, and the desired/ recommended casting method Production organization and capabilities (design and engineering support, testing and evaluation capabilities) Quality-assurance procedures and capabilities Quality, delivery, and cost (goals and reliability) Based on the purchasing objectives, the defined strategy, and the detailed deliverables and requirements, the purchasing plan will specify the criteria in detail along with weighting factors for evaluating the bids. Depending on the complexity, critical value, and risk for a given cast part, the evaluation approach for bids may take two general forms: For low-risk castings, qualify bidders as go/no
go against the different criteria, eliminate unqualified bidders, and then rank qualified bidders by quoted cost. For complex, critical components, grade bidders against the designated criteria (not qualified, qualified, or exceptionally qualified), use weighting factors for the different criteria, and then rank bidders high/low by summed and weighted criteria. The OEM buyer has a responsibility to assess the capability of a metalcasting facility to meet the purchase requirement. Assessing the capability of a metalcasting facility to produce high-strength structural cast components is difficult when the metalcasting facility is new to the buyer and does not have a track record based on previous purchases. A critical review of a metalcasting facility facilities, organizational structure, and process documentation is necessary to judge its capabilities. The capability survey is usually combined with qualitycontrol and business assessments before any effort to use the metalcasting facility as a supplier occurs. A questionnaire to address these issues is given in Fig. 5.
20 / Applications and Specification Guidelines The reviewer/assessor should write a capability report on the facility and the type of castings it is capable of providing. In considering and evaluating the five selection factors, the assessor should use the following guidelines to focus inquiries and determine the capabilities of the metalcasting facility. Metalcasting facilities making structural castings must have the proper equipment to attain the necessary quality. The production capability (and available capacity) of the facility should be consistent with the production quantity and rate required by the contract. A capability assessment should be made before giving any metalcasting facility a purchase order. Costs and schedules may be adversely affected if the capability assessment is not performed conscientiously. A user survey team approach often is necessary; the team usually is composed of engineering, quality-control, and purchasing personnel familiar with metalcasting facility operations. The survey usually is not repeated unless there are changes in the metalcasting facility, personnel, or organization or if a significant time has elapsed since the metalcasting facility last supplied components to the user. Metalcasting facility records should be available for similar castings that were successfully produced to comparable requirements. This demonstrated experience is an important factor in qualifying and selecting a metalcasting facility that can meet the quality and production requirements for cast metal parts. Metalcasting facility records should include mechanical test results and photographs or drawings of cast components. These results should be verified by periodic destructive tests. Separate test bar and appendage data are not adequate proof of capability. Metalcasting facilities that have not had experience producing the size, complexity, or quality of the required castings will need time to develop this expertise. Often, metalcasting facilities without experience are not aware of the requirements to cast a particular type of component. To gain the experience needed, they should start with small, less complex parts. Complex, high-performance metal castings must have high radiographic quality to be capable of exhibiting the necessary mechanical properties. Experience, combined with the quality-assurance equipment and personnel of the metalcasting facility, indicates if the metalcasting facility can achieve high-quality cast components. The ability of the metalcasting facility to hold tight dimensions regarding the location of critical features and the location of contoured mold lines is important. Special equipment, techniques, and procedures to mold, trim, heat treat, straighten, machine, and inspect are necessary to accomplish this on a production basis. The organizational structure of a metalcasting facility varies to some degree with the size of the facility. However, the organizational structure of most metalcasting facilities is basically the same:
senior management, production, quality assurance, marketing/sales, and maintenance/operations. The larger facilities usually have a greater depth of personnel in each department. Smaller metalcasting facilities are under direct day-byday control of fewer people. Sometimes, only the owner can make decisions. Large organizations tend to delegate authority, and the owner or chief executive officer is involved only in business decisions on an administrative level. Metalcasting facilities that produce quality cast parts need the technical capability to develop metallurgical and dimensional procedures and to solve metallurgical and dimensional problems as they occur. A degreed metallurgist with several years of on-site experience should be a permanent part of the organizational structure. Each cast component is produced in a manner that is specific to that part. Casting procedures must be adequately documented, and methods should be in place to monitor and record the actual casting process. The details of each operation should be maintained in a manner that permits them to be retrieved easily each time production is required. Photographs can aid considerably in the process description. Details on rigging, chilling, coring, molding and mold assembly, melting, pouring, processing, welding, heat treatment, straightening, and final inspection should be recorded during processing and should be available for review. A quality-control department should be staffed to implement the necessary inspections and testing throughout the entire metalcasting facility process (e.g., Fig. 6), from the receipt of raw material through the shipment of the finished cast component. There should be a production-control department that has an efficient system of processing each purchase order, from receipt of the order to final shipment, and that can retrieve daily input from the floor regarding the production status of each order. The reliability of a metalcasting facility in meeting customer requirements is best determined by direct experience and demonstrated reliability on previous orders. If a history has not been directly established, references and verified performance on other contracts should be provided. However, performance must be considered at two levels: a demonstrated record of performance across all orders as well as the capability and expertise to meet the performance, cost, and schedule requirements of a specific cast component in terms of the required alloy, complexity, molding technique, and quality control. Purchasing Tasks, Responsibilities, and Schedule. The final step in developing the purchasing plan is to lay out the specific tasks, responsibilities, and schedules that must be achieved to meet the required production and delivery times. Schedule layout is an essential element to integrating and coordinating the various tasks that must be accomplished in purchasing metal castings. One of the largest frustrations felt by a casting user is when a metalcasting facility does not meet expectations. It is important to
accurately and completely specify requirements, because inferior specifications will lead to poor performance or unnecessary costs due to overspecification (Table 1). Ideally, cast components are ordered after the metalcasting facility quotes on the job and all specification requirements are given in the purchase order (PO) for accurate pricing. If the metalcasting facility has taken exception to any of the specifications, these must be noted in the PO.
Fig. 5
Questionnaire for foundries—short form
Fig. 6
Gas testing for quality assurance
Metalcasting Technology and the Purchasing Process / 21 Table 1 Sample cost differences attributed to overspecification Overspecified procedure
Additional costs, $/lb
Melting Heat treatment for riser removal Weld repair, preheat Increased cracking susceptibility
0.05 0.08 0.04 0.04
Total additional cost
0.21
Guidelines on Specifications for Accurate Pricing. The following is a guide for cast component purchasers to navigate their way through the pitfalls of over- or underspecification and toward accurate pricing to receive the most from their cast parts. For new jobs (first-timeordered castings), the following information must be included in the RFQ and PO. Parent number and engineering drawing with the correct revision for both casting and machining drawings must be supplied. The drawings must be legible, with tolerances clearly indicated. The referenced drawing number and revision in the PO must coincide with the actual number and revision of the drawing. Quantity, price, and delivery must be clearly defined, with no misunderstandings. For example, an order for 5000 ductile iron automotive connecting rods can read, “5000 castings— 1000 pieces by October 1999, 500/month thereafter.” Castings can be supplied in various conditions, such as blasted (shot or grit), ground, machined (rough or finished), heat treated, pickled, and/or painted. They can be shipped in open or closed boxes, crates and metal tubs, or individually packaged for protection. For every metalcasting specification, there is a particular grade or class of alloy. Examples of terminology in material specifications are as follows: In low-alloy steel, the grade “low” alloy
refers to the minimum mechanical properties required. For example, ASTM A148 grade 105-85 means that the test bars must meet or exceed 105,000 psi tensile strength and 85,000 psi yield strength. Although not directly noted by the grade, there also are minimum values for ductility, as measured by the percent elongation and reduction of area from a tensile test. In ductile iron castings, the tensile yield and elongation minimum requirements are specified in the grade. For example, the ductile iron specification ASTM A 536 class 7050-05 means that the minimum requirements are 70,000 psi tensile strength, 50,000 psi yield strength, and 5% elongation. In gray iron castings, the class refers to the minimum tensile strength and the size test coupon from which properties are determined. For example, ASTM A48 class 3DB requires a 30,000 psi tensile strength determined from a 30.5 mm (1.2 in.) diameter test bar. Class A requires properties to be determined from a 22.4 mm (0.88 in.) test bar.
Aluminum cast components, similar to cast
iron, can be supplied in the as-cast or heat treated condition. The most critical aspect of ordering cast components is quality requirements. How does the casting purchaser assure that the quality requirement is met? It is difficult to assign quality requirements to the casting, especially nondestructive testing (NDT), but without sufficient knowledge of existing NDT methods and specifications, the result is either an overly conservative approach or one that is too stringent. Therefore, some knowledge of the test methods and quality levels is an asset for casting purchasers to ensure that they receive what they need. In addition to the standard requirements, the purchaser often has supplementary requirements that are not covered under the designated specification. For example, a steel casting may require good weldability, and as such, the composition may be adjusted to ensure the carbon equivalent is below 0.5. On the other hand, an application in which the casting is subjected to impact loading at colder temperatures may require that Charpy V-notch tests be performed at the metalcasting facility. A critical aluminum alloy casting may require traceability to the heat from which it was produced. This can be accomplished by identification of the casting by serial number. Stainless steel castings, which require superior corrosion properties, can have a maximum allowed carbon content. Certifications are required only when they are critical to the cast component performance. Common specifications are chemistry (usually obtained from an analysis of the heat but also from test bars). Tensile properties also are commonly specified from a test bar but less commonly from a test block of equivalent section thickness as the casting. Other specifications may include liquid penetrant, radiography, ultrasonic, and magnetic particle testing. It is sometimes necessary to record and certify heat treatment data. In more critical applications, the actual heat treatment chart may be required. Expectations for the sample casting must be clearly defined. Are they required to be liquid penetrant inspected? Does each component require an x-ray? Is 100% layout required? Is written sample approval required prior to production? Are castings subject to third-party inspection? If sample requirements are not listed, the metalcasting facility must assume that sample approval is not required prior to production. For repeat orders with no changes, it may be sufficient to reference the requirements from the original PO. For repeat orders with an engineering drawing revision, this must be noted clearly on the PO, with the new drawing attached.
Requesting and Evaluating Quotations/Bids The second step of the purchasing process is requesting and evaluating quotations/bids from
metalcasting facilities. This section describes how to locate and identify potential metalcasting facilities, request and receive quotations, and then evaluate and rank the offer for best value. Locate and Identify Suitable Metalcasting Facilities. There are many ways to locate suitable metalcasting facility suppliers: current casting suppliers, names selected from a buyer’s guide, advertising material and magazines, and recommendations from other buyers and engineering staff. In all cases, the buyer must actively seek reliable information about sources of supply and try to match metalcasting facility capabilities and capacity to purchasing needs. Replies should be studied carefully because they will provide the guidelines for future purchasing decisions. For example, the buyer will know whether the supplier can produce the grade of material required in sufficient quantity. Attitudes toward quality can be assessed from the percentage of total production and qualitycontrol personnel directly involved in the quality-control program. The buyer should recognize that the strong emphasis of a metalcasting facility on quality control indicates it is probably producing high-grade cast components. The price structure will reflect this commitment to quality. If the component is very complex, look for a supplier who specializes in this type of cast component. Discuss capacity, to avoid disappointment after the order is placed. Buyers would do well to make buying decisions on the realities of the situation: price, quality, and inspection standards required for any given set of physical and mechanical properties. Perhaps the most important point, when all other conditions have been satisfied, is to look at the current pattern of production at the metalcasting facility. Important factors can affect prices and quality of the cast part, such as whether or not the metalcasting facility is high volume, casts small components, is a jobbing shop, and so on. An initial survey of casting sources will provide a list of candidate suppliers, based on capabilities and expertise. With a source list, the next purchasing step is to request bids. Maintaining a Good Relationship. The buyer-supplier relationship does not end when the appropriate metalcaster is found. A close, mutually beneficial relationship must be established between the casting user and the metalcasting facility to produce acceptable, cost-effective cast components. The metalcasting facility must be encouraged to learn more about the casting application, including high-stress or critical sections. With this knowledge, it can suggest additional safeguards for good performance. The metalcasting facility also can suggest various quality requirements to assure that purchasers receive what they want. In some situations, a redesigned casting can have greater integrity at the critical sections. The casting purchaser (and user) also is encouraged to learn more about castings and processes, including newer technologies for producing high-quality
22 / Applications and Specification Guidelines cast parts. Cleaner casting practices allow weight savings in many types of cast components without sacrificing quality. Send Out the Bid Package and Review Quotations. The bid package consists of an RFQ form along with supporting technical drawings/information. The bid package should include contact information for the cast component buyer and the production designer as well as the closing dates for quotations. Requests for quotations may be distributed by both paper and electronic means, such as fax and e-mail. During the quotation process, metalcasting facility engineers often call the buyer and/or the designer for additional information on technical, quality, quantity, cost, and schedule and service issues. The buyer and the design engineer should be prepared to discuss these issues with the responding metalcasting facilities. When quotations are received, they should be reviewed for completeness and responsiveness. Incomplete quotations should be returned to the sender for review and resubmission. Grade and Rank the Bids/Quotations. When all the quotations have been received, they must be graded against the selection criteria defined in the purchasing plan. This may be an informal process or a detailed review and evaluation of the bids, based on the size, value, complexity, and risk involved in the component purchases. Evaluation of the supplied quotations consists of applying the selection/criteria defined in the original purchasing plan: Assess the quote against the specific quality,
quantity, schedule, technical service, and cost targets for the casting. Determine the overall reliability and value of the individual bids against the supplychain management strategy. Rank/classify the quotations in terms of the best value.
Select a Supplier and Negotiate Contract Terms After the initial bids have been received, evaluated, and ranked, the third step is selecting the qualifying and best-value bids. If the component is well defined in terms of simple shape/complexity and of relatively low risk, the quotation may be sufficient for closing the contract and moving into production. However, for the more complex, high-value cast components, additional engineering design and contract negotiation may be required to reach final agreement. Specifically, it may be necessary to: Perform a final review and refinement of the
design for manufacturability with the casting engineer Finalize technical, quality, and schedule requirements Review and resolve nontechnical contractual issues
In contract negotiations, all parties should consider the following factors:
Find a better way to phrase existing clauses Improve the agreement by adding clarifying
clauses
Level and scope of technical service, includ
ing design consultation, production, quality control, and product improvement Detailed instructions and guidance on specifications and quality-assurance requirements Assigned responsibilities for and integration of supporting engineering activities among the various production activities. These include the patternmaker, metalcasting facility, machine shop, and heat treater. Cost of quality assessment (rejects, delays, rework, etc.) Scheduling and lead-time planning Product liability
After these factors are considered, the next steps are to review, negotiate, and resolve cost details, performance clauses, and conditions of sale and then finalize and sign the contract. Figure 7 is a sample template on conditions of sale that are considered in casting purchase contracts. Terms and Conditions for Sale Agreements in the Metalcasting Industry. Over the years, there have been constant inquiries about what constitutes typical terms and conditions of sale in the metalcasting industry. There is no association-endorsed standard agreement. However, the North American Die Casting Association compiled terms and conditions of sale information from dozens of metalcasting operations around the country in 2002. Each agreement was separated into various clauses to constitute a comprehensive agreement between the seller and the buyer. If there was a significant redundancy from one agreement to the next, then the exact text for a particular passage was not repeated. One section identifies the common points for each clause that creates the agreement. In addition, the percent of agreements that contain each common point is reported. Finally, a list of atypical clauses is discussed in the section about unique or special terms. This section should be given special attention, since several adept clauses may protect the interest of a foundry. The Report on Terms and Condition of Sale Agreements in the Die Casting Industry—2002 is meant to provide the foundry with information on various terms and conditions currently being supported by the metalcasting industry. (American Foundry Society (AFS), American Metalcasting Association (AMC), Advanced Technology Institute (ATI), Defense Logistics Agency (DLA), North American Die Casting Association (NADCA), Non-Ferrous Founders Society (NFFS), and Steel Founders’ Society of America (SFSA) neither endorse nor recommend the adaptation of any of the terms or conditions documented in this report, which is available through NADCA at http://www.diecasting.org/publications/default.htm.) These terms are available electronically for cutting and pasting to help those purchasing castings to:
The principal sections include:
Alloy and tolerances Cancellation Clerical errors Credit in returns Delivery Design changes Dispute resolution Engineering disclaimer Errors in quantity Inserts Overage allowances Packaging Patents Patterns, foundry, tooling, and dies Payment and credit terms Preliminary agreement Pricing Samples Taxes Unique or special terms Warranty
Contract Fulfillment and Continuous Improvement A signed contract is not the completion of the purchasing process. The contract must be completed by two actions: on-time delivery of a fully qualified product at the agreed cost by the metalcasting facility, and prompt and timely payment by the buyer. During the life of the contract, all parties have a valuable stake in continuous improvement in quality, delivery, and best value. Specific issues that must be considered during the production cycle include: Open and timely communication (by visits,
phone calls, e-mail, etc.) between customer and supplier on project status, shortfalls, potential problems, forecasted changes, and opportunities for value improvement Monitoring of quality, delivery, cost, and performance by both the supplier and the customer Identification of opportunities to improve quality and increase value to the final customer by the customer and supplier
Considerations in Purchasing Metal Castings The metalcasting process requires designers to consider a specific list or requirements prior to purchasing that they may not otherwise be familiar with. The following is a list of the most significant considerations when attempting to determine the overall cost and design requirements of a metal casting.
Metalcasting Technology and the Purchasing Process / 23
Fig. 7
Conditions of sale—sample contract
24 / Applications and Specification Guidelines Production Cost Factors. Efficient and effective purchasing of cast components requires the use of a considerable amount of specialized information. With the increased emphasis on value-analysis techniques, purchasing agents must evaluate the many important factors that directly influence the total cost of a component part. The total cost factors are usually grouped under the general headings of quality, service, price, and delivery. Each buyer will determine the relative importance of these factors, based on particular circumstances. Since a universal formula is not possible, the following information will assist purchasing managers by covering the important criteria that affect the cost of a cast component. It is important that the casting supplier be included during the early stages of the design and purchase planning. Regardless of the type of metalcasting facility, dollar costs will break down into three main categories: materials, direct labor, and expenses. Materials may be subdivided into raw materials (primary metal and alloying constituents) and process materials (coke, sand, molding materials, bonding agents, shot, paint, etc.). Automation has markedly improved productivity in the metalcasting facility industry, but labor remains a high percentage of costs. Expenses are extremely variable, depending on the metalcasting facility facilities and the cost-accounting procedures. Although there is no magic formula that can be universally applied to metalcasting facility costs, it is possible for the buyer to develop a basic understanding of the cost factors for any metalcasting facility or particular cast component. Numerous factors influence the total cost of a finished cast component. These are considered in two groups: factors directly affecting the cost of a cast component and indirect cost factors for a finished casting. Dimensional tolerances and allowances are also discussed, since they have an important bearing on casting costs. Size, weight, and complexity of a metal casting are generally the most important cost factors, because they directly influence the materials and labor required for its production. Weight of the metal in a cast part also is important, but it no longer dominates the cost, as was true at one time. Overall size of a casting also is important, because several tons of mold material and other expendables are involved in the casting process for each ton of cast components produced. A large, thin casting that must be made in a large mold can be much more expensive than a heavier, simpler part made in a small mold. When designing a cast component, complexity is also an important factor. In general, the more complex a design, the more difficult it will be to manufacture. This may lead to higher costs. However, this is not an ironclad rule. With careful design, complexity can be molded into the metal casting without increased production cost (Fig. 8). To minimize the cost of complex designs, the engineer should call on the metalcasting facility to make manufacturing
suggestions at the early design stage. The metalcasting facility can help to minimize costs and problems with such things as draft, fillet radius, parting line locations, coreprints, wall thickness, and surface finish. The type of pattern equipment provided is an important factor in the total cost of producing a metal casting. The term pattern equipment includes coreboxes, if they are necessary, and any jigs or fixtures used for checking or assembling cores or for checking the dimensions of the part. Note that part quality is only as good as the pattern quality. Geometric complexity of a design raises costs because the mold must be much larger in order to contain external coring to form complex outer configurations. Cored castings are more costly. Cores must be made ahead of time, transported to the molding department, and placed carefully in the mold. However, equally important is the extra cleaning and trimming that is usually necessary on cored cast components. At the same time, the cores make the metalcasting process unique since they allow for the creation of internal cavities while the external surfaces of the part are produced. Wall thicknesses are an important design feature in metal castings, with the part configuration and alloy selection playing a major role in determining what wall thicknesses can be cast. As an example, in aluminum castings, the normal minimum wall thicknesses will range from 3.0 to 6.4 mm (0.12 to 0.250 in.). Depending on the part size and the area involved, thinner wall sections, down to 2.0 mm (0.080 in.), are possible over limited areas (65 cm2, or 10 in.2). Type of Metal. The cost of a casting will vary in three ways, based on the alloy used: Inherent cost of the alloy: Cost depends on
both the baseline cost of the metal along with the alloying elements in the alloy. Purity of the alloy: Premium prices must be paid by the metalcasting facilities to obtain alloys with fewer impurities. The lower the impurity level requirements of the alloy, the fewer the gating and risering returns allowed in melt charges. Castability of the alloy. In steel and aluminum alloys, various alloy compositions have differing levels of fluidity, weldability, and castability. With reduced castability, additional metal must be poured into the mold in the form of flow channels and feed heads to assure a sound casting. For such parts, the weight of the actual metal in the finished cast component may be less than half of the total weight poured. This total-metalpoured factor is especially important when considering alloys with costly alloying elements, since the alloying elements must be added to all of the metal poured into a mold. Of these three considerations, castability of the alloy has the most effect on the cost of a cast component.
Fig. 8
Comparison of (a) fabricated agricultural component and (b) redesign to a one-piece casting. Metalcasting allows production of geometrically complex shapes, consolidating multiple welded components into one casting. This casting provided a cost-savings while improving functionality.
Dimensional Allowances and Tolerances. Because dimensional allowances and tolerances affect both the cost and delivery of metal castings, it is important that part buyers have a working knowledge of the major factors influencing them. Realistic tolerance requirements are of prime importance in minimizing overall costs. Tolerances that are more rigid than necessary increase metalcasting facility production costs and lengthen delivery schedules. Those that are too loose result in castings that require more expensive machining and usually have a heavier section size than required. The following is a brief review of the more important dimensional allowances and tolerances that should be considered. Shrinkage Allowance. When metal is poured into and takes shape in the mold, the metal is still at or above its liquidus temperature. As the metal solidifies and cools to room temperature, the metal contracts. Thus, to obtain a casting of desired size, the pattern is made slightly oversized. This addition to the dimensions is called a patternmaker’s shrinkage allowance. Another type of shrinkage that occurs while the liquid metal solidifies in the mold is called solidification shrinkage. This aspect of metal solidification is corrected by feed heads, risers, or shrink-bobs, which are attached to the casting when they are poured. This is a separate phenomenon and should not be confused with pattern shrinkage allowance. Draft Allowance. This is the term describing taper on the vertical faces of a pattern that allows the pattern to be removed without binding or tearing the mold. Draft will normally vary between 1 and 3 of taper. Small or machinedrawn patterns require minimum draft. Large patterns, inside pockets on patterns, and handdrawn patterns generally require a greater draft. Under ideal conditions, some patterns as deep as 150 mm (6 in or more may be drawn with almost no draft, using special techniques. Maximum draft should be provided where it does not interfere with desired casting geometry.
Metalcasting Technology and the Purchasing Process / 25 The metalcasting facility should be consulted so that draft allowance may be minimized or even eliminated entirely on surfaces where it must be machined off. For a no-draft condition, a loose piece in the pattern or corebox can be used. This will increase the cost of the part due to an expensive pattern and higher labor costs in molding and coremaking. Alternatively, the buyer should consider the lost foam process. Cast tooling or datum points should be fixed on designated no-draft surfaces. Machining Allowances. Additional stock that is added to the surface of a metal casting so that it can be finished to size by machining is called machine-finish allowance. Thickness of this allowance will depend on size and type of the cast component, the type of surface, how it is to be machined, and how accurately the casting can be made. Large castings and those made on a low-volume basis will generally have larger machining allowances than smaller castings or components produced in large quantities. With newly designed cast parts, a full machining allowance is usually added. After some experience is gained with the pattern, machining allowances may be reduced to the minimum necessary. To avoid local surface hard spots, a ground finish can be applied directly to a cast surface, with just enough stock so that the surface will clean up well. A general guide to machining allowances is given in Table 2. Distortion Allowance. When one part of a casting cools more rapidly than another, the part may become distorted. The first line of defense is to pay attention to design details, such as avoiding sharp changes in adjacent sections. When this is not possible, distortion often can be minimized by special metalcasting facility practice to control heat extraction or by followup heat treatment. In some cases, the pattern may be intentionally distorted so that the casting will be true. This is called distortion allowance. Tolerances. Because castings are individually designed and process details are engineered for the characteristics of each pattern, casting requirements such as surface smoothness, dimensional tolerances, or machining allowances cannot be stated in a general specification. These requirements usually are established through mutual agreement between the casting user and the metalcasting facility. Ordinarily, variation in casting dimensions can be held to within half of the patternmaker’s Table 2
Machining allowance guide On bore
mm
in.
mm
Up to 305 330–610 635–1067 1092–1524 1549–2032 2057–3048 Over 3048
Up to 12 13–24 25–42 43–60 61–80 81–120 Over 120
3.2 4.8 6.3 7.9 4.8 11.1
in.
Outside surface mm
1=8 2.4 3=16 3.2 1=4 4.8 5=16 6.3 3=16 7.9 7=16 9.5 Special instructions
Use the proper shrink factor in pattern
Guide to pattern machine-finish allowance Pattern size
shrinkage allowance. On a simple 305 mm (12 in.) casting dimension, this would be a tolerance of þ2.1 mm (þ 0.0825 in.). Closer tolerances often are obtained on repetitive production where corrective adjustments can be made. For such cast parts, a tolerance of 1.3 mm in a 203 to 305 mm (þ0.05 in. in an 8 to 12 in.) dimension may be con sidered normal. Special procedures have been developed to produce some cast components with dimensions as close as 0.13 mm (0.005 in.). However, note that tolerances do not vary proportionally with the component dimensions. Dimensions within each half of the mold are generally more accurate than those across the parting plane or those related to a cored surface. The type of pattern equipment used directly affects dimensional tolerances. This is particularly true when machine molding or coremaking methods are employed. Metal patterns and coreboxes that are machined to size often will provide closer tolerances than will cast-to-size metal or wooden patterns. High-strength cast components frequently demand tight dimensional tolerances or thin-wall sections. In either, higher production costs are incurred if a systems approach is not used by the OEM to work with the metalcasting facility to recognize the true requirements of the component. Frequently, components are overspecified. In other cases, the design can be altered to take advantage of the capabilities of the many casting processes that are commercially available to meet true design requirements. The factors that most affect dimensional tolerances are:
in.
3=32 1=8 3=16 1=4 5=16 3=8
design. This is less of a problem on smaller castings. For larger castings, the metalcasting facility must have experience with the type of part required so it can make proper judgments about what shrinkage factors to use. Variability in the pattern equipment and coreboxes from nominal blueprint dimensions. To minimize this, quality tooling is required. Variability in the molds and cores themselves due to the people, equipment, and materials used in their production. Quality metalcasting facilities are aware of this and have controls and gages in place to minimize the effect. Variability in cleaning and finishing operations performed on the part. The metalcasting facility will perform inspections after this process to assure that finishing operations are done properly. Distortions due to the heat treatment of the castings. The metalcasting facility should plan to use heat treatment fixtures and special quenchants where the potential for this problem exists.
The quality of both the tooling and the production equipment is of great importance in dimensional tolerances. Even if only a small quantity of castings is required, the tooling
must be of prime quality if the dimensional tolerances are to be satisfied. Designers must be aware that there are additional costs incurred when tight tolerances are specified. Linear tolerances often must be increased for larger casting sizes. Critical dimensions should he identified to the metalcasting facility so that proper tooling will be built. Cast Surface Finish. There is no standard finish specification for cast components. Although root mean square values for cast surfaces are sometimes given on drawings, they are of questionable value because there is no dependable method of quantitative surface roughness measurement for unmachined cast surfaces. However, sample castings of the type required provide a satisfactory way of designating desired surface finish and observing what can be produced. In some casting applications, the degree of roughness of the cast surface is critical to successful use. Dimensions that are held within close tolerances may be ineffective if the surface is not smooth enough for accurate measurement or proper positioning in a fixture. While a textured surface is desirable or even necessary for some finishing or coating procedures, a surface that is too rough can be detrimental to finish quality and may require extra finishing, with increased costs. Although a smooth surface on a cast part often is considered an aspect of its quality, this is not an accurate indicator of the overall quality of the casting. Surface quality can involve several characteristics that may not be evident or even measurable. Superficial roughness on the casting can be caused by duplicating the rough surface of a pattern in metal or by an improperly made mold. However, most of the surface roughness of iron castings is the direct result of the highenergy abrasive cleaning methods in common use. Accurately evaluating a cast component surface roughness is difficult. Instruments and methods that are used successfully for measuring the roughness of machined surfaces are unsuitable for cast component surfaces. A sand cast surface is not composed of regularly spaced surface features, as found on a machined surface. The size distribution of the sand mix produces a much more complex and richer pattern of surface features. A finish that varies from location to location on a casting (due to inherent variations in sand compaction around a complex engineered geometry) further complicates the problem of measurement. The surface texture or roughness is generally specified by the term root mean square (rms) or arithmetic average (AA). The rms yields approximately 11% higher values than AA, but the difference is not significant. The surface finish may be specified as 300 rms maximum, 115 AA maximum, or C50 maximum on a visual comparator. Finer finish requirements will require special molding media or secondary operations,
26 / Applications and Specification Guidelines such as grinding or machining. This will add to the cost. Measurement should be made by surface comparator. The profilometer probe diameter acts as a filter for the input data. Fine surface areas should be identified and specified on drawings. Order Quantity. An important factor in selecting appropriate tooling is order quantity. Even with a given pattern, the order quantity may have an important influence on casting cost. One reason for this is that there are a number of standard operations in the metalcasting facility that must be performed for each part order, regardless of the quantity involved. Some of these are: Quoting, acknowledgment, and recording of
the order
Scheduling production and issuing shop
work orders
Retrieving pattern equipment from storage
and inspecting and transporting it to production departments Setting up tooling for production (i.e., laying out or attaching to machines and lining up proper flasks) If cores are required, delivering core equipment to coreroom, making cores, and transporting them to the molding line Shipping and billing Efficiency of production operations is also greatly influenced by quantity. When an order is issued subject to release, the total quantity is important in the selection of metalcasting facility tooling, but each release is normally considered as a separate order by the metalcasting facility, since each production order requires separate processing. The purchasing manager can determine the most economical lot size for cast parts when he knows what his casting requirements will be for some time in advance. Due to usual billing practices, less than a one-month supply is seldom practical. Savings realized by larger releases should be weighed against inventory costs, which include investment, space, and records. Possible changes in product demand and casting design or prices may not be equated directly in numbers but are incorporated as a matter of judgment by experienced buyers. The production quantity and the production rate are major factors in selecting the casting process and the type of mold pattern/tooling for the casting run. Various metalcasting processes have different production rate capabilities, different tooling costs, and different design change flexibility. Sand cast components often are preferred, because design changes can be made with relative ease. Two factors play important roles in tool materials selection. The first factor is the total quantity of cast parts needed. The second factor depends on the delivery period involved. Quality wood tools will generally produce up to 100 parts within approximately a year. Over a longer period of time, the tool may distort and
produce castings with dimensional discrepancies. Wood and plastic tooling should be usable for up to 500 parts in a moderate time frame (one to five years) with periodic refurbishment. Whenever small lot quantities are required for more than three years, metal or plastic tooling should be used to assure only minor refurbishment costs. Metal tooling is recommended when more than 2000 castings are required annually. Additional Operations. Often, the metalcasting facility can perform additional production operations that may increase the direct cost of the casting but will reduce the total cost of the finished product. Heat treating, painting or coating, and inspections that are more sophisticated and testing procedures are typical of additional operations that may be efficiently performed by the metalcasting facility. For quantity production runs, castings may be justified or targeted in a gaging fixture to provide accurate locating points for subsequent machining operations. This assures that the correct amount of stock is available to be removed on each critical finish surface. Technical Services. Close cooperation through concurrent engineering between the metalcasting facility and its customer is the approach required to yield the greatest benefit for buyer and supplier. However, for the more complex engineered cast components, an additional amount of service or technical assistance may be required from the metalcasting facility. This commonly occurs in development of new products. Assistance with component design, information on properties and metallurgy, troubleshooting, pattern assistance, or pattern models and experimental castings may be provided by the metalcasting facility. These services may be billed to the customer at cost or included in the price of the parts. However, in either case, quotations from a metalcasting facility that stands ready to provide its customers with this assistance cannot be compared directly with quotations from a metalcasting facility that does not. The latter type of metalcasting facility should not be ignored, because it can be a good, economical source for castings. This is particularly true where special problems or requirements do not exist and where additional services are not necessary. When a complex casting will be produced in large quantities, a model casting may be made to assist in planning production equipment. The model will assist in establishing the most efficient metalcasting method by providing an opportunity to establish critical factors such as the parting line and core prints on the actual shape. This also can be of assistance in planning for subsequent processes, such as machining. Types of Price Quotations. At one time, metal was expensive compared with labor, and castings were commonly sold by the pound. This was a simple buying method and worked well because the weight of a cast part was a more important part of total cost than was the
workman’s time to make it. This is no longer true. With the tremendous increase in cost of labor and mechanized equipment in a modern metalcasting facility, the value of the metal in most castings has become a secondary factor. Accurate cost accounting must include cost of the worker’s time, value of materials used, and overhead costs. This provides price quotations that are based on cost per casting and a ready reference for the purchaser to arrive at unit cost. The cost-per-piece can easily be converted to cost-per-pound, but such a figure cannot be justly compared to other castings that may be ordered in different quantities or require more or less man-hours to be produced. Some metalcasting facilities, however, still retain the price-per-pound method of quoting for specific types of parts. Thus, equivalent quotations based on price-per-pound and price-per-piece may show variations when compared for individual castings. This is particularly true when the quotation is based on an estimated casting weight. Variations of actual from estimated weight still affect the cost more on a price-per-pound basis than on a price-perpiece basis. Lead Time and First-Article Delivery. Essentially, the metalcasting process involves pouring molten metal into a cavity close to the final dimensions of a desired component. It is the most direct and simplest forming method available for metals and is often an essential production factor in rapidly manufacturing finished metal parts. Lead time is commonly defined as the amount of time between contract agreement and delivery of the first article, either prototype or initial production. Lead time is reduced with metalcasting compared to other manufacturing methods by eliminating and/or shortening the time required for tool production, parts ordering and delivery, assembly, finishing, and machining. In the metalcasting arena, new rapid prototyping (RP) tools (Fig. 9) are being developed to produce patterns and tools for castings. The RP tools can dramatically reduce lead times and production costs required for patterns and tools, particularly for designs that have complex detail that would be difficult to produce by conventional machining. An example of an investment casting produced with rapid tooling is shown in Fig. 9. Quality Assurance and Specifications. Quality assurance has benefits for both the buyer and the seller of metal castings. The buyer is assured of a cast component that meets performance requirements and the designated specifications. The metalcasting facility has confidence in its production process, the documentation of the process methods, and the application of quality standards in its day-to-day operations. All parties benefit because a high-quality product is delivered on schedule without cost overruns. Standard specifications for various types and grades of alloys and castings are promulgated by several national organizations and
Metalcasting Technology and the Purchasing Process / 27
Fig. 10
Fig. 9
Example of rapid prototyping process with investment casting. (a) Valve body prototype prepared for investment casting manufacturing without hard tooling. (b) Soft tooling for production investment casting mold. (c) Final casting. With rapid tooling, metal casting has developed to the point of producing castings in days, with no tooling.
government agencies. The most commonly used specifications for alloys and metal castings in the United States and Canada are those of the ASTM International, which is composed of many volunteer committees that write and maintain the standard specifications. Each committee is assigned jurisdiction over standards for one type of material, application, or testing methodology. The Society of Automotive Engineers (SAE) also publishes commonly used specifications for cast components. Other organizations, such as the American Society of Mechanical Engineers and the American Petroleum Institute, have specifications concerning castings for specific applications, but the material portion of these specifications typically refers to an appropriate ASTM International specification. The engineering departments of some companies that design and use a large number of castings write their own specifications. Some casting applications require an unusual combination of
Indentation profiles of three types of hardness testing methods
properties that would not be provided by cast parts made to standard specifications. Types of Specifications. There are several methods of specifying cast components. Specifications can be written to designate the properties of the metal that are specifically needed (strength, hardness, finish, accuracy, or special characteristics, such as resistance to heat or corrosion). In addition, there are many casting applications in which the ordinary properties of the designated type of alloy are quite adequate. In purchasing metal castings for noncritical applications, there may be an advantage in economy and availability to ignore standard metal specifications. The purchase order simply could designate the generic type of alloy by name (gray iron or malleable iron). Other requirements, such as for dimensional tolerances, may be added. When the mechanical properties of the metal are important, there are two commonly used methods to designate this requirement. The more frequently used procedure for engineering metal castings uses a standard test bar that is poured along with the specified lot of parts. The required minimum tensile properties should be obtained in testing a sample from this test bar. Most ASTM International specifications apply this method to qualify the metal used to pour the casting. The actual properties of the metal in the part will depend on its characteristics and the metal cooling rate in the various sections of the casting. The relation between properties in the casting and the test bar will depend on their relative cooling rates. For example, the properties of gray iron are especially sensitive to the section thickness in which they are cast, so a series of standard test bars of increasing size have been established specifically for gray iron and related to various casting thicknesses. One of these standard test bars, or, if necessary, a larger- or smaller-sized one, should be selected to approximate the cooling rate in the critical sections of the cast component. Malleable iron is not poured into heavy sections, and since all malleable castings are heat treated, a single-sized test bar is used. Ductile iron has some variation in properties with metal section when not heat treated.
However, a single-sized test bar is satisfactory, except for very large castings. In another approach, a commonly used method of specifying iron for castings produced in larger quantities is based on the microstructure and Brinell hardness of the metal in the part itself at a designated location. This method was originally developed for automotive and agricultural implement castings but has proven satisfactory for many other types of iron castings produced in quantity. When the microstructure of the iron is normal, there is a good general relationship between Brinell hardness and tensile properties; however, this relationship is not part of the specification, nor are the tensile properties of the iron stipulated. The suitability of a casting for its intended use as produced in a particular type of iron at a specified hardness level must be established for each application. This is typically done by trial use or by simulated load tests. Some castings purchased under SAE specifications for critical applications are tested for a specified Brinell hardness (Fig. 10). Standard Specifications. Many casting purchases are referenced to standard specifications, primarily to those of ASTM International and SAE. In addition, there are also cast parts purchased without metal specifications other than simply stating the generic type of alloy wanted. This purchasing procedure is satisfactory when the requirements of the casting are quite ordinary and typical for that type of alloy. Some metalcasting facilities will not accept orders that require metal property tests but specialize in providing castings with other specified requirements. Where metal properties are to be specified, the use of a standard specification provides the buyer with a maximum number of sources and possible economic advantages. The use of a familiar specification saves time and reduces the possibility of mistakes. Even when purchasing cast parts with unusual requirements, it is preferable to use an appropriate standard specification as a base and then add additional requirements. Of course, the additional requirements should be clearly stated and given with the reference number of the standard specification on the purchase order.
28 / Applications and Specification Guidelines Quality-Management Systems. A buyer may ask a caster, “How good is your quality, and how do you control it?” The caster has two possible replies: (1) The caster can provide a detailed description of the quality-management system and wait for the buyer to review and qualify the methods and procedures, or (2) the reply can be, “My quality system has been reviewed and certified against a recognized quality-management standard. Here is the certification.” A certified quality-management program has two basic features:
Definition and documentation of the policies
and procedures required that meet the specified quality standards Application and practice of the quality procedures on a day-to-day basis, supported by traceable records There are several quality-management standards and certifications that are widely accepted in the manufacturing community and are applicable to the metalcasting market: ISO 9000, QS 9000, AS 9000, and NQS 9000.
SELECTED REFERENCES Cost-Effective Casting Design: A Designer’s
and Purchaser’s Reference, American Foundry Society, Inc. Engineered Casting Solutions, American Foundry Society, Inc., www.castsolutions. com Principles of Purchasing Castings, American Foundry Society, Inc. Website for casting buyers and designers: www.metalcastingdesign.com
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 31-40 DOI: 10.1361/asmhba0005189
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Chemical Thermodynamics and Kinetics Revised by Seymour Katz
TWO SUBJECTS that underlie all foundry processes are chemical thermodynamics and chemical kinetics. Relationships derived from these two disciplines of physical chemistry are extremely useful in providing a quantitative framework for explaining and predicting chemical behavior. This includes such presolidification chemical processes as preparing liquid metal for casting (going back to the winning of ores, if desired), melt purification from scrap recycling, and alloy additions. Chemical processes are also involved in making molds and cores, the casting and solidification processes, the microstructural development of alloys, and the phase changes that occur after casting (for example, by heat treatment). This article introduces the fundamental concepts of chemical thermodynamics and chemical kinetics in describing presolidification phenomena. Thermodynamics defines the most stable chemical system (equilibrium system) that can be produced, given a set of starting conditions (compositions, temperature, pressure, and so on). An example is the phase diagram. It gives the stable phases that coexist at equilibrium, given the composition, temperature, and pressure. Chemical thermodynamics is only concerned with the initial and final states of a system, while chemical kinetics is concerned with how and at what rate the system is transformed from the initial to the final state. The practical importance of chemical kinetics implies that systems often fail to reach the most stable condition and that the path taken plays a major role in determining the speed of the reaction. In general, kinetics assumes its greatest importance when atomic movement is restricted, such as at low temperatures or in solids. Thermodynamics most often applies when atoms have high mobility, as in gases and liquids and at high temperatures. Another factor that determines the relative applicability of thermodynamics and kinetics is time. The less time available for a process, the more likely that kinetics will govern the chemical state of the system. The thermodynamics and kinetics are fundamental in the composition control of liquid metal. Composition control consists of two areas: alloy addition and melt purification. Both are of vital importance in producing quality
castings. The chemical and physical processes involved in alloy addition do not vary much with base metal; therefore, this subject can be treated generically in this article. On the other hand, melt purification processes tend to be material-specific and vary according to base metal. For example, metal purification may involve processing of scrap, which is a common source material used for iron and aluminum castings. By its nature, scrap is of uncertain composition, both with respect to concentrations of desirable alloy elements and the presence of deleterious elements. Thus, despite all precautions in melting, adjustments in melt composition may be required. Following the introduction to chemical thermodynamics and kinetics in this article, the next two articles provide pertinent thermochemical data for ferrous and nonferrous (aluminum and copper) casting alloys. Casting defects with origins in the liquid state (gas defects and inclusions) are covered in the articles “Gases in Metals” and “Inclusion-Forming Reactions.” Nomenclature of symbols used in physical chemistry is given in Table 1.
mass. Common units of H are joules/mole, calories/gram, or Btu/pound. Thus, the total enthalpy of ni moles of material i is: ðHi Þtotal ¼ ni Hi
(Eq 1)
For a mixture of materials, the total enthalpy is the sum of the individual enthalpies: Htotal ¼
X
ni Hi
(Eq 2)
i
The heat capacity, CP, given in joules/mole K, and so on is a measure of how H varies with temperature and is defined as: CP ¼
@H @T P
(Eq 3)
where the subscript P indicates that the values refer to constant pressure conditions. Enthalpy and heat capacity govern the thermal state of a system and the heat generated or absorbed by chemical reactions. These variables are particularly important for establishing the energy requirements for processes, for determining rates of solidification, and for heat treatment processes. Enthalpy and heat capacity data are tabulated in Ref 1.
Chemical Thermodynamics
Gibbs Free Energy
There are three thermodynamic laws that provide the basis for the relationships that describe chemical systems. These laws hold that perpetual motion and absolute zero temperature can never be achieved. The relationship between the thermodynamic laws and, for example, phase diagrams is not immediately apparent. However, the connection exists and is testimony to a great deal of brilliant theoretical work carried out over the last 100 years. For metallurgical systems, the most important thermodynamic variables are enthalpy and Gibbs free energy. The former governs the disposition of heat, and the latter governs chemical equilibrium.
Conditions for Equilibrium. Chemical equilibrium is a condition in which the chemical components of a system have no tendency to change. In mechanical systems, equilibrium is understood as the condition in which the potential energy is at a minimum. For chemical systems at constant temperature and pressure, an analogous statement for chemical equilibrium is that the Gibbs free energy is at a minimum, or equivalently from calculus:
Enthalpy and Heat Capacity The heat content of a material is termed its enthalpy, H. Because the value of the enthalpy is proportional to the amount of material present, H is usually tabulated as enthalpy per unit
dGðP; T; ni . . .Þ ¼ 0
(Eq 4)
For systems in which chemical change can occur, dG can be expressed in partial differential form:
@G @G dT þ @T P;ni ... @P T;nj @G dP þ dni þ @ni T;P;nj
dG ¼
(Eq 5)
where ni stands for moles of component i. The term ð@G=@ni ÞT;P;nj ; is called the partial
32 / Principles of Liquid Metal Processing Table 1 Physical chemistry nomenclature ai aoi A A0 Ci Cp Cs CE dp D DR eii E fi F* F Gi DG Gi Goi xs G Goi Goi ð%iÞ Goi ðXi Þ h hi Hi Hio H i
Activity of component i relative to pure material standard state Activity of component i in the pure material standard state Interfacial surface area Surface area of an alloy particle Concentration of component i Heat capacity Sulfide capacity Carbon equivalent Particle diameter Diffusion coefficient Desulfurization ratio Interaction coefficient for 1 wt% standard state Cell potential Activity coefficient for 1 wt% standard state Integral free energy of solution Faraday’s constant Gibbs free energy of component i Integral molar free energy Partial molar free energy of component i Molar free energy of component i Excess partial molar free energy of component i Free energy of component i in solution at the 1 wt% standard state Free energy of solution from pure material to 1 wt% standard state Free energy of solution from the pure material to the hypothetical pure material standard state Convective heat transfer coefficient Activity of component i relative to the 1 wt% standard state Molar enthalpy of component i Standard enthalpy of formation Partial molar enthalpy of component i
molar free energy, Gi , or the chemical potential. It is clear from Eq 4 and 5 that, at equilibrium and at constant temperature and pressure, the following relationship exists: 0 ¼ Gi dni þ Gj dnj þ . . .
(Eq 6)
When the change in composition is dictated by a chemical reaction, stoichiometry provides a relationship between dni, dnj, and so on, and Eq 6 can suitably be transformed to: 0 ¼ ni Gi þ nj Gj þ . . .
(Eq 7)
where ni, nj, and so on are the stoichiometric coefficients in the chemical reaction equation. For products and reactants, the stoichiometric coefficients are positive and negative numbers, respectively. As an example, for the chemical reaction: 4Al þ 3O2 ¼ 2Al2 O3
(Eq 8)
Equation 7 becomes: 0 ¼ 2GAl2 O3 4GAl 3GO2
(Eq 9)
Tabulation of Thermodynamic Data. To make thermodynamic calculations convenient, a system was developed for tabulating thermochemical data so that relatively small amounts of data could describe a wide range of chemical conditions. This was done by defining special
xs
H i J k, k0 , kB KA kH km KmN K m,m0 Mi ni Ni pi poi P P S, P M
Q Q r, r0 R S Sc Sr S i xs S i Sio
Excess partial molar enthalpy of component i Mass flux Rate constants Constant related to the absorption of an element on a surface Constant related to hydrogen solubility in iron Mass transfer coefficient Liquid-phase mass transfer coefficient Equilibrium constant Solubility factors giving change in carbon solubility in Fe-C-X from unit mass addition of a third element Molecular weight of component i Moles of component i Mole fraction of component i Pressure of component i in the gas phase Pressure of component i in its standard state Constant pressure conditions Partition coefficient. Ratio alloy element concentration between liquid and solid phases in equilibrium. The superscripts S and M refer to stable and metastable products of Fe-C solidification, respectively. Heat flux Gas flow rate Radius of a rod Universal gas constant Ratio of surface area to volume Ratio carbon concentration: iron alloy melt to eutectic Rectified saturation degree, that is, weight fraction Fe-C eutectic formed on solidification of hypoeutectic iron Partial molar entropy of component i Excess partial molar entropy of component i Molar entropy of component i in the pure material standard state
conditions called standard states. Thus, for a chemical compound i that is an ideal gas at pressure pi,Gi is now written: Gi ¼ Goi þ RT ln
pi poi
(Eq 10)
where Goi is the free energy per mole of i in its standard state and poi is the pressure of i in its standard state. The second term on the right-hand side of Eq 10 is the free energy change to take a mole of i from its standard state poi to pi. For gases, the usual standard state is poi ¼ 1 atm. Thus, poi is usually omitted in writing, Eq 10. Both Gi and Goi are absolute values that cannot be measured. Changes in these variables can be measured from an arbitrary baseline. By convention, the baseline is taken as the elements in their standard states for which ðGoi Þelements ¼ 0 is defined at all temperatures. To indicate that the values are relative to a baseline, Eq 6, 7, and 10 are rewritten by replacing Gi and Goi with Gi and Goi , respectively. Values of Goi for many compounds (called the standard free energy of formation) and phase changes have been tabulated (Ref 1, 2). Because Goi is a function of temperature, the most common tabulations are based on another thermodynamic definition of Goi , that is: Goi ¼ Hio T Sio
(Eq 11)
t t tm T V V0 V0 w W Xi
Time Celsius temperature Melting time Kelvin temperature Gas flow rate Gas volume per tonne of metal Volume of an alloy particle Mass fraction: desulfurization to iron Weight of melt Mole fraction of component i
Greek symbols a Phase designation aij Alpha activity coefficient function for Darken’s quadratic formalism gi Activity coefficient of component i relative to pure material standard state io Henry’s law constant d Boundary layer thickness Interaction coefficient; pure material standard eji state y Fraction of surface sites covered by surface-active atoms k Empirical constant indicating iron compositions that have the same microstructure at a given cooling rate l Particle shape factor ni Stoichiometric coefficient for component i in a reaction r Density Superscripts B e, equil f ‘ s S *
Bulk property Equilibrium condition Final condition Liquidus concentration Solidus concentration Surface condition Interface condition
where Hio and Sio are the standard enthalpy and the entropy of formation, respectively. The tabulations provide Hio and Sio data. The value of Goi is calculated for any temperature using Eq 11. Equilibrium Constant. Substituting Eq 10 into Eq 7 gives the most important relationship governing reaction equilibrium: X
ni Goi ¼ RT
i
X
ln pni i
(Eq 12)
i
Here, the thermodynamic data on the left are related to actual pressures (concentrations) on the right. Equation 12 is ordinarily written in abbreviated form as: Go ¼ RT ln K
(Eq 13)
where K is the equilibrium constant. The form of K is illustrated in the following. Given the gas phase chemical reaction: 2H2 þ O2 ¼ 2H2 O
(Eq 14)
the equilibrium constant is: K14 ¼
p2H2 O p2H2 pO2
(Eq 15)
Activity and Condensed Phase Equilibrium. Metallurgical applications are primarily concerned with condensed phases (solids and liquids), which do not behave in the manner of an ideal gas. That is, for a condensed
Chemical Thermodynamics and Kinetics / 33 material, Goi does not vary with the logarithm of the concentration as may be inferred from Eq 12, where, for an ideal gas, Goi varies directly with ln pi. To preserve the useful equilibrium constant expression for nonideal, condensed phases, Eq 10 and 12 are retained in form but rewritten to enable the use of data expressing nonideality. Thus: ! ai aoj
(Eq 16)
ln ani i
(Eq 17)
Gi ¼ Goi þ RT ln X
ni Goi ¼ RT
i
X i
The difference between Eq 16 and 10 and Eq 17 and 12 is that a term ai, called activity, replaces the pressure term. The activity is a function of concentration but usually not a simple function of concentration. Analogous to the case with gases, the activity in the standard state aoi is taken as unity. The equilibrium constant is now defined in terms of activity: ln K ¼
X
ln ani i
(Eq 18)
i
Activity Coefficients—Departure from Ideality. Common expressions for activity are: ai ¼ gi Xi
(Eq 19)
hi ¼ fi ð%iÞ
(Eq 20)
where the concentration terms Xi and (%i) are the mole fraction and weight percent, respectively, of the solute i, and gi and fi are the activity coefficients that measure departure from ideality, that is, generally gi and fi 6¼ 1. The terms ai and hi are often called the Raoultian activity and the Henrian activity, respectively, to distinguish the standard state being used. The Raoultian activity is usually used in conjunction with the pure material standard state, while the Henrian activity is used with the 1 wt% standard state (to be explained). The latter standard state and Eq 20 are used to describe dilute solutions and have been extensively applied to ferrous systems, especially steel. 1.0 A
T = 1873 K
Activity, ai
0.6
aNi aFe
0.4
γ⬚
Ni
0.2 B 0 0 (Fe)
Activity coefficient, γi
Raoult’s law
0.8
Henry’s law 0.2
0.4
0.6
0.8
1.0 (Ni)
Mole fraction, XNi
Fig. 1
Activities of iron and nickel in their binary solutions at 1873 K. Source: Ref 1, 3
The pure material standard state and Eq 20 are used in most other cases. The two standard states, their activities, and their activity coefficients are illustrated with Fig. 1, which gives the activity composition relationships for the liquid iron-nickel system. The point “A” represents the pure material standard state for nickel (pure nickel in its most stable form—liquid—at the temperature of concern, 1873 K). Activity, given on the ordinate, is based on Eq 19. The ideal solution line (Raoult’s law line), where aNi = XNi, is labeled. The change in the partial molar free energy of nickel in transferring nickel to solution is given by Eq 21, which was obtained by substituting Eq 19 into Eq 16: GNi ¼ GoNi þ RT ln XNi þ RT ln gNi
(Eq 21)
In this case, Goi ¼ 0 because the standard state and the baseline conditions are identical. The term RT ln XNi is the change in GNi ideal for an ideal solution GNi . The term RT ln gNi provides the correction for the real system. As seen, gNi < 1. Thus, GNi is more negative ideal than GNi . Larger negative free energy values indicate greater stability. Thus, real solutions of nickel in iron are more stable than ideal solutions of nickel in iron. A common characteristic of the activity of a component in solution is the relatively linear portion of the curve as Xi ! 0. This limiting slope, designated goi (Fig. 1), is known as the Henry’s law constant. The 1 wt% solution standard state is marked point “B” in Fig. 1. It falls within or near the range in which the activity coefficient is a constant ðgoi Þ. By defining the standard state at point “B”, fi = 1 in the linear region and conveniently, hi = (%i). The expression that corresponds to Eq 21 for the 1 wt% standard state is: ^o þ RT lnð%NiÞ GNi ¼G Ni þ RT lnðfNi Þ
(Eq 22)
^o to indiIn Eq 22, GoNi is designated G Ni cate that it refers to the 1 wt% standard state. Clearly, for any iron-nickel composition, GNi is the same whether it is calculated by Eq 21 or 22. Therefore, the choice is one of ^o 6¼ 0. It is the convenience. In this case, G i change in free energy in taking pure liquid nickel to point “B” on the Henry’s law line. The difference in Goi between standard states A and B is: ^o ¼ Go þ RT ln aB G i i aA Ms loi ¼ Goi þ RT ln Ms þ 99Mi
(Eq 23)
where the activities for composition A and B are Raoultian, and Ms and Mi are the molecular weights, respectively, of solvent and solute. The equilibrium constant expressions (Eq 17) for the pure material and 1 wt% standard states are: X i
ni Goi ¼ RT
X
lnðgi Xi Þni
(Eq 24)
X
^o ¼ RT ni G i
X
ln½fl ð%iÞni
(Eq 25)
i
o
Very often Gi for components in the equilibrium constant are expressed in different standard states. There is no problem with this as long as appropriate values Goi or ^o are used. G i Tabulations of Activity Coefficient Data. Tabulations of Goi data were already discussed. From Eq 23 to 25, it is seen that data for goi , gi, and fi are also needed. Tabulations of goi and ^o for elements in solution in liquid iron, G i cobalt, and copper are given in Ref 2. Similar data for aluminum are included in the article “Thermodynamic Properties of AluminumBase and Copper-Base Alloys” in this Section. Data for gi and fi for iron, copper, and aluminum multicomponent alloys are available in Ref 4 to 6. These data are based on a formalism (Ref 7) that is discussed in the article “Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys” in this Section. Another formalism (Ref 8, 9), found in the article “Thermodynamic Properties of Iron-Base Alloys” in this Section, is used to describe the activities of carbon and silicon in multicomponent ferrous alloys. Most of the tabulated values were obtained from studies in dilute solution (Henry’s law region). From Fig. 1, it can be seen that the shapes of the activity curves are complex. Therefore, data for the Henry’s law region have only limited applicability outside the region. To describe Gi over the entire composition range, tables giving Gi as a function of composition are used. These data are usually available only for binary alloy systems. Tables for important aluminum and copper binary systems are given in the article “Thermodynamic Properties of Aluminum-Base and CopperBase Alloys” in this Section. Comparable data for iron are available in Ref 10. Because Gi is temperature dependent, the data are commonly given as H i and S i , where Gi is obtained by: Gi ¼ H i T S i
(Eq 26)
Examples of the use of Eq 1 to 26 are found in other articles in this Section.
Phase Diagrams Phase diagrams describe the phases that coexist at equilibrium. Because chemical equilibrium is involved, phase diagrams have a thermodynamic basis. The complexity and variety of relationships demonstrated by phase diagrams place a detailed description beyond the scope of this article. A qualitative demonstration of the interrelationship between phase diagrams and thermodynamics is given. For a multiphase system, a condition of equilibrium is that the partial molar free energy or chemical potential of each component must be the same in all phases (a, b, . . .):
34 / Principles of Liquid Metal Processing a
b
Gi ¼ Gi
(Eq 27)
Thus, for example, the chemical potential of carbon in carbon-saturated iron is the same as graphite, despite there being less than 5 wt% C in saturated iron. A similar situation exists with respect to silicon/aluminum-silicon alloy equilibrium, where the chemical potential of pure 0
silicon is matched by less than 15 wt% Si in solution. Phase diagrams describe stability relationships for compounds and solutions (solid and liquid). To understand the relationships based on thermodynamics, it is necessary to compare the phase diagram (temperature-composition plot) with the equivalent free energy-composition plot. To construct the free energy-composition plot, it must be recognized that for a binary solution the integral molar free energy is:
Free energy (F *), cal
G ¼ X1 G1 þ X2 G2 −500
−1000
F ¼ G ðX1 G1 þ X2 G2 Þ
(ΔG1 – ΔG1⬚)
−1500
−2000
(Eq 28)
Substituting Eq 16 into Eq 28 and assuming ideal solutions (ai = Xi) yields F*, the free energy change in forming 1 mole of an ideal, binary solution from components in the pure standard state (taken as the pure liquid in this case): ¼ RT ðX1 ln X1 þ X2 ln X2 Þ (ΔG2 – ΔG2⬚) 0
0.75
0.50
0.25
1.0
Mole fraction (X2)
Fig. 2
Free energy of formation of an ideal solution from liquid components in the pure material standard state. Also illustrated is the graphic method of tangents to determine the partial molar free energies of solution.
(Eq 29)
A plot of F* as a function of X2 is given in Fig. 2. Clearly, F* < 0 with a minimum at X2 = 0.5. For 6 1), the curve can be higher nonideal solutions (gi ¼ or lower than the ideal line, and it can be asymmetrical. It can be shown that the intercepts of a tangent line drawn at any point on the curve will give ðG1 G1 Þ and ðG2 G2 Þ (Ref 1). These are the partial molar free energies of solution from the liquid standard state.
0
For metallic solutions, Gi ¼ 0 ; therefore, GI is obtained directly from this plot. The free energy of a phase of definite composition is represented as a point on the free energy-composition diagram. Clearly, if such a phase were in equilibrium with a solution, the equilibrium solution and the compound would have to lie along the same tangent line in order for Eq 27 to hold. This condition is illustrated in Fig. 3(a). Figure 3(b) shows the phase diagram that results from the thermodynamic constraints. In Fig. 3(a), F* < 0 for pure solid phase 1 because the liquid was chosen as the standard state, and at the temperature chosen, the liquid is less stable than the solid. The equilibrium between a liquid and a primary solid solution is illustrated in Fig. 4. From Fig. 4(a), the respective equilibrium concentrations lie along a common tangent, thus satisfying the equilibrium conditions expressed by Eq 27. The phase diagram that results from these thermodynamic conditions is given in Fig. 4(b). For small values of X2, either solution can theoretically exist. The solid phase is expected to be present, however, because it is more stable (lower F*). If, by virtue of such factors as rapid cooling, the liquid phase is retained in a region where the solid is more stable, a metastable condition would exist. This is not uncommon in systems of metallurgical interest. Common examples are cementite (Fe3C), which is metastable with respect to iron and graphite, and the glass (supercooled liquid) condition of slags, which is metastable with respect to crystalline compounds.
Solid Solution
B
Composition of solution in equilibrium with pure solid
Solid + solution
α solid solution
Free energy (F *), cal
Free energy (F *), cal
0 A
Solution
Liquid solution
(a)
Temperature (T )
(a)
Temperature (T )
This line corresponds to temperature of free energy chart above
T1 L
α+L
β+L
α
0 0
(b)
Fig. 3
1.0 Mole fraction (X2) Relationship between (a) free energy-composition and (b) temperature-composition diagrams for the equilibrium between a solution and a pure solid component. Source: Ref 11
β
α+β
(b)
Fig. 4
1 Mole fraction (X2)
Relationship between (a) free energycomposition and (b) temperaturecomposition diagrams for the equilibrium between a liquid and a primary solid solution. Source: Ref 12
Chemical Thermodynamics and Kinetics / 35
Chemical Kinetics The rates of chemical processes depend so much on local conditions that it is difficult to describe the processes in the universal terms used in thermodynamics. The rates of chemical change in solid cast metals are usually governed by solid diffusion. This subject is examined in the Section “Principles of Solidification” in this Volume. For processes involving fluids at high temperatures, the rates of chemical reactions are inherently rapid and usually not rate limiting. The processes that generally limit the rate are nucleation of the product phase and interphase mass transport (moving material between immiscible phases).
Nucleation The problem of nucleation is universal to the kinetics of solidification, and the subject is treated more extensively in the article “Nucleation Kinetics” in this Volume. However, a few general comments are in order. Nucleation of a phase in an inclusion-free metal requires supersaturation with respect to the material being precipitated. For this condition, the chemical potential of the material in the melt exceeds that in the precipitate, which is the requirement for spontaneous reaction. Initially, the diameter of the precipitated phase is in the atomic range. Thermodynamics dictates that the activity of particles increases with decreasing radius of curvature. Therefore, the activity of atomic-sized particles is very high. This in turn requires very high melt supersaturation to meet the requirement for spontaneous nucleation. This barrier to the generation of nuclei from a homogeneous melt, imposed by thermodynamics, makes it unlikely that appreciable nucleation occurs by this mechanism. Instead, nucleation is initiated on particles purposely introduced to the melt before solidification. For example, TiB2 particles are in use to nucleate aluminum melts (Ref 13), and complex oxysulfide particles nucleate graphite in gray and ductile iron (Ref 14, 15).
Interphase Mass Transport Because high-temperature reactions are rapid, the rate-limiting step in many cases of heterogeneous reaction is the diffusion of reactants to and from the reaction interface. This is particularly the case when the reactants are dilute. Because the reaction rate in these cases is rapid relative to rates of diffusion, the concentration of a reacting species at the reaction interface is either above or below the bulk concentration. The condition that exists depends on whether the activity of the reactive species in the second phase is higher or lower than the activity in the bulk. The model, frequently used in quantifying mass transfer limited kinetics, assumes that there is a thin stagnant film at the interface
across which a concentration gradient exists. The model is schematically represented in Fig. 5. Transport across the boundary layer is assumed to occur only by diffusion. Applying Fick’s first law of diffusion, the flux of species i across the boundary layer is: J¼
D C CiB ¼ km Ci CiB d i
(Eq 30)
where D is the diffusion coefficient, d is the boundary layer thickness, and Ci and CiB are concentrations of the diffusion species, respectively, at the interface and in the bulk. The term D/d is seldom used; instead, the mass transfer coefficient km is used to represent D/d. When phase 2 is a fluid, d decreases with increasing fluid velocity along the interface (that is, increased convection). Thus, km is a function of the fluid dynamic condition. To determine the rate of reaction, that is, J, it is necessary to know Ci . This is difficult to determine experimentally. To circumvent the problem, one of two assumptions is commonly made. For extremely fast reactions in which the equilibrium constant is very large, Ci can be assumed to be 0. Thus, Eq 30 reduces to: J ¼ km CiB
J ¼ km Ciequil CiB
(Eq 32)
The value of Ciequil is determined from the equilibrium constant expression, knowing the activities of all the reactants and products except Ciequil . A number of examples of this treatment are found in succeeding articles in this Section. There are cases in which the reaction rate (rate constant k) at the interface, unimpeded by diffusion, is of the same order as the mass transfer limited rate. In this case, the observed rate and the rate constant k0 are affected by both processes. This relationship is illustrated as follows. Because diffusion of i through the boundary layer and the reaction of i at the interface occur sequentially, the rates of both processes must be the same at steady state. Thus: J ¼ km Ci CiB ¼ k Ci Ciequil
1 ¼ k0
1 1 þ km k
δ
Stagnant boundary layer
Fig. 5
Boundary layer model used for mass transfer limited kinetics
It is clear from Eq 34 and 35 that for the condition k km, Eq 34 reduces to Eq 32. Applications in which mixed control is important are nitrogen absorption in iron alloy melts (Ref 17) and coke gasification at intermediate temperature (Ref 16).
Kinetics of Alloy Additions As previously noted, the chemical and physical processes involved in alloy addition do not vary much with base metal and can be treated generically. There are two fundamental mechanisms governing the kinetics of alloy additions, depending on whether the alloy is liquid or solid at the temperature of the metal bath. These processes are schematically illustrated in Fig. 6.
Dissolution of Alloy with Melting Point Lower than Bath Temperature
(Eq 33)
where the last term gives the first-order rate of reaction occurring at the interface. Eliminating Ci in the last two terms in Eq 33 yields: J ¼ k0 Ciequil CiB
C Bi
Phase 1
(Eq 31)
This assumption has been used to describe the high-temperature gasification of coke (Ref 16). The second assumption is that equilibrium between the two phases is established at the reaction interface. In this case, Eq 30 is written as:
where:
Phase 2 C *i
(Eq 34)
(Eq 35)
For the case where the alloy addition is liquid at the melt temperature (Fig. 6, top row sequence), addition of the cold alloy at time t0 is quickly followed by formation of a layer of frozen base metal on the alloy surface (t1) due to heat extraction by the alloy. At longer times (t2), the shell thickness increases, while melting occurs at the alloy surface. As the interior of the alloy heats up, the thermal demand decreases. As a result, the frozen base metal shell begins to redissolve while the thickness of the liquid
36 / Principles of Liquid Metal Processing Superheat, ⬚F
Time t1
t0
t2
t3
t4
tm
840
Alloy liquid at bath temperature
32
48
64
80
96
112
60
70
720
Melting time, s
600
Alloy solid at bath temperature
480
360 C 240
Solid alloy Liquid alloy
B 120
Solid base metal
Fig. 6
A
Two kinetic paths for melting and/or dissolving alloy additions in liquid metal baths. Source: Ref 18
0
20
30
40
50
Superheat, ⬚C
Table 2 Data for melting time calculations for a 100 mm (4 in.) diameter sphere of silicomanganese in a steel melt Condition
Fig.
DH, J/g
r0, g/cm3
V0/A0, cm
h, J/g K
TL , K
ðTBL TL Þ, K
t m, s
387 307
5.6 5.6
1.67 1.67
0.15 0.15
1803 1395
20 418
1200 45
Frozen shell No frozen shell Source: Ref 18
7
Predicted melting/dissolution times for silicomanganese spheres in quiescent steel baths versus steel superheat temperature. Particle diameters of the spheres are as follows: A, 25 mm (1 in.); B, 50 mm (2 in.); C, 100 mm (4 in.). Source: Ref 18
requirement for melting, divided by the rate at which heat is supplied from the bath (Ref 18), that is:
Table 3 Data for melting time calculations for a 10 mm (0.4 in.) diameter aluminum sphere in different liquid metal melts Base metal
DH, J/g
r0, g/cm3
V0/A0, cm
h, J/g K
TL , K
ðTBL TL Þ, K
tm, s
Steel Cast iron Copper Aluminum
576 576 576 576
2.7 2.7 2.7 2.7
0.167 0.167 0.167 0.167
0.15 0.15 0.15 0.15
1823 1723 1323 1023
890 790 290 90
1.9 2.2 6.0 19.2
annulus increases (t3). In the final stage, beginning at t4, the redissolution of the base metal shell is complete, and as a result, the liquid alloy held in the annular space is released to the melt. Any alloy that remains solid is exposed directly to the liquid metal bath and is progressively incorporated into the melt by the surface melting and convection. Dissolution is completed at tm (melting time). The length of the incubation period (t0 t4) and the length of the alloy dissolution period (t4 tm) are governed by heat transfer. Conditions favoring rapid alloy dissolution are high superheat temperatures, small alloy particle size (large surface for heat transfer), and low alloy thermal conductivity. The heat flux to the alloy particle is given by: Q ¼ h TLB TL
(Eq 36) 2
1
where Q is the heat flux (J m s ), h is the convective heat transfer coefficient (J m2 s1 K1), TLB is the temperature of
the bulk liquid (in degrees Kelvin), TL is the melting point temperature of the alloy at the solid/liquid interface (in degrees Kelvin), and TLB TL is the superheat temperature (in degrees Kelvin). During the incubation period, the superheat temperature is less than that in the alloy dissolution period because TL corresponds to the melting point of the base metal and alloy in the former and latter cases, respectively. For steel and aluminum melts, superheat temperatures in the incubation period are generally smaller (20 to 100 K) than for cast iron (200 to 300 K). In the alloy dissolution period, the superheat temperature for a given alloy addition increases with increasing melting point of the base metal. Thus, superheat temperatures, and therefore dissolution rates, increase in the following order: aluminum < copper < cast iron < steel. Melting times, tm, can be estimated by assuming that tm is equal to the total heat
tm ¼
Hr0 V0 QA0
(Eq 37)
where DH is the heat supplied to 1 g of alloy at the time it dissolves, and r0, V0, and A0 are the density, volume, and surface area of the alloy particle. The variable Q is defined by Eq 36. Equation 37 is most accurate when applied at the extremes of possible superheat temperatures. For small superheat temperatures, it can be assumed that by the time the frozen shell redissolves, the alloy is a liquid at the melting point of the base metal. At high superheat temperatures, it can be assumed that no frozen layer forms and that the alloy liquefies at its melting point. These two cases have been examined for the dissolution of a 100 mm (4 in.) particle diameter B(dp) sphere of silicomanganese in a steel melt TL ¼ 1823 K . Data for the calculations are given in Table 2. The very long melting time of 1200 s obtained by restricting the calculations to a small superheat temperature (20 K) is in good agreement with the 900 s obtained with a more sophisticated model (Fig. 7). The difference in tm for the two extreme cases considered in Table 3 underscores the importance of superheat temperature in alloy dissolution processes. This is clearly illustrated in Fig. 7, in which the role of particle diameter is also demonstrated. It is apparent from Fig. 7 that the large value of tm in Table 2 for the case
Chemical Thermodynamics and Kinetics / 37 a common foundry practice that is necessary for control of cast iron composition. The rate at which carbon dissolves in cast iron is controlled by the rate of diffusion of carbon from carbon-saturated iron at the carbon/iron interface into the bulk liquid (Ref 21–23). Thus, Eq 38 can be written:
Density ratio 40
0
0.2
0.4
0.6
0.8
1.0
Liquid iron-carbon alloy
Carbon
C sat * Carbon concentration
20
10
Total immersion time, s
6 4
d 2 Distance C
1
Fig. 9
0.6
B
0.4
A
0.2
0.1 0
1
2
3
4
5
6
7
Addition density, g • cm-3
Fig. 8
Total immersion times versus density ratio for alloys entering a stagnant steel bath at 7.67 m s1 (25.2 ft s1) (drop height: 3.0 m, or 10 ft). Particle diameters are as follows: A, 10 mm (0.4 in.); B, 50 mm (2 in.); C, 250 mm (10 in.). Source: Ref 18
J ¼ km ðCsat CLB Þ
C BL
TLB TL ¼ 20 and dp = 100 mm can be substantially B reduced with moderate changes in TL TL and dp. The range of melting times obtained with the addition of an alloy to different base metals is examined subsequently. Consider the addition of pure aluminum to melts of steel, cast iron, copper, and aluminum under conditions where no frozen shell is formed. In this case, TL is the same (933 K), but TLB varies. Pertinent data for the calculation and the calculated melting times are given in Table 3. It is clear from Table 3 that the duration of alloy melting becomes a more serious problem as the melting point of the base metal decreases. The short melting times predicted for steel and cast iron melts in Table 3 highlight another important consideration in alloy addition, namely, immersion time due to alloy buoyancy. Because iron is denser than common alloy additions such as carbon, silicon, and aluminum, buoyant forces prevent these alloys from being easily submerged for times approaching tm. This is illustrated in Fig. 8, which plots immersion times versus the density ratio, ralloy/ rbase metal. For 10 mm (0.4 in.) aluminum particles (density ratio: 0.39) that are dropped
Carbon concentration profiles for a carbon rod dissolving in an iron-carbon melt
a distance of 3 m (10 ft) to the surface of an iron bath, the immersion time is less than 1 s, while the melting time (from Table 3) is 1.9 s. Immersion times can be increased by entraining particles in strong recirculating flows or by wire and rod feeding (Ref 19, 20). These conditions have higher associated heat transfer coefficients, which further facilitate dissolution.
Dissolution of Alloy That Is Solid at Bath Temperatures The processes involved in the dissolution of this type of addition are schematically illustrated in the bottom row sequence in Fig. 6. As in the previous case, heat transfer controls the duration of the incubation period (t0 t4). However, in this case, mass transfer is the controlling factor in the alloy dissolution period (t4 tm). In general, mass transfer controlled dissolution is slower than heat transfer controlled dissolution, and the number of variables controlling the rate is greater. For mass transfer control, melt temperature and fluid dynamic considerations are important, as in the heat transfer controlled case. However, melt composition and diffusion variables are also important. The flux of alloy element to the melt is given by: J ¼ km ðCL CLB Þ
(Eq 38)
where J is the mass flux of alloy element (g cm2 s1), km is the mass transfer coefficient (cm s1), CLB is the concentration of the alloy element in the bulk melt (g cm3), and CL is the concentration of the alloy element in the melt at the solid/liquid interface (g cm3). To illustrate mass transfer limited dissolution processes, two industrially important examples are examined as follows. Carbon Dissolution in Cast Iron. Carbon is the most important alloy element in cast iron. The addition of carbon to iron-carbon melts is
(Eq 39)
where Csat indicates carbon-saturated conditions at the carbon/melt interface. The conditions in the melt are qualitatively described in Fig. 9. Referring to Fig. 9 and Eq 39, the flux of carbon, and therefore the dissolution rate, increases with decreasing bulk carbon concentration CLB , increasing carbon concentration at saturation Csat (increasing bath temperature), and decreasing boundary layer thickness, d. The boundary layer thickness is a function of the fluid velocity, and it affects the value of km. CLB ) on The effect of the difference (Csat the dissolution rate of carbon is illustrated in Fig. 10 with data for the dissolution of graphite granules in induction-stirred cast iron melts. Theoretical curves are drawn for different rate constants k and particle shape factors l, where l is defined as:
¼
Sparticle Ssphere
(Eq 40)
where Sparticle is the surface-to-volume ratio of the graphite particles, and Ssphere is the equivalent surface-to-volume ratio if the particles are spherical. The role of boundary layer thickness (fluid dynamics) on dissolution kinetics is illustrated in Fig. 11, which indicates the variation of the rate constant (dissolution rate) with the peripheral velocity (rotational speed) of a cylindrical graphite rod. The line drawn through the data has a slope of 0.7, which is expected from earlier correlations (Ref 24). Dissolution of Steel in Iron-Carbon Melts. Another well-established case of mass transfer controlled dissolution is the dissolution of steel in iron-carbon melts (Ref 25–28). This process is critical in the electric melting of steel and cast iron. For this case, the rate of steel dissolution increases with melt carbon concentration (Fig. 12), melt temperature (Ref 27), and bath velocity (Ref 25, 26, 28) (affecting the boundary layer thickness). The data can be accounted for on the basis that carbon diffuses from the bulk of the melt to the steel/liquid iron interface, where steel is isothermally converted to a liquid with a liquidus composition, CL‘ . The conditions existing in the bath are illustrated in Fig. 13. Figure 13(a) shows the carbon concentration profiles in the liquid bath in the vicinity of a dissolving steel bar. The meaning of the carbon concentration notation is indicated in Fig. 13(b). As indicated in Fig. 13(a), carbon diffuses across the liquid boundary layer, while equilibrium conditions are maintained at the solid/liquid interface. Thus, the carbon concentrations in the solid and the liquid
38 / Principles of Liquid Metal Processing 1.0
1000 15
800
λ = 2.0
10
600 λ = 2.5
0.8
Time to dissolve, s
200
3 λ = 1.4 2
100
λ = 1.6
80 60
k = 0.007 λ = 1.6
1
k = 0.007
Time to dissolve, min
5
0.6
0.4
0.2
k = 0.02
40
Fractional thickness, r /r0
400
B A
k = 0.02 20
0 k = 0.03
0
10
20
30
40
Immersion time, min 10 0
0.5
1.0
2.0
1.5
2.5
3.0
Fig. 12
Melting curves for 75 mm (3.0 in.) diameter steel rods immersed in iron-carbon baths. A, high (4.6% C) iron-carbon bath; B, low (2.1% C) ironcarbon bath. Plotted is the fractional thickness (the fraction of the rod diameter remaining undissolved) versus immersion time. Source: Ref 27
3.5
2 Weight fraction of solution (C*sat – CB L ) ⫻ 10
Fig. 10
Time of dissolution of carbon particles (10 + 14 mesh) versus (Csat CLB ). Source: Ref 23
J ¼ km ðCLB CL‘ Þ
0.1
(Eq 41)
The carbon flux can also be defined in terms of a carbon balance: B ‘ r C dr J ¼ S ‘L CSB dt rL
Mass transfer coefficient (km), cm • s−1
0.05 0.7 k = (const) V i
where rBS and r‘ L are the respective densities of solid steel and liquid iron at the liquidus composition, CSB is the bulk carbon concentration of solid steel, and dr/dt is the rate of change of the radius or thickness of the steel with time. By combining Eq 41 and 42 and integrating between the limits r = r0 at t = 0 and r = 0 at t = tm, the isothermal melting time tm is given as:
0.02
0.01
(C *sat – C BL )
tm ¼
0.8–2.2 wt % 2.2–3.7 wt % 3.7–5.0 wt %
0.005
0.002 2
5
10
20
50
100
200
400
Peripheral velocity (Vi), cm • S−1
Fig. 11
Experimental mass transfer coefficient versus peripheral velocity for the dissolution of a rotating carbon rod in an iron-carbon melt. Source: Ref 22
at the interface are the isothermal solidus Css and the liquidus CL‘ concentrations, respectively. Dissolution takes place by isothermal conversion of the steel into a liquid, that is, by a chemical melting process. It is clear from
(Eq 42)
Fig. 13(b) that as temperature increases, CL‘ decreases significantly. Thus, the driving force for diffusion increases with temperature. Based on Eq 38 and Fig. 9, the flux of carbon entering the steel is:
B ðrBS CL‘ =r‘ L Þ CS r0 ‘ B km ðCL CL Þ
(Eq 43)
In laboratory studies, Eq 43 accurately predicted steel dissolution rates in a stagnant bath (Ref 27). Research has suggested that melting times are reduced by approximately a factor of 2 in induction-stirred melts (Ref 26). Other work has indicated that tm can be reduced by an order of magnitude with strong convection (Ref 28). On the basis of Eq 43 and the mass transfer correlation for rapidly dissolving rods (Ref 29), correlations were developed relating tm to bath temperature and to carbon concentration (Ref 27). This is given in Fig. 14 for a steel thickness of 2.5 mm (0.098 in.). Because tm is directly proportional to the thickness of the steel, Fig. 14 can be used to estimate tm for a wide range of steel thickness by making the appropriate thickness correction. Checks of the
Chemical Thermodynamics and Kinetics / 39
Steel
Liquid ironcarbon alloy
4. 5.
C BL
Carbon concentration
6.
7.
C *Ll
8. 9.
C *Ss δ
10. C BS
11. Distance (a)
12. 13.
1600 δ+L
2600
C BL
15. 2400
1300 γ+L
1200
2200
1100
Temperature, ⬚F
C *Ll
C *Ss
C BS
1400 Temperature, ⬚C
14.
2800
Liquid alloy
δ
1500
2000 γ + Fe3C
1000
16. 17.
1800
900 0
0.5
Fig. 13
1
2
3
4
5
18.
Carbon, wt%
(b)
(a) Carbon concentration profiles for a steel rod dissolving in an iron-carbon melt. (b) Iron-carbon phase diagram defining the carbon concentrations noted in (a). Source: Ref 27
predictions from Fig. 14 in production operations have indicated that predicted tm values are realistic (Ref 27). ACKNOWLEDGMENT Revised from the articles “Principles of Physical Chemistry” and “Composition Control” by Seymour Katz in Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 50–55 and 71–81
REFERENCES 1. Y.K. Rao, Stoichiometry and Thermodynamics of Metallurgical Processes, Cambridge University Press, 1985, p 285–287, 298, 383–394, 880–891 2. E.T. Turkdogan, Physical Chemistry of High Temperature Technology, Academic Press, 1980, p 5–26, 81, 82 3. G.R. Belton and R.J. Fruehan, The Determination of Activities by Mass Spectrometry. I. The Liquid Metallic Systems Iron-
19. 20.
21.
22.
Nickel and Iron-Cobalt, J. Phys. Chem., Vol 71, 1967, p 1403–1409 G.K. Sigworth and J.F. Elliott, The Thermodynamics of Liquid Dilute Iron Alloys, Met. Sci., Vol 8, 1974, p 298–310 G.K. Sigworth and J.F. Elliott, The Thermodynamics of Dilute Liquid Copper Alloys, Can. Metall. Q., Vol 13, 1974, p 455–461 G.K. Sigworth and T.A. Engh, Refining of Liquid Aluminum—A Review of Important Chemical Factors, Scand. J. Metall., Vol 11, 1982, p 143–149 C. Wagner, Thermodynamics of Alloys, Addison-Wesley, 1962, p 51 L.S. Darken, Thermodynamics of Binary Metallic Solutions, Trans. Metall. Soc. AIME, Vol 239, 1967, p 80–89 L.S. Darken, Thermodynamics of Ternary Metallic Solutions, Trans. Metall. Soc. AIME, Vol 239, 1967, p 90 R. Hultgren et al., Selected Values of the Thermodynamic Properties of Binary Alloys, American Society for Metals, 1973 L.S. Darken, Thermodynamics and Physical Metallurgy, American Society for Metals, 1950 L.S. Darken and R.W. Gurry, Physical Chemistry of Metals, McGraw-Hill, 1953, p 235 Aluminum Casting Technology, American Foundrymen’s Society, 1986, p 21 M.J. Lalich and J.R. Hitchings, Characterization of Inclusions as Nuclei for Spheroidal Graphite in Ductile Cast Irons, Trans. AFS, Vol 84, 1976, p 653–664 B. Francis, Heterogeneous Nuclei and Graphite Chemistry in Flake and Nodular Cast Irons, Metall. Trans. A, Vol 10, 1979, p 21–31 S. Katz, The Properties of Coke Affecting Cupola Performance, Trans. AFS, Vol 90, 1982, p 825–833 R.J. Fruehan, B. Lally, and P.C. Glaws, A Model for Nitrogen Absorption in Iron Alloy Melts, Fifth International Iron and Steel Congress Process Technology Proceedings, Vol 6, The Iron and Steel Society of AIME, 1986, p 339–346 R.I.L. Guthrie, Addition Kinetics in Steelmaking, Proceedings of the 35th Electric Furnace Conference, Iron and Steel Society of AIME, 1978, p 30–41 F.A. Mucciardi and R.I.L. Guthrie, Aluminum Wire Feeding in Steelmaking, Trans. ISS, Vol 3, 1983, p 53–59 J.W. Robison, Jr., “Ladle Treatment with Steel-Clad Metallic Calcium Wire,” Paper 35, presented at Scaninject III, Part II, MEFOS, Lulea, Sweden, 1983 L. Kalvelage, J. Markert, and J. Po¨tschke, Measurement of the Dissolution of Graphite in Liquid Iron by Following the Buoyancy, Arch. Eisenhu¨ttenwes., Vol 50, 1979, p 107–110 R.G. Olsson, V. Koump, and T.F. Perzak, Rate of Dissolution of Carbon in Molten
40 / Principles of Liquid Metal Processing Bath temperature, ⬚F 2400
2300 3.5
2500
3.0
2600
2700
2% C bath 4% C bath
1.5% C bath
3% C bath
Melting time, min
2.5
2.0
1.5
1.0
5% C bath 0.5
0 1260
1315
1370
1425
1480
Bath temperature, ⬚C
Fig. 14
Predicted melting times for vertical steel plates and cylinders, 2.5 mm (0.098 in.) scrap thickness and 0.1% C composition, [(%S)i = 0.03] immersed in stagnant iron-carbon baths at different bath temperatures. Source: Ref 27
23. 24.
25.
26.
Fe-C Alloys, Trans. Met. Soc. AIME, Vol 236, 1966, p 426–429 O. Angeles, G.H. Geiger, and C.R. Loper, Jr., Factors Influencing Carbon Pickup in Cast Iron, Trans. AFS, Vol 74, 1968, p 3–11 M. Eisenberg, C.W. Tobias, and C.R. Wilke, Mass Transfer at Rotation Cylinders, Chem. Eng. Prog. Symp. Series, Vol 51, 1955, p 1–16 R.G. Olsson, V. Koump, and T.F. Perzak, Rate of Dissolution of Carbon Steel in Molten Iron-Carbon Alloys, Trans. Met. Soc. AIME, Vol 233, 1965, p 1654–1657 R.D. Pehlke, P.D. Goodell, and R.W. Dunlap, Kinetics of Steel Dissolution in Molten
Pig Iron, Trans. Met. Soc. AIME, Vol 233, 1965, p 1420–1427 27. R.I.L. Guthrie and P. Stubbs, Kinetics of Scrap Melting in Baths of Molten Pig Iron, Can. Metall. Q., Vol 12, 1973, p 465–473 28. K. Mori and T. Sakuraya, Rate of Dissolution of Solid Iron in Carbon-Saturated Liquid Iron Alloys with Evolution of CO, J. Iron Steel Inst. Jpn., Vol 22, 1982, p 964–990 29. P.T.L. Brian and H.B. Hales, Effects of Transpiration and Changing Diameter on Heat and Mass Transfer to Spheres, AICHE J., Vol 15, 1969, p 419–425
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 41-55 DOI: 10.1361/asmhba0005190
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Thermodynamic Properties of Iron-Base Alloys Doru M. Stefanescu, Ohio State University Seymour Katz, S. Katz Associates
THE STRUCTURE AND PROPERTIES of cast metals are sensitive to numerous impurities. Because purification of melts generally adds considerable cost to castings, the lowest cost and surest defense against contamination is careful selection of scrap. Purification is generally reserved for elements that are so pervasive that avoidance is impossible. This is exemplified by sulfur and oxygen removal from cast iron and steel, respectively, and removal of alkali and alkaline earth elements from aluminum. Because of their immediate importance, aspects of the physical chemistry of these processes are reviewed. The structures and properties of steel and cast iron also depend on the control of the solidification structure in castings. To accomplish this, it is necessary to understand and control the thermodynamics of the liquid and solid phases and the kinetics of solidification (nucleation and growth of various phases). The information linked to the practical interests that thermodynamics can provide when considering iron-base alloys encompasses a rather wide range. Only two issues are addressed in sections on thermodynamics in this article. The first is the calculation of solubility lines, relevant to the construction of phase diagrams, and the second is the calculation of the activity of various components, which allows for the determination of probability of formation and relative stability of various phases. Alloying elements are then discussed in terms of their influence on the activity of carbon, which provides information on the stability of the main carbonrich phases of iron-carbon alloys, that is, graphite and cementite. Information on the iron-oxygen system can be found in the article “InclusionForming Reactions” in this Volume.
the level of inclusions, leading to stronger and more fatigue-resistant steels. For cast iron, desulfurization is practiced in the manufacture of ductile iron castings in order to develop spherical graphite morphology. Ductile iron is used in applications where high fracture toughness is needed. Sulfur is removed from iron and steel when the metals are liquid. Although a variety of reagents are employed to remove sulfur, namely, calcium, magnesium, sodium, and rare earths, the most important is calcium. Common forms of calcium include the metal; alloys such as calcium silicon (CaSi); the oxide, calcium oxide (CaO); and the carbide, calcium carbide (CaC2). Despite the use of various forms of calcium, the governing chemical reaction in all cases appears to be the same (Ref 1, 2): CaO þ S ¼ CaS þ O
(Eq 1)
The equilibrium constant for the reaction (Eq 1) is: aCaS hO K9 ¼ aCaO hS
(Eq 2)
For both cast iron and steel, target sulfur concentrations after desulfurization are in the range of 0.006 to 0.010% S. In electric-melted cast iron and steel, the sulfur levels before desulfurization are 0.02 to 0.03% S, while input sulfur levels to the cupola are generally much higher—0.1 to 0.2% S. Requirements for Desulfurization. For reasons related to reaction kinetics and thermodynamics, the final sulfur (%S)f concentration achieved in desulfurization processes depends on three factors:
Purification of Ferrous Melts
Initial sulfur concentrations (%S)i Amount of desulfurizer used, usually expressed
One of the most important processes involved in cast iron and steel production is desulfurization. For steels, desulfurization is necessary to reduce
as the weight fraction of desulfurizer to metal, W Extraction efficiency of the desulfurizer, which is measured by the desulfurization
ratio (DR), that is, the ratio of sulfur concentrations: desulfurizer to metal The three variables can be related as follows through a mass balance on sulfur: ð%SÞf ¼
ð%SÞi 1 þ W ðDRÞ
(Eq 3)
For liquid-desulfurizing slags that are not saturated with respect to calcium sulfide (CaS), the maximum value of DR, that is, the equilibrium value, can be predicted from thermodynamic considerations: ðDRÞmax ¼
ð%SÞslag ð%SÞT
¼ max
CS fS K14 K15 hO
(Eq 4)
where CS is the slag sulfide capacity, defined as (Ref 3): CS ¼ ð%SÞslag
1=2 pO2 pS2
(Eq 5)
and K14 and K15 are the respective equilibrium constants for the reactions: 1 O2 ¼ O 2
(Eq 6)
1 S2 ¼ S 2
(Eq 7)
and fS is the activity coefficient for 1 wt% S in the standard state. Equation 4 can be derived from Eq 2. Using known input and desired output sulfur values, Eq 3 gives the required desulfurization ratios as a function of weight fraction desulfurizer. These data are plotted in Fig. 1 for two cases. In the electric-melting case, initial and final sulfur concentrations were assumed to be 0.03 and 0.008% S, respectively. In the cupola-melting
42 / Principles of Liquid Metal Processing
Weight ratio, Wdesulfurizer/Wiron
1
Table 1 Theoretical weight ratios: desulfurizer to iron to achieve 0.008% S in various systems T, K
CS ( 104)(a)
fS
ho ( 104)(b)
44 CaO-15MgO-5Al2O3-36SiO2 CaOsat-Al2O3
1773 1773
2.7 59
5 5
1.3 5.0
90 417
CaOsat-Al2O3 CaOsat-Al2O3
1773 1873
59 316
5 1
5.0 0.45
417 5280
Type of melt
A. cast iron—cupola B. cast iron—cupola (ladle desulfurized) C. cast iron—electric D. steel—electric
0.1 B A
100
(DR)max(c)
W
0.13 0.027 0.0065 0.00052
(a) CS is obtained from optical basicity correlations (Ref 3). fs is based on iron composition (Ref 4). (b) Case A: ho is governed by Si/SiO2 equilibrium based on respective concentration in iron and slag (Ref 5). Cases B and C: ho is governed by Si/SiO2 equilibrium with aSiO2 ¼ 1 due to ladle exposure to air (Ref 6). Case D: ho is governed by Al/Al2O3 equilibrium with %Al = 0.05 (Ref 1), assumed no contact with air. (c) K14 and K15 data from Ref 7
0.01
10-3 10
Slag composition
103
104
104
Desulfurization ratio (DR), (% S)slag/(%S)iron
Fig. 1
The best desulfurizing cupola slags have lower
sulfide capacity and desulfurization ratio than slags used in ladle desulfurization; the consequence is the need for much larger quantities of slag*. Compared to steel, cast iron desulfurization benefits from higher fS because of the presence of relatively high concentrations of carbon and silicon in cast iron. Ladle desulfurization systems that are exposed to air suffer higher hO and, as a result, poorer desulfurization than may otherwise be anticipated. For the cupola slag case in Table 1, the thermodynamically predicted value of (DR)max = 90 is in good agreement with measured data (Fig. 2). This indicates that the cupola desulfurization process operates close to equilibrium levels. Further evidence for this is given Fig. 3, which plots desulfurization data for a cupola operated with varying amounts of municipal ferrous refuse in the charge. The upper portion of Fig. 3 plots CS and oxygen activity. The latter is expressed as the partial pressure of oxygen. *The actual differences are less than those indicated. Calcium sulfide has only limited solubility in CaO-Al2O3 slags. As a result, a minimum W = 0.01 is needed to maintain the CaS-unsaturated condition.
103 Desulfurization ratio, aCaS/as
case, the equivalent concentrations were 0.10 and 0.008% S. The cupola line applies for both cupola iron and ladle-desulfurized cupola iron. The relationships illustrated in Fig. 1 are useful in defining systems that will provide the necessary desulfurization conditions. Four systems are compared in Table 1. These cover cupola- and electric-melted cast iron and electric-melted steel. Also examined are two liquid slag systems: a basic cupola slag (dicalcium silicate saturated) and a CaO-saturated CaOAl2O3 slag. Table 1 shows that:
100
Desulfurization ratio, (%S)slag /(%S)iron
Weight fraction desulfurizer required to achieve given desulfurization ratios. Plots for electricmelted iron and steel (curve A) assume initial sulfur concentrations [%(S)i = 0.03] different from those of cupola-melted iron [(%S)i = 0.10] (curve B).
100
10
10 0.7
0.9
1.1
1.3
1.7
1.5
1.9
2.1
2.3
Slag basicity
Fig. 2
Desulfurization ratio-basicity comparisons for cupola (closed circles) and laboratory data (open circles). Equilibrium values are indicated by the angled line. The vertical line indicates the slag basicity above which slags are saturated at 1500 C (2730 F), with respect to dicalcium silicate. This is the point at which the observed DR should equal (DR)max. Source: Ref 2
The lower portion of Fig. 3 is a comparison of (DR), measured and calculated. The good agreement found is evidence that near-equilibrium conditions existed. For the cases discussed previously, the oxygen activity in cupola iron has been found to be governed by the reaction (Ref 5, 8):
Mn þ O ¼ MnO
(Eq 8)
Therefore, the overall cupola desulfurization reaction is: CaO þ S þ Mn ¼ CaS þ MnO
(Eq 9)
Thermodynamic Properties of Iron-Base Alloys / 43
106
0.01
%C ⬚C 3.7 1550 4.0 3.7 1500 4.0
10−3
10−5 100 Desulfurization ratio
Inlet S 0.015 – 0.03% Inlet S 0.06 – 0.08%
(6 ⫻ 10−4)
Measured Calculated
50
0
0.5
1.0
1.5
3.7 4.0
1450
3.7 4.0
1400
2.0
2.5
3.0
Silicon concentration in iron, wt %
20 10
Fig. 4
Equilibrium sulfur concentrations for CaO desulfurization calculated with Eq 11 and 12, assuming aCaO ¼ aCaS ¼ aSiO2 ¼ 1. Source: Ref 2
5
Measured sulfur concentration, wt %
10−4
Sulfur concentration in iron, wt %
107
10
(po2)−1/2
Slag sulfide capacity
Cs (p o2)−1/2 Mn −3
0.1
(0.02)
108
0.01
0.01
2 1
0
20
40
60
80
100
Municipal refuse in charge, %
Fig. 3
Slag sulfide capacities, oxygen activities, and desulfurization ratios, measured and calculated, for a cupola operated with municipal ferrous refuse as a charge material. Source: Ref 8
Considerably lower sulfur could be achieved (Ref 2) if hO were governed by: C þ O ¼ CO
(Eq 10)
However, equilibrium for this reaction has not been observed. This discussion has concerned CaS-unsaturated slags. However, a desulfurizing slag, saturated with CaS, can in many cases continue to desulfurize as long CaO is present. In this case, the ultimate sulfur levels are not as low as those for CaS-unsaturated slags. Nevertheless, CaSsaturated slags can produce sulfur levels that are more than sufficient for cast iron and steel applications. In fact, CaS-saturated conditions are often employed in industrial applications because much less desulfurizer is required. In addition to the CaS-saturated liquid slags, totally solid slags such as CaO and CaC2 are in the same category, because CaS does not form solid solutions with these materials. Also in this category are desulfurizers containing small amounts of liquid phase. These materials possess the desirable properties of liquid and solid slags. That is, they combine the fast reaction rates of liquid slags with the large desulfurizing capacities (high CaO concentration) of solid slags. The final sulfur concentrations attainable under CaS-saturated conditions can be obtained from Eq 2 by setting aCaS = 1. For ladle desulfurization, the most important industrial desulfurizers are also CaO-saturated, that is, aCaO = 1.
Applying the condition of double saturation to Eq 2 yields:
1 10−3
ð%SÞr ¼
hO K9 fS
(Eq 11)
In the ladle desulfurization of cast iron, hO is governed by the reaction (Ref 2): 1 1 Si þ O ¼ SiO2 2 2
1
(Eq 12)
where aSiO2 ¼ 1. This is attributed to the exposure of the iron to air (Ref 2). Sulfur concentrations obtained with Eq 11, assuming siliconsilicon dioxide (Si-SiO2) equilibrium and aSiO2 ¼ 1, are given in Fig. 4. In continuous ladle-desulfurization processes, equilibrium sulfur levels are achieved when input sulfur levels are low but are only approached when input levels are high. This is illustrated in Fig. 5 with desulfurization data for CaC2 (Ref 2). Other desulfurizers such as CaO-CaF2 behave similarly. High desulfurization rates are needed for effective desulfurization in the short times required. For CaO-base desulfurizers, the presence of relatively small amounts of liquid phase (<25 vol%) significantly increases the rate of desulfurization. This is illustrated in Fig. 6, in which the rates of desulfurization by CaO with varying amounts of calcium fluoride (CaF2) are compared. The faster rates obtained with CaF2 additions were due to the formation of a CaOCaF2 liquid phase that provided a path for CaS diffusion from the reaction interface into the porous CaO particle. This prevented the normal rapid development of an impervious CaS coating on the CaO surface (Ref 10, 11). A similar explanation, involving liquid-phase formation, was used to account for the higher
(6 ⫻ 10−4) 0.0006
0.001
0.002
0.004
Theoretical sulfur concentration, wt%
Fig. 5
Comparisons of the sulfur concentrations in production-continuous desulfurization using CaC2 with the sulfur concentrations from Eq 11 and 12. Open circles indicate electric-melted iron. Closed circles indicate cupola iron. Source: Ref 2
CaO desulfurization rates of steel when aluminum was concurrently added (Ref 12, 13). Another rate-controlling variable in ladle desulfurization is bath agitation. For gas-stirred melts, the desulfurization rate constant k is _ that is, k / Q_ n . a function of gas flow rate Q, For well-dispersed solid-liquid mixtures, the rate of diffusion-controlled interphase reaction _ that is, n is a relatively weak function of Q, 0.2 to 0.4 (Ref 14–16). For a poorly dispersed system, such as a ladle of iron with a cover slag of desulfurizer, agitation has a much greater influence on the rate constant, with n typically in the range of 1.0 to 2.5 (Ref 14–16). This is due to the entrainment of increasing amounts of top slag into the liquid metal with increasing _ This effect is illustrated in Fig. 7, which Q. plots the apparent mass transfer coefficient for desulfurization versus the gas flow rate. Deoxidation. In the manufacture of steel, FeO-saturated conditions (0.2%O) are approached as a result of oxygen blowing to remove carbon and silicon as the respective oxides. To make a steel product, the oxygen levels are subsequently lowered to 0.005 to 0.02% O by reactive alloy additions of
44 / Principles of Liquid Metal Processing 1
0% CaF2
0.5
2%
0.2
%S/ %So
8%
0.1
5% CaF2
0.05
20%
10% 0.02
0.01 0
10
20
30
40
50
60
Time, min
Fig. 6
Desulfurization rates of carbon-saturated iron, containing 0.4% Si, with CaO and varying amounts of CaF2 at 1450 C (2640 F). Source: Ref 9
For single-element deoxidation, expressed by:
Desulfurization rate constant (k ), s -1 ⫻ 10-3
10
xM þ yO ¼ Mx Oy
5
(Eq 13)
the equilibrium oxygen concentration in the liquid steel, obtained by rearranging the equilibrium constant expression, is:
2 1
•
k a Q0.25
•
k a Q 2.1
0.5
ð%OÞ ¼
0.2 0.05 0.1
0.2
0.5
1
2
5
10
Gas flow rate (Q ) , m3 • kg-1 • s-1 ⫻ 10-6
Fig. 7
Dependence of CaO desulfurization rate constant on the rate of gas flow through a reactor. Source: Ref 16
electropositive elements. The low oxygen concentrations are necessary to maintain the number of oxide inclusions at a suitable low level, to prevent the formation of CO blowholes, and to ensure effective desulfurization.
a
Mx Oy x y K22 ð%MÞx fM fO
y1 (Eq 14)
The calculated equilibrium values of oxygen solubility in liquid iron at 1600 C (2910 F), based on Eq 14, are given in Fig. 8 for several elements as a function of the concentration of the element, assuming aMx Oy ¼1 . As seen among the common deoxidizers, aluminum produces the lowest oxygen. Although not shown, rare earth metals produce even lower oxygen (Ref 17). Figure 8 plots data for oxygen concentration and oxygen activity. In all cases, oxygen activity decreases with increasing levels of deoxidizer, but the oxygen concentration can go through a minimum. This occurs in cases where the
interaction coefficient eM O is a large negative number, and as a result, fO decreases significantly even at relatively low concentrations of M. To obtain exact values for [%O] or hO, equations for K22 can be found in Ref 1 and 7. As indicated by Eq 14, the oxygen concentration can be reduced if a complex deoxidation product is formed so that aMx Oy <1. Common deoxidation systems of this type are Si-Mn, Al-Si-Mn, or Al-CaO. An advantage, in addition to lower oxygen, is that less deoxidizer is required in solution to achieve a given level of oxygen. This is illustrated in Fig. 9, which plots the oxygen activity of iron as a function of the concentration of CaO in the calcium aluminate inclusion. Separate lines are given for aluminum concentrations ranging from 0.001 to 0.05% Al. Also indicated is the oxygen activity when pure aluminum oxide (Al2O3) is the reaction product. The data in Fig. 9 show that an oxygen activity of 4 ppm can be produced at three aluminum levels—0.002, 0.005, or 0.01%—depending on whether the respective deoxidation product was CaO-saturated calcium aluminate, CaAl2O4-saturated calcium aluminate, or Al2O3. The levels of total oxygen measured in steel melts considerably exceed equilibrium levels (Ref 7). Two possible explanations are the kinetic limitations in the deoxidation reaction and the ineffective separation of deoxidation products from the melt. Consideration of the first possibility suggests that the most important kinetic limitation would be oxygen and/or deoxidant solute diffusion to inclusions in the melt (Ref 7). Considering the presence of 105 to 107 inclusion particles in a cubic centimeter of melt, deoxidation to equilibrium levels of 10 to 20 ppm would require a few minutes. Appreciably longer times would be needed to reduce oxygen to less than 1 ppm, as with rare earth additions. For most cases, oxide formation appears faster than oxide particle separation from the bath. Thus, for applications that depend on dissolved oxygen, such as desulfurization or CO blowhole formation, equilibrium oxygen values can be assumed to exist after relatively short reaction periods. For considerations of the inclusion content of steel, the kinetics of particle separation need to be considered. The limiting case of particle flotation at velocities obtained from Stokes’ law calculations is shown in Fig. 10. Inclusion size versus time of flotation is plotted. The particles produced during deoxidation are generally small (<5 mm, or 200 min.); accordingly, separation times are long. The times indicated in Fig. 10 are representative of conditions in a quiescent bath. Times are considerably shorter for stirred melts because articles grow by collision and coalescence to sizes of 30 to 100 mm (Ref 7). The efficiency of coalescence of particles is material dependent. The energy of adhesion of particles is determined by the wetting characteristics of the particle, measured by wetting angle. For example, the force of adhesion of Al2O3
Thermodynamic Properties of Iron-Base Alloys / 45 20
1
0.002 % Al 0.001
Oxygen concentration Oxygen activity ao ∫ wt% O for %X Æ 0 Iron oxide saturation
Oxygen activity of iron, ppm 0
V2O4
0.1
MnO Cr2O4
SiO2 Oxygen is solution, wt %
16
B2O3
CO
Cr2O3
B2O3
V 2O 3
0.01
Ti2O3
10-3
0.005
0.01
0 30
0.01
CaO saturated
0.05 40
50
60
70
CaO concentration, mass%
ZrO2
Fig. 8
8
0.05
Al2O3
Fig. 9
Ti2O3 0.01
0.002 0.005
4 CaAl2O4 saturated
Ti3O5
10-4 -3 10
12
0.1 Deoxidizer in solution, wt %
1
10
Concentration of oxygen and aluminum in liquid steel in equilibrium with calcium aluminate at 1600 C (2910 F); arrows indicate oxygen values when the reaction product is pure Al2O3. Source: Ref 18
Deoxidation equilibria in liquid iron alloys at 1600 C (2910 F). Source: Ref 17
107
106
10
15 105
20 30 40 0.1
A
B
1
C
10
Z, number of nuclei • cm−3
Inclusion size for 0.05%O removal, µm
5
100
Flotation time, min
Fig. 10
Calculated time of flotation of inclusions in stagnant melts as a function of inclusion size. Melt depth: A, 50 mm (2 in.); B, 500 mm (20 in.); C, 2000 mm (80 in.). Source: Ref 7
particles with a wetting angle of 140 is twice that for SiO2 with a wetting angle of 115 . On this basis, Al2O3 inclusions are expected to cluster more easily than SiO2 inclusions, a phenomenon that has been observed in practice (Ref 7).
Thermodynamics of Ferrous Systems For binary systems, only the iron-carbon and iron-silicon systems are discussed in this article. Information on the iron-oxygen system is available in the article “Inclusion-Forming Reactions” in this Volume.
Fig. 11
Phase diagram for the binary iron-carbon system. Source: Ref 19
The Iron-Carbon System The phase diagram of the binary iron-carbon system includes the stable (iron-graphite) and
metastable (Fe-Fe3C) equilibria. The most recent diagram is shown in Fig. 11. The solubility of carbon in iron is described in Eq 15–20, Eq 21. (Ref 20, 21).
46 / Principles of Liquid Metal Processing For the stable system, that is, with graphite as the equilibrium high-carbon phase, Eq 15–17 are used. For the liquid:
The general equation used to describe the activity of a component 2 in a component 1 according to the Darken formalism (Ref 22) is:
wt%Cmax ¼ 1:3 þ 2:57 103 t
log
(Eq 15)
in the interval 1152 to 2000 C (2106 to 3632 F), with t being the temperature in degrees Celsius. Using thermodynamic quantities, Eq 15 can be written as: log XCmax
12:728 ¼ T þ 0:727 log T 3:049
(Eq 16)
where XCmax is the maximum solubility of graphite in liquid iron in mole fraction, and T is the temperature in degrees Kelvin. Equation 16 is valid for the interval 1425 to 2300 K. For the austenite: wt%Cmax ¼ 0:435 þ 0:355 103 t þ 1:61 106 t2
wt%Cmax
(Eq 18)
(Eq 19)
For the austenite:
(Eq 23)
1180 0:87 T
(Eq 24)
Using the natural logarithm (base e) instead of the common logarithm (base 10), Eq 22 can be written as follows for the iron-carbon systems: 2714 ln gC ¼ 2 T 2920 þ þ 4:01 2XC XC2 T
(Eq 25)
153 aFe-Si ¼ þ 3:02 T
log goSi ¼ 0:914
The Iron-Silicon System
6863 T
(Eq 26)
(Eq 27)
As shown in Fig. 13, calculations with Eq 26 and 27 are in good agreement with experimental data (note that the equation in Fig. 13 results from introducing Eq 26 and 27 into Eq 22 and using a natural logarithm rather than a common logarithm).
Ternary Iron-Base Alloys Using a Taylor series expansion for the excess partial molar free energy, one researcher derived an expression for the activity coefficient of component 2 in a multicomponent solution, which for dilute solutions can be written as (Ref 24): ln f2 ¼ ln
g2 X2 e22 þ X3 e23 þ X4 e24 ; . . . go2
(Eq 28)
where g2 is the activity coefficient of component 2 dissolved in the three-component system, go2 is the activity coefficient of component 2 in the range of infinitely dilute solution, and e22, e23, e24, . . . are coefficients defined as: @ ln f2 @X2 limit @ ln f2 ¼ X1 !1 @X3 limit
e22 ¼ X1 !1 e23
(Eq 20)
For the ferrite: 10; 908 wt%Cmax ¼ 1:8 103 exp T
1270 aFe-C ¼ þ 1:74 T
Equation 25 is in good agreement with experimental data, as shown in Fig. 12.
wt%Cmax ¼ 0:628 þ 1:222 103 t þ 1:045 106 t2
where g2 is the activity coefficient of component 2 = a2/X2 (a2 being the activity of component 2 and X2 the mole fraction of component 2); go2 is the activity coefficient of component 2 in the range of infinitely dilute solutions, that is, when %component 2 ffi 0; it is only a function of temperature (g2 ! go2 as X2 ! 0); and a12 is the function of temperature, usually of the shape A/T + B, where A and B are constants. For the activity of carbon in an iron melt, Eq 22 is valid for component 1 (iron) and component 2 (carbon) with (Ref 23):
log goC ¼
For the metastable system, that is, with Fe3C as the equilibrium high-carbon phase, Eq 19–21 apply. For the liquid: wt%Cmax ¼ 4:34 þ 0:1874ðt 1150Þ t 200 ln 1150
(Eq 22)
(Eq 17)
For the ferrite: 11; 460 ¼ 2:46 103 exp T
g2 ¼ a12 ðX22 2X2 Þ go2
with the following remarks (Ref 23): XSi 0.2; component 1 is iron; component 2 is silicon:
(Eq 29)
Equation 22 can be used to calculate the activity of silicon in the iron-silicon system,
For infinitely dilute solutions, that is, X2 ! 0, X3 ! 0: e23 ¼ e32
Fig. 12
Fig. 13
(Eq 21)
Equation 15–21 can be used to calculate the maximum solubility of carbon in various phases (liquid, austenite, ferrite) as a function of temperature. For example, to calculate the carbon content of the stable eutectic, Eq 15 is used with t = 1152 C (2106 F) to obtain %Cmax = 4.26%. Prediction of phase stability relies heavily on knowledge of the activity of various elements. In the specific case of iron-carbon alloys, an increase in the activity of carbon in the liquid parallels a greater ability of carbon to separate as graphite; that is, the activity of carbon is a measure of an increased tendency of the alloy to solidify as the stable system. On the other hand, a decreased activity of carbon reflects the carbide-promoting behavior of the system, that is, the metastable solidification tendency.
Dependence of the activity coefficient of carbon on the mole fraction of carbon in liquid iron at 1550 C (2820 F). Source: Ref 23
Dependence of the activity coefficient of silicon on the mole fraction of silicon in liquid iron at 1600 C (2910 F). Source: Ref 23
Thermodynamic Properties of Iron-Base Alloys / 47 Simplification of the Darken formalism for infinite dilutions gives the interrelationship between e and a, as follows (Ref 22): e22 ¼ 2a12 2:303 e33 ¼ 2a13 2:303 e23 e32 ¼ ¼ a23 a12 a13 2:303 2:303
ðEq 30Þ
The solute concentration is often given in weight percent. For a ternary system, Eq 28 is written: log fi ¼ eii ½%i þ eij ½%j
(Eq 31)
where the interaction coefficient e is related to E by: eii ¼ 230ðMi =M1 Þeii þ ½ðM1 Mi Þ=M1 eij ¼ 230ðMj =M1 Þeij þ ½ðM1 Mj Þ=M1
log ðEq 32Þ
Data on interaction coefficients in liquid iron for carbon, hydrogen, nitrogen, oxygen, and sulfur are given in Table 2. For metallurgical reactions involving dilute solutions, it is often necessary to change the standard state from pure component to that of 1 wt% in solution. The free energy change is given by: Gs ¼ RT ln
0:5585goi Mi
available in Ref 25. Only the Fe-C-Si, Fe-C-P, and Fe-C-Mn systems are discussed in this article. The Fe-C-Si System. The addition of silicon to a binary iron-carbon alloy decreases the stability of Fe3C, which is already metastable, and increases the stability of ferrite (the a field is enlarged, and the g field is constricted). The equilibrium diagrams in Fig. 14 show that as the silicon content in the Fe-C-Si system increases, the carbon contents of the eutectic and eutectoid decrease (%Cmax decreases), while the eutectic and eutectoid temperatures increase (Ref 26). The activity coefficient, g2, of component 2, dissolved in a three-component system as a function of the concentration X2 and X3 of dissolved components 2 and 3, is described by Eq 34 (Ref 22):
(Eq 33)
Free energies of solution DGs of selected elements in liquid iron are given in Table 3.
g2 ¼ 2a12 X2 þ ða23 a12 a13 ÞX3 go2 þ a12 X22 þ a13 X32 þ ða12 þ a13 a23 ÞX2 X3
ðEq 34Þ
where a12 and a13 are the modified interaction parameters of the two-component systems, a23 is the modified interaction parameter that describes the activity of both dissolved components, and go2 is the activity coefficient of component 2 in the range of infinitely dilute solution. Starting from Eq 34, the activity of carbon and silicon in an Fe-C-Si melt can be calculated as a function of temperature and concentration with the following relationships (Ref 23): log aC ¼ð1180=T Þ 0:87 þ ½ð2540=T Þ þ 3:48 XC
Ternary Fe-C-X Systems
½ð1270=T Þ þ 1:74 XC2 þ log XC
Detailed equilibrium diagrams of the Fe-C-B, Fe-C-Cr, Fe-C-Cu, Fe-C-Mn, Fe-C-Mo, Fe-CN, Fe-C-Ni, Fe-C-Si, and Fe-C-W systems are
þ ½ð1423=T Þ þ 4:75 XSi 2 ½ð153=T Þ þ 3:02 XSi
½ð1423=T Þ þ 4:76 XC XSi
ðEq 35Þ
log aSi ¼ ð6863=T Þ þ 0:914 þ ½ð306=T Þ þ 6:04 XSi 2 ½ð153=T Þ þ 3:02 XSi þ log XSi
þ ½ð1423=T Þ þ 4:76 XC ½ð1270=T Þ þ 1:74 XC2 ½ð1423=T Þ þ 4:76 XC XSi
ðEq 36Þ
The Fe-C-P System. Because the partition coefficient of phosphorus between solid and liquid iron is low, phosphorus segregates into the liquid during solidification, and the melt becomes supersaturated in phosphorus, which results in the formation of Fe3P. Therefore, the four possible solid phases of the stable and metastable systems are austenite (g), Fe3P, Fe3C, and graphite. The calculated liquidus surfaces for the four phases and the eutectic lines formed by their interactions are shown in Fig. 15, which is a superimposition of the g-Fe3P-Fe3C and the g-Fe3P-graphite (Gr) diagrams. Two ternary eutectic points are shown in Fig. 15. To the left is L ! g + Fe3P + Gr at 954 C (1749 F) and yC = 0.099, yP = 0.123. To the right is L ! g + Fe3P + Fe3C at 948 C (1738 F) and yC = 0.106, yP = 0.123. Here, yP = XP/(1 XC), and yC = XC/(1 XC), where XP is the weight fraction of phosphorus and XC is the weight fraction of carbon. In general, it is agreed that phosphorus does not have a strong graphiteor carbide-stabilizing influence. The Fe-C-Mn System. The influence of increasing manganese contents in the Fe-C-Mn system over the equilibrium phase diagram is shown in Fig. 16. It is evident that manganese: Decreases the eutectoid temperature Increases the eutectoid interval, that is, the
range of temperature and carbon contents over which a, g, and carbide can coexist Decreases the carbon content in the eutectoid and in the eutectic Increases the eutectic temperature (by approximately 3 C, or 5 F, for each 1% Mn)
Table 2 Interaction coefficients in ternary iron-base alloys for carbon, hydrogen, nitrogen, oxygen, and sulfur at 1600 C (2910 F) Element j
Aluminum Boron Carbon Cobalt Chromium Copper Manganese Molybdenum Nitrogen Niobium Nickel Oxygen Phosphorus Sulfur Silicon Tin Titanium Vanadium Tungsten Zirconium Source: Ref 4, 7
eCj
eHj
eNj
eOj
eSj
0.043 0.24 0.14 0.008 0.024 0.016 0.012 0.008 0.11 0.06 0.012 0.34 0.051 0.046 0.08 0.041 ... 0.077 0.006 ...
0.013 0.05 0.06 0.002 0.002 <0.001 0.001 0.002 ... 0.002 0 0.19 0.011 0.008 0.027 0.005 0.019 0.007 0.005 ...
0.028 0.094 0.13 0.011 0.047 0.009 0.02 0.011 0 0.06 0.01 0.05 0.045 0.007 0.047 0.007 0.53 0.093 0.002 0.63
3.9 2.6 0.13 0.008 0.04 0.013 0.021 0.004 0.057 0.14 0.006 0.20 0.07 0.133 0.131 0.011 0.31 0.14 0.009 (3.0)
0.035 0.13 0.11 0.003 0.011 0.008 0.026 0.003 0.01 0.013 0 0.27 0.29 0.028 0.063 0.004 0.072 0.016 0.01 0.052
Table 3 Free energies of solution in liquid iron at 1 wt% Mi, liquid iron at 1600 C (2910 F) Element i
Aluminum (l) Carbon (g) Cobalt (l) Chromium (s) Copper (l) ½ H2(g) Manganese (l) ½N2 (g) Nickel (l) ½O2 (g) ½P2(g) ½ S2(g) Silicon (l) Titanium (s) Vanadium (s) Tungsten (s) Zirconium (s)
goi
D Gs (cal g atom1)
0.029 0.57 1.07 1.14 8.6 ... 1.3 ... 0.66 ... ... ... 0.0013 0.038 0.1 1.2 0.043
15,100 6.67T 5400 10.10T 240 9.26T 4600 11.20T 8000 9.41T 8720 + 7.28T 976 9.12T 860 + 5.71T 5000 7.42T 28,000 0.69T 29,200 4.60T 32,280 + 5.60T 31,430 4.12T 7440 10.75T 4950 10.90T 7500 15.20T 8300 11.95T
48 / Principles of Liquid Metal Processing
Fig. 14
Influence of silicon content on the solubility lines and equilibrium temperatures of the iron-carbon system. (a) to (c) Source: Ref 26, (d) to (g) Source: Ref 25
The carbide M3C, which is (FeMn)3C, is stable over a wide range of manganese and iron contents, extending at 1000 C (1832 F) from Mn3C to nearly pure Fe3C. At high manganese contents (for example, above 40% Mn), other manganese carbides such as M5C2, M7C3, M23C6, and M4C1+x(e) can form (Ref 25). Influence of a Third Element on the Equilibrium Temperatures. Various elements, dissolving in liquid and solid iron phases, change the equilibrium temperatures on the iron-carbon diagram, as indicated in Table 4. This results in increased or decreased a and g fields.
Of particular interest for the case of cast iron is the influence of third elements on the stablemetastable eutectic interval (Tst Tmet). In general, elements that increase the Tst Tmet interval promote graphite formation, while those that decrease the interval promote carbide formation. Based on their specific influence on the Tst Tmet interval, third elements can be classified into the following four groups (Fig. 17): Strong graphitizers that increase Tst and
decrease Tmet, such as silicon, aluminum, nickel, and copper (Fig. 17a)
Weak graphitizers that decrease both Tst and
Tmet but increase the Tst Tmet interval overall, such as phosphorus and arsenic (Fig. 17b) Strong carbide stabilizers that decrease Tst but increase Tmet, such as chromium, vanadium, and manganese (Fig. 17c) Weak carbide stabilizers that decrease both Tst and Tmet, such as molybdenum and tungsten (Fig. 17d) st ¼ Tst Tmet interval Real values for the Tmet for a number of elements can be calculated with
Thermodynamic Properties of Iron-Base Alloys / 49 the data given in Table 4, which also includes data on the influence of some elements on the eutectoid temperature and on the maximum solubility of carbon in austenite. Two examples of the use of data in Table 4 are provided as follows.
having 2% Si, 0.5% Mn, and 1% Cu by weight: Tst ¼ 1154 C þ 4% Si 2% Mn
Example 1: Calculation of the value of the Tst Tmet eutectic interval for a cast iron
TFe-C-Mn-Cu ¼ 1166 1117:2 ¼ 48:8 C
þ 5% Cu ¼ 1166 C Tmet ¼ 1148 C 15% Si þ 3% Mn 2:3% Cu ¼ 117:2 C TFe- C 1154 1148 ¼ 6 C
Fig. 15
Calculated liquidus surfaces in the Fe-C-P phase diagram. Source: Ref 27
Fig. 16
Vertical sections through the Fe-C-Mn ternary phase diagram at (a) 4.92, (b) 12.8, and (c) 19.7 wt% Mn
Example 2: Calculation of the temperatures of the point of maximum carbon solubility in austenite, TE and TE0 , for the aforementioned iron: TE ¼ 1154 C 10%Si þ 3:2% Mn 2%Cu ¼ 1133:5 C TE0 ¼ 1148 C þ 2:5% Si 2% Mn þ 5:2%Cu ¼ 1157:2 C Some typical experimental results for silicon, chromium, and vanadium are shown in Fig. 18. These results are in agreement with theoretical predictions. Partition of a Third Element Between Various Phases in the Fe-C-X System. It has been established from the thermodynamic stability of phases involved in the eutectic solidification of Fe-C-X alloy that the differences in the equilibrium partition coefficients DP (= PM PS, where PM and PS are the partition coefficients of element X between liquid and eutectic metastable or stable solid), the partition of element X between austenite g=L and liquid, PX , and the partition of element X Fe C=g between cementite and austenite, PX 3 , are related to one another and have similar effects on graphitization during eutectic solidification (Ref 28). Calculated and experimental data for partition coefficients of various elements in phases of iron-carbon alloys are given in Table 5. Figure 19 illustrates the correlation between the graphitizing influence of a third element (expressed by the normalized addition of each element required to increase or decrease the chill depth by a given amount) and the DP value. It is quite obvious that a definite correlation exists, although silicon, aluminum, and titanium deviate from the line. These three elements have a strong affinity for nitrogen; this may explain their higher-than-expected graphitizing influence. Indeed, it is accepted that the nitrides may act as
50 / Principles of Liquid Metal Processing nuclei for the stable eutectic, promoting gray solidification. This is why experiments show that titanium, for example, can act either as a carbide or graphite promoter, depending on the nitrogen content. Influence of a Third Element on Carbon Solubility in the Fe-C-X System. The influence of a third element on the solubility of carbon can be expressed by (Ref 20): XCX
¼
XCXmax
XCmax ðmole fractionÞ
ðEq 37aÞ
concentrations. In the region of low concentration of the third element (%X = 0 to 5%), the change in the solubility of carbon can be represented by the temperature-independent linear equation: XCX ¼ m XX
ðEq 38aÞ
or
%CEX ¼ m0 g %X 0
%C ¼ m %X X
ðEq 38bÞ
%C ¼
X %Cmax
%Cmax %Cmax ðwt%Þ
ðEq 37bÞ
where XCX and D%CX represent the increase or decrease of carbon solubility in mole fraction or weight percent, respectively; XCmax and % Cmax represent the saturation concentration in the iron-carbon system calculated from Eq 15 or 16; and XCXmax and %CX max are the saturation concentration in the Fe-C-X system. The changes in carbon solubility resulting from additions of third elements, X, are shown in Fig. 20. These data are valid in the range from 1200 to 1700 C (2190 to 3090 F), with the exception of silicon and sulfur at high
where m and m are solubility factors. In Eq 38a and 38b, m and m0 as well as XCX and D% CX are positive for an increase in the solubility of carbon and negative in the other case. It is also generally accepted that carbide-promoting elements increase the solubility of carbon (activity coefficient of carbon in solution is decreased), while graphite-promoting elements decrease it (activity coefficient of carbon in solution is increased). Therefore, third elements that have a negative solubility factor promote graphitization, while elements that have a positive solubility factor promote carbide formation. The value of the solubility factor is proportional to their effect. It must be noted, however, that the accuracy of the prediction of the behavior of elements based on solubility factor is questioned in
Table 4 Influence of a third element on the temperature change of critical points on the iron-carbon diagram +, increase; , decrease. Maximum solubility of carbon in austenite, C/wt%X Element
Metastable (E)
Silicon Copper Aluminum Nickel Cobalt Chromium Manganese Molybdenum Tungsten Boron Nickel Titanium Vanadium Phosphorus Sulfur
10 to 15 2 14 4.8 ... +7.3 +3.2 ... ... ... +6–8 180 ...
0
Stable (E )
+2.5 +5.2 +8 +4 ... 2 ... ... ... ... 180 ...
ðEq 39Þ
Data for m0 g are also given in Table 6.
0
X
Ref 28 for elements that have a strong influence on the austenite liquidus. As indicated in Table 6, experimental values are close to the theoretical values for many elements. Equations similar to Eq 38a and 38b can be used to calculate the change in the maximum solubility of carbon in austenite upon the addition of a third element:
Multicomponent Iron-Carbon Systems Carbon Solubility in Multicomponent Systems. Assuming that the values in Table 6 for m and m0 , which were determined for the ternary Fe-C-X system, can be extended to multicomponent systems (in this case, the effect of alloying elements on the solubility of carbon can be considered to be additive, as shown in Eq 41, at least in the region of low concentrations), the saturation concentration of carbon in multicomponent systems can be calculated as (Ref 20): %CSi;Mn;P;S;... ¼ %Cmax þ %CSi;Mn;P;S;... max
ðEq 40Þ
which is Eq 37b rewritten for multicomponent systems. In turn: %CSi;Mn;P;S;... ¼ %CSi þ %CMn þ %CP max
Eutectoid, C/wt%X
Eutectic, C/wt%X 0
Metastable (S) Stable (S ) Metastable C Stable C
+8 ... +10 20 ... +15 9.5 + + ... ... ... +15 + ...
0–30 10 +10 30 ... +8 3.5 + + ... ... ... + +6 ...
10 to 20 2.3 15 6 ... +7 +3 ... ... ... 6–8 37 ...
+4 +5 +8 +4 ... 2 ... ... ... 30 ...
þ %CS þ . . .
Eutectic, C/at. %X(a) 0
Metastable C Stable C
3.25 3.32 7.74 1.33 2.23 4.97 2.33 8.90 12.13 15.47 19.67 16.91 ... 16.05 .. . .
ðEq 41Þ
0
10.91 8.41 0.91 6.50 1.48 10.23 7.15 12.36 15.55 18.53 16.87 18.91 ... 16.98 18.53
(a )Calculated results. Source: Ref 21, 28
Then, using Eq 15 and data from Table 6, Eq 40 becomes: %CSi;Mn;P;S;... ¼ 1:3 þ 2:57 103 t 0:31% Si max þ 0:027% Mn 0:33% P 0:4% S . . .
ðEq 42Þ
Saturation Degree and Carbon Equivalent. The iron-carbon diagram with letter notations for critical points (Fig. 21) is used in this section. It is well known that mechanical properties of cast iron strongly depend on the amount and shape of graphite. In turn, for hypoeutectic gray irons, the amount of graphite depends on the amount of eutectic. The amount of eutectic can be calculated with the lever rule. For hypoeutectic irons, one can write: Amount of eutectic Amount of eutectic þ amount of austenite %Canal %C0E ¼ %Canal %C0E þ %C0C %Canal %Canal %C0E ¼ ðEq 43Þ %C0C %C0E
Sr ¼
Fig. 17
Classification of the influence of a third element in the Fe-C-X system on the graphite- or carbide-promoting tendency in cast iron, based on their influence on the eutectic stable and metastable temperatures. (a) Strong graphitizers. (b) Weak graphitizers. (c) Strong carbide stabilizers. (d) Weak carbide stabilizers
where Sr is the rectified saturation degree, % Canal is the analyzed carbon content of cast iron, and %C0 C and %C0 E represent the carbon content
Thermodynamic Properties of Iron-Base Alloys / 51
Fig. 18
Influence of (a) chromium, (b) silicon, and (c)
Table 5
Equilibrium partition coefficients of a third element X in the Fe-C-X system g=L
Fe C=L
PX (a)
PX 3
(a)
Element X
calc
exp
calc
exp
PS
PM
DP = PM PS
Silicon Copper Aluminum Nickel Cobalt Chromium Manganese Molybdenum Tungsten Boron Nitrogen Titanium Phosphorus Sulfur
1.71 1.57 ... 1.46 1.18 0.53 0.70 0.41 0.23 0.06 2.04 0.04 0.15 0.06
1.72 1.62 1.15 1.61 1.13 0.55 0.75 0.38 0.42 ... 2.04 ... ... ...
0 0.12 ... 0.43 0.59 1.96 1.03 0.60 0.42 0.22 2.12 0.09 0.08 ...
0.05 0.08 0.03 0.32 0.60 1.95 1.21 0.84 0.88 ... ... 0.27 0.09 ...
1.55 1.43 1.05 1.33 1.07 0.48 0.64 0.37 0.21 0.06 1.86 0.04 0.14 0.06
0.78 0.78 0.55 0.90 0.85 1.32 0.90 0.52 0.33 0.15 2.09 0.07 0.11 ...
0.77 0.65 0.50 0.43 0.21 0.84 0.26 0.15 0.12 0.09 0.23 0.03 0.03 ...
(a) calc, calculated; exp, experimental. Source: Ref 28
temperature of the multicomponent system. Further simplification gives: SC ¼
Fig. 19
Correlation between the partition coefficient, DP, and the graphitizing influence of a third element in the Fe-C-X system. Source: Ref 28
of the eutectic and of the austenite, respectively, at the eutectic temperature in the multicomponent system (Fig. 21). Assuming the solubility factors can be extrapolated to the hypoeutectic region, %C0 C can be calculated with Eq. 42. A similar approach can be used to calculate %C0 E using data in Table 6 for m0 g: %C0E ¼ 2:11 0:11% Si 0:35%P þ 0:006% Mn 0:08% S
ðEq 44Þ
In foundry practice, the following simplified form of Eq 43 is used: SC ¼ ¼
%Canal %C0C %Canal 4:26 0:31% Si 0:33% P þ . . .
ðEq 45Þ
For more accuracy, %Cmax as a function of temperature rather than %Cmax = 4.26 can be used, where the temperature is the eutectic
%Canal 4:25 0:3ðSi þ PÞ
ðEq 46Þ
where SC < 1 for hypoeutectic irons, SC = 1 for eutectic irons, and SC > 1 for hypereutectic irons. Saturation degree is used in most of the European literature. In the Anglo-Saxon literature and foundry practice, carbon equivalent rather than saturation degree is used. Carbon equivalent (CE) can be calculated as: CE ¼ %Canal %CSi %CMn %CP %CS . . . ¼ %Canal þ 0:31%Si þ 0:33% P 0:027% Mn þ 0:4% S
ðEq 47Þ
A eutectic iron has CE = 4.26%. It must be noted that although Sr gives directly the amount of eutectic in the structure (for example, Sr = 0.9 means 90% eutectic), SC and CE, although easier to calculate, do not allow for exact estimation of the amount of eutectic.
Structural Diagrams For the practicing metallurgist involved in cast iron production, one of the main applications of the thermodynamics of the iron-carbon system is the calculation of structure-composition
correlations. The so-called structural diagrams are used for this purpose. The first and simplest structural diagram is the Maurer diagram (Fig. 22a), in which the as-cast structure is considered to be solely the product of the carbon and silicon contents of cast iron. Because it was recognized long ago that solidification rate, as well as chemical composition, plays a significant role in the formation of the as-cast structure of casting, especially for cast iron, Laplanche (Ref 29) further developed the Maurer diagram to include the influence of cooling rate through the section size (or bar diameter) of the casting (Fig. 22b). The k lines on the diagram are lines of same structure (same degree of graphitization) at a given cooling rate. The value k is an empirical function of composition, as follows: k ¼ 4=3 Si 1
5 3C þ Si
ðEq 48Þ
The correlation between k and cooling rate is given in Table 7. Thus, if one wants to produce a pearlitic iron in casting with equivalent section sizes of 10 to 30 mm (0.4 to 1.2 in.) bar diameters, Table 7 would suggest a k of 0.85 to 2.35. For example, k = 1.8 is then chosen. A saturation degree, SC, is also selected depending on the required mechanical properties (for example, SC = 0.8), and the intersection of the k = 1.8 with SC = 0.8 will then give the carbon and silicon required for the iron.
52 / Principles of Liquid Metal Processing A further contribution to structural diagrams was made by Patterson and Doepp (Ref 30). This diagram (Fig. 23) encompasses not only structure, composition, and cooling rate but also a broad range of mechanical properties. Part of this diagram is based on Laplanche’s k lines. As an example of the way this diagram is used, assume it is desired to pour a 50 mm (2 in.) diameter cast iron bar (or equivalent section size) with a pearlitic-ferritic structure. A heavy horizontal line is drawn from the y-axis at the 50 mm (2 in.) diameter bar location to the middle of zone IIb. This corresponds to a determination of k = 2.2. This k line can be followed to the lower diagram until it intersects a preselected saturation degree—in this case, SC = 0.6. The composition is now read as C = 2.6% and Si = 3%, together with the equilibrium tem! perature for the ðSiO2 Þ þ 2½C ½Si þ fCOg reaction, which is 1460 C (2660 F). An additional explanation of the significance of this reaction can be found in the article “Solidification of Eutectic Alloys: Cast Iron” in this Volume. It must be pointed out that foundry variables such as superheating and holding during melting, charge composition (raw materials), and inoculation can cause significant deviation from these calculations.
ACKNOWLEDGMENTS Reviewed and adapted from Seymour Katz, Composition Control, Casting, Vol 15, Metals Handbook, 9th Edition, ASM International, 1988, p71–82; and D.M. Stefanescu, Thermodynamic Properties of Iron-Base Alloys, Casting, Vol 15, Metals Handbook, 9th Edition, ASM International, 1988, p 61–70 REFERENCES 1. R.J. Fruehan, Ladle Metallurgy: Principles and Practices, Iron and Steel Society of AIME, 1985, p 7–8 2. S. Katz and C.F. Landefeld, Cupola Desulfurization, Cupola Handbook, American Foundrymen’s Society, 1984, p 351–363 3. I.D. Sommerville, The Capacities and Refining Capabilities of Metallurgical Slags, Foundry Processes: Their Chemistry and Physics, S. Katz and C.F. Landefeld, Ed., Plenum Press, 1988, p 101–133 4. G.K. Sigworth and J.F. Elliott, The Thermodynamics of Liquid Dilute Iron Alloys, Met. Sci., Vol 8, 1974, p 298–310 5. S. Katz and H.C. Rezeau, The Cupola Desulfurization Process, Trans. AFS, Vol 87, 1979, p 367–376 6. S. Katz, D.E. McInnis, D.L. Brink, and G.A. Wilkinson, The Determination of Aluminum in Malleable Iron from Measured Oxygen, Trans. AFS, Vol 88, 1980, p 835–844 7. E.T. Turkdogan, Physical Chemistry of High Temperature Technology, Academic Press, 1980 Table 6
Solubility factors of various third elements for carbon-saturated Fe-C-X melts
Influence of third elements, X, on the solubility of carbon in molten iron. (a) In mole fraction. (b) In weight percent. Source: Ref 20
Fig. 20
5 6 13 14 15 16 22 23 24 25 26 27 28 29 32 33 41 42 44 50 51 52 73 74 76 78 79 92
Symbol
B C Al Si P S Ti V Cr Mn Fe Co Ni Cu Ge As Nb Mo Ru Sn Sb Te Ta W Os Pt Au U
D %CX = m0 %X
DXX C ¼ m XX (a)
Third element X Atomic No.
8. S. Katz and V.R. Spironello, Effect of Charged Aluminum on Iron Temperature, Silicon Recovery and Desulfurization in an Iron-Producing Cupola, Trans. AFS, Vol 92, 1984, p 161–172 9. C.F. Landefeld and S. Katz, Kinetics of Iron Desulfurization by CaO-CaF2, Proceedings of the Fifth International Iron and Steel Congress, Vol 6, Iron and Steel Society of AIME, 1986, p 429–439 10. S. Katz and C.F. Landefeld, Plant Studies of Continuous Desulfurization with CaOCaF2-C, Trans. AFS, Vol 93, 1985, p 215–228 11. S. Katz and B.L. Tiwari, A Critical Overview of Liquid Metal Processing in the Foundry, Foundry Processes: Their Chemistry and Physics, S. Katz and C.F. Landefeld, Ed., Plenum Press, 1988, p 1–52 12. J. Niederinghaus and R.J. Fruehan, Desulfurization Mechanisms for CaO-Al and CaO-CaS in Carbon Saturated Iron, Metall. Trans. B, to be published 13. E.T. Turkdogan, Physiochemical Phenomena of Mechanisms and Rates of Reaction in Melting, Refining and Casting of Foundry Irons, Foundry Processes: Their Chemistry and Physics, S. Katz and C.F. Landefeld, Ed., Plenum Press, 1988, p 53–100 14. S. Asai and I. Muchi, Fluid Flow and Mass Transfer in Gas Stirred Ladles, Foundry Processes: Their Chemistry and Physics, S. Katz and C.F. Landefeld, Ed., Plenum Press, 1988, p 261–292 15. S.-H. Kim and R.J. Fruehan, Physical Modelling of Liquid/Liquid Mass Transfer in Gas Stirred Ladles, Metall. Trans. B, Vol 18, 1987, p 381–390
0 D%CX C ¼ mg %X
m exp
m calc
Validity
m exp
m calc
Validity
m0 g
Validity
0.575 0.67 0.52 0.71 0.81 1.00 +0.508 +0.359 +0.215 +0.105 0.0 0.10 0.20 0.313 0.677 0.734 +0.51 +0.24 0.0 0.70 0.79 0.85 +0.49 +0.256 0.0 0.224 0.333 +0.53
0.51 0.68 0.51 0.68 0.85 1.02 +0.44 +0.33 +0.22 +0.11 0.0 0.11 0.22 0.33 0.66 0.77 +0.33 +0.22 0.0 0.66 0.77 0.88 +0.33 +0.22 0.0 0.22 0.33 +0.55
<0.04 ... <0.05 <0.007 <0.048 <0.004 <0.013 <0.1 <0.13 <0.30 ... <0.20 <0.10 <0.02 <0.1 <0.07 <0.045 <0.14 <0.008 <0.02 <0.05 <0.005 <0.03 <0.05 <0.013 <0.015 <0.024 <0.01
0.539 0.610 0.220 0.310 0.331 0.405 +0.159 +0.105 +0.064 +0.028 0.0 0.027 0.051 0.076 0.144 0.151 +0.058 +0.014 0.020 0.110 0.119 0.122 +0.004 0.015 0.035 0.052 0.060 0.006
0.462 0.622 0.215 0.294 0.349 0.414 +0.138 +0.097 +0.062 +0.029 0.0 0.029 0.055 0.080 0.141 0.159 +0.030 +0.012 0.023 0.104 0.117 0.126 0.009 0.01 0.035 0.052 0.060 0.005
<1.0 ... <2.0 <5.5 <3.0 <0.4 <1.6 <11.0 <10.0 <25.0 ... <40.0 <8.0 <3.8 <14.0 <10.0 <9.0 <2.0 <1.5 <4.5 <15.0 <1.5 <11.0 <18.0 <4.0 <6.0 <10.0 <5.0
... ... +0.04 0.11 0.35 0.08 2.08 +0.18 0.07 +0.006 ... 0.017 0.07 +0.014 ... ... ... 0.3 ... ... ... ... ... 0.12 ... ... ... ...
... ... <2 <6 <0.4 <5 <0.8 <1 <20 <60 ... <20 <6 <5 ... ... ... <4 ... ... ... ... ... <12 ... ... ... ...
(a) exp, experimental; calc, calculated. Source: Ref 20
0
Thermodynamic Properties of Iron-Base Alloys / 53 16. J. Ishida, K. Yamaguchi, S. Sugiura, N. Demukai, and A. Notoh, Denki Seiko, Vol 52, 1981, p 2–8 17. E.T. Turkdogan, Ladle Deoxidation, Desulfurization and Inclusions in Steel—Part I: Fundamentals, Arch. Eisenhu¨ttenwes., Vol 54, 1983, p 1–10 18. E.T. Turkdogan, Slags and Fluxes for Ferrous Ladle Metallurgy, Ironmaking Steelmaking, Vol 12, 1985, p 64–78 19. T.B. Massalski et al., Ed., Binary Alloy Phase Diagrams, Vol 1, American Society for Metals, 1986 20. F. Neumann, The Influence of Additional Elements on the Physico-Chemical Behavior of Carbon in Carbon Saturated Molten Iron, Recent Research on Cast Iron, H.D. Merchant, Ed., Gordon and Breach, 1968, p 659
Fig. 21
21. N.G. Girsovitch, Ed., Spravotchnik po tchugunomu litja (Cast Iron Handbook), Mashinostrojenie, 1978 22. L.S. Darken, The Thermodynamics of Ternary Metallic Solutions, Trans. TMS, Vol 239, 1967, p 80, 90 23. F. Neumann and E. Do¨tsch, Thermodynamics of Fe-C-Si Melts with Particular Emphasis on the Oxidation Behavior of Carbon and Silicon, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 31 24. C. Wagner, Thermodynamics of Alloys, Addison-Wesley, 1952 25. Metallography, Structures and Phase Diagrams, Vol 8, Metals Handbook, 8th ed., American Society for Metals, 1973, p 400–416 26. E. Piwowarsky, Hochwertiger Gusseisen, 2nd ed., Springer-Verlag, 1958
The iron-carbon diagram with letter notations. S.S., soild solution
27. M. Hillert and P.O. So¨derholm, White and Gray Solidification of the Fe-C-P Eutectic, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 197 28. A. Kagawa and T. Okamoto, Partition of Alloying Elements on Eutectic Solidification of Cast Iron, The Physical Metallurgy of Cast Iron, H. Fredriksson and M. Hillert, Ed., North-Holland, 1985, p 201 29. H. Laplanche, The Maurer Diagram and Its Evolution and a New Structural Diagram for Cast Iron, Foundry Trade J., No. 1669–1671, 1948, p 191, 225, 249 30. W. Patterson and R. Doepp, Giessereiforschung, Vol 21 (No. 2), 1969, p 91
54 / Principles of Liquid Metal Processing
Fig. 22
Structural diagrams for cast iron. (a) Maurer diagram. (b) Laplanche diagram. See also Fig. 23. Source: Ref 29
Table 7
k lines for various structural ranges
Bar diameter mm
in.
k for mottled structure
k for pearlitic structure
k for pearlitic-ferritic structure
30 20 10
1.2 0.8 0.4
0.65–0.85 0.75–1.10 1.05–1.50
0.85–2.05 1.10–2.25 1.50–2.35
2.05–3.10 2.25–3.40 2.235–3.50
Thermodynamic Properties of Iron-Base Alloys / 55
Fig. 23
Patterson and Doepp structural diagram for cast iron. See the corresponding text for details. Source: Ref 30
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 56-63 DOI: 10.1361/asmhba0005191
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys* THIS ARTICLE provides accessible information on the thermodynamic properties of liquid aluminum-base and copper-base alloys. Three alternative means have been used to report the thermodynamic data: Activities and activity coefficients Partial and integral molar thermal properties Interaction coefficients
symbol Gxsi . It represents the contribution to Gi resulting from the departure of solution from ideal behavior. Thus, one may write: xs
Gi ¼ RT ln gi xs Gi
can also be
Gi ¼ H i T S i ¼ RT ln gi
(Eq 4)
Corresponding to Eq 1, expressed as: xs
In addition, the utility of phase diagrams is briefly discussed. Knowledge of activities and activity coefficients is necessary in describing solution behavior and in solving problems that involve chemical equilibria. The thermal properties are useful in understanding the liquid state and in correlating data on solution behavior. The interaction coefficients provide a simple means of calculating activity coefficients in dilute solutions and are also used to correlate experimental data on dilute solutions.
Activities and Thermal Properties The activity a of constituent i in a solution is given by the relationship: Gi ¼ H i T S i ¼ RT ln ai
(Eq 1)
where Gi , H i , and Si are the partial molar free energy, enthalpy, and entropy, respectively, of constituent i in solution relative to the pure constituent i at the solution temperature as the standard state; T is the temperature (in degrees Kelvin); and R is the universal gas constant. Thus, from the knowledge of H i and Si as functions of temperature and solution composition, one may calculate Gi , ai, and gi (activity coefficient of component i relative to pure material standard state) for all conditions of solution. It is important to realize, however, that thermal property values vary with solution composition, unlike the properties of pure substances. Replacing ai in Eq 1 with Xi gi yields: Gi ¼ RT ln Xi þ RT ln gi
(Eq 2)
where Xi is the mole fraction of component i. In Eq 2, the term RT ln gi is called partial molar excess free energy and is expressed by the
(Eq 3)
xs
xs
However, because H i for the ideal solution is 0 by definition, H xsi and are H i are identical. Therefore, Eq 4 can be rewritten as: xs
xs
Gi ¼ H i T S i ¼ RT ln gi
(Eq 5)
Equations 1 and 5 are commonly used to calculate activities and activity coefficients from the thermal properties of the solution. Although the partial molar thermal properties of individual constituents in a solution are most useful in solving chemical process problems, the integral properties are valuable in the discussion of the nature of solutions and also serve in some cases as the source of data on the partial properties. The integral molar free energy is related to the partial molar free energy by the following relationship: Gi ¼ Xi Gi þ X j Gj þ . . . Xn Gn
(Eq 6)
Similar equations can also be written to express other integral thermal properties. The following tables compile, thermodynamic, data in the form of activities, activity coefficients, partial molar thermal properties, and integral molar properties for selected aluminum-base and copperbase alloys. The alloy systems have been chosen based on their commercial importance and on the concentration of the major alloying elements. In general, the concentrations of the alloying elements in the commercial alloys are below 10 at.%, and under this condition the interaction coefficient method described as follows provides a convenient means of calculating the activities of the constituents in solution. When the concentrations are close to or greater than 10 at.%, the interaction coefficient method does not yield accurate values. With this factor in mind, thermodynamic data have
been selected for the following alloy systems: aluminum-magnesium, aluminum-silicon, copperaluminum, copper-nickel, copper-tin, and copper-zinc. Aluminum-magnesium thermodynamic data:
Tables 1 and 2
Aluminum-silicon
thermodynamic data: Tables 3 and 4 Copper-aluminum thermodynamic data: Tables 5 and 6 Copper-nickel thermodynamic data: Tables 7 and 8 Copper-tin thermodynamic data: Tables 9 to 17 Copper-zinc thermodynamic data: Tables 18 and 19
Interaction Coefficients The use of interaction coefficients, first suggested by Wagner (Ref 4), provides a convenient means of organizing thermodynamic data on dilute solutions. Wagner proposed a Taylor series expansion for the excess partial molar free energy in order to express the logarithm of the activity coefficient of a dilute constituent in a multicomponent solution. The values of the activity coefficients calculated from this method are as accurate as the original data from which the interaction coefficients are determined, provided it is applied to solutions that are quite dilute (Xi < 0.1). Consider a dilute solution of i, j, and k dissolved in a common solvent, s. If all but the first-order terms of the Taylor series expansion are neglected, the activity coefficient of solute i is expressed by the relationship: @ ln gi @Xi @ ln gi @ ln gi þ Xk þ Xj @Xj @Xk
ln gi ¼ ln goi ¼ Xi
(Eq 7)
In Eq 7, the partial derivatives are called interaction coefficients and are expressed by the symbol eji , where the superscript denotes the constituent that affects the activity coefficient of the subscript constituent. Thus:
*Updated from B.L. Tiwari, Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 55–60.
Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys / 57 Table 2 Integral thermal properties for liquid aluminum-magnesium alloys at 1073 K
Table 1 Activities, activity coefficients, and partial molar thermal properties for liquid aluminum-magnesium alloys at 1073 K Al component Alð‘Þ ¼ Alðin alloyÞð‘Þ XAl
aAl
0.9997 0.9992 0.9989 0.9945 0.9886 0.9785 0.9553 0.8870 0.8095 0.6600 0.4175 0.2510 0.1792 0.0450
gAl
DGAl , cal/g-atom
DSAl , cal/K-g-atom
0.9 1.7 2.3 12.8 26.6 45.0 93.5 251.6 473.7 934.7 1893.4 2874.9 3839.5 7988.8
0.0091 0.0016 0.0105 0.0203 0.0417 0.0420 0.0958 0.2157 0.3790 0.7806 1.1548 1.2240 0.2920 3.8840
DGAl , cal/g-atom
0.2 0.0 0.0 0.9 1.9 1.3 4.3 4.3 22.7 48.5 30.3 74.0 173.4 1373.6
0.9996 0.9992 0.9989 0.9940 0.9876 0.9791 0.9571 0.8887 0.8008 0.6451 0.4115 0.2597 0.1652 0.236
0.9999 1.0000 1.0000 0.9996 0.9991 1.0006 1.0020 1.0020 0.9894 0.9775 0.9859 1.0353 0.9219 0.5251
9 0 9 9 18 0 9 20 67 97 654 1561 4153 12157
aMg
gMg
DGMg, cal/g-atom
DSMg, cal/K-g-atom
DHMg , cal/g-atom
DGxs Mg, cal/g-atom
0.00022 0.00063 0.00081 0.0048 0.0107 0.0183 0.0366 0.0929 0.1679 0.3097 0.5176 0.6367 0.7184 0.9198
0.8060 0.7974 0.7814 0.8780 0.9353 0.8535 0.8188 0.8221 0.8813 0.9109 0.8885 0.8500 0.8752 0.9631
18,017 15,696 15,171 11,358 9683 8530 7051 5069 3806 2500 1405 963 705 178
15.4180 13.9463 12.6176 9.6420 7.4045 6.9708 5.4069 3.6123 2.7957 1.6331 1.2733 1.3102 1.7577 1.1672
1470.8 729.8 1630.2 1010.9 1736.7 1049.0 1248.7 1192.1 805.9 747.6 38.2 443.2 1180.8 1074.3
459.73 482.71 525.87 227.42 142.53 337.68 424.12 418.20 269.73 198.80 251.88 346.49 284.20 79.98
Mg component Mgð‘Þ ¼ Mg ðin alloyÞð‘Þ XMg 0.00027 0.00079 0.00104 0.0055 0.0114 0.0215 0.0447 0.1130 0.1905 0.3400 0.5825 0.7490 0.8208 0.9550
XMg
DG, cal/g-atom
DS, cal/K-g-atom
DH, cal/g-atom
DGxs, cal/g-atom
0.00027 0.00079 0.00104 0.0055 0.0114 0.0215 0.0447 0.1130 0.1905 0.3400 0.5825 0.7490 0.8208 0.9550
5.7 14.2 18.1 75.7 136.9 227.3 404.8 796.1 1108.7 1466.9 1608.9 1442.6 1266.9 529.7
0.01320 0.01274 0.02365 0.07366 0.12577 0.19086 0.33346 0.59962 0.83951 1.07043 1.22383 1.28856 1.39046 0.93986
8.61 0.58 7.29 3.34 2.04 22.54 47.28 152.48 207.80 318.19 295.28 59.77 225.25 478.87
0.34 0.39 0.55 2.39 3.53 6.00 14.91 43.49 69.79 99.61 159.37 240.99 264.35 138.20
Source: Ref 1
Table 4 Integral thermal properties for liquid aluminum-silicon alloys at 1100 K ð1 xÞAlð‘Þ þ xSið‘Þ ¼ Alð1xÞ Sixð‘Þ XSi
0.1 0.2 0.3 0.31(a)
DG, cal/g-atom
DGxs, cal/g-atom
1210 1931 2403 2430 (þ100)
499 837 1067 1076 (þ100)
(a) Phase boundary. Source: Ref 2
Source: Ref 1
If the solution of i in s obeys Henry’s law over a small range of concentration, then gi is a constant over this range, and eii ¼ 0. It can be shown by using the Gibbs-Duhem equation that:
Table 3 Activities, activity coefficients, and partial molar thermal properties for liquid aluminum-silicon alloys at 1100 K Al component Alð‘Þ ¼ Alðin alloyÞð‘Þ XAl
aAl
1.00 0.90 0.80 0.70 0.69(a)
ð1 xÞAlð‘Þ þ xMgð‘Þ ¼ Alð1xÞ Mgxð‘Þ
xs
DHAl , cal/g-atom
gAl
1.000 0.859 0.698 0.503 0.473 (þ0.01)
xs
DGAl , cal/g-atom
DGAl , cal/g-atom
0 333 786 1500 1638 (þ50)
0 102 298 720 826 (þ50)
1.000 0.954 0.872 0.719 0.685 (þ0.02)
Si component Sið‘Þ ¼ Siðin alloyÞð‘Þ DGSi, cal/g-atom
XSi
aSi
gSi
0.00 0.10 0.20 0.30 0.31(a)
0.000 0.016 0.051 0.127 0.147 (þ 0.02)
0.040 0.155 0.255 0.424 0.473 (þ 0.60)
1 9106 6510 4509 4194 (þ 300)
xs DGSi , cal/g-atom
7060 4073 2991 1877 1634 (þ 300)
DHSi, cal/g-atom 2502 ... ... ... ... (þ 50)
DSSi, cal/ K-g-atom 1 ... ... ... ...
xs DSSi , cal/ K-g-atom
4.144 ... ... ... ...
(a) Phase boundary. Source: Ref 2
@ ln gi Xk ¼ 0; Xi ! 0 @Xi Xj @ ln gi ¼ 0; Xi ; Xj ! 0 eji ¼ @Xj Xk
In Eq 8, goi is the limiting value of gi at infinite dilution:
eji ¼
ln gi ¼ ln goi þ Xj eii þ
Xj eji
þ Xk eki
goi ¼ lim gi (Eq 8)
Xi ! 0
(Eq 9)
eji ¼ eij
(Eq 10)
and this reciprocal relation is most useful in determining values of e from activity data. Values of eii are determined from experimental data on the binary system i-s, and values of eji are determined from experimental data on the ternary system i-j-s. Values of go and e for aluminum-base and copper-base alloys are given in Tables 20 to 24. The standard state for and e and go described previously is the pure material at the temperature of the solution, and the concentration is expressed in mole fraction. In dealing with dilute solutions, however, it is common to use a hypothetical 1 wt% solution as the standard state. Under this condition, the Taylor series expansion corresponding to Eq 8 is: log fi ¼ eii ð%iÞ þ eji ð%jÞ eki ð%kÞ
(Eq 11)
where fi, which is the activity coefficient for 1 wt% standard state, equals ai/%i and: eji ¼
@ log fi @%j %k¼0; %i; %j!0
(Eq 12)
58 / Principles of Liquid Metal Processing Table 5 Activities, activity coefficients, and partial molar thermal properties for liquid copper-aluminum alloys at 1373 K
Table 6 Integral thermal properties for liquid copper-aluminum alloys at 1373 K
Al component Alð‘Þ ¼ Alðin alloyÞð‘Þ XAl
aAl
1.0 0.9 0.8 0.7 0.6 0.5
gAl
1.000 0.889 0.759 0.609 0.441 0.266 (þ0.02) 0.116 0.028 0.006 0.001 0.000
0.4 0.3 0.2 0.1 0.0
1.000 0.988 0.949 0.870 0.735 0.532 (þ0.04) 0.290 0.095 0.029 0.008 0.002
(1x)Al‘ þ xCu‘ ¼ Al(1
DGAl, cal/g-atom
xs DGAl , cal/g-atom
DHAl, cal/g-atom
0 320 753 1354 2235 3611 (þ150) 5873 9709 14,062 19,307 1
0 32 144 381 842 1720 (þ150) 3373 6424 9670 13,025 16,640
0 29 44 60 391 1055 (þ300) 2878 5012 5864 7415 8625
xs DSAl , cal/K-g-atom
DSAl, cal/K-g-atom
0.000 0.254 0.580 0.942 1.343 1.862 (þ0.25) 2.181 3.421 5.971 8.661 1
0.000 0.044 0.137 0.234 0.328 0.484 (þ0.25) 0.361 1.028 2.772 4.086 5.838
XCu
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9
x)Cux(‘)
DG, cal/ g-atom
DH, cal/ g-atom
DS, cal/ K-g-atom
DGxs, cal/ g-atom
DSxs, cal/ K-g-atom
1723 2969 3955 4700 5170 (þ 150) 5289 4906 3927 2361
459 960 1445 1863 2163 (þ 150) 2254 1979 1380 800
0.921 1.463 1.828 2.066 2.190 (þ 0.1) 2.210 2.132 1.855 1.137
836 1604 2289 2863 3278 (þ 150) 3453 3240 2562 1474
0.275 0.469 0.615 0.728 0.812 (þ 0.1) 0.873 0.918 0.861 0.491
Source: Ref 2
Cu component Cuð‘Þ ¼ Cuðin alloyÞð‘Þ XCu
aCu
gCu
0.0 0.1 0.2 0.3 0.4 0.5
0.000 0.005 0.013 0.025 0.046 0.085 (þ0.005) 0.166 0.352 0.600 0.839 1.000
0.042 0.052 0.065 0.084 0.115 0.170 (þ0.01) 0.277 0.503 0.750 0.932 1.000
0.6 0.7 0.8 0.9 1.0
DGCu, cal/g-atom
xs DGCu , cal/g-atom
DHCu, cal/g-atom
1 14349 11833 10025 8396 6728 (þ150) 4900 2848 1394 478 0
8671 8067 7442 6741 5896 4837 (þ150) 3506 1875 785 191 0
4225 4849 4975 4675 4071 3271 (þ300) 1838 679 259 65 0
DSCu, cal/K-g-atom
xs DSCu , cal/K-g-atom
1 6.919 4.995 3.897 3.150 2.518 (þ0.25) 2.230 1.580 0.827 0.301 0.000
3.238 2.344 1.797 1.505 1.329 1.141 (þ0.25) 1.215 0.871 0.383 0.092 0.000
Source: Ref 2
Table 7 Activities, activity coefficients, and partial molar thermal properties for liquid copper-nickel alloys at 1823 K Cu component Cuð‘Þ ¼ Cuðin alloyÞð‘Þ XNi
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
aCu
1.000 0.902 0.814 0.740 0.677 0.611 (þ 0.02) 0.534 0.444 0.334 0.191 0.000
gCu
1.000 1.002 1.017 1.058 1.128 1.222 (þ 0.03) 1.334 1.480 1.669 1.912 2.227
DGCu, cal/g-atom
xs DGCu ,
0 374 747 1088 1414 1786 (þ 100) 2275 2941 3974 5992 1
cal/g-atom
0 8 61 204 436 725 (þ 100) 1044 1421 1856 2349 2900
Table 8 Integral thermal properties for liquid copper-nickel alloys at 1823 K (1 x)Cu(‘) þ xNi(‘) ¼ Cu(1 XNi
0.1 0.15 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9
DG, cal/ g-atom
948 1195 1377 1617 1745 1786 (þ100) 1742 1604 1349 917
x)Nix(‘)
DH, cal/ g-atom
DS, cal/ K-g-atom
DGxs, cal/ g-atom
DSxs, cal/ K-g-atom
240 341
0.652 0.842
... ... ...
... ... ...
0.006 0.003 ... ... ... ...
... ... ... ...
... ... ... ...
229 336 435 596 693 725 (þ100) 696 609 464 261
... ... ... ...
Source: Ref 2
Ni component Ni(‘) ¼ Ni (in alloy)(‘) aNi
0.000 0.185 0.341 0.455 0.539 0.611 (þ 0.02) 0.682 0.752 0.826 0.907 1.000 Source: Ref 2
gNi
1.906 1.864 1.704 1.517 1.347 1.222 (þ 0.03) 1.136 1.075 1.032 1.008 1.000
DGNi, cal/g-atom
1 6120 3899 2851 2241 1786 (þ 100) 1337 1031 692 353 0
xs DGNi , cal/g-atom
2336 2222 1931 1510 1079 725 (þ 100) 464 261 116 29 0
In Eq 11, fi is the activity coefficient, and the zeroeth-order term, log fio , disappears because the activity coefficient at infinite dilution fio is equal to 1. The reciprocal relationship, corresponding to Eq 10, is: eji ¼
Mi i e Mj j
(Eq 13)
where Mi is the atomic weight of i. The relationship between the two types of the interaction coefficients is: eji ¼
Ms ej ð2:303Þð100ÞMj i
(Eq 14)
where Ms is the atomic weight of the solvent (aluminum, copper).
Thermal Properties for Hypothetical Standard State It is a common practice in solving problems on chemical equilibria to use the hypothetical pure component (Xi) or hypothetical weight percent (%i) as the standard state. For this purpose, thermodynamic properties for the hypothetical standard states are needed. The values of go can be used to determine Gibbs free energies of mixing for elements in liquid base metal, using Eq 15 and 16: Goi ðXi Þ ¼ RT ln goi
(Eq 15)
and Goi ð%iÞ ¼ RT ln
goi Ms 100Mi
(Eq 16)
The values of Goi ðXi Þ and Goi ð%iÞ thus calculated for aluminum-base and copper-base alloys are also given in Tables 20 and 22. The principal advantage of using the hypothetical standard states is that at low concentrations, the activity of i can be set equal to its atom fraction or weight percent, depending on the composition scale used to express the standard state.
Phase Diagrams A phase diagram is essentially a map that shows the relative stabilities of various phases present at equilibrium under the varying conditions of temperature and composition. Therefore, for a given composition of an alloy, a phase diagram can be used to determine the phases, which will be present at equilibrium as the melt
Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys / 59 Table 9 Activities, activity coefficients, and partial molar thermal properties for liquid copper-tin alloys at 1400 K
Table 10 Partial molar thermal properties for liquid copper-tin alloys at 633 K
Cu component Cuð‘Þ ¼ Cuðin alloyÞð‘Þ XCu
1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0
aCu
gCu
1.000 0.802 0.539 0.389 0.284 0.220 (þ0.025) 0.169 0.125 0.082 0.038 0.000
Sn(‘) ¼ Sn (in alloy)(‘)
DGCu, cal/g-atom
xs DGCu , cal/g-atom
DHCu, cal/g-atom
0 613 1717 2626 3497 4215 (þ300) 4949 5781 6971 9104 1
0 320 1096 1633 2076 2287 (þ300) 2400 2432 2493 2698 3197
0 159 749 1443 1662 1631 (þ250) 1418 1049 640 81 1050
1.000 0.891 0.674 0.556 0.474 0.440 (þ0.05) 0.422 0.417 0.408 0.379 0.317
DSCu, cal/K-g-atom
xs DSCu , cal/K-g-atom
0.000 0.325 0.691 0.845 1.311 1.845 (þ0.3) 2.522 3.380 4.522 6.561 1
0.000 0.115 0.248 0.136 0.295 0.468 (þ0.3) 0.701 0.988 1.324 1.985 3.034
XSn
0.98 0.99
aSn
gSn
0.965 0.985 (þ0.005)
0.985 0.995 (þ0.005)
DGSn, cal/g-atom
47 13 (þ5)
DGxs Sn, cal/g-atom
20 7 (þ5)
Source: Ref 2
Sn component Sn(‘) ¼ Sn (in alloy)(‘) XSn
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
aSn
0.000 0.007 0.072 0.197 0.340 0.467 (þ 0.05) 0.580 0.681 0.784 0.892 1.000
gSn
DGSn, cal/g-atom
DGxs Sn, cal/g-atom
DHSn, cal/g-atom
DSSn, cal/K-g-atom
DSxs Sn, cal/K-g-atom
1 13,706 7304 4523 3003 2119 (þ 300) 1516 1069 678 317 0
13,609 7301 2827 1173 454 190 (þ 300) 95 77 57 24 0
8000 5233 1901 252 706 681 (þ 250) 506 311 176 49 0
1 6.053 3.860 3.411 2.649 2.000 (þ 0.3) 1.444 0.986 0.610 0.261 0.000
4.006 1.477 0.661 1.018 0.828 0.622 (þ 0.3) 0.429 0.277 0.167 0.052 0.000
0.007 0.072 0.362 0.656 0.849 0.934 (þ 0.1) 0.966 0.973 0.979 0.991 1.000
Source: Ref 2
Table 11
Heats of formation of solid and liquid copper-tin alloys at 723 K ð1 xÞCuðsÞ þ xSnðsÞ ¼ Cuð1xÞ SnxðsÞ ð1 xÞCuð‘Þ þ xSnð‘Þ ¼ Cuð1xÞ Snxð‘Þ Cuð‘Þ ¼ Cuðin alloyÞð‘Þ
XSn
Phase
0.091(b)
(Cu)
DH, cal/g-atom
xSn
Phase
DCp(a), cal/K-g-atom
DH Cu , cal/g-atom
280
0.825(b)
‘
1105
1.7
...
...
... ... ...
990 760 530 (þ 50) 0
... 1.2 ...
... ... ...
... ... ...
260
... ...
... ...
1.9 (þ2) ... ...
0.204(b) 0.209
d d
1300 1260
0.850 0.900 0.950
0.244(b)
e
1686
1.000
e
1800 1920 (þ100)
... ...
0.250 0.255(b)
‘
DH, cal/g-atom
... ...
(a) Cp ¼ heat capacity. (b) Phase boundary. Source: Ref 2
Table 12
Integral thermal properties for liquid copper-tin alloys at 1400 K ð1 xÞCuð‘Þ þ xSnð‘Þ ¼ Cuð1xÞ Snxð‘Þ
XSn
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 Source: Ref 2
DG, cal/g-atom
1922 2834 3195 3300 3167 (þ 300) 2889 2483 1937 1196
DH, cal/g-atom
666 979 934 715 475 (þ 150) 264 97 13 52
DS, cal/K-g-atom
0.897 1.325 1.614 1.846 1.923 (þ 0.24) 1.875 1.705 1.392 0.891
DGxs, cal/g-atom
1018 1442 1495 1427 1238 (þ 300) 1017 784 545 291
DSxs, cal/K-g-atom
0.251 0.331 0.400 0.509 0.545 (þ 0.24) 0.538 0.491 0.398 0.245
... ...
DC pCuð‘Þ , cal/K-g-atom
60 / Principles of Liquid Metal Processing Table 13 Measured activities of tin in liquid copper-tin alloys Temperature, K
Table 15 Recommended partial molar thermodynamic quantities for liquid alloys at 1400 K Cu component Cu‘ ¼ Cu(in alloy)‘
Activity
Composition, mole fraction of Sn
Sn
Cu
XCu
aCu
gCu
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 0.1 0.2 0.3 0.4 0.5 0.65 0.8 0.9 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9
0.02 0.082 0.22 0.373 0.499 0.614 0.712 0.811 0.905 0.1 0.055 0.198 0.307 0.466 0.652 0.808 0.919 0.007 0.055 0.214 0.353 0.474 0.575 0.683 0.805 0.903
0.856 0.678 0.498 0.373 0.297 0.227 0.17 0.113 0.064 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...
1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0
1.000 0.822 0.611 0.448 0.336 0.258 0.197 0.143 0.092 0.045 0.000
XSn
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
1573
1373
1073
DGCu, J/mol
Gex Cu , J/mol
1.000 0.913 0.763 0.640 0.560 0.516 0.492 0.475 0.460 0.450 0.468
0 2282 5744 9356 12,693 15,777 18,929 22,674 27,780 36,100 1
0 1055 3146 5204 6747 7708 8262 8658 9044 9294 8842
aSn
gSn
DGSn, J/mol
Gex Sn , J/mol
0.000 0.013 0.076 0.196 0.336 0.465 0.581 0.690 0.798 0.901 1.000
0.021 0.132 0.378 0.652 0.839 0.930 0.968 0.986 0.997 1.001 1.0000
1 50,370 30,058 19,002 12,716 8913 6326 4316 2633 1213 0
45,239 23,564 11,322 4986 2049 844 380 164 35 13 0
Source: Ref 3
Table 14 Recommended integral thermodynamic quantities for liquid alloys at 1400 K
Table 16 Recommended enthalpies of formation of solid and liquid alloys at 723 K
DG
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9
7091 10,607 12,250 12,702 12,345 11,368 9824 7662 4702
DH
3360 4374 4010 2996 1838 857 210 79 94
Gex
DS
2.655 4.452 5.886 6.933 7.505 7.508 6.867 5.529 3.426
3307 4782 5139 4868 4277 3534 2713 1837 918
Note: Values in J/mol or J/mol K. Source: Ref 3
solidifies. In addition, a phase diagram can also be used to estimate the activity of the components in liquid solution and to understand the behavior of the liquid solution (Ref 9). For example, if the diagram is for a simple eutectic system, it is likely that no strong intermetallic compound exists. If, however, a strong compound in the solid state appears in the diagram, then association of dissimilar atoms is probably occurring in the liquid. Therefore, one can predict whether a liquid solution will exhibit strong negative or positive deviations from the ideal behavior from the shapes of the liquids and azeotropes. The phase diagrams of the aluminum-silicon and copper-zinc systems, which are chosen based on their commercial importance, are shown in Fig. 1 and 2. The aluminum-silicon diagram (Fig. 1) represents a simple eutectic system, and the alloys containing 8.5 to 13% Si are widely used
Sex Cu , J/mol K
0.000 0.661 1.243 2.123 3.628 5.999 9.360 13.736 19.152 26.317 1
0.000 0.215 0.612 0.843 0.619 0.235 1.742 3.725 5.770 7.172 6.891
Phase
0.090(a)
DH
XSn
(Cu)
1165
0.825(a)
0.205
d
5350
0.85
0.240
e
7650
0.90
0.435
Z
7030(b)
0.95
Phase ‘ ‘ ‘ ‘
49,228 21,394 5854 1526 3930 3704 2455 1176 341 22 0.000
DSSn, J/mol K
Sex Sn , J/mol K
1 20.697 17.289 14.663 11.890 9.012 6.272 3.923 2.124 0.882 0.000
2.849 1.552 3.907 4.652 4.272 3.249 2.025 0.957 0.269 0.006
XSn
Phase
GSn, J/mol
671
0.008 0.012 0.021 0.043 0.069
(Cu) (Cu) (Cu) (Cu) (Cu)
54,350 50,354 45,606 38,618 33,600
515
Data given relative to pure liquid Sn. Source: Ref 3
(1 X) Cu‘ þ XSn‘ ¼ Cu(1x)Snx‘ XSn
DHSn, J/mol
Table 17 Recommended partial Gibbs energy of tin in solid alloys at 1000 K
(1 X) Cu(s) þ XSn(s) ¼ Cu(1x)Snx(s)
Sex
0.038 0.291 0.807 1.337 1.742 1.912 1.788 1.368 0.723
0 1356 4004 6384 7614 7379 5825 3444 967 744 806
DSCu, J/mol K
Sn component Sn‘ ¼ Sn(in alloy)‘
Source: Ref 3
XSn
DHCu, J/mol
DH
263 104
(a) Phase boundary. (b) DH at 237 K. Source: Ref 3
to produce automotive parts. Generally, the higher the silicon content, up to the eutectic composition (12.6% Si), the greater the fluidity and, consequently, the easier an alloy is to cast. The copperzinc diagram (Fig. 2) however, shows a high solubility of zinc in solid copper-zinc solution and the formation of several other phases over a wide range of zinc concentrations. Copper alloys containing up to 35% Zn and significant amounts of other alloying elements are quite common. The copper-zinc phase diagram is used to adjust the compositions in order to produce alloys with the desired properties. REFERENCES 1. B.L. Tiwari, Metall. Trans. A, Vol 18, 1987, p 1645–1651
2. R. Hultgren et al., Selected Values of the Thermodynamic Properties of Binary Alloys, American Society for Metals, 1973 3. P.R. Subramanian, D.J. Chakrabarti, and D. E. Laughlin, Ed., Phase Diagrams of Binary Copper Alloys, ASM International, 1994 4. C. Wagner, Thermodynamics of Alloys, Addison-Wesley, 1952 5. G.K. Sigworth and T.A. Engh, Scand. J. Metall., Vol 11, 1982, p 143–149 6. J.M. Dealy and R.D. Pehlke, Trans. Met. Soc. AIME, Vol 227, Feb 1963, p 88–94 7. G.K. Sigworth and J.F. Elliott, Can. Met. Q., Vol 13, 1974, p 455–461 8. J.M. Dealy and R.D. Pehlke, Trans. Met. Soc. AIME, Vol 227, Aug 1963, p 1030–1032 9. L.S. Darken and R.W. Gurry, Physical Chemistry of Metals, McGraw-Hill, 1953 10. T.B. Massalski et al., Binary Alloy Phase Diagrams, Vol 1 and 2, American Society for Metals, 1986
Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys / 61 Table 18 Activities, activity coefficients, and partial molar thermal properties for liquid copper-zinc alloys at 1200 K Cu component Cu(‘) ¼ Cu (in alloy)(‘) XZn
aCu
0.334(a) 0.4 0.5
gCu
0.432 0.334 0.219 (þ0.05) 0.139 0.085 0.049 0.023 0.000
0.6 0.7 0.8 0.9 1.0
xs DGCu ,
DGCu , cal/g-atom
2003 2614 3621 (þ400) 4705 5869 7187 9010 1
0.648 0.557 0.438 (þ0.1) 0.348 0.284 0.245 0.229 0.235
cal/g-atom
1034 1396 1968 (þ400) 2520 2998 3349 3519 3454
Zn component Zn(‘) ¼ Zn (in alloy)(‘)
aZn 0.132 0.207 0.347 (þ 10.05) 0.505 0.657 0.789 0.900 1.000
xs
gZn
DGZn , cal/g-atom
DGZn , cal/g-atom
0.398 0.517 0.695 (þ 0.1) 0.841 0.939 0.986 1.000 1.000
4814 3757 2521 (þ 400) 1630 1002 564 250 0
2199 1572 868 (þ 400) 412 151 32 1 0
(a) Phase boundary. Source: Ref 2
Table 19
Integral thermal properties for liquid copper-zinc alloys at 1200 K ð1 xÞCuð‘Þ þ xZnð‘Þ ¼ Cuð1xÞ Znxð‘Þ
XZn
DG
0.334(a) 0.4 0.5
2942 3071 3071 (þ 400)
DGxs, cal/g-atom
1423 1466 1418 (þ 400)
xZn
0.6 0.7 0.8 0.9
DG, cal/g-atom
2860 2462 1889 1126
DGxs, cal/g-atom
1255 1005 695 351
(a) Phase boundary. Source: Ref 2
Table 20
Standard Gibbs free energies for solution of elements in liquid aluminum
Solution reaction
Be(s)¼Be Bi(l)¼Bi Ca(l)¼Ca Cd(l)¼Cd Cu(s)¼Cu Fe(s)¼Fe Ga(l)¼Ga Ge(l)¼Ge ½H2 ¼ H In(l)¼In Li(l)¼Li Mg(l)¼Mg Na(l)¼Na Ni(s)¼Ni Pb(l)¼Pb Sb(l)¼Sb Si(s)¼Si Sn(l)¼Sn Zn(l)¼Zn Source: Ref 5
g0, 1100 K
19.6 24.5 0.0086 19.9 0.037 1.6 104 1.1 0.16 ... 12.3 0.40 0.18 293 0.7 106 115 3.4 0.27 4.68 1.92
DGoi for i (X), cal/g-atom
9404 2.639T 5309 þ 1.524T 10,400 7100 0.518T 1135 5.538T 27,000 þ 7.18T 832 0.52T 2761 1.176T ... 6800 1.201T 5800 þ 3.435T 3478 0.30T 8230 þ 3.798T 28,280 2.442T 9970 þ 0.363T 13,100 9.45T 9598 11.291T 5845 2.245T 2538 1.007T
DGoi for i(%), cal/g-atom
9404 9.607T 5309 11.688T 10,400 9.93T 7100 12.498T 1135 16.385T 27,000 3.411T 832 11.551T 2761 12.288T 11,664 þ 6.523T 6800 13.223T 5800 3.014T 3478 9.24T 8230 5.03T 28,280 13.132T 9970 12.832T 13,100 21.589T 9598 20.517T 5845 14.333T 2538 11.911T
Table 21 Interaction coefficients for elements in liquid aluminum i
j
«ji
Temperature, K
Ag Cu H H H H H H H H H Si Cd Cu Ga Ge In Mg Si Sn Zn Zn Li Na Mg Sn H H H H H H H H H H H
Ag H Cu Cu Cu Cu H Si Si Si Si H Cd Cu Ga Ge In Mg Si Sn Zn Si Sn Si Si Pb Ce Cu Cr Fe Mg Mn Ni Th Ti Si Sn
3.1 see eCu H 39.0 16.6 20.1 4.3 0 11.5 6.2 4.2 1.8 see eSiH 5.0 2.2 0.3 þ3.0 4.5 3.0 16.0 6.0 0.9 2.2 16.0 12 9 1.5 100 15 1 1 2 27 19 20 42 7.1 0.6
1273 973–1273 973 1073 1173 1273 973–1273 973 1073 1173 1273 973–1273 1373 1373 1023 1200 1173 1073 1100 1100 1000 963–1053 949 973 973–1073 973–1073 800 700–800 800–900 800–900 700–800 800 700–800 800–900 800–900 700–800 800–900
Source: Ref 5, 6
62 / Principles of Liquid Metal Processing Table 22
Standard Gibbs free energies for solution of elements in liquid copper
Table 24 Activity coefficients at infinite dilution in liquid metals
Temperature Element(a), i
goi , 1200 C (2190 F)
Ag(l) Al(l) As(v) Au(l) Bi(l) C(graphite) Ca(l) Cd(v) Cd(l) Co(s) Cr(s) Fe(s) Fe(l) Ga(l) Ge(l)
3.23 0.0028 4.8 104 0.14 1.25 1.4 105 5.1 104 15.6 0.53 15.4 43 24.1 19.5 0.034 0.009 ... 0.41 0.08 0.044 0.51 0.53 2.22 2.66 ... 5.27 1.3 0.05 ... 0.014 0.002 0.006 0.01 0.048 0.0328 8.5 130 0.146
1=2 H2 ðgÞ
In(l) Mg(v) Mg(l) Mn(l) Mn(s) Ni(l) Ni(s) 1 2 O2 ðgÞ
Pb(l) Pd(s) Pt(s) 1=2 Ss ðgÞ
Sb(l) Se(v) Si(l) Si(s) Sn(l) Te(v) Ti(l) V(s) Zn(l)
DGoi ðXÞ, cal/g-atom
DGoi (wt%), cal/g-atom
3900 0.32T 8630 5.84T 22,350 4630 0.73T 5960 3.6T 8550 þ 17.8T 22,200 25,700 þ 22.9T 1860 8000 11,000 12,970 2.48T 9300 0.41T 10,800 þ 0.61T 16,000 þ 1.52T 10,400 þ 8.4T 9550 þ 4.71T 40,200 þ 22.3T 8670 0.31T 1950 1550 2.31T 2340 6500 2.5T 20,400 þ 10.8T 8620 2.55T 800 10,200 þ 0.87T 28,600 þ 13.79T 12,500 18,200 15,000 2900 7.18T 8900 10,000 6730 0.31T 28,100 9.4T 5640
3900 10.52T 8630 13.84T 22,350 9.44T 4630 12.09T 5960 15.1T 8550 þ 12.0T 22,200 25,700 þ 12.7T 1860 9.0T 8000 9.0T 11,000 8.72T 12,970 11.34T 9300 9.27T 10,800 8.68T 16,000 7.5T 10,400 7.5T 9550 5.58T 40,200 15.1T 8670 7.53T 1950 8.83T 1550 11.14T 2340 9.0T 6550 11.5T 20,400 4.43T 8620 14.01T 800 10.1T 10,200 10.47T 28,600 6.03T 12,500 10.4T 18,200 9.5T 15,000 7.5T 2900 14.68T 8900 10.4T 10,000 10.53T 6730 11.74T 28,100 18.1T 5640 9.1T
C
1100–1200 1100 1000 1175–1325 1100–1300 1100–1300 800–925 ... ... ... ... 1460–1580 1460–1580 1100–1280 1255–1545 1100–1300 700–1000 650–927 650–927 1244 1244 ... ... 1100–1300 1000–1300 1500–1600 ... 1050–1250 1000–1200 1200 1550 1550 1100–1300 1200 1000–1300 ... 1150
(a) l, liquid; v, vapor; s, solid; g, gas. Source: Ref 7
Table 23 Interaction coefficients for elements in liquid copper alloys Temperature i
j
H H H H H H H H H H H H H H H H H O O O O
Ag Al Au Co Cr Fe Mn Ni P Pb Pt S Sb Si Sn Te Zn Ag Au Co Fe
O O O O O O
Ni P Pb Pt S Si
«ij
0.5 6.2 1.9 3.1 1.6 2.9 1.1 5.5 10.0 21.0 8.0 9.0 13.0 4.8 6.0 6.6 6.8 0.7 8.6 68 4.04 106/T 2183 36,000/T þ 17 700,000/T þ 385 7.4 38 19 6300
Source: Ref 5, 6
C
Temperature
F
1225 1225 1225 1150 1550 1150–1550 1150 1150–1240 1150 1100 1225 1150 1150 1150 1100–1300 1150 1150 1100–1200 1200–1550 1200 1200–1350
2235 2235 2235 2100 2820 2100–2820 2100 2100–2260 2100 2010 2235 2100 2100 2100 2010–2370 2100 2100 2010–2190 2190–2820 2190 2190–2460
1200–1300 1150–1300 1100 1200 1206 1250
2190–2370 2100–2370 2010 2190 2203 2280
i
O S S S S S S Ag Al Au Bi Ca Fe Ga Ge H Mg Mn O Pb S Sb Sn Tl Zn Zn Zn Zn Zn
j
Sn Au Co Fe Ni Pt Si Ag Al Au Bi Ca Fe Ga Ge H Mg Mn O Pb S Sb Sn Tl An Zn Zn Zn Zn
«ij
4.6 6.7 4.8 25,400/T þ 8.7 29,800/T þ 13 11.5 6.9 2.5 14 3.7 6800/T þ 1.65 20 5.7 7 13.4 1.0 9.8 6 24,000/T þ 7.8 2.7 20,800/T 15 10 4.8 4 0.38 0.72 1.185 1.40
C
1100 1115–1200 1300–1500 1300–1500 1300–1500 1200–1500 1200 1150 1100 1277 1000–1200 877 1550 1280 1255 1123 927 1244 1100–1300 1200 1050–1250 1000–1200 1300–1320 1300 1150 902 727 653 604
Temperature
F
2010–2190 2010 1830 2150–2415 2010–2370 2010–2370 1470–1700 ... ... ... ... 2660–2880 2660–2880 2010–2340 2290–2815 2010–2370 1290–1830 1200–1700 1200–1700 2271 2271 ... ... 2010–2370 1830–2370 2730–2910 ... 1920–2280 1830–2190 2190 2820 2820 2010–2370 2190 1830–2370 ... 2100
F
2010 2040–2190 2370–2730 2370–2730 2370–2730 2190–2730 2190 2100 2010 2331 1830–2190 1611 2820 2335 2290 2053 1701 2271 2010–2370 2190 1920–2280 1830–2190 2370–2410 2370 2100 1656 1341 1207 1119
Solvent
Solute
Aluminum
Silver
Copper
Magnesium Zinc
Source: Ref 2, 8
g
0.38 0.47 0.53 0.88 0.14 0.17 0.21 0.24
C
700 900 1000 800 604 653 727 802
F
1290 1650 1830 1470 1119 1207 1341 1476
Thermodynamic Properties of Aluminum-Base and Copper-Base Alloys / 63
Silicon, at.% 1500
10
0
20
30
40
50
60
70
80
90
100 2700 1414 ⬚C
2400
1300
1100
2000
900
1700
700 660.452 ⬚C
1300 577 ± 1 ⬚C
12.6
300
(Si)
(Al)
500
0
Al
10
20
30
40
50 Silicon, wt%
Fig. 1
The aluminum-silicon phase diagram. Source: Ref 9
Fig. 2
The copper-zinc phase diagram. Source: Ref 10
60
70
80
90
900
600 100 Si
Temperature, ⬚F
Temperature, ⬚C
L
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 64-73 DOI: 10.1361/asmhba0005192
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Gases in Metals Richard J. Fruehan, Carnegie Mellon University Prince N. Anyalebechi, Grand Valley Sate University
(Eq 1)
where the underlining of aluminum (Al) indicates that it is dissolved in the metal. The free energy for reaction in Eq 1 is highly negative even for very low concentrations of aluminum, indicating that the thermodynamics for the reaction are highly favorable. Similar reactions can be written for silicon, carbon, and iron. For carbon, carbon monoxide is also evolved, according to the reaction: H2 O þ C ! H2 þ CO
(Eq 2)
For the casting of aluminum and its alloys, the reaction is even more favorable. Whether or not the reaction will actually occur depends on many complex factors, including liquid-phase mass transfer and surface tension; therefore, its occurrence is difficult to predict. However, if there is excessive moisture in the sand, the reaction is highly likely to occur. The other major cause of gas porosity is the evolution of dissolved gases during casting. For example, liquid cast iron may have dissolved hydrogen and nitrogen. The solubility of these gases in the solid may be less than that in the liquid, and the gases may therefore be evolved during solidification. Whether or not this will occur depends on the amount of hydrogen and nitrogen present, the alloy being cast, chemical kinetics, and the surface tension of the alloy. Similarly, hydrogen can dissolve in liquid aluminum, but its solubility is much lower in solid aluminum
Cast Iron (Ref 1) Solubility of Hydrogen and Nitrogen. The gases in cast iron that can cause porosity are hydrogen and nitrogen. Castings from iron produced in a cupola are nearly saturated with nitrogen; therefore, the presence of nitrogen is a major concern (Ref 2). Hydrogen and nitrogen dissolve in liquid iron alloys and other metals as atomic species according to: 1
/2 H2 ðgasÞ ! H
(Eq 3)
/ N2 ðgasÞ ! N
(Eq 4)
12
The underlined symbols H and N denote hydrogen dissolved in liquid aluminum and its alloys. The solubilities of both elements in liquid and solid iron alloys have been extensively measured, and the results have been adequately summarized (Ref 3). The solubility of nitrogen and hydrogen in liquid iron is decreased by carbon. The nitrogen and hydrogen contents in iron-carbon alloys in equilibrium with 1 atm pressure of the respective gases are shown in Fig. 1 as a function of carbon content. The solubility of nitrogen at 1823 K decreases from 450 ppm for pure iron to approximately 80 ppm for an alloy containing 4.5 wt% C. The solubilities of hydrogen for the same metals are 24 and 10 ppm, respectively. The hydrogen solubility decreases with temperature and at
eutectic temperature is 6.5 ppm; temperature has only a small effect on the solubility of nitrogen. Silicon decreases the solubility of nitrogen even further. For a Fe-4wt%C-1wt% Si melt, the solubility of nitrogen is approximately 75 ppm. When cast iron solidifies, it forms austenite and graphite or cementite. The solubility of nitrogen in austenite (Ref 3) is also shown in Fig. 1. For austenite containing approximately 1.8 wt% C, the solubility is approximately 150 ppm. Therefore, when the liquid cast iron freezes, the solubility in the solid is higher than that in the liquid; this interesting phenomenon is discussed later in detail. There is some solubility of nitrogen in cementite (Ref 4). The exact amount is not known but is approximately the same as that in austenite. The solubility of hydrogen in austenite is approximately 7 ppm; therefore, the solubility of hydrogen in the liquid and solid are about equal, and there is little segregation of hydrogen during solidification (Ref 3). Kinetics of Gas-Liquid Reactions. The absorption of hydrogen and nitrogen is controlled by one of the following steps:
500
50
400
40 Hydrogen in liquid
300
30 Nitrogen in liquid
200
20 Nitrogen in austenite
100
0
Fig. 1
Hydrogen, ppm
2Al þ 3H2 O ! Al2 O3 þ 3H2
and can cause porosity. For copper, the evolution of hydrogen, water vapor, or carbon monoxide could cause gas porosity. In this article, the solubilities of the common gases present in ferrous metals, such as cast irons, and nonferrous metals, such as aluminum, copper, and magnesium, are reviewed. The kinetics of the relevant reactions, reactions during solidification, and possible methods of control or removal of the dissolved gases are discussed. The discussion is primarily focused on cast iron because more is known about the thermodynamics and kinetics of iron than other elements. However, th e same basic principles also apply to aluminum, copper, and magnesium.
Nitrogen, ppm
GAS POROSITY is one of the most serious problems in the casting of both ferrous and nonferrous metals. It is generally caused by the evolution of gases during the casting and solidification process. The gases may be the result of a reaction between the casting sand mold and the metal, or they may result from the evolution of gases dissolved in the liquid metal during solidification. An example of a chemical reaction evolving gas and causing porosity is the reaction of moisture (H2O) in a sand mold with elements in the cast iron, such as carbon, silicon, aluminum, or iron itself. For example, aluminum in cast iron often causes porosity problems, and the reaction responsible for the gas evolution is:
10 0
0
1
2 3 Carbon, %
4
5
Effect of carbon on the solubility of hydrogen and nitrogen in the iron-carbon alloys at 1823 K (i.e., liquid) and 1500 K (i.e., austenite) and 1 atm pressure
Gases in Metals / 65
(Eq 5)
where km is the mass transfer coefficient, pBN2 is the pressure of nitrogen in the bulk gas, and pSN2 is the surface pressure of nitrogen in equilibrium with the melt where: pSN2 ¼ KN2 fN2 ðwt%NÞ2
(Eq 6)
where KN is the equilibrium constant for Eq 4, and fN is the activity coefficient of nitrogen with respect to 1 wt%. However, in general, gas diffusion for these reactions is faster than the chemical reaction or liquid-phase diffusion and can be neglected in most cases. The rate of the chemical reaction is controlled by the dissociation of the nitrogen molecule on the surface, and the rate is given by: dnN2 ¼ kB Að1 yÞ pN2 peN2 dt
(Eq 7)
where kB is the rate constant for pure iron, A is the surface area, peN2 is the equilibrium nitrogen pressure given by an expression similar to Eq 6, and (1 y) is the fraction of vacant sites not occupied by surface-active elements. Certain elements are surface active on liquid iron in that they lower the surface tension of iron by covering most of the surface. For example, oxygen and sulfur are surface active on iron, and for bulk concentrations of 0.03% O or 0.03% S, over 90% of the surface sites will be covered by oxygen or sulfur. These elements therefore retard the rate of chemical reaction. At high coverage, the surface-active element (1 y) is inversely proportional to the activity, a, of the element. Therefore, the rate is given by: dnN2 kB kA A ¼ pN2 peN2 dt a
(Eq 9)
where kmN is the liquid-phase mass transfer coefficient of nitrogen, CNB is the bulk concentration of nitrogen, and CNS is the surface concentration of nitrogen. The integrated form of Eq 7 is: %N %Ne ArkmN t ¼ ln %No %Ne W
(Eq 10)
where r is the density of iron, W is the weight of the metal, %Ne is the equilibrium of the nitrogen contents, and %No is the initial nitrogen contents. The rate is often controlled by two processes in series—usually the chemical reaction and liquidphase mass transfer. The rate in this case can be obtained by equating the fluxes given by Eq 8 and 9 and solving for the surface concentration, as indicated by Pehlke and Elliot (Ref 5). 3.0 2.5
1.0
2.0 1.5
1.0 1450⬚C 1600⬚C
0.5
0.75 0.5 0.25 0
0
100
0 0
10
20
30
40
50
Fig. 2
Fig. 4
Effect of sulfur in the rate of reaction of nitrogen with Fe-CSAT-S alloys
1.0 0.75 0.5 0.25 0
Fig. 3
0
2.5
5 1/aBi
7.5
300
400
500
600
Effect of tellurium on the reaction of nitrogen with Fe-CSAT-Te alloys at 1450 C (2640 F)
1.5
1.0
0.5 as = 0.11 Present study Ref 9 0
10
Effect of bismuth on the reaction of nitrogen with Fe-CSAT-Bi alloys at 1450 C (2640 F)
200
1/[% Te]
60
1/as
(Eq 8)
where kA is related to the adsorption coefficient of the element on the surface, and the quantity kBkA/a is the overall rate constant k.
Rate constant, mol/cm2 • s • atm) ⫻ 106
JN2
km S p pBN2 ¼ RT N2
JN ¼ kmN CNB CNS
The rate of the nitrogen reaction with iron alloys containing oxygen, sulfur, chromium, and other elements has been measured by many investigators, and there is good agreement in most cases (Ref 5–10). However, the rate for carbon-saturated iron containing other elements was not investigated until recently (Ref 11). In the study, an isotope exchange technique was used to measure the rate of dissociation of the nitrogen molecule (N2). In particular, the effects of carbon, sulfur, phosphorus, lead, tin, bismuth, and tellurium on the rate were investigated. Sulfur, phosphorus, lead, bismuth, and tellurium decreased the rate, while carbon and tin had no significant effect. For example, the effect of sulfur is shown in Fig. 2 as the rate constant versus the reciprocal of the activity of sulfur (Ref 10). The activity of sulfur is in weight percent, and the activity coefficient of sulfur in carbonsaturated iron is 6.3. The rate for carbon-saturated iron with no sulfur is approximately 105 mol/ cm2 s atm at 1450 C (2640 F). For example, for 0.009 wt% S (1/as = 18), the rate is decreased to 8 107 mol/cm2 s atm. Bismuth and tellurium had an even larger effect (Ref 11), as shown in Fig. 3 and 4. The activities are relative to pure bismuth and tellurium, respectively, which have larger deviations from ideal behavior. As little as 50 ppm Te reduced the rate by 90%, and 50 ppm Bi decreased it by 80%. On the other hand, carbon is not surface active and does not affect the rate at moderate sulfur contents (0.017 wt% S) up to 4.5 wt% C (Ref 11), as shown in Fig. 5. Therefore, the
Rate constant, mol/(cm2 · s · atm) ⫻ 106
For the desorption of a gas, the same three steps are important but occur in reverse order. The process can be controlled by two or more of the steps in series. For liquid cast iron, nitrogen is the more important of the gases, and it is considered in detail. Gas diffusion is usually not rate controlling for nitrogen absorption from the atmosphere because of the pressure of N2; consequently, the driving force for diffusion is high. As a result, one of the other steps is slower and rate controlling. However, for nitrogen or hydrogen removal by an inert gas, diffusion on N2 or H2 must be considered. The flux of nitrogen,JN2 , away from the surface is given by:
Rate constant, mol/(cm2 ·s · atm) ⫻ 106
from the surface
For dilute solutions, the activity of the surfaceactive element is proportional to its concentration. Therefore, the overall rate constant is inversely proportional to the weight percent of the surface-active element in iron. Liquid-Phase Mass Transfer. If the rate is controlled by liquid-phase mass transfer, the flux of nitrogen atoms is given by:
Rate constant, mol/cm2 ·s· atm) ⫻ 106
Diffusion of the gas to the surface Chemical reaction on the surface Diffusion of the element in the liquid away
Fig. 5
0
1
2 3 4 Carbon content, wt%
5
6
Effect of carbon on the reaction of nitrogen with Fe-C-S alloys at constant sulfur activity
66 / Principles of Liquid Metal Processing pT ¼ KH2 fH2 ð%HÞ2 þKN2 fN2 ð%NÞ2
(Eq 11)
where Ki is the equilibrium constant for the gas reaction, and fi is the activity coefficient. At the end of solidification, the concentrations will be as indicated in Eq 11. Even a small amount of hydrogen may be important. For example, in the case of an Fe-3.8wt%C alloy containing only 4 ppm H, the hydrogen content in the liquid will increase slightly, because of the enrichment, to approximately 4.5 ppm just prior to the eutectic reaction. However, even for this small amount of hydrogen, its thermodynamic pressure is 0.4 atm. Therefore, if the thermodynamic pressure of nitrogen exceeds 0.6 atm, a gas pinhole may result. Calculations indicate that iron with as little as 80 ppm N and 4 ppm H may have a gas pressure exceeding 1 atm. Figure 6 shows a plot of the total pressure of hydrogen and nitrogen just prior to final solidification equal to 1 atm as a function of bulk nitrogen and hydrogen contents in an Fe-3.8C alloy. If the hydrogen and nitrogen contents are such that they are below the line, the pressure will not exceed 1 atm, and a pinhole due to gas evolution will not occur; if above the line and the pressure exceeds 1 atm, a pinhole from this source is possible. For pinholes to develop, it is also necessary to nucleate a bubble. Bubble nucleation is a complex phenomenon and is influenced by surface tension. For melts with low surface tension, bubble nucleation is favored. On the other hand, elements that reduce the surface tension also reduce the rate of N2 formation. Therefore, it is difficult to predict the net effect on the probability of bubble formation of an alloying element that reduces the surface tension of iron. Nitrogen and Hydrogen Removal by Inert Gas Flushing. The preceding discussion indicates that nitrogen and hydrogen evolution during solidification may be a cause of pinholes in castings. Therefore, it would be desirable to 120
remove a portion of these gases if possible. One method is by argon bubbling in the desulfurization reactor or ladle. The desulfurization reactor is a continuous reactor with metal flowing in and out almost continuously. However, for the present calculation, it is assumed to be a batch reactor. Because this is only an order of magnitude calculation, this assumption is not unreasonable. When argon is bubbled through the melt, nitrogen atoms combine on the bubble surface to form N2. The rate is controlled by the chemical reaction on the surface and the liquid-phase mass transfer of nitrogen to the surface in series. For the purpose of the present calculations, the following were assumed: 10 Mg (11 ton) reactor 1 m (3.3 ft) deep Argon flow rate of 0.005 m3/s (10 scfm)
through the melt
Initial nitrogen content of 100 ppm 60 mm (2.4 in.) diameter bubbles
The rate constant for iron containing a relatively small amount of sulfur is approximately 1.5 106 mol/cm2 s atm, and the mass transfer coefficient, km, is approximately 0.1 cm/s. The gas velocity, retention time, and surface area were estimated as done previously (Ref 13). Equations 8 and 9 were solved simultaneously, and the results are given in Fig. 7. The calculations demonstrate that the rate is truly mixed control, as indicated by the surface concentration being between zero and the bulk concentration. The results indicate that it would take approximately 40 min to remove 20 ppm; if the argon flow were doubled, it would still take over 20 min to remove 20 ppm. Although these calculations are rather crude, they indicate that it would be difficult to remove nitrogen by argon bubbling. For hydrogen, the chemical reaction and mass transfer are considerably faster than for nitrogen. If equilibrium between the metal and the gas bubbles leaving the system is assumed, it is possible to calculate the amount of hydrogen that can be removed by gas bubbling from thermodynamic considerations alone. The
100 0.012 80 0.010
p N + pH = 1 2
2
60
Nitrogen, %
Nitrogen, ppm
initial rate of nitrogen formation is primarily controlled by chemical kinetics at the surface, and surface-active elements, such as sulfur and tellurium, can reduce the rate. This information can be helpful in controlling the rate of the reaction. For example, if it is desired to remove nitrogen by argon gas flushing, the concentrations of these elements should be as low as possible. On the other hand, it may be possible to retard nitrogen evolution during solidification by the deliberate addition of these elements. Reactions During Solidification. During the solidification of most simple iron alloys, there is enrichment of the alloying element because there is greater solubility in the liquid than in the solid. The difference in the solubilities is expressed by the partition ratio kS/L, which is determined from the slopes of the solidus to liquidus lines on the phase diagram. For most simple alloys, the partition ratio is less than 1, resulting in liquid enrichment during solidification. For example, for an Fe-0.02wt% N alloy, the last portion of liquid to solidify will be enriched and will have a concentration of 0.045 wt% N, which is the solubility of nitrogen in liquid iron at 1 atm (Ref 12). Therefore, even though the liquid alloy is far from being saturated with nitrogen, during solidification, the concentration will increase to the point where the solubility is exceeded and nitrogen gas will be evolved. Researchers have measured the partition ratio for nitrogen in high-carbon iron alloys (Ref 4). The partition ratio was found to be 1.9 and 2.2 for stable and metastable eutectic solidification, indicating that there is greater solubility in the solid than in the liquid. Therefore, in this case, as the alloy solidifies, nitrogen is actually enriched in the solid, and the nitrogen content decreases in the liquid. It may appear that, because there is no enrichment during solidification, the thermodynamic pressure of nitrogen cannot increase. The thermodynamic pressure is defined by an expression similar to Eq 6. Carbon greatly increases the activity coefficient of nitrogen in the liquid and decreases its solubility; therefore, the thermodynamic pressure may increase because the carbon content of the liquid is increasing during solidification. For example, for an alloy containing 3.8 wt% C that is saturated at 1500 C (2730 F) with 110 ppm N just prior to eutectic solidification, the concentration in the liquid will decrease to 97.5 ppm because of enrichment in the austenite. However, the solubility of nitrogen in the liquid has decreased to 90 ppm primarily because of the increase in the carbon content. Consequently, the thermodynamic pressure of nitrogen exceeds 1 atm, and the evolution of nitrogen may cause a pinhole. When determining the possibility of pinhole formation, one must consider the total thermodynamic pressure of all the gases, and if the total pressure, pT, exceeds 1 atm, a gas pinhole may result. For example, for hydrogen and nitrogen, the total pressure is given by:
40
0.008
0.006
20
0 0
Fig. 6
2
6 4 Hydrogen, ppm
8
10
Thermodynamic pressure of nitrogen and hydrogen at the end of solidification equal to 1 atm for an Fe-3.8wt%C alloy
0.004 0
Fig. 7
20
60 40 Time, min
80
100
Rate of removal of nitrogen from an Fe-3.8wt% C alloy with argon bubbling at 0.005 m3/s (10 scfm) in a 10 Mg (11 ton) reactor
Gases in Metals / 67 equations are developed similarly to those for steel (Ref 12), taking the form: 1 1 ¼ kH V ½H ½Ho
(Eq 12)
where [H] is the hydrogen content after bubbling, [Ho] is the initial hydrogen content, and V is the total volume of gas used per ton of metal. The constant kH is related to the solubility of hydrogen and therefore depends on the alloy composition. For example, 10 Mg (11 tons) of an Fe-4.5wt%C melt at 1773 K containing 4 ppm H bubbling argon for 20 min at 0.005 m3/s (10 scfm) will reduce the hydrogen content to 1.6 ppm. The previous calculation assumes equilibrium and therefore the fastest rate possible and may be an overestimation of the amount of hydrogen removed. Generally, the efficiency of the flushing gas would be roughly 50% for the conditions encountered in practice. However, the calculation does indicate that it is possible to remove significant amounts of hydrogen by gas bubbling.
Aluminum and Its Alloys Compared to other nonferrous metals, aluminum is the most susceptible to hydrogeninduced defects. This has been attributed to two primary reasons. The first reason is that molten aluminum and its alloys readily absorb hydrogen by the reaction with water vapor: 2AlðlÞ þ 3H2 OðvÞ ! Al2 O3 ðsÞ þ 3H2 ðin AlÞ (Eq 13)
The standard free energy change, DGo, for this reaction is extremely high: Go ¼ 979; 100 719T log10 T þ 413T
(Eq 14)
where T is the temperature in K. Inserting this value of free energy change into the van’t Hoff reaction isotherm yields an activity coefficient that is exceedingly high and, for practical purposes, implies complete conversion to hydrogen of all traces of moisture contacting the metal. That is: Go ¼ RT ln Kp
(Eq 15)
where DGo is the standard free energy change, T is the temperature in K, R is the gas constant, and Kp is the activity coefficient. The activity coefficient is given by the equation: Kp ¼
pH2 aH2 O
(Eq 16)
where pH2 is the partial pressure of hydrogen, and aH2 O is the activity of the water vapor. The second reason is that, unlike other nonferrous metals, the partition coefficient, k, that is, the ratio of hydrogen solubility in the solid state to that in the liquid state, is relatively small. That is, the solubility of hydrogen in liquid aluminum is more than an order of magnitude greater than
that in the solid. The first reason is not unique since, for a similar reaction for magnesium and lithium and their alloys, the free energy change is even higher. This explains the greater propensity for hydrogen absorption and the high solubility of hydrogen in magnesium, lithium, and their alloys. However, the low partition coefficient is unique to aluminum and its alloys and partly explains why, despite the comparatively lower solubility of hydrogen in aluminum, it is still the most susceptible of the nonferrous metals to hydrogen-induced defects. Hydrogen Solubility and Reactions During Solidification. Hydrogen solubility is often a direct indication of the propensity for porosity formation and its evolution during solidification of a casting. It indicates the nature of portioning of hydrogen in both cast and wrought products and the effects of alloying elements on the susceptibility of aluminum alloys to hydrogen-induced defects. It is generally agreed that during solidification, hydrogen in the melt is rejected at the solid-liquid interface (at the liquid side) in the same manner in which alloying elements are rejected from the solid in alloys having an equilibrium partition coefficient k less than unity: k¼
Ss <1 Sl
(Eq 17)
where k is the equilibrium partition coefficient for hydrogen in aluminum, and Ss and Sl are the solubility of hydrogen in the solid and liquid, respectively (in the case of alloys, the solubility of hydrogen at the solidus/eutectic and liquidus temperatures, respectively). The solubility of hydrogen in both the liquid and solid states of aluminum has been measured by several researchers. Of the more than 26 sets of results (Ref 14) reported on the liquid state of the aluminum-hydrogen system, only 9 are assessed to be reliable (Ref 15–25). Combined regression analysis of these reliable sets of results yields the following recommended empirical equation for hydrogen solubility as a function of temperature, T (K), at 1 atm hydrogen partial pressure (Ref 26): Liquid aluminum-hydrogen : 2692 log10 S; cm3=100 g ¼ þ 2:7292 T
Of the 14 sets of results reported on the solubility of hydrogen in solid pure aluminum, only 7 are assessed to be reliable (Ref 14, 15, 26–33). Combined linear regression analysis of these results yielded the following empirical equation for the solubility of hydrogen in pure sold aluminum as a function of temperature at 1 atm hydrogen partial pressure: Solid aluminum-hydrogen: 2965:8 þ 1:8032 log10 S; cm3 =100 g ¼ T
where S is the hydrogen solubility in cm3/100 g, and T is the temperature in K. Hydrogen solubility in both liquid and solid pure aluminum follows Sievert’s law. That is, the solubility of hydrogen is directly proportional to the square root of the hydrogen pressure. Equations 18 and 19 yield values for the solubility of hydrogen in liquid and solid pure aluminum of 0.70 and 0.042 cm3/100 g, respectively, at the melting point, 933 K. This confirms the fact that the solubility of hydrogen in liquid pure aluminum is more than 16 times the solubility of hydrogen in solid aluminum, and the equilibrium partition coefficient of hydrogen in aluminum is significantly less than 1. Both recommended sets of results of hydrogen solubility in liquid and solid pure aluminum are presented graphically in Fig. 8. Alloying elements affect the solubility of hydrogen in aluminum alloys. Unfortunately, there are a limited number of published data on the solubility of hydrogen in aluminum alloys. This is especially the case for multicomponent and commercial aluminum alloys. However, alloying elements have a strong and varying influence on the solubility of hydrogen in aluminum. For example, in liquid aluminum, silicon, copper, zinc, and iron decrease hydrogen solubility, whereas lithium, magnesium, and titanium increase it (Fig. 9). The effects of temperature on the solubility of hydrogen in liquid binary aluminum alloys can be represented by the general equation of the same form as Eq 1 and 2: log10 S; cm3=100 g ¼
(Eq 18) 3
where S is the hydrogen solubility in cm /100 g, and T is the temperature in K. Unlike in the liquid state, published values of hydrogen solubility in solid pure aluminum are limited (Ref 14). This is because the solubility of hydrogen in the solid state of pure aluminum is significantly lower than that in the liquid state. Consequently, determination of the solubility of hydrogen in solid aluminum is a more formidable experimental task than the determination of the solubility of hydrogen in the liquid state. This problem is compounded by the fact that diffusion of hydrogen into aluminum in the solid state is several orders of magnitude slower than that into the liquid.
(Eq 19)
A þB T
(Eq 20)
where S is the hydrogen solubility in cm3/100 g, T is the temperature in K, and A and B are the solubility constants or coefficients that depend on alloy composition. The coefficients for the effects of temperature on the solubility of hydrogen in some liquid binary aluminum alloys, obtained from published data assessed to be reliable, are given in Table 1. In solid aluminum, published results show that silicon, magnesium, copper, lithium, and zinc increase the solubility of hydrogen (Fig. 10). The addition of lithium increases the solubility of hydrogen in solid aluminum by 1 to 2 orders of magnitude and that in liquid aluminum by two- to threefolds. Consequently, the partition coefficient of hydrogen in aluminum-lithium
68 / Principles of Liquid Metal Processing alloys is significantly greater than that in pure aluminum. This explains why aluminum-lithium alloys are less susceptible to gas porosity than other aluminum alloys. This is despite the comparatively higher hydrogen concentration in
lithium-containing aluminum alloys. Figure 11 shows the comparatively higher hydrogen solubility in the lithium-containing aluminum alloy 2090 than in other conventional aluminum alloys.
Reactions During Solidification of Aluminum and Its Alloys. Because the solubility of hydrogen is significantly higher in the liquid aluminum as compared with the solid, there will be enrichment of the liquid during solidification. Assuming no solid diffusion and complete liquid diffusion, at the end of solidification there will be a large increase in the hydrogen concentration. Therefore, even when the hydrogen content is far below the solubility for the bulk liquid during solidification, the concentration in the enriched liquid will increase, and the solubility limit may be exceeded. For example, consider pure aluminum containing a hydrogen concentration of 0.45 cm3/100 g at 750 C (1380 F). By the time the alloy is 90% solidified, the concentration in the remaining liquid will be approximately 4.03 cm3/100 g H, which will exceed the hydrogen solubility for 1 atm, resulting in the formation of hydrogen bubbles. Inert Gas Flushing. It is possible to remove hydrogen from liquid aluminum by inert gas (argon or nitrogen) flushing. As discussed pre viously for cast iron, it is possible to calculate the fastest possible rate, or the minimum amount of purge gas needed to remove hydrogen. The general form of the equation is: 1 1 kH Vt ¼ ½H ½Ho W
Fig. 8
where [H] is the hydrogen content in ppm at time t, [Ho] is the initial hydrogen content, W is the weight (in kilograms), t is the time (in minutes), kH is a constant related to the solubility, and V is the flow rate (in m s1). In practical units, at 750 C (1380 F):
Hydrogen solubility in liquid and solid pure aluminum at 1 atm hydrogen partial pressure
Temperature, ⬚F 1290 3.5
1380
1470
1560
1650
1740
Table 1 Empirical constants for the calculation of hydrogen solubility (in cm3/ 100 g) in the liquid state of some binary aluminum alloys at 1 atm hydrogen partial pressure
3.0
Hydrogen solubility, cm3/100 g
3.2 wt% Mg
4 wt% Cu Pure AI
2.5
(Eq 21)
Solubility constants, log10 S; cm3 =100 g ¼ A T þB A
B
Al-Mg alloys Al-3.2wt%Mg Al-6wt%Mg
2735.6 2695.5
2.9752 3.1452
Al-Zn alloys Al-1.8wt%Zn Al-3.5wt%Zn Al-6wt%Zn
3193.9 3058.1 2957.7
3.1392 3.0382 2.8602
Al-Si alloys Al-2wt%Si Al-4wt%Si Al-7wt%Si
2777.8 2941.3 2731.0
2.7702 2.9032 2.6562
Al-Cu alloys Al-2wt%Cu Al-4wt%Cu
2955.3 3047.7
2.9060 2.9392
Al-Li alloys Al-1wt%Li Al-2wt%Li Al-3wt%Li
2141.4 2850.7 2869.7
2.5952 3.3832 3.4862
Alloy
2.0
1.5 4.4 wt% Zn
1.0 7 wt% Si 0.5 700
Fig. 9
750
800 850 Temperature, ⬚C
900
950
Hydrogen solubility in liquid pure aluminum and binary aluminum alloys at 1 atm hydrogen partial pressure
Source: Ref 14
Temperature, ⬚F 570 0.10
660
750
840
930
1020
1110
Hydrogen solubility, cm3/100 g
0.08
0.06 7 wt% Si 5.3 wt% Mg 0.04
0.00 300
300
200
100
0
0
350
400
450 Temperature, ⬚C
500
550
1380
1470
Temperature, ⬚F 1560
1650
1740
600
1830
14.0 AA2090 12.0
10.0
ln 6.0
Pure Al AA7050
2.0 Duralumin B 0.0 700
750
800
850 Temperature, ⬚C
900
950
Hydrogen solubility in liquid multicomponent aluminum alloys at 1 atm hydrogen partial pressure
0.2
0.3
0.4
Gas removal ratio for equilibrium in aluminum at 760 C (1400 F)
1000
kg m3ðstpÞ
ðppmÞ s
(Eq 22)
For example, for 10 Mg (2.2 104 lb) of aluminum, a gas purge rate of 0.01 m3/s (0.35 ft3/s) for 100 s could reduce the hydrogen from 0.2 to 0.09 ppm. As the hydrogen content decreases, it becomes increasingly difficult to remove hydrogen (Ref 34), as indicated in Fig. 12. As the hydrogen content decreases, the ratio R of the purge gas to the hydrogen gas removed increases from approximately 15 at 0.4 ppm to over 500 at 0.1 ppm (1 cm3/100 g = 0.9 ppm). It should be emphasized that this calculation indicates the theoretical minimum amount of purge gas required. Due to kinetic factors, the amount of purge gas can be significantly higher. There are two limiting cases with respect to hydrogen removal: the thermodynamic limit discussed previously and the diffusion in the liquid film boundary layer surrounding the gas bubble. For the limiting case of diffusion control, the rate equations are similar to those discussed previously for nitrogen in iron, and the rate can be expressed by:
8.0
4.0
0.1
kH ¼ 62; 000
Hydrogen solubility in solid pure aluminum and binary aluminum alloys at 1 atm hydrogen partial pressure
1290
Fig. 11
400
Fig. 12 4 wt% Cu
Fig. 10
500
Hydrogen dissolved in aluminum (H), cm3/100 g
Pure Al 0.02
Hydrogen solubility, cm3/100 g
Gas removal ratio (R ), (standard ft3 N2)/(standard ft3 H2)
Gases in Metals / 69
%H km At ¼ H %Hi W
(Eq 23)
where %Hi is the initial hydrogen content, r is the liquid density, kmH is the mass transfer coefficient for hydrogen, A is the bubble surface area, and W is the weight of the melt. Both A and kmH increase with decreasing bubble size; consequently, the rate is sensitive to bubble size, with hydrogen removal favored by small bubbles. In general, the rate will be controlled by both diffusion and the thermodynamic limit in series. In an analysis of the case of mixed control, it was found that the purging efficiency depends on the bubble size and hydrogen content (Ref 35), as shown in Fig. 13. Purging efficiency is the ratio of the actual amount of gas required to the theoretical minimums given by Eq 23. Both
70 / Principles of Liquid Metal Processing Temperature, ⬚F
100
1110 10.0
1290
1470
1650
1830
2010
2190
2370
2550
0.05 cm3/100 g
Pure Cu
50
8.0
25 0.1 cm3/100 g 0
2 1 Bubble diameter (d B), cm
0
Fig. 13
3
Purging gas efficiency as a function of bubble size for 0.05 and 0.1 cm3/100 g of hydrogen
the interfacial area and the mass transfer coefficient increase with decreasing bubble size; consequently, the rate of hydrogen removal is enhanced by having small bubbles. The use of porous plugs helps improve the rate by providing small bubbles, but the bubbles generally coalesce in the bath, reducing the beneficial effect. The use of an impeller with a porous plug improves the rate further by dispersing the bubbles. It has also been found that the addition of chlorine (Ref 36) or Freon (E.I. Du Pont de Nemours, Inc.) (Ref 37) improves the rate in some cases. These halogen gases increase the mass transfer coefficient km and the rate. However, if the rate is being primarily controlled by the thermodynamic limit, the addition of these gases will have little effect. It is also possible that the halogen is actually taking part in a hydrogen removal reaction: 2H þ Cl2 ¼ HCl
(Eq 24)
Equation 24, under some circumstances, is thermodynamically favorable. This will increase the amount of hydrogen that can be present in the inert gas bubble; that is, hydrogen can be present as H2 or hydrogen chloride gas, thus increasing the rate of removal.
Copper and Its Alloys As in aluminum and its alloys, gas porosity is a problem in the casting of copper and copper alloys. Hydrogen is the most important dissolved gas in copper and its alloys. It forms dilute solutions in oxygen-free copper. However, in impure copper there are complex interactions among the gas-forming elements hydrogen, oxygen, and sulfur. Thus, in addition to hydrogen, gaseous compounds such as water vapor, carbon monoxide, and sulfur dioxide can also evolve during solidification. Hydrogen Solubility and Reactions. The solubility of hydrogen in pure copper has been experimentally determined by various investigators (Ref 38–47). Combined linear regression
Hydrogen solubility, cm3/100 g
Efficiency, %
75
6.0 5.9 wt% Sn
4.0 11.5 wt% Sn 21.7 wt% Sn 54.8 wt% Sn
2.0
40.2 wt% Sn
0.0 600
Fig. 14
700
800
900
1000 1100 Temperature, ⬚C
1200
1300
1400
Effect of temperature on the solubility of hydrogen in pure copper and copper-tin alloys at 1 atm hydrogen partial pressure
analyses of the sets of reported data for hydrogen solubility in liquid and solid pure copper are given by the following equations: Liquid copper-hydrogen: 2327:59 þ 2:4649 log10 S; cm3=100 g ¼ T Solid copper-hydrogen: 2057:76 þ 1:6829 log10 S; cm3=100 g ¼ T
(Eq 25)
(Eq 26)
Figure 14 is a graphical representation of the effect of temperature on the solubility of hydrogen in copper and copper-tin alloys. The difference in solubility between the liquid and solid pure copper is not as marked as for pure aluminum. At the melting point, 1083 C (1981 F), the concentrations of hydrogen in equilibrium with the gas at atmospheric pressure are 5.6 and 1.46 cm3/100 g for the liquid and solid, respectively. Thus, the partition coefficient at 0.26 for hydrogen in pure copper is comparatively higher than that in pure aluminum, where the partition coefficient is extremely low at 0.04. Hydrogen solubility in both liquid and solid pure copper is directly proportional to the square root of the hydrogen partial pressure, in accordance with Sievert’s law (Ref 38, 45). Published results on the effect of alloying elements on the solubility of hydrogen in copper are sparse. Nickel increases the solubility of
hydrogen in liquid pure copper. For example, at 1200 C (2190 F), the solubility of hydrogen in Cu-10wt%Ni is 11.2 cm3/100 g compared to 7.7 cm3/100 g in pure copper. However, large additions of aluminum, zinc, and tin decrease the solubility of hydrogen in pure copper. For example, at 1200 C (2190 F), the solubility of hydrogen in Cu-5.9wt%Sn is 6.2 cm3/100 g compared to 7.7 cm3/100 g in pure copper. The coefficients for the effects of temperature on the solubility of hydrogen in some liquid binary copper-tin alloys, obtained from data reported by Bever and Floe (Ref 38), are given in Table 2. Hydrogen can enter the copper directly from the hydrogen in the atmosphere, but most likely it enters according to: 2Cu þ H2 O ¼ Cu2 O þ 2H
(Eq 27)
H2 O ¼ 2H þ O
(Eq 28)
Reactions given in Eq 27 and 28 are favored by copper with low oxygen contents. If there are high concentrations of alloying elements with more stable oxides than cuprous oxide, such as tin, the alloying element may react to put hydrogen into solution, for example: Sn þ H2 O ¼ SnO þ 2H
(Eq 29)
The main cause of porosity in the casting results from the evolution of hydrogen or water
Gases in Metals / 71 Table 2 Empirical constants for calculating hydrogen solubility (cm3/100 g) in the liquid state of copper-tin alloys at 1 atm partial pressure Solubility constants, log10 S; cm3 =100 g ¼ A T þB Alloy
Cu-5.9wt%Sn Cu-11.5wt%Sn Cu-21.7wt%Sn Cu-40.2wt%Sn Cu-54.8wt%Sn
A
B
2284.8 2304.0 2551.0 4239.0 3410.8
2.3465 2.2962 2.3289 3.0562 2.3737
Temperature range
C
1100–1300 1000–1300 1000–1300 1000–1300 1000–1300
F
2010–2370 1830–2370 1830–2370 1830–2370 1830–2370
Source: Ref 38
vapor during solidification. The hydrogen and oxygen contents increase because of enrichment during solidification, as described previously for cast iron. As the concentrations increase toward the end of the solidification, hydrogen or water vapor is evolved by the reverse reactions of Eq 27 and 28, which become thermodynamically favorable because the hydrogen and oxygen contents in the last liquid to solidify are high. This mechanism is supported by the observation that the voids in the casting are normally distributed along the grain boundaries, which are the last parts to solidify. Other gases that can cause porosity are sulfur dioxide and carbon monoxide, which also form during solidification according to Eq 30 and 31, respectively: S þ 2O ¼ SO2
(Eq 30)
C þ O ¼ CO
(Eq 31)
The thermodynamic pressure of carbon monoxide gas can be very high, even at moderate oxygen and carbon contents, because the solubility of the carbon in copper is limited, and therefore, its chemical activity is high. Porosity in Copper Alloys. In copper ingots, gas porosity usually results from H2 and water evolution. The evolution of as little as 1.12 cm3/100 g of hydrogen results in a gas evolution in 44% of the volume of the metal. If the sulfur content exceeds 0.05 wt%, the evolution of sulfur dioxide may also contribute to gas porosity. Only if the oxygen content exceeds 0.01% is carbon monoxide considered to be the cause of porosity (Ref 48). Because phosphorus deoxidizes copper significantly, the oxygen content of phosphorized copper is very low, and water, carbon monoxide, and sulfur dioxide do not contribute to porosity; only H2 need be of concern. Copper-zinc alloys rarely have problems associated with gas porosity, primarily because zinc deoxidizes the metal. Consequently, brasses do not dissolve significant quantities of oxygen. Zinc may also remove dissolved gases because of its high vapor pressure, which results in the zinc vapor flushing hydrogen out of the melt. During casting of copper-zinc alloys, degassing is not usually required. This is because the quantities of hydrogen absorbed by industrially produced copper-zinc melts are usually insufficient to cause problems in castings. The lowest hydrogen contents are usually
found in brasses with the highest zinc contents. However, the copper-tin alloys are the copper alloys that are most susceptible to gas porosity. This is because tin significantly increases the solubility of hydrogen in copper, tin mildly deoxidizes copper, and copper-tin alloys have long freezing ranges. Degassing of Copper Alloys. The two primary methods of removing dissolved gases from copper alloys are oxidation-reduction and inert gas flushing. Although vacuum degassing is theoretically possible, it is rarely used because it is not cost-effective. In oxidation-reduction, the first step is to remove the hydrogen by oxidation using an oxidizing slag or oxygen-rich copper. This involves increasing the activity of oxygen in the melt by melting it under an oxidizing flame or covering it when molten with an oxidizing flux such as a mixture of cupric oxide-boric acid-sodium chloride-sodium fluoride (Ref 49). This drives off the hydrogen in the form of steam. Then, just prior to pouring, the melt is deoxidized by adding phosphorus or other deoxidizers. The phosphorus deoxidizer produces cuprous phosphate, a molten deoxidation product that separates fairly easily from the molten copper metal. Generally, an addition of 0.02 to 0.03 wt% P is sufficient to deoxidize bronzes and gun metals. Calcium boride, boron carbide, and lithium have also been used for deoxidation (Ref 50). A charcoal cover is often used for deoxidation and protection of the melt from reoxidation. For copper alloys that contain strong deoxidizers, such as phosphorus, zinc, and tin, it is not possible to remove hydrogen by oxidation, because these elements will form stable oxides. For these alloys, inert gas flushing with nitrogen is commonly used. The theoretical minimum amount of inert gas is determined by using relationships similar to those derived for iron and aluminum. In practice, 150 to 250 L (5.3 to 8.8 ft3) of nitrogen per 1000 kg (2200 lb) of copper alloy is generally recommended (Ref 51).
Magnesium and Its Alloys Magnesium and its alloys are susceptible to gas-induced porosity. Of the nonferrous metals considered in this article, magnesium dissolves more hydrogen and has the potential for forming hydrides. Magnesium hydride, MgH2, is stable only for very high hydrogen activities
that are not normally encountered in metallurgical operations. At comparable temperature and hydrogen partial pressure, the solubility of hydrogen in magnesium and its alloys is an order of magnitude of that in aluminum and copper and their alloys. However, the propensity for gas porosity formation in magnesium is less than that in aluminum and copper. The principal source of hydrogen in magnesium and its alloys is moisture in the surrounding atmosphere or introduced into the melt when corroded materials are charged or damp fluxes are used. Magnesium and its alloys readily absorb hydrogen by the reaction with water vapor: MgðlÞ þ H2 OðvÞ ! MgOðsÞ þ H2 ðin MgÞ
(Eq 32)
The standard free energy change, DGo, for this reaction is extremely high: Go ¼ 368; 600 19:8T log10 T þ 122T J=mol
(Eq 33)
where T is the temperature in K. Application of the van’t Hoff reaction isotherm yields a value for aH2 =aH2 O that is even higher than the corresponding quotient for aluminum. In the absence of hydride formation, hydrogen dissolution in magnesium and its alloys is in the atomic form. Hydrogen Solubility and Reactions During Solidification. The solubility of hydrogen in pure magnesium has been experimentally determined by several investigators (Ref 52–59). However, the measurement of the solubility of hydrogen in magnesium and its alloys is an experimentally difficult task. This is because magnesium is too volatile and reactive toward traces of water vapor for direct measurement with the Sievert’s method. Combined linear regression analyses of the sets of reported data considered to be reliable yielded the following empirical equations for hydrogen solubility in liquid and solid pure magnesium as a function of temperature: Liquid magnesium-hydrogen: log10 S; cm3=100 g 1290:5 þ 3:042 ¼ T =K
(Eq 34)
Solid magnesium-hydrogen: log10 S; cm3=100 g 1062:47 þ 2:6254 ¼ T =K
(Eq 35)
At the melting point, 651 C (1204 F), the concentrations of hydrogen in equilibrium with the gas at atmospheric pressure are 44.2 and 29.9 cm3/100 g for the liquid and solid, respectively. Thus, the partition coefficient at 0.68 compared to 0.27 and 0.04 for hydrogen in pure magnesium is much higher than that in copper and aluminum, respectively. Hydrogen solubility in both the liquid and solid pure magnesium
72 / Principles of Liquid Metal Processing Temperature, ⬚F 750
570 120.0
930
1110
1290
1470
1650
1830
Pure Mg
Table 3 Coefficients for calculating hydrogen solubility (cm3/100 g) in the liquid state of magnesium alloys at 1 atm partial pressure Solubility constants, log10 S; cm3 =100 g ¼ A T þB
5.3 wt% Al
2 wt% Sr 1.3 wt% Zn Hydrogen solubility, cm3/100 g
A
B
Mg-Al alloys Mg-5.3wt%Al Mg-8.3wt%Al
1127 1201.5
2.8219 2.8532
Mg-Zn alloys Mg-1.3wt%Zn Mg-3.2wt%Zn Mg-6.8wt%Zn
1253.6 1074.7 1262
2.9733 2.7958 2.9577
Mg-Sr alloys Mg-0.5wt%Sr Mg-2wt%Sr Mg-3.4wt%Sr
1198.9 1028.3 886.9
2.9473 2.8323 2.7487
Mg-Si alloy Mg-1wt%Si
1126.1
2.875
Alloy
1 wt% Si
100.0
80.0
60.0
Derived from results given in Ref 38
40.0
that in aluminum and copper alloys. This can be attributed to the following reasons: 20.0
The significantly higher partition coefficient
0.0 300
400
500
600
700
800
900
1000
Temperature, ⬚C
Fig. 15
Solubility of hydrogen in magnesium and magnesium alloys as a function of temperature and at 1 atm hydrogen partial pressure
Hydrogen solubility, cm3/100 g
60
55
50
Ag Sn Cu Zn Sb
45
40 0
2
6 8 4 Alloying element, wt%
10
12
Fig. 16
Effect of alloying elements on the solubility of hydrogen in binary magnesium alloys at 700 C (1290 F) and 1 atm hydrogen partial pressure
is directly proportional to the square root of the hydrogen partial pressure, in accordance with Sievert’s law (Ref 52, 53). Alloying elements affect the solubility of hydrogen in magnesium (Fig. 15, 16). With the exception of nickel, zirconium, and strontium, the addition of alloying elements to magnesium generally results in a decrease in the solubility of hydrogen in magnesium. For example, the solubility of hydrogen in both liquid and solid
magnesium decreases with an increase in the addition of zinc, silicon, aluminum, lead, silver, tin, manganese, antimony, and copper (Ref 57, 60). However, because it apparently forms hydrides in the molten aluminum, strontium increases the solubility of hydrogen in liquid magnesium (Ref 57). Although zirconium increases the solubility of hydrogen significantly in solid magnesium, at 700 C (1290 F) it does not have any discernible effect on the solubility of hydrogen in liquid magnesium (Ref 60). In liquid magnesium-zirconium alloys, zirconium apparently acts as a scavenger for hydrogen, precipitating as zirconium hydride, ZrH2, and consequently obviating formation of gasinduced porosity in the castings. In the solid state, the absorbed hydrogen is also precipitated within the structure as zirconium hydride. At above 1 wt%, the addition of nickel significantly increases the solubility of hydrogen in liquid magnesium. The coefficients for calculating the solubility of hydrogen as a function of temperature in some binary magnesium alloys are given in Table 3. Porosity in Magnesium and Its Alloys. Castings of magnesium alloys are susceptible to gas porosity. The prevalence of gas porosity apparently increases with an increase in the hydrogen content of the molten metal. However, the susceptibility of magnesium alloys to gas porosity formation is significantly less than
of hydrogen in magnesium and its alloys reduces hydrogen evolution during solidification. The nominally higher solubility of hydrogen in liquid magnesium alloys makes the nucleation of gas porosity during solidification rather difficult. Nucleation of the gas bubble requires that the hydrogen concentration in the solidifying melt should exceed the equilibrium solubility of hydrogen in the melt. The comparatively higher partition coefficient of hydrogen also makes the characteristic of gas porosity in magnesium alloys different from that in aluminum and copper alloys. The gas porosity in magnesium alloys is more interdendritic in nature, more like a shrinkageinduced porosity. This is due to the proportionately much lower decrease in hydrogen solubility during solidification, so that a hydrogen content in the residual liquid sufficient to nucleate gas bubbles is not reached until at a later stage in solidification, if at all. Degassing of Magnesium Alloys. Degassing of magnesium and its alloys is usually accomplished by bubbling nitrogen, helium, nitrogen-chlorine mixture, or chlorine through the molten metal (Ref 61). The most effective of the gas-flushing methods is the bubbling of chlorine through the molten metal at 720 to 780 C (1330 to 1440 F). It can be used to reduce hydrogen content to 2 to 3 cm3/100 g. Bubbling of nitrogen reduces hydrogen to approximately 4 cm3/100 g. Other degassing methods that have been shown to be effective in reducing the hydrogen content of magnesium and its alloys include vacuum degassing and treatment with a flux. Treatment with a flux consisting of 55 wt% KCl, 34 wt% MgCl2, 9 wt% BaCl2, and 2 wt% CaCl2 apparently decreases the
Gases in Metals / 73 hydrogen content to 3 to 4 cm3/100 g. Although more environmentally friendly, vacuum degassing is not as effective as the gas-bubbling methods.
Overcoming Gas Porosity Gas in cast iron, aluminum, copper, and magnesium can be a major problem. The porosity results from reactions with the environment or the evolution of gases during solidification. In determining if gas evolution can cause porosity, the thermodynamics for the reactions must be favorable. In this article, the solubilities of the common gases and the reactions during solidification were discussed. Because of enrichment during solidification, the concentration of the elements increases to the extent that the thermodynamic pressure of the dissolved gases may exceed 1 atm. In addition to the thermodynamics, kinetic factors and interfacial energies determine if the gases will evolve. The most common method for removing hydrogen from aluminum, copper, and magnesium, which can also be used for removing hydrogen and nitrogen from cast iron, is inert gas flushing. The theoretical minimum amount of flushing gas required for removing hydrogen and nitrogen can be computed. In practice, the amount of inert gas required is somewhat higher because of kinetic factors. The kinetics of inert gas flushing can be improved by the dispersion of small bubbles throughout the melt. ACKNOWLEDGMENT Updated from Richard J. Fruehan, Gases in Metals, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 82–87. REFERENCES 1. R.J. Fruehan, in Proceedings of the Physical Chemistry of Foundry Processes Symposium (Warren, MI), General Motors Corporation, 1986 2. S. Katz and C. Landefeld, General Motors Research, private communication, 1986 3. Making, Shaping and Treating of Steel, 10th ed., United States Steel Corporation, 1985 4. A. Kagawa and T. Okamoto, Trans. Jpn. Inst. Met., Vol 22 (No. 2), 1981, p 137 5. R.D. Pehlke and J. Elliot, Trans. TMSAIME, Vol 227, 1963, p 894 6. M. Inouye and T. Choh, Trans. JISI, Vol 8, 1968, p 134 7. R.J. Fruehan and L.J. Martonik, Metall. Trans. B, Vol 11, 1980, p 615 8. P.C. Glaws and R.J. Fruehan, Metall. Trans. B, Vol 16, 1985, p 551 9. P.C. Glaws and R.J. Fruehan, Metall. Trans. B, Vol 17, 1986, p 317
10. M. Byrne and G.R. Belton, Metall. Trans. B, Vol 14, 1983, p 441 11. F. Tsukihashi and R.J. Fruehan, Trans. JISI, Vol 27, 1987, p 858 12. R.J. Fruehan, Ladle Metallurgy: Principles and Practices, Iron and Steel Society of AIME, 1985 13. R.J. Fruehan, B. Lally, and P.C. Glaws, in Proceedings of the Fifth International Iron and Steel Congress (Washington, D.C.), Iron and Steel Society of AIME, 1986 14. P.N. Anyalebechi, “A Comprehensive and Critical Review of Reported Hydrogen Limits in Aluminum and Its Alloys,” unpublished research, 2000 15. C.E. Ransley and H. Neufeld, J. Inst. Met., Vol 74, 1947–1948, p 599–620 16. W. Baukloh and F. Oesterlen, Z. Metallkd., Vol 30 (No. 11), 1938, p 386–389 17. W.R. Opie and N.J. Grant, J. Met. (Trans. AIME), Vol 188, 1950, p 1237–1241 18. C.E. Ransley and D.E.J. Talbot, Z. Metallkd., Vol 46 (No. 5), 1955, p 328–337 19. D.E.J. Talbot and P.N. Anyalebechi, Mater. Sci. Technol., Vol 4 (No.1), 1988, p 1–4 20. P.N. Anyalebechi, Ph.D. thesis, Brunel University, Middlesex, England, 1985 21. D. Stephenson, Ph.D. thesis, Brunel University, Middlesex, England, 1978 22. H. Shahani, Ph.D. thesis, The Royal Institute of Technology, Stockholm, Sweden, 1984 23. J. Kocur, K. Tomasek, and L. Rabatin, Hutn. Listy, Vol 44 (No. 4), 1989, p 269– 275 24. H. Liu, L. Zhiang, and B. Bouchard, Recent Developments in Light Metals, M. Gilbert, P. Tremblay, and E. Ozberk, Ed., The Metallurgical Society of Canadian Institute of Metallurgists, 1994, p 257–268 25. H. Liu, B. Bouchard, and L. Zhiang, Light Metals 1995, J. Evans, Ed., The Minerals, Metals and Materials Society, 1995, p 1285–1291 26. P.N. Anyalebechi, Light Metals 1998, B. Welch, Ed., TMS, 1998, p 827–842 27. W. Eichenauer and A. Pebler, Z. Metallkd., Vol 48, (No. 7), 1957, p 373–378 28. W. Eichenauer, K. Hattenbach, and A. Pebler, Z. Metallkd., Vol 52, (No. 10), 1961, p 682–684 29. W. Eichenauer, Z. Metallkd., Vol 59, 1968, p 613–616 30. M. Ichimura, M. Imabayashi, and M. Hayakawa, Nippon Kinzoku Gakkaishi, Vol 43, (No. 9), 1979, p 876–883 31. E. Hashimoto and T. Kino, J. Phys. F., Met. Phys., Vol 13, 1983, p 1157–1165 32. M. Ichimura, Y. Sasajima, and M. Imabayashi, Mater. Trans., JIM, Vol 32 (No. 12), 1991, p 1109–1114 33. M. Ichimura, Y. Sasajima, and M. Imabayashi, Mater. Trans., JIM, Vol 33 (No. 5), 1992, p 449–453 34. G.K. Sigworth, Trans. AFS, Vol 95, 1987, p 73–78
35. G.K. Sigworth and T.A. Engh, Metall. Trans. B, Vol 13, 1982, p 447–460 36. J. Botor, Metal. Odlev, Vol 6, 1980, p 21–27 37. J. Botor, Aluminum, Vol 56 (No. 8), 1980, p 519–522 38. M.B. Bever and C.F. Floe, Trans. AIME, Vol 156, 1944, p 149–159 39. A. Sieverts, Z. Phys. Chem., Vol 77, 1911, p 11 40. P. Rotgen and F. Moller, Metallwirtschaft, Vol 13, 1934, p 81, 97 41. E. Kato, H. Ueno, and T. Orimo, Trans. Jpn. Inst. Met., Vol 11, 1970, p 351 42. R.B. McLellan, J. Phys. Chem. Solids, Vol 34, 1973, p 1137–1141 43. C. Thomas, Trans. Met. Soc. AIME, Vol 239, 1967, p 485–490 44. F.G. Jones and R.D. Pehkle, Met. Trans., Vol 2, 1971, p 2655–2663 45. M. Mokaram, Ph.D. thesis, Brunel University, London, 1974 46. W. Eichenauer and A. Pebler, Z. Metallkd., Vol 48, 1957, p 373–378 47. W. Eichenauer et al., Z. Metallkd., Vol 56, 1965, p 287 48. R.H. Waddington, J. Inst. Met., 1948– 1949, p 311 49. W.T. Pell-Walpole, J. Inst. Met., Vol 70, 1944, p 127 50. E.R. Thews, Metall., June 1956, p 431 51. W.A. Baker and F.C. Child, J. Inst. Met., Vol 70, 1944, p 349 52. J. Koeneman and A.G. Metcalfe, Trans. ASM, Vol 51, 1959, p 1072–1082 53. T. Watanabe, Y.C. Huang, and R. Komatsu, J. Jpn. Inst. Light Met., Vol 46, (No. 2), 1975, p 76–81 54. Y.C. Huang, T. Watanabe, and R. Komatsu, Proc. Fourth Int. Conf. Vacuum Metallurgy, 1973, p 173–176 55. D. F. Chernega, Gotvjanskij and Prisjasznjokm Bestn. Kiev. Politexn. In-Ta, Kiev, Vol 17, 1959, p 55–58 56. V. Shapavalov, N.P. Serdyuk, and O.P. Semik, Dop. Akad. Nauk Ukr. RSR. Ser. A. Fiz.-Mat. Tekh. Nauki, Vol 6, 1981, p 99–101 57. E. Overlid, P. Bakke, and T.A. Engh, Light Metals 1997, Metaux Legers, Canadian Inst. of Metallurgists, 1997, p 141–154 58. Z.D. Popovic and G.R. Piercy, Met. Trans. A, Vol 6, 1975, p 1915–1917 59. M.V. Sharov and V.V. Serebryakov, Nauchnyye Doklady Vvyshe Shkoly, Metallurgiya, (No. 2), 1958, p 37 60. T. Watanabe, Y. Tachihara, Y.C. Huang, and R. Komatsu, Kei Kinzoku, Vol 26 (No. 4), 1976, p 167–174 61. T. Watanabe, Y.C. Huang, R. Komatsu, and T. Ito, J. Jpn. Inst. Light Met., Vol 26 (No. 6), 1976, p 266–272
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 74-83 DOI: 10.1361/asmhba0005193
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Inclusion-Forming Reactions Revised by Paul K. Trojan, University of Michigan—Dearborn; and Richard Fruehan, Carnegie Mellon University
INCLUSIONS can be defined as nonmetallic and sometimes intermetallic phases embedded in a metallic matrix. They are usually simple oxides, sulfides, nitrides, or their complexes in ferrous alloys and can include intermetallic phases in nonferrous alloys. In almost all instances of metal casting, they are considered to be detrimental to the performance of the cast component. For example, mechanical properties can be adversely influenced by inclusions, which act as stress raisers. There is no set pattern for the effect, but some properties are more sensitive to the presence of inclusions than others. Elongation or reduction in area is usually modified more significantly than ultimate tensile strength. As a result, ductility specifications are common quality-control indices in cast products. Similar observations are found with porosity due to gas or occluded shrinkage. The loss in casting properties measured by a tensile test may then reflect a combination of defects; of these defects, inclusions are only a single but very important source of poor performance. Because imperfections become areas of higher stress concentration, the percentage of property loss becomes greater when the strength requirement is higher. An important example is the decrease in fracture toughness when inclusions are present in higher-strength lower-ductility alloys. Similar pronounced property degradation is observed in tests that reflect slow, rapid, or cyclic strain rates, such as creep, impact, and fatigue testing. In all cases, the degree of property loss will depend on the parameters that control the characteristics of all microstructures. These metallurgical phase controls are the nature, amount, size, shape, distribution, and orientation of the phases. If inclusions are to be controlled, these characteristics must be suitably modified to minimize their impact on casting performance. Although inclusions are generally considered to be detrimental, their intentional introduction in larger quantities can lead to unique dispersion-strengthened materials. An example would be the dispersion of insoluble silicon carbide whiskers or alumina (Al2O3) fibers in a molten aluminum alloy that is subsequently cast into a final shape by squeeze casting methods (Ref 1).
Inclusion Types There are essentially two classifications for all inclusions: Exogenous: those derived from external
causes
Indigenous: those that are native, innate, or
inherent in the molten metal treatment process Exogenous Inclusions. Slag, dross, entrapped mold materials, and refractories are examples of inclusions that would be classified as exogenous. In most cases, these inclusions would be macroscopic or visible to the naked eye at the casting surface (Fig. 1). When the casting is sectioned, they may also appear beneath the external casting surface if they have had insufficient time to float out or settle due to
1
Fig. 1
density differences with respect to the molten metal. Indigenous inclusions include sulfides, nitrides, and oxides derived from the chemical reaction of the molten metal with the local environment. Such inclusions are usually small and require microscopic magnification for identification. They are often uniformly distributed within the microstructure. However, a grainboundary distribution of indigenous inclusions (Fig. 2) can be particularly damaging to mechanical properties. There are cases in which indigenous inclusions such as oxides may react with mold materials to promote an exogenous inclusion variety that otherwise may not appear in the absence of the indigenous oxide. To circumvent the strict definition as indigenous or exogenous, a number of investigators prefer to catalog all inclusions as either macroinclusions of microinclusions.
2
Ductile iron crankshaft segment essentially free of exogenous inclusions (1, left) and with numerous exogenous inclusions (2, right). Low pouring temperature and poor mold-filling practice were the cause of the inclusions in part 2.
Inclusion-Forming Reactions / 75 Oxygen, at.% 49 50
51
52
53
54
55
56
L2
L1 + L 2
(δ-Fe) 1394⬚ 1373⬚
59
60
61
62
63
64
L2 + O2 (g; 1atm) 28.30
3000
1585⬚
28.07 Fe O + O (g; 1atm) 3 4 2 1459⬚
22.60
1500 1400
58
1599⬚
1600 1529⬚
57
25.31
22.84
1426⬚
28.36
25.60 27.64
2800
30.04
2600
22.91 23.16
2400
1300 FeO (Wustite)
Temperature, ⬚C
1200
2200 Fe3O4 (Magnetite)
1100
2000 (γ-Fe)
Fe2O3 + O2 (g; 1atm)
1000 912⬚
23.10
900
1800
Fe2O3 (Hematite)
Temperature, ⬚F
1700
47 48
1600 Curie temperature
800
Fig. 2
Type II indigenous sulfide inclusions in a 0.25% C steel. Etched using 2% nital. Original Magnification: 500
600
Although such definitions describe the anticipated effect on mechanical properties, there can be a loss in source identification of the inclusion. Whatever definition is used, control of inclusions requires an understanding of their origin and the associated physical chemistry.
770⬚
1400
(α-Fe)
1200
700
560⬚ 23.26
1000
500 400 20
800
30.06 21
22
23
24
25
26
27
28
29
30
31
32
33
34
Oxygen, wt%
Physical Chemistry
Fig. 3
Phase diagrams, thermochemical relationships, and reaction rates must be considered for proper understanding of inclusion formation. The general concepts of each, along with their practical significance, are established in the following sections in this article. Phase Equilibria. Phase diagrams map the occurrence of phases to be expected under equilibrium conditions. Equilibrium also requires long periods of time to achieve. This is certainly not the case in castings, in which even moderate cooling rates may inhibit equilibrium phases. For example, in the iron-oxygen system shown in Fig. 3, wustite (FeO) produced at high temperatures should break down to a-iron and magnetite (Fe3O4) at 560 C (1040 F). However, iron scales may show considerable amounts of retained FeO at room temperature. A further limitation is that pure elements seldom exist in the casting process, and more than two elemental species are usually present. In Fig. 3, pure oxygen is one component, while air (20% O2) is more likely to be the oxidizing source during casting. Furthermore, a cast steel will contain carbon and other elements, such as silicon or aluminum, that will have a greater affinity for oxygen than iron. Therefore, phase diagrams become a tool for understanding, but the predicted equilibrium for inclusion occurrence may not always be observed
for the casting cooled under nonequilibrium conditions. Thermochemistry. Any chemical reaction will occur if the products are at a lower energy level than the reactants. The energy measurement is the chemical free energy, which can be treated simply as the energy of chemical position. Free energy values can be obtained directly from numerous sources or can be calculated from tabulated enthalpy, entropy, and heat capacity data. An equilibrium is then established between the products and reactants such that very little remaining reactant indicates a more or less complete reaction. A more thorough discussion of chemical free energy is provided in standard textbooks on chemical thermodynamics. By way of example, assume pure molten iron contains dissolved oxygen at 1600 C (2910 F). A small amount of aluminum is added to the melt, and the chemical reaction is: 2Al þ 3O ! Al2 O3 ðsÞ where the underlined indicates the elements to be dilute and dissolved in iron while the (s) indicates that Al2O3 is solid at 1600 C (2910 F). (The melting point of Al2O3 is 2072 C, or 3762 F.) For the reaction to occur, the standard free energy exchange at 1600 C (2910 F) must be negative. However, the quantity of each component in the reaction is also important, and the chemical activity of each species must
Iron-oxygen phase diagram
be known. An overabundance of aluminum will react stoichiometrically with oxygen, and the remainder may go into liquid solution or react with the surroundings. In the previous example, the Al2O3 product is insoluble in the molten iron and can result as inclusions trapped in the solidified casting. It is then possible to calculate the free energy exchange for the occurrence of various inclusion chemistries. Their relative stability is dictated by the most negative value of free energy. In the previous example, iron oxide would not occur, because Al2O3 has a lower free energy of formation. In fact, iron oxide would be predictably reduced by aluminum to form Al2O3. As with phase equilibria, thermochemical calculations are again based on equilibrium data, and castings may not be at equilibrium throughout the solidification process. Kinetics. There are several reasons why inclusion-forming reactions are usually incomplete and why reaction rates therefore become the most important practical consideration. Much of the understanding of reaction rates is associated with the following observations: Not all species diffuse at the same rate; when
coupled with the dependence of diffusion rates on temperature, nonuniform inclusion formation and distribution may be anticipated.
76 / Principles of Liquid Metal Processing Chemical free energy is also temperature-
Experimental kinetic data can be presented in several ways, including mathematical relationships associated with activation energy and first-order kinetic reaction rate theory. Perhaps more familiar are isothermal and continuous cooling transformation graphical interpretations consistent with nucleation and growth-controlled phenomena.
3990
2200
Liquid 3630
2000
1840⬚ 3270
1800 Cristobalite + liquid
Corundum + mullite solid solutiion
Mullite + liquid 1595⬚
1600
2910 Mullite solid solution
Cristobalite + mullite
1400 0
20
Temperature, ⬚F
1850⬚
40
60
SiO2
80 Mullite 3Al2O3.2SiO2
2550 100 Al2O3
Al2O3, wt%
Control of Inclusions The previous discussion suggests several practical means of controlling the occurrence of inclusions in any cast metal. All of the control methods can be cataloged as chemical, mechanical, or a combination of these. Several of the more common techniques are discussed separately as follows. Chemistry Control. Although control of alloy chemistry may be obvious, composition cannot always be modified to minimize inclusions. The metal compositions may be inherently reactive with the local environment. For example, the presence of magnesium is required in nodular cast iron; the propensity for inclusions is greater than for the same nominal chemistry of a gray cast iron. Similarly, the existence of inclusions is accepted because the alternative is worse. Aluminum added to steels may result in Al2O3 inclusions, but the lack of deoxidation if aluminum were not used would result in gas holes with a greater negative effect on casting performance. However, this does not imply that if a small addition is good then more is better. Phosphorus added to a copper-base alloy can be an effective deoxidizer, but excess phosphorus can be very aggressive and can react with mold materials commonly used with this alloy family. Therefore, the importance of chemistry control cannot be overemphasized if inclusions are to be controlled. Mold/Metal Interface and Refractory Control. Another chemical variable that can lead to inclusions is that of the mold materials and refractories. The SiO2-Al2O3 phase diagram is shown in Fig. 4. It is obvious that the solidus temperature increases with Al2O3
Corundum + liquid
Mullite + liquid Temperature, ⬚C
dependent and may not vary in the same fashion for all of the species under consideration. Available phase equilibrium data may predict the relative solubility of inclusion-forming species in the liquid state but may not treat decreasing solubility in the cooling liquid or the effects of the solidification process on the equilibrium. The surface energy of species will change with temperature and will influence the ability of an inclusion to float, sink, agglomerate, or react secondarily with its surroundings. Reaction products at a surface can become a barrier to further reaction and can block completion of the chemical reaction.
Fig. 4
SiO2-Al2O3 phase diagram. Source: Ref 2
content. Even if a refractory is to be used well below the solidus temperature, a high-Al2O3 refractory may still be preferred because Al2O3 plus mullite would be present rather than the silica (SiO2) plus mullite that would be present at compositions less than 72% Al2O3. The SiO2 does not possess the same thermochemical stability as Al2O3; therefore, aggressive liquid metals or oxides may react more readily with SiO2. Higher temperatures and longer holding times at temperature increase the possibility for reaction and the production of inclusions associated with refractory wear and erosion. Therefore, melt overheating, especially for long periods of time, is to be avoided. The scale on scrap surfaces can also lead to refractory erosion and exogenous inclusions. Rust on ferrous scrap can be considered a hydrated iron oxide scale that, when heated and dehydrated, can form low-melting complex oxides with the refractories. Inclusion difficulties have also been known to increase when adhering mold materials remain on charged revert scrap. The severity depends on the nature of the mold bonding agent; bentonite and organically bonded mold materials would display different tendencies toward refractory attack. The aforementioned observations can be applied to reaction between the mold materials and the molten metal. This is not to be confused with metal penetration, in which molten metal merely fills the void spaces between grains of
the mold material. However, the solution to the problem—for example, a mold wash— may be the same. In this case, the wash fills the voids between sand grains and isolates the mold material from contact with a potentially aggressive molten metal or oxide. Longer solidification times merely increase the possibility for mold-metal reaction and the accompanying inclusion-forming tendency. Separation Techniques. The choice for inclusion separation is between removal before the molten metal enters the mold or use of a removal system as a component part of the mold. Gating systems are discussed in the next section in this article. In both cases, inclusion control is basically mechanical. External methods include simple skimming of ladle slag or dross to minimize exogenous inclusion sources before pouring of the castings. A teapot ladle can be very effective for decanting clean metal from beneath a layer of floating slag or dross. Screens or sieves have been used in gating systems to filter inclusions. In general, metallic filters suffer from the tendency to alloy with the poured metal, and they are difficult to separate from the charge returns. Ceramic foam filters with consistent pore size and structural integrity are being increasingly used for those alloys from which definite cost benefits can be derived (Ref 3–5). Gating Methods. Gating and pouring practice has been known to create and to solve inclusion problems in castings. If exogenous
Inclusion-Forming Reactions / 77
500
475 68
64 425 60 400 56
Alternating stress, ksi
Alternating stress, MPa
450
375 Conventional 52 Calcium treated
350
48
325 104
105
106
107
Cycles to failure
Fig. 5
Comparison of axial fatigue data for untreated and calcium-treated rolled ASTM A516 steel. 51 mm (2 in.) thick plates tested with alternating stress ratio of 0.1. Source: Ref 8
Yield strength, ksi 30 320
50
70
90
110
130 240
280
180
200 160
120
120 Calcium treated 80
60
40 Conventional
0 −40 200
Charpy V-notch USE, ft · lb
240 Charpy V-notch USE, J
inclusions leave the ladle because of inadequate upstream control, the gating system should be designed to remove such inclusions in the pouring basin or in the runner system. Too often, a pouring basin is not designed to fill easily or to remain filled, and inclusions that may float are forced into the casting cavity. Whirl gates can be used to collect the inclusions and to ensure a better opportunity to have clean metal enter the casting proper (Ref 6). In addition, runners can have dead-end extensions that collect the first metal to enter the mold, which is the most likely to contain slag or dross (Ref 7). Pressurized gating systems are often recommended for preventing the aspiration of air that can cause the formation of inclusions during the pouring process. This technique is more important when the metal or alloy has a great tendency toward inclusion formation because of so-called secondary oxidation. More detailed information on correct gating is available in the article “Gating Design” in this Volume; many of the gating practices discussed in that article are used to minimize inclusions. Inclusion Shape Control. As pointed out previously, mechanical properties suffer when inclusions are present, with a more pronounced degradation in strain-rate-sensitive applications. For example, an inclusion at the surface of a casting subjected to cyclic strain can become a source for fracture initiation in fatigue. The most conservative approach in design is to determine the smallest inclusion at the surface that may be present without casting rejection. With no variation in inclusion continuity or size, a spherical shape (lower microstress concentration) would be less damaging than more angular shapes. Inclusions beneath the surface would exhibit the same dependence on shape for fracture propagation rather than initiation. As shown in Fig. 5, calcium treatment of a steel provides longer life under cyclic loading due to inclusion shape modification. Although the data in Fig. 5 are for a rolled constructional steel, similar results would be anticipated for a cast steel, where anisotropy also exists due to directional solidification. The advantage of a spherical inclusion shape is possibly more evident in Charpy V-notch impact data for constructional steels (Fig. 6). SEM fractographs show a pronounced difference in deformation mode with an inclusion shape modification (Fig. 7). Local ductility increases adjacent to the inclusions when the shape is spherical. This would then suggest that if inclusions are to occur and cannot be completely removed, a spherical shape is the least detrimental. In these examples, the MnS inclusions assume a planar shape that allows easier fracture propagation between inclusions or lower observed alloy ductility. The addition of calcium converts the inclusion to essentially CaS with a much higher melting point, spherical shape, and hence, higher alloy ductility. Inclusion shape control, therefore, is a useful technique that is often applied to more demanding casting applications in which added
300
400
500
600
700
800
900
0 1000
Yield strength, MPa
Fig. 6
Comparison of Charpy V-notch upper-shelf energies (USEs) for several grades and thicknesses of untreated and calcium-treated steel. Source: Ref 9
processing costs can be justified. The principle is a control of surface energy between the inclusion and the metal—normally through suitable chemistry modification of the metal. By analogy, the technique may not be unlike the addition of magnesium to cast iron to change the graphite shape from flakes to nodules.
available for modifying the viscosity of slag or dross, thus facilitating their ease of removal. On the other hand, treatment of steels and cast irons is likely to be different, and a few examples, along with the application of appropriate physical chemistry, are given as follows.
Steels
Inclusions in Ferrous Alloys Some applications for inclusion minimization or shape control are immediately obvious. For example, proprietary addition agents are
Oxides. The amount of carbon in iron controls the amount of dissolved oxygen (Fig. 8). The curve has been calculated from thermochemical data with variable CO/CO2 ratios. Experimental points would generally lie above
78 / Principles of Liquid Metal Processing
(a)
(b)
Fig. 7
SEM fractographs showing the effect of calcium treatment on the fracture morphology of ASTM A633C steel impact specimens. (a) Untreated steel with type II manganese sulfide inclusions showing evidence of brittle fracture. (b) Calcium-treated steel with spherical inclusions; the fracture is ductile. Courtesy of A.D. Wilson, Lukens Steel
Company
1.0
0.10
0.5 1705 ⬚C 1600 ⬚C 1515 ⬚C
0.2
0.08
2Fe
O+
0.1
Si
SiO
0.06
2 +2 Fe
0.02 FeO, %
Oxygen, wt%
0.05
0.04
0.01
3Fe
O+
0.005
2A
l Al
0.002 0.02
2O 3 +3 F
e
0.001
Equilibrium with CO + CO2 at 1 atm pressure and 1600 ⬚C
0.0005 0.0002
0
0
0.2
0.4
0.6
0.8
0.0001 0.001 0.002
0.005 0.01 0.02
0.05
0.1
0.2
0.5
1.0
2.0
5.0
10
20
50
Carbon, wt% Alloy remaining in steel, %
Fig. 8
Carbon-oxygen equilibrium at 1600 C (2910 F). Source: Ref 10
the theoretical curve, but the relationship between carbon and oxygen would be the same. Low carbon in solution, as in many cast steels, results in higher dissolved oxygen. The solid solubility of oxygen in iron is low, and during solidification, oxygen is rejected to the remaining liquid, where it can react with carbon to form CO and result in gas porosity. Although the porosity could occur in rimmed steel ingots (induced porosity that is “healed” during subsequent processing), the amount of dissolved oxygen in steel castings is reduced by additions of aluminum or silicon. The resultant oxides are insoluble and form inclusions.
Fig. 9
Equilibrium between FeO and dissolved aluminum and silicon. Source: Ref 11
Because Al2O3 has a lower free energy of formation than SiO2, aluminum is a more effective deoxidizer than silicon (Fig. 9). The FeO content is an index of the amount of dissolved oxygen in the steel where the oxygen can also be reduced by carbon, as shown in Fig. 8. At a given carbon content, the oxygen can then be reduced by aluminum or silicon to still lower levels, because the free energies for SiO2 and Al2O3 are lower than those for CO or CO2 at the reaction temperature. Once the oxide forms, the inclusion removal rate depends on the nucleation of particles, their growth, ultimate agglomeration, and rise to the surface. Although there is a significant
difference in density between the solid particles and the liquid metal, a simple application of Stoke’s law is difficult. The problem is dynamic and is influenced by the presence of other indigenous or exogenous inclusion species. Furthermore, gravitational effects are different while holding the treated metal in the ladle, pouring through air with an opportunity for the metal to reoxidize, and during solidification with the associated thermal convective currents. Therefore, it should not be surprising that oxide inclusions can become trapped in the final casting. Sulfides. Sulfur is also often present in steels and is controlled by the nature of the
Inclusion-Forming Reactions / 79 50 Reduction in area 40
54.2 Izod
40.7
30 Elongation 20
27.1 0.06% S
10 60 Reduction in area
50
67.8
Izod
40
54.2
Elongation 40.7
30 0.04% S
20 60 Reduction in area 50
67.8 Izod 54.2
40 Elongation
40.7
30
(b)
Izod impact strength, J
Izod impact strength, ft · lb Reduction of area and elongation, %
(a)
0.03% S
20 70 Izod
81.3
60 Reduction in area
67.8
50 40
54.2 Elongation
30 20
40.7 0.02% S 0
0.04
0.08
0.12
0.16
0.20
0.24
0.28
Aluminum added to heat, %
(c)
Fig. 10
Illustrations of common inclusion shapes and distributions found in steel. (a) Type I: globular silicates and oxides (no aluminum deoxidizer). (b) Type II: continuous eutectic sulfides and Al2O3 at grain boundaries (low aluminum addition). (c) Type III: duplex irregularly shaped sulfides and Al2O3 (large aluminum additions). Source: Ref 12, 13
slag-refining processes in the furnace. In general, as the process becomes more reducing (that is, the oxygen content of the metal is lower) and the slag more basic, lower sulfur will be present. Because the sulfur content is not reduced to zero, manganese is always present in steels to produce the more desirable manganese sulfide (MnS) inclusions rather than FeS inclusions with lower melting point. The resulting inclusions, although usually identified as manganese sulfides, may have several distributions and shapes, depending on the presence of other elements and the deoxidation procedure. Complex Inclusions. The various elements compete to form compounds with oxygen and sulfur. Furthermore, alloying elements in excess of those necessary to form compounds will modify the surface energy of the resultant inclusions.
Fig. 11
Combined effect of aluminum and sulfur on the ductility and impact strength of cast and normalized medium-carbon steel. Source: Ref 11–13
The shape and distribution of the inclusions will then also vary, depending on the variables discussed at the beginning of this article. Historically, inclusions have been classified into three general types, as shown in Fig. 10. Of these, type II has been demonstrated as particularly damaging to mechanical properties. The argument has been that sulfur, oxygen, and deoxidizer play a combined role in the inclusion type. Figure 11 shows this effect and the resulting decrease in structure-sensitive properties such as ductility and impact strength. The minimum in the curves is attributed to the formation of type II inclusions. The role of aluminum in the modification of inclusion shape and distribution also depends on the oxygen-sulfur equilibrium shift when the aluminum is added. Therefore, complex inclusions such as oxy-sulfides may be present in a cast steel but can be controlled by the quantity of oxide-sulfide-forming elements. A more complete discussion of oxide inclusion formation mechanisms and the relationship to mechanical properties is given in Ref 14.
Inclusion shape control becomes a consideration especially if fracture toughness, impact strength, or fatigue resistance must be optimized. Magnesium and calcium additions can be successfully used for shape control, that is, to produce more spherical shapes (Ref 15). It is noteworthy that both magnesium and calcium form stable sulfides and oxides as indexed by the free energy of formation of the compounds. Several production methods have been employed for calcium or magnesium additions, either as the metallic elements or as compounds. Because of the high vapor pressure of the pure metals and, to a lesser extent, the compounds at the addition temperatures, special procedures have evolved, including injection at the bottom of a filled ladle. In this case, the ferrostatic head provides a pressure that suppresses the boiling action and gives a longer reaction dwell time. The most common practice is to deoxidize the steel with aluminum followed by calcium for inclusion shape control. Solid Al2O3 particles are formed upon deoxidation, occasionally
80 / Principles of Liquid Metal Processing with complex aluminates such as spinels (Al2MgO4). These solids prefer their own kind and can therefore agglomerate into larger, more harmful inclusions. The addition of calcium lowers the melting range to form liquid calcium aluminates (40 to 60% CaO) that do not agglomerate but appear as small and more desirable spherical oxide inclusions. Rare earth elements added in conjunction with calcium, magnesium, and aluminum have also been reported to enhance the spheroidization and fineness of the inclusions by breaking up the Al2O3 galaxies (Ref 16, 17). However, difficulty in removal of the rare earth oxide inclusions, especially in scrap remelt, has limited the practice in present-day steel casting processes. Nitrides. Just as there are elements that preferentially form oxides or sulfides, there are differences in the free energy of formation of the nitrides. These elements control the amount of nitrogen that can be dissolved in both the solid 0.020 1260 ⬚C
Nitrogen, %
0.015 980 ⬚C
1150 ⬚C
0.010 1065 ⬚C 900 ⬚C
0.005
0 0
0.02
0.04
815 ⬚C
0.08 0.06 Aluminum, %
0.10
Fig. 12
0.12
Isothermal precipitation of aluminum nitride. Compositions to the right or above each isotherm cause precipitation. Source: Ref 12
and liquid states. For example, vanadium and manganese increase nitrogen solubility, while carbon and phosphorus decrease nitrogen solubility (Ref 12). The effect of aluminum on nitrogen solubility has been extensively studied, and criteria that result in the precipitation of aluminum nitrides have been developed (Fig. 12). Other nitride-forming elements, such as titanium, zirconium, and vanadium, would show similar but numerically different results. It is not uncommon to consider these nitride inclusions to be beneficial; they peg austenite grain boundaries, and this results in grain refinement. On the other hand, brittleness results when the nitrides become more or less continuous. Exogenous Inclusions. Numerous examples of inclusions occur because of the reaction between the molten steel and the adjacent environment. A few of these are considered as follows. When a steel is killed with aluminum, residual aluminum remains in solution. If the metal contacts a refractory that contains free SiO2, the aluminum can reduce the SiO2 because of the lower free energy of formation of the Al2O3. The eroded or partially dissolved refractory becomes a source of inclusions, as does the Al2O3 that is produced. Careful control of refractory chemistry and good pouring practice minimize this problem. Similar reactions can occur at the mold/metal interface, where the first equilibrium of importance is iron-oxygen-silicon. When conditions are oxidizing, FeO is present, and the corresponding form for silicon is SiO2. The FeOSiO2 phase diagram is shown in Fig. 13. The most significant observation here is the formation of
1800
3270 2 liquids ~1700⬚
1600
2910 SiO2 (Cristobalite) + L
1470⬚ 1400
2550
1371⬚ SiO2 (Tridymite) + L
2FeO·SiO2 + L FeO + L
1208⬚
1200
2190 1180⬚
1175⬚
FeO + 2FeO•SiO2 1000 FeO
20
2FeO·SiO2 (Fayalite) + Tridymite 40
60 SiO2, wt%
Fig. 13
FeO-SiO2 phase diagram. Source: Ref 10
80
1830 SiO2
Temperature, ⬚F
Temperature, ⬚C
Liquid
lower-melting liquid phases as the amount of FeO is increased. With minimal oxidation, burned-on sand and metal penetration will occur. However, when the oxidation level increases, local macroinclusions may form, especially if poor pouring practice, segregated bentonite, nonuniform mold ramming, and excessive mold moisture exist. A number of the problems and solutions associated with exogenous inclusions in steel have been summarized in Ref 18.
Cast Irons Gray Cast Iron. The earlier discussion of the carbon-oxygen equilibrium (Fig. 8) suggests that dissolved oxygen should not be a problem in cast irons because of their high carbon contents. Therefore, oxide inclusions should not occur in gray cast iron. However, because the silicon content of this material is also significant, the carbon-silicon-oxygen equilibrium is of interest. The temperature dependencies of the free energies of formation for SiO2 and CO are different for the two species. At temperatures above approximately 1535 C (2795 F), CO is the more stable compound, while below this temperature, SiO2 is more stable. The exact temperature depends on the carbon and silicon content of the metal. The phenomenon can be observed on the surface of a cast iron melt that appears clean at high temperatures and yet develops a slag covering at lower temperatures. The significance of the equilibrium shift is that skimming the slag from a melt surface may be unnecessary at high temperatures, because little may exist. Inclusions are then found in the casting because of pouring at lower temperatures and exposure of the metal to reoxidation during pouring. These inclusions are often silicates derived from the normal equilibrium change with temperature. Once the molten iron is in the mold cavity, there is less opportunity to form the silicates because conditions are usually less oxidizing. Nitride and sulfide inclusions also exist in gray cast iron but are not as severe a problem as in steel because of the inherently lower dissolved oxygen content of the cast iron. Special gray cast iron chemistries containing alloying elements with an affinity for nitrogen, oxygen, or sulfur may result in inclusions by further reaction with the refractories, mold, or atmosphere. The remedies are usually similar to those given for steel. Ductile Cast Iron. A number of foundries erroneously assumed that ductile (nodular) cast iron would be as easy to process as gray cast iron because the major difference was merely a lower sulfur plus a modest magnesium addition. Although many characteristics were found to be different, one of the more perplexing was the formation of dross usually not found in gray cast iron. The result was the so-called cope-side defect that was very prevalent in early ductile iron castings.
Inclusion-Forming Reactions / 81 Magnesium forms a very stable oxide and can therefore reduce oxides of lower stability, as determined by their free energy of formation. Furthermore, the alloys used to produce the nodular graphite found in ductile iron contain considerable amounts of silicon and are low in density, which allows them to float on ladle surfaces rather than reacting completely. These observations help to explain how dross particles form and can become entrapped in castings. In general, there are three types of dross inclusions in ductile cast iron (Ref 19). Type I consists of large dross particles generally derived from oxidation of the silicon and magnesium used in the nodularizing alloys. The dross particles are present on the treated liquid-metal surface and can be removed by skimming. Silica does not exist in the pure form; rather, magnesia (MgO) and forsterite (Mg2SiO4) are the predominant species. This suggests a driving force from the high-MgO end of the phase diagram shown in Fig. 14. Under some circumstances, especially after a late silicon inoculation, the dross also contains fayalite (2FeOSiO2 or Fe2SiO4) (Fig. 13). Fayalite and forsterite form a complete series of solid solutions known as the olivines and are used as molding materials. Type II inclusions are in a more or less stringlike configuration in conjunction with decayed graphite shapes. Chemically, they are predominantly MgO with some iron and silicon atom replacement for the magnesium in the MgO lattice. Reaction with the oxidizing components in the mold material is one source for this type of inclusion, as is an excess of nodularizing agent. Their position is usually just beneath the casting surface. Control of the mold
atmosphere and the amount of nodularizer usually minimizes the defect. Type III inclusions are fine dispersions of MgO and MgS, occasionally mixed with Mg2Si. These components would normally be expected only if the oxygen content in the immediate vicinity is very low. Their normal uniform distribution does not usually detract from the overall performance of the casting. However, undissolved nodularizer may be a source of type III inclusions, because Mg2Si is the active magnesium source in the nodularizing agent. Finally, the examples provided in this section are not intended to represent all of the inclusions to be anticipated in every ferrous alloy. On the other hand, they are sufficiently representative so that analogies for occurrence and remedy can be derived for the alloys not discussed.
Inclusions in Nonferrous Alloys Several important cases are given subsequently for aluminum, copper, and magnesium alloys. The principles discussed in this section can be extended to include other alloy systems. In the alloys discussed, both nonmetallic and intermetallic compounds are considered deleterious.
Aluminum Alloys Oxides. It is nearly impossible to prevent dross formation in the melting of aluminum alloys. The most common form of dross contains large quantities of Al2O3 and, when agglomerated, can be removed from the melt surface by skimming. Pouring operations should employ 3450
1900 1890⬚
MgO + L 1850⬚
Liquid 3270
1800
skimming ladles to prevent larger dross particles from entering the mold. Furthermore, a gating practice must be used that minimizes Al2O3 formation during mold filling. The fact that Al2O3 forms is not surprising, because the free energy of formation of Al2O3 is very low. Furthermore, the specific gravity of Al2O3 is not greatly different from that of the molten aluminum alloys. Although more dense, Al2O3 forms at an outside surface first if the melt is exposed to the atmosphere. The tenacious skin will float unless it is broken up into small particles, and it can become quite thick when melts are held for extended periods in the furnace. When small oxide particles are formed, they may not sink but may become suspended in the melt. In high-quality castings, these particles must be removed; ceramic filters have been used to remove the particles, as mentioned earlier in this article (Ref 3). The degradation of mechanical properties also depends on the size and distribution of oxide inclusions. Therefore, solidification rate is very important for minimizing property loss. When the casting solidifies rapidly, as in thin sections or in permanent molds, the oxide inclusions are smaller and more uniformly dispersed. This observation is consistent with the general constraints of nucleation and growth, where less time during solidification not only creates smaller particles but also prevents their agglomeration and growth. Complex Oxides. Other elements in the aluminum alloys may react with oxygen sources and produce complex oxides, with Al2O3 as one of the components. This is the case in magnesium-containing alloys, in which the spinel MgAl2O4 may form (Fig. 15). Here, MgO has a lower free energy of formation than Al2O3; the spinel results when they form together. Again, precautions in pouring, gating, and minimizing the reoxidation of the metal are important in inclusion control. The use of filtering systems is also advantageous, especially in high-strength castings in which fracture toughness characteristics are important.
1695⬚
3090
SiO2 + L
1600
2910 1557⬚
20
40
Fig. 14
MgO-SiO2 phase diagram. Source: Ref 10
Liquid + MgAl2O4
2600
2200
MgO + liquid
MgAl2O4 1800
1400
MgO·SiO2 + SiO2
4710
Liquid + spinel
Liquid + Al2O3
Spinel
MgO + MgAl2O4
3990
3270
Al2O3 + spinel
2550
(Enstatite)
60 SiO2, wt%
Liquid
1543⬚
2MgO·SiO2 + MgO·SiO2 1500 MgO
5430
3000
80
2730 SiO2
1000 0 MgO
Fig. 15
20
60 40 Al2O3, mol%
80
1830 100 Al2O3
MgO-Al2O3 phase diagram. Source: Ref 2
Temperature, ⬚F
1700
Temperature, ⬚C
2 liquids
Temperature, ⬚F
(Forsterite)
2MgO·SiO2 + L
Temperature, ⬚C
MgO + 2MgO·SiO2
2030
1100
2012
Temperature, ⬚C
1110
1090
L
1976
1080 1070 1060
1050 0.001
Fig. 16
1994
1958 α
α + Cu2O 0.01 0.1 Oxygen, wt%
Temperature, ⬚F
82 / Principles of Liquid Metal Processing
1940 1922 (0.7)
A portion of the copper-oxygen equilibrium diagram. Source: Ref 20
Other Inclusion Types. Inclusions in aluminum alloys can occur from grain-refining operations in which the grain-refining additions may not be adequately controlled. Titanium, vanadium, and boron are used in various combinations to produce grain refinement, and the diborides can become undesirable inclusions. Again, the fact that a small addition is beneficial does not mean that larger additions will provide greater rewards. Many of the commercial aluminum alloy compositions have been developed to exploit their age-hardening characteristics in conjunction with an alloyed matrix. They then become susceptible to the presence of seemingly minor elements such that undesirable intermetallic particles can be formed during cooling. In those cases of precipitation during solidification, the particles are placed in the category of inclusions. The intermetallic inclusions are particularly troublesome in aluminum alloys containing large quantities of zinc or magnesium. Complexes can be formed with the transition elements. These observations suggest that the remedy is close control of the metal chemistry. There are many examples in which control is inadequate and the problem is not identified until mechanical properties are marginally out of specification. Even aluminum carbide and iron-aluminum compounds have degraded properties by the formation of inclusions. Although the solubilities are low, the problem is additive in that scrap/revert procedures can cause a gradual increase in the impurity levels, and inclusion problems become more severe. The hazards associated with recycling then become apparent, especially when the final product is sensitive to the elements present and when strict mechanical property specifications are imposed. The development of new alloys, such as the aluminum-lithium alloys, brings its own set of metal-processing difficulties that begin with making ingots for forming. The question of casting more complex shapes then becomes a natural evolution. Here, the reaction of lithium with air requires melting under liquid salts that themselves are aggressive with the surroundings, such as refractories or mold materials. The opportunity for trapped inclusions is apparent.
Copper Alloys
Magnesium Alloys
A portion of the copper-oxygen system is shown in Fig. 16. The eutectic shows that the last liquid to solidify may contain 0.39% O, even though the initial melt may contain 0.01% O. Therefore, Cu2O inclusions may be more predominant internal to the casting (liquid enrichment in oxygen due to slow diffusion) and may also be present at grain boundaries. The key to control is to add elements that combine with dissolved oxygen (that is, to form oxides with free energies lower than that of Cu2O). Examples are lithium, boron, magnesium, or phosphorus, where control of the excess addition becomes important to modification of the electrical conductivity in the high-copper alloys. The excess deoxidizer goes into solid solution. A more thorough discussion of gas solubility is available in Ref 20 (see also the article “Gases in Metals” in this Volume). A number of the copper alloys are susceptible to dross formation, and the normal precautions in pouring and gating are absolutely necessary to minimize exogenous inclusions. The drosses are usually complex oxides of copper, zinc, tin, lead, or aluminum in the aluminum bronze family. Iron is added as both a dispersion strengthener and as a grain refiner in aluminum and manganese bronzes. However, iron can become an undesirable inclusion in the other copper alloys. When a foundry pours several different alloy chemistries, scrap segregation becomes important to inclusion control. Consideration can also be given to alternative grain-refining agents such as zirconium. In some cases, the intermetallic inclusions are rich only in iron, and complexes are formed. This is especially true when alloys are high in zinc, such as the high-strength yellow brasses. The problem becomes more severe when a melt is held close to its liquidus for extended periods. In this case, the solubility of the complex intermetallics is lower, and their formation is enhanced. High-zinc alloys offer fewer problems with dissolved oxygen because of the vaporization of zinc and the purging action of the melt. Oxides as inclusions are then a lesser problem. However, zinc oxide (ZnO) can react with refractories and mold materials to produce inclusions during pouring. In addition to careful pouring and gating practices, exogenous inclusions can be removed by filtering processes. However, the practice has not been as well received for copper alloys as for aluminum alloys. Although not strictly an inclusion from the previous definition, lead in many of the copper alloys can form oxides that will react with mold bonding agents such as bentonite. These lead complexes can become leachable in standard tests, and waste sand streams that contain lead compounds may be classified as toxic, with the attendant difficulty in disposal (Ref 21).
Of the nonferrous alloys considered in this section, magnesium has the lowest free energy of formation for its oxide. Therefore, the complete prevention of oxide inclusions is impossible, and a number of methods have been developed to minimize oxide inclusion formation. Agents that evolve gas to partially insulate the melt from the atmosphere and fluxes that agglomerate the oxide (plus other impurities that are to be removed) are two examples of control of the melting practice. Again, great care must be exercised in gating to minimize reoxidation. The magnesium alloys are light in weight, and many of the potential exogenous inclusions are collected as sludge in the bottom of the furnace and ladle. Desludging operations are then necessary. The use of filtering systems, whether metallic or ceramic, is common to all of the higher-quality magnesium castings. As with aluminum alloys, excess grain-refining agents can result in undesirable inclusions that can have complex chemistries because of secondary reactions. The fluxes themselves must be reactive in order to perform their intended function. Excess flux can then react with the atmosphere, refractories, or the mold materials to produce inclusions. On the other hand, insufficient flux causes metal reaction with the atmosphere to produce oxides and nitrides that may be difficult to remove if not in an agglomerated form. Temperature, time at temperature, alloy composition, pouring practice, and gating procedures dictate the optimal fluxing process. So-called fluxless melting has been suggested for magnesium alloys in order to overcome the environmental and metallurgical problems associated with fluxes (Ref 22). Fluxless melting uses nonreactive gaseous atmospheres above the melt. The most effective protective atmosphere is CO2 plus the addition of trace inhibitors such as SO2 and SF6. On the other hand, the potential addition of fluorine-containing gases to a gas waste stream must also be of concern. Finally, the general observations and processing techniques in this article have assumed the presence of air to generate the bulk of the inclusions. Alternative melt processing can exclude air for the primary purpose of inclusion prevention. Both vacuum and inert gases are used to process alloys for more demanding applications such as metal bearings, superalloys, and metal alloy prosthetics. Further discussion is included under specific alloy systems and melt methods elsewhere in this Volume.
ACKNOWLEDGMENT From: P.K. Trojan, Inclusion-Forming Reactions, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 88–97
Inclusion-Forming Reactions / 83 REFERENCES 1. L. Ackermann, J. Charbonnier, G. Desplanches, and H. Kaslowski, Properties of Reinforced Foundry Alloys, Trans. AFS, Vol 94, 1986, p 285–290 2. R.A. Flinn and P.K. Trojan, Engineering Materials and Their Applications, 3rd ed., Houghton-Mifflin, 1986 3. F.R. Mollard and N. Davidson, Experience with Ceramic Foam Filtration of Aluminum Castings, Trans. AFS, Vol 88, 1980, p 595–600 4. L.A. Aubrey, J.W. Brockmeyer, and P.F. Wieser, Dross Removal from Ductile Iron with Ceramic Foam Filters, Trans. AFS, Vol 93, 1985, p 171–176 5. L.A. Aubrey, J.W. Brockmeyer, P.F. Wieser, I. Dutta, and A. Ilhan, Cast Steel Improvement by Filtration with Ceramic Foam Filters, Trans. AFS, Vol 93, 1985, p 177–182 6. P.K. Trojan, P.J. Guichelaar, and R.A. Flinn, An Investigation of Entrapment of Dross and Inclusions Using Transparent Whirl Gate Models, Trans. AFS, Vol 74, 1966, p 462–469
7. D.G. Schmidt, Gating of Copper Base Alloys, Trans. AFS, Vol 88, 1980, p 805–816 8. A.D. Wilson, Calcium Treatment of Plate Steels and Its Effect on Fatigue and Toughness Properties, Proc. 11th Annual Offshore Technology Conference, OTC 3465, May 1979, p 939–948 9. A.D. Wilson, Effect of Calcium Treatment on Inclusions in Constructional Steels, Met. Prog., Vol 121 (No. 5), April 1982, p 41–46 10. The Making, Shaping and Treating of Steel, 9th ed., United States Steel Corporation, 1971 11. C.W. Briggs, The Metallurgy of Steel Castings, McGraw-Hill, 1946 12. R.A. Flinn, Fundamentals of Metal Casting, Addison-Wesley, 1963 13. C.E. Sims and F.B. Dahle, The Effect of Aluminum on the Properties of Medium Carbon Cast Steel, Trans. AFS, Vol 46, 1938, p 65–132 14. W.O. Philbrook, M.C. Flemings, L.H. Van Vlack, A. Gittins, and J. Lankford, Oxide Inclusions in Steel, A series of five papers, Int. Met. Rev., Sept 1977, p 187–228 15. Low Sulfur Steel, Symposium proceedings, AMAX Inc., 1986
16. S.K. Paul, A.K. Chakrabarty, and S. Basu, Effect of Rare Earth Additions on the Inclusions and Properties of Ca-Al Deoxidized Steel, Metall. Trans. B, Vol 13 (No. 2), June 1982, p 185–192 17. M. Olette and G. Gatellier, Effect of Calcium, Magnesium or Rare Earth Additions on Cleanness of Steel, Rev. Metall. Cah. Inf. Tech., Vol 78 (No. 12), Dec 1981, p 961–973 18. R.A. Flinn, L.H. Van Vlack, and G.A. Colligan, Mold-Metal Reactions in Ferrous and Nonferrous Alloys, Trans. AFS, Vol 94, 1986, p 29–46 19. D.R. Askeland and P.K. Trojan, The Approach to Equilibrium and Dross Formation in Nodular Cast Iron, Trans. AFS, Vol 77, 1969, p 344–352 20. Casting Copper Base Alloys, American Foundrymen’s Society, 1984 21. T.R. Ostrom, B.P. Winter, and P.K. Trojan, Lead Transfer from Copper Base Alloys into Molding Sand, Trans. AFS, Vol 93, 1985, p 757–762 22. J.W. Fruehling and J.D. Hanawalt, Protective Atmospheres for Melting Magnesium Alloys, Trans. AFS, Vol 77, 1969, p 159–164
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 87-98 DOI: 10.1361/asmhba0005195
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Electric Arc Furnace Melting*
THE MELTING OF STEEL is performed in both arc furnaces and induction furnaces. The arc furnace is the primary melting method and is used by over 80% of steel foundries (Ref 1). However, arc furnaces are not good holders or efficient at superheating metal, so it is not uncommon to use induction furnaces to superheat steel after the arc melters. Induction furnaces used to melt steel require cleaner charge materials to meet metallurgical requirements and therefore require low-carbon charge to produce lower slag levels. The basic concept of an electric arc melting furnace is a refractory-lined steel vessel having a bowl-shaped hearth and a domed-shaped refractory roof. The hearth and sidewalls are typically lined with bricks, monolithic materials, or a combination of both. Acid refractories (silica-based), basic refractories (magnesitebased), and, in some cases, neutral refractories (alumina-based) are used to line the furnace sidewalls, hearth, and roof. All major additions are made to the furnace through the roof by swinging the roof out of the way and charging scrap, fluxes, and carbon from loaded buckets. The furnace has two or three electrodes mounted vertically through the roof of the furnace that float just above the surface of the cold ferrous scrap or molten bath. Over 80% of arc furnaces used for steel castings are alternating current furnaces with three electrodes. These electrodes strike an arc on the cold metal scrap or molten metal to impart energy to the bath and melt or superheat the furnace. Control of the electrode level above the bath is critical to the melting process efficiency. Arc melting furnaces are a widely used method for melting steel scrap. They can melt a wide variety of ferrous scrap materials and can accept certain levels of dirty scrap. They also use chemical energy to assist in the melting process and to lower the carbon levels of the molten metal. Since steel castings require low-carbon-level metallurgy, carbon from scrap metal feedstock and carbon electrodes must be removed from the molten metal. Injecting oxygen into the molten bath oxidizes carbon and uses the exothermic reaction to assist in melting and superheating the metal. This is referred to
as supplying chemical energy to the furnace. In some instances, different forms of oxyfuels are used to provide additional energy to speed up the melting process. Although refractories line the furnace shell and roof, significant energy losses occur through the surfaces of the furnace and from exposed surfaces during charging and slagging operations. Energy losses from the surface of the bath are due to the large surface-area-to-volume ratio of the furnace, resulting in significant energy losses during delays in the operation. In addition, the large surface area of the furnace results in large losses through the refractory walls, hearth, and roof. Nearly all steel foundries blow oxygen into the bath after most of the scrap is melted, which generates chemical energy from exothermic decarburization, assisting in the melting process. Some operations use additional sources of chemical energy to decrease tap-to-tap time. Electric arc furnaces are used as melters and holders in duplex operations and as melting and refining units. Since its first appearance as a production tool at the beginning of the 20th century, the electric arc furnace(EAF) has become one of the primary melting tools used by foundries and steel mills. The early furnaces had capacities of 910 to 14,000 kg (1 to 15 tons), but new technology has vastly increased EAF productivity. There are currently approximately 100 minimills in the United States capable of producing 50 million or more tons per year, with numerous furnaces capable of producing in excess of 100 tons per hour (Ref 2). This article focuses on the construction and operation of these furnaces and their auxiliary equipment in the steel foundry (shape casting) industry. Primary steelmaking practices differ from that of smaller foundry furnaces, and most primary steelmaking EAF operations have adopted separate practices since the 1980s. For example, methods of improving the efficiency of arc melting furnaces in the primary steel industry are not directly transferable to the smaller furnaces used by steel shape casting producers. Oxygenenriched burners and foaming slag methods are not considered effective on the smaller furnaces typical of steel foundry operations. For primary
melting, the reader is referred to The Making, Shaping, and Treating of Steel, 11th ed., published by The AIST Foundation. More information on ladle metallurgy is also contained in the article “Steel Melt Processing” in this Volume.
Power Supply The current applied to the EAF is supplied by a local electric utility company and passes through an electrical substation that is designed specifically for the furnace(s) to be supplied. A simplified diagram is shown in Fig. 1. The power supplied to the furnace during melting is provided by an electrical arc
Fig. 1
Schematic of electrical network for the electric arc furnace
*Revised from Melting Furnaces: Electric Arc Furnaces, Casting, by Nick Wukovich, in Vol 15, ASM Handbook, ASM International, 1988, p 356–368 from Reviews by Peter Glaws and Craig Darragh, The Timken Co.
88 / Melting and Remelting established from three carbon or graphite electrodes. During the meltdown portion of the heat, the three arcs or phases act as a single phase. Two arcs can strike the furnace charge and draw current without a current going to the third arc or electrode. Because of the type of scrap used and the arc lengthening and shortening that take place, there is a great fluctuation in the current during meltdown, which causes a significant variation in the electrical supply system. If the same power source is the supply for other plant power, a flickering of lights and voltage fluctuation will be noticed on machinery and electrical equipment. During the refining cycle or after meltdown, the arcs tend to stabilize, partially because of a slag cover on the liquid metal and the flatness of the bath. In addition, the arcs have been shortened to direct their energy in a smaller area. A great amount of energy is produced during this meltdown and refining of ferrous metals. Controls for harnessing and directing the energy of the arcs are required in order to produce molten iron and steel in the electric furnace without destroying the furnace refractories. The energy requirements for melting various carbon levels in iron or steel are shown in Fig. 2. Power Factor. The efficiency with which power is transferred to the melt is called the power factor. The power factor, PF, is the ratio of watts, W, divided by volt-amperes, VA: PF ¼
W VA
have been chosen. The electrical system for the furnace includes the transformer and the control panel. Very small furnaces (<910 kg, or 1 ton) are not used in current production schedules; heats of these sizes are normally made in induction furnaces. Furnaces are described as supplying so many tons or pounds per heat under normal conditions, but larger heats can be made by modifying
the furnace. This is done by stopping off the tap hole and raising the breast on the working and charge door (if the furnace has one). The furnace is lined with refractories that determine the type of melting practice (acid or basic) that will be used (Fig. 3). Table 1 lists various refractory furnace materials used in acid and basic melting practices. Depending
100
One method of measuring the average power factor over a period of time requires the use of two separate meters: a watt-hour meter to accumulate the useful power and a var-hour meter to accumulate the reactive power. The var-hours (in a given time) divided by the watt-hours (in the same time) will give the tangent of the power factor angle. The cosine of this angle is the power factor. The local utility company can be contacted for this information and for recommendations on the effect on energy usage.
Fig. 2
Power consumed in melting iron and steel in the electric arc furnace. Values will vary depending on scrap, transformer, lining, and so on. The melting point of pure iron (0.0% C) is 1535 C (2795 F); of iron containing 4.3% C, 1130 C (2066 F).
Arc Furnace Components The list of equipment associated with an EAF can be extensive. This list is assembled as if a proposal for the furnace has been made. The locations for the furnace foundation and pits or elevated platforms are selected. Factors that also must be considered include the location of the furnaces in the plant, flow and access to raw materials, storage of melt materials, ladle construction, ladle preheating, crane runways, water, air, electrical lines, transformers, laboratory, and temperature equipment. This list is not all-inclusive. Fume and dust collection equipment can be added, as well as used-refractory and slag removal equipment. For the melting portion, a supply of oxygen, hoses, pipe for oxygen blowing, gages, valves, thermocouples, and so on will be needed. The furnace itself is the primary concern after the location, cement work, and electrical supply
Fig. 3
Cutaway view of electric arc furnace showing the typical refractories used
Electric Arc Furnace Melting / 89 on whether the process is basic or acid, slag color may be sufficient to determine the iron oxide content of the slag, as shown: Color
FeO, %
CaO/SiO2
Acid slag Gray-black Dark green Streaked dark green Blue-green (jade green)
30–40 20–30 15–20 <15
... ... ... ...
Basic slag Black Dark brown-gray Light brown Gray White White
14.85 11.50 1.05 1.05 0.56 0.51
3.33 2.19 1.91 1.88 2.01 2.53
The silica content of basic slag can profoundly affect the melt. If present in sufficient (quantities, silica can acid constituent) can significantly reduce the slag’s ability to effectively desulfurize the steel. The size of the furnace can vary from 450 to 91,000 kg (0.5 to 100 tons) or larger. Typical furnace data are given in Table 2.
Table 1
delta section of the roof and can cause excessive refractory wear in this area. The use of water-cooling rings around the electrode ports can help to reduce this wear. The electrodes should be vertical from all sightings around the furnace. When the electrodes are not vertical, the distance and arc flare on the sidewalls are not uniform, and a hot spot may cause the refractory wear to be greater in the areas closer to the electrodes. In addition, because the arc works between the electrodes and the bath, the current per phase can fluctuate (Fig. 5). The roof structure is arched to be selfsupporting and to keep the greatest distance between the arc action in the delta section. Refractory roofs start with a roof ring that is designed as a channel or an angle beam that is rolled or cast into a ring. The ring generally has a water-cooling chamber on all or a portion of its height (Fig. 6). Newer designs of watercooled roofs are cast in gray iron except for the delta section and the exhaust and oxygen lance holes on large roofs, and they are cooled by a network of steel pipes running throughout the structure. The delta section and electrode ports are
The electrode can be graphite or carbon. The graphite electrode has greater current capacity than the carbon unit for the same size electrode. Table 3 lists some data on electrode sizes, weights, and current capacities. The electrodes are supported on three separate arms and held to the electrode arm by a spring-clamped, pneumatically released electrode holder. On older furnaces, the electrode holder had a mechanical wedge-type clamp; on newer, larger furnaces, power-operated clamping devices are used. The electrode arms support the electrodes, holders, bus tubes, and transformer power cables. Smaller bus tubes and power cables can be used if they are air or water cooled. The electrodes are raised and lowered by manual or automatic controls. These controls can be operated manually when the bottoms of the electrodes are raised to the roof or when electrodes are shifted. The electrodes are raised and lowered automatically during the melting and refining process (Fig. 4). The furnace designer will make the electrode circle as small as possible to increase the speed of melting. This will intensify the heat in the
Refractories as a source of slag based on typical composition Composition, %
Type
Fireclay Super duty Medium duty Semi silica Alumina type (high) 50% 60% 70% 80% 90% Mullite Corundum Silica type Silica super duty Conventional Basic type Chrome Magnesite Magnesite high periclase Chrome-magnesite
Table 2
Softening point
SiO2
Al2O3
Cr2O3
TiO2
CaO
MgO
Fe2O3
Other
C
F
49–56 57–70 72–80
40–44 25–38 18–26
... ... ...
1.5–2.5 1.3–2.1 1.0–1.5
... ... ...
... ... ...
... ... ...
2.5–4.0 4.0–7.0 1.0–3.0
1745–1765 1660–1685 1640–1685
3175–3210 3020–3065 2985–3065
41–47 31–37 20–26 11–15 7½–9 18–34 0.2–1.0
47½–52½ 57½–62½ 67½–72½ 77½–82½ 89–91 60–78 98–99.5
... ... ... ... ... ... ...
2.0–2.8 2.0–3.3 3.0–4.0 3.0–4.0 0.4–0.8 0.5–3.1 Trace
... ... ... ... ... ... ...
... ... ... ... ... ... ...
... ... ... ... ... ... ...
3.0–4.0 3.0–4.0 3.0–4.0 3.0–4.0 1.0–2.0 1.0–3.0 0.3–1.0
1785 1805–1820 1820–1850 1865 1930 1850 2000
3245 3280–3310 3310–3360 3390 3505 3360 3630
95–97 94–97
0.15–0.35 0.45–1.20
... ...
... ...
2.5–3.5 1.8–3.5
... ...
0.3–2.2 0.3–0.9
0.02–0.10 0.10–0.30
1680–1700 1635–1665
3060–3090 2975–3025
3.0–6.0 0.7–1.0 0.5–5.0 4–8
15–34 0.3–1.5 0.2–1.0 16–27
28–33 ... ... 18–28
... ... ... ...
... 1.0–3.5 0.5–1.5 0.7–1.5
14–19 85–93 92–98 27–53
11–17 0.3–7.0 0.2–1.0 ...
1.0–2.0 0.5–1.0 0.0–0.6 ...
1290–1425 1480–1675 1595–1705 1595–1675
2350–2600 2700–3050 2900–3100 2900–3050
Typical dimensions and capacities of foundry-sized electric arc furnace components Refractory
Shell dimensions (inside) Diameter m
1.5 1.8 2.20 2.4 2.67 2.74 3.05 3.35 3.81
Dimensions
Depth
Melting rate
ft
m
ft
5.0 6.0 7.25 8.0 8.75 9.0 10.0 11.0 12.5
1.14 1.4 1.7 1.8 2.01 2.06 2.3 2.51 2.84
3.75 4.5 5.5 6.0 6.58 6.75 7.5 8.25 9.33
kg/h
450–640 910–1,270 1,360–1,910 1,810–2,540 2,270–3,180 2,720–3,810 3,630–5,440 5,440–8,160 7,260–10,900
Sidewall
Liquid metal capacity
Bottom
Normal
Maximum
Electrode diameter
lb/h
mm
in.
mm
in.
kg
lb
kg
lb
mm
in.
Transformer capacity, kVA
1,000–1,400 2,000–2,800 3,000–4,200 4,000–5,600 5,000–7,000 6,000–8,400 8,000–12,000 12,000–18,000 16,000–24,000
254 254 254 254 368 368 368 368 368
10 10 10 10 14.5 14.5 14.5 14.5 14.5
305 381 432 457 483 483 508 508 533
12 15 17 18 19 19 20 20 21
1,360 2,810 4,260 6,080 9,070 9,840 14,500 22,700 29,900
3,000 6,200 9,400 13,400 20,000 21,700 32,000 50,000 66,000
1,810 3,750 5,670 8,120 12,100 13,200 19,100 29,900 39,900
4,000 8,260 12,500 17,900 26,700 29,000 42,000 66,000 88,000
75 152 178 203 229 229 254 305 356
3 6 7 8 9 9 10 12 14
1,000–1,500 1,500–2,000 2,000–3,000 3,000–5,000 4,000–6,000 4,000–6,000 5,000–9,000 7,500–12,500 10,000–15,000
90 / Melting and Remelting
Table 3 Electrode sizes, weights, and current capacities Diameter length mm
Approximate weight of electrode and nipple in.
Graphite electrodes 102 1015 4 40 152 1220 6 48 152 1520 6 60 178 1220 7 48 178 1520 7 60 203 1220 8 48 203 1520 8 60 229 1520 9 60 254 1220 10 48 254 1520 10 60 305 1520 12 60 305 1830 12 72 356 1520 14 60 356 1830 14 72 Carbon electrodes 203 1520 254 1520 305 1520 356 1520 356 1830
8 60 10 60 12 60 14 60 14 72
Approximate current capacity, A
kg
lb
1,800–3,300 3,500–5,800 3,500–5,800 4,400–7,500 4,400–7,500 5,500–9,300 5,500–9,300 6,700–11,300 8,000–13,300 8,000–13,300 11,300–17,000 11,300–17,000 18,000–25,000 18,000–25,000
13.6 37.6 47.1 46.3 59.9 62.1 77.1 96.6 97.1 118 171 201 228 273
30 83 104 102 132 137 170 213 214 261 378 444 502 602
2,500–4,500 3,100–6,300 4,500–7,900 5,400–10,000 5,400–10,000
79.8 125 178 240 293
176 275 392 528 647
usually constructed of a high-alumina refractory brick, rammed material, or precast shapes. The roofs of all modern electric furnaces are designed for top charging. The roof and electrodes are raised several inches above the top of the furnace shell and swung to one side. The charge is dropped into the furnace, and the roof is swung back into place. The raising of the roof is done by electrohydraulic systems. Schematics of roof-lifting devices are shown in Fig. 7. Tilt Control. The electric furnace is tapped out by tilting the furnace into a position in which the metal and slag will run out of a hole above the slag line in the furnace wall down into a spout and into a ladle. Tapping of the furnace requires that the roof and electrodes remain attached to the furnace shell; therefore, safety locks or stops are found welded to the shell and roof ring. The tilting mechanism can be mechanical for very small furnaces or can be motor driven (either by a screw pitman or a rack-and-pinion mechanism) or hydraulic, and it has rockers attached to the furnace and rails attached to the foundation. Emergency equipment is usually provided to pour the heat from the furnace in case of a power failure. The furnace shell is typically made of carbon steel sections that are rolled and welded to the diameters listed in Table 2. The diameter and height define the furnace size. The furnace bottom is formed and welded to complete the pot. The shell will have door openings cut into it and a cooling frame welded around the door. The door-lifting supports will also be welded onto the shell. The furnace rockers and other lifting equipment attachments are welded onto the shell. With the rockers and door aligned, the spout opening is cut out of the shell and the spout frame attachment welded onto
the shell. To ensure that the furnace is properly grounded, straps, rods, or mesh are attached to the bottom on the inside and ground on the outside prior to lining of the furnace. The choice of refractories will depend on the melting practice selected (Fig. 3 and Table 1). Furnaces built in the 1980s and beyond have used water-cooled panels in the furnace sidewalls above the slag line (Fig. 8). Spout and Tap Hole. The length and shape of the furnace spout have historically created design issues. The melter and metallurgist want the shortest possible spout, while the design engineer is confronted with a network of electrodes, crane cables, pit dimensions, and ladle and tap hole dimensions needed for positioning the ladle in order to obtain a fast and smooth tap from the furnace, using an acceptable spout. New designs and changing the locations of tap holes may reduce the physical and mechanical problems encountered with the spout. One of these alternatives is the side slide gate tap hole (Fig. 9), and another is eccentric bottom tapping with the use of a slide gate (Fig. 10). These alternative devices are primarily used to eliminate furnace slag carryover. In order to do so, they must use a hot heel practice. Eccentric bottom tapping does a very good job of preventing slag carryover at the beginning of pouring. Both do a good job when an effective hot heel practice is employed. Tap-Out and Back-Slagging Pits. The design of the furnace pit should include sufficient depth for the slag boxes, and along with this depth, provision must be made for preventing ground water from entering the pit. Sump pumps are needed to minimize the presence of water and to ensure an environment that is safe from steam explosions. The length and width of the pit should be sufficient to allow movement
of the ladle; this is necessary if the furnace slag is to be kept out of the ladle so that ladle metallurgy or secondary refining can be done following tap out. Water-Cooling System. Water-cooled sidewall panels have become common practice. Even before this innovation, water cooling was used on: Electrode holders, arms, and clamps Roof rings and electrode rings in the delta
section
Around the door Cooling panels positioned on the outside of
the furnace to cool the furnace wall refractories Preheat and Furnace Scrap Burners. Preheating scrap with natural gas or flue gases helps reduce the electrical cost of melting a charge. On the electric furnace, this preheating may also help stagger furnace operation to keep electrical demand peaks to a minimum. Furnace oxygen/natural gas burners can help melt the charge and reduce electrical charges and electrode usage while shortening the melt time from tap to tap. Improper use of these burners can cause excessive oxidation of the heat and the electrodes and can reduce the life of the refractory lining. Fume and dust collection at the electric furnace has been a required practice in melting during the last 25 to 30 years. The complexity of the fume and dust collection equipment varies from a simple fourth-hole connector to an entire separate room for the melting furnace in which the fumes and dust are collected and where the furnace is tapped into a ladle that is brought to the furnace and then removed through a dust protection door. Figure 11 shows typical dust and fume collectors. The charge bucket is sized to fit within the refractory walls of the furnace and is open at the top, with a lifting bale that folds back or is removable from two trunions. The bottom of the charge bucket can be designed to drop the charge in one of several ways: The clam shell bottom is actuated with a trip
lever.
The orange-peel bottom, which is set in a
stand to fold the leaves back and is tied with a rope, burns off when the charge bucket is placed into a hot furnace. The preheat charge bucket uses a trip lever and, after being charged with scrap, is preheated to some temperature above 260 C (500 F) to up to 815 C (1500 F). Figure 12 shows a typical charge loaded in two different types of charge buckets. Ladles may be part of the equipment assigned to the electric furnace, or they may be assigned to the pouring crew or the mold department. Nevertheless, the heat is not finished until the metal is removed from the furnace.
Electric Arc Furnace Melting / 91 slag will attack the refractory, as in Hadfield manganese steel. In foundry operations, the type of ladle used varies. For holding ladles, bottom-pour, teapot, and lip-pour ladles are used. For pouring ladles, bottom-pour, teapot shanks, lip-pour shanks, and drum ladles are used. The ladle linings have changed from refractory lining to the one-heat type with insulating linings made of insulating board or of one-piece designs on shank ladles to eliminate preheating. The repeated use of the refractory reduces cast steel cleanliness and quality, but the board ladles have overcome this problem (Fig. 13).
Weighing and Chemical Analysis Scales. Charge and alloy scales are required for weighing the charge and alloy additions. The charge scale can be the standard beam scale type or the newer load cell type. These scales should be checked daily or weekly to ensure accurate weighing of the furnace melt. The alloy scale can be a small spring or a large balance beam type capable of measuring in pounds and ounces. These should be periodically checked for accuracy. Quality-Control Testing Equipment. The demands for faster and more accurate chemical analysis methods have spawned advancements in the spectrographic, x-ray defraction, and scanning electron microscope analysis of metals. Carbon, sulfur, alloy, and residual analysis is quick and accurate. Gas analysis, oxygen probes, and nitrogen and hydrogen analysis equipment are within the reach of the average melt shop to assist in quality control. Temperature-control immersion thermocouples that are easy to use and recorders that are accurate to within 3 or 6 C (5 or 10 F) are available.
Scrap and Alloy Storage Purchase scrap, shop returns, and alloys should be stored in a covered area. Wet scrap will produce small errors in weight and increase the gas content of the melt. Purchase scrap should be stored by size and the density of its bulk. Scrap should be stored according to specification or alloy content. Alloys and melt additions, that is, slag formers, fluxes, and deoxidizers, should be organized by usage.
Acid Melting Practice Fig. 4
Schematics of apparatuses for raising and lowering electrodes. (a) Arrangement of electrode and mast. (b) Electrode arm arrangement with power-operated electrode holder in which clamp shoe is held in place by a push rod mechanism. (c) Typical arrangement of electrode arm using power-operated electrode holders; pressure on clamp shoe is exerted by spring-operated pull rod.
In foundries and steel mills, the ladles are refractory lined—usually with an alumina lining. More than 72% alumina is used if a clean
steel low in inclusions due to ladle refractory is desired. In one application, magnesite or olivine is used in situations where the metal and
The melting operation from charge to tap out will vary, depending on the type of charge materials and alloy additions used to produce steel in the acid melting practice. The furnace lining will be made from silica or alumina refractories. These refractories determine the type of slag generated during melting, along with the dirt or silica sand included in the charge. The acid practice is a
92 / Melting and Remelting
Fig. 5
Effect of arcs on furnace refractories due to electrodes not being vertical (a) on section left of center and (b) running too high a tap after meltdown
Lower refractory costs due to the initial cost
and the refractory life
Lower power consumption as compared to a
two-slag basic practice
Greater tonnage per unit of time
Acid Steelmaking Steelmaking practices can be classified as partial-oxidation or complete-oxidation operations. These practices are differentiated by the amount of carbon removal during oxidation compared to the final desired carbon content of the melt. The advantages of acid melting practices using partial oxidation are: Rapid steel production Good refractory life due to lower iron oxide
content in the slag
Lower power input Lower electrode consumption No recarburization required
The disadvantages include: Insufficient removal of gases Occasional poor fluidity due to oxides Occasional poor mechanical properties
The advantages of acid melting practices using complete oxidation are:
Fig. 6
Typical roof ring with cooling face. (a) Designed for use with special-shaped or cut starter brick. (b) Designed for use with skewback starter brick. (c) Designed for use with special-shaped or cut starter brick
fast oxidizing operation that does not remove phosphorus or sulfur from the melt. Therefore, charge materials with low phosphorus and sulfur levels should be selected in order to meet the final chemical analysis of the steels produced. The acid melting practice is selected for the following reasons: Flexibility in producing small quantities of
steel to various chemical compositions
Low gas content Good fluidity Good mechanical properties Elimination of nearly all silicon and most of the manganese to a uniform base metal to which additions can be calculated more closely
The disadvantages are: Slower steel production Shorter refractory life due to higher iron
oxide content in the slag
Higher production cost
Fig. 7
Roof-lifting devices and swing mechanisms used on top-charging electric furnaces. (a) Roof-lifting device and swing arm mounted separately from furnace. (b) Roof-lifting device and swing mechanism attached to furnace. (c) Roof lifting by cylinder and cable/swing system by rotating roof with a kingpin arrangement
In partial oxidation, the charge carbon content is 0.05 to 0.15% above the finishing carbon content. If the carbon is above this level, a short blow with the oxygen lance is required. No slag-forming materials are added to the bath or the charge, and a fast melting practice is used. The rust and scale in the charge, along with the oxidation of the charge during melting, are sufficient to produce a light boil. When melting is complete, the carbon content will be approximately 0.25 to 0.35%, and the silicon content will be 0.15 to 0.20%; the actual value depends on the amount of oxidation from the charge and the oxidation during meltdown. By the time the laboratory report on the melteddown carbon and manganese is prepared, the carbon content will have dropped another two
Electric Arc Furnace Melting / 93 or three points, and by the time the tap temperature is reached, the carbon will have dropped so that any alloy additions will bring the carbon content to the desired finishing range. Slag additions are seldom required. The slag will be 3 to 5% of the charge weight because of silica dirt and refractory erosion. Foundries using partial oxidation have minimal mechanical test properties to meet. They
use a good grade of scrap, and their primary concern is the high production of molten steel from their furnace capacity. This practice is losing favor because of quality standards and product liability laws affecting the finished product. Complete Oxidation. Most steel foundries using an acid melting method employ the complete-oxidation practice. This is because nearly all steel castings must meet minimum
mechanical test requirements, along with magnetic-particle, visual, and radiographic inspection. In this process, rapid melting is necessary to start a boil during meltdown, followed by oxidation with the use of the oxygen lance. Some melters add iron ore or mill scale to the charge to obtain a partial boil during meltdown. This boiling action or oxidation is accomplished with the following reaction: FeO þ C ! Fe þ CO
Fig. 8
Fig. 9
Furnace design incorporating a water-cooled roof and upper sidewalls. The location of the water-cooling passages depends on manufacturer’s suggestions.
Furnace slide valve used to hold back slag
Fig. 10
Slag-forming materials are not added to the charge or the bath. Slag is formed from silica and dirt in the charge and erosion from the furnace refractories. A fluid black slag should be obtained by prudent additions of mill scale to ensure consistent silicon and manganese contents at meltdown. Meltdown Period. The meltdown carbon content should be 0.20 to 0.40% above the aimed carbon level. The meltdown carbon level will drop 20 to 40 points after oxygen lancing to approximately 0.20% C. The silicon and manganese contents at meltdown will be approximately 0.13 to 0.20% and 0.25 to 0.35%, respectively. After oxygen lancing, the silicon content will be 0.02 to 0.10%, and the manganese content will be 0.07 to 0.15%. A vigorous boil (oxygen lancing) will eliminate carbon, most of the hydrogen, and small amounts of nitrogen from the bath; the elimination of oxygen depends on the carbon-iron oxide reaction during this period. The slag should be green when the boil is completed. After the boil, silica sand can be added to increase the slag volume and to reduce its iron oxide content. Lime is sometimes added to the slag to reduce some of the iron oxide in the slag. Most slag formers are added to the bath during the refining period. Lime is sometimes added to thin the slag. To minimize the loss of oxidizable elements—silicon and manganese—during
Eccentric bottom tap hole furnace (a) with slide gate attached to furnace bottom (b)
94 / Melting and Remelting melting, the percentages of these elements in the charge should be low. Overoxidizing of the heat should be avoided, and alloying elements should be added immediately before tap out. The Refining Period. After decarburization, the boil will slow to a gentle roll, and recarburization and slag additions can be made if required. The recarburizers may be low- silicon, low-phosphorus pig iron, sorel pig iron, highcarbon steel, broken electrodes, petroleum coke, and so on, and occasional dipping of the carbon electrodes. The addition of carbon gives a carbon boil: O ¼ COðgÞ
This helps keep the gas content down. Small additions of ferrosilicon and ferromanganese or silicomanganese are made to the bath at the end of this boil to hold the bath at the desired carbon level. This is known as a block. The heat should be tapped as soon as possible after this procedure and after the addition of any required alloying elements. Deoxidation is required for preventing the reaction of dissolved oxygen and carbon in producing gaseous products such as CO or CO2, into which hydrogen can find a nuclei and into which it can accumulate and form pinholes in the cast product. The first of these additions is silicon in the form of ferrosilicon or silicomanganese and then aluminum, a powerful deoxidizer. Special deoxidizers such as titanium and
Fig. 11
zirconium help tie up nitrogen, while complex deoxidizers containing silicon, aluminum, calcium, and other elements help modify inclusions.
Melting Heats and Raw Materials The preceding discussion is a general description of the acid melting process. This section discusses the heat and the materials used. The charge will consist of 30 to 50% shop returns (heads, gates, and scrap castings) and 50 to 70% purchase scrap (plate, clean punchings, railroad wheels, and so on). The weight of each group of materials is recorded on the furnace log. The charge distribution into the furnace will normally be heavy scrap and casting on the bottom, followed by light scrap and purchased plate. Nonoxidizable materials such as nickel and molybdenum can also be added with the scrap charge. Carbon raiser additions to the furnace, if desired, are normally done before the charge is placed into the furnace. The roof is swung into place after an inspection of the furnace top to ensure that the scrap does not contact the roof ring, and the roof is then locked down. Melting should proceed as soon as the charge is in the furnace. The No. 1 tap will normally be used for meltdown. To protect the roof from excessive erosion, the No. 2 tap can be used to start the electrodes into the charge, followed by a change back to the No. 1 tap. This procedure will reduce roof wear in the delta section.
Four methods of dust and fume control in electric furnaces. (a) Prepollution control ventilation for dust and fume removal. (b) Direct furnace dust and fume collection (both front view and top view are shown). (c) Total furnace hood for fume and dust collection. (d) Canopy hood for fume and dust collection
A pool of metal will form on the bottom of the furnace as the electrodes burrow through the charge. Once the electrodes have reached this pool of metal, the remainder of the charge will be melted by radiation from the pool and the arc splash off of the electrodes. Rust and any iron ore added to the charge will help in forming a watery slag with the sand, eroded refractory, and slag remaining from the previous heat to form an oxidizing slag on the pool of metal produced from the melting portion of the heat. When the charge is completely melted, the tap is changed to No. 2 or lower to reduce arc splash on the furnace walls and roof. A laboratory sample is taken to determine the carbon level of the bath in preparation for the blowdown, or oxygen lancing. The oxidation following meltdown should not begin until a temperature of at least 1540 C (2800 F) is reached. When heats containing chromium are to be made, temperatures of 1595 to 1620 C (2900 to 2950 F) should be reached before lancing to minimize the loss of chromium to the slag, from which it cannot be recovered. Depending on the amount of carbon, the bath temperature, and the amount of oxygen present in the bath, a decrease of 0.08 to 0.10% C will occur for each minute of oxygen injection. Generally, a steel pipe is used as the lance (ceramic-coated steel pipes are also available). The pipe must be clean and free of grease and oil. The blow time is recorded on the furnace log. After the oxygen blow, a lab sample is taken, and the heat is prepared for finishing and tap out. When the lab test is returned to the melter, adjustments to the carbon level will be made by recarburizing or by additional lancing to the required carbon level. A final check on carbon is made, and the heat is blocked with ferrosilicon. This silicon addition will bring the bath silicon content up to approximately 0.10% or slightly higher to stop any remaining carbon boil. Final additions are made to the furnace to adjust the bath to the final chemistry. These additions may include ferrosilicon, ferromanganese, ferrochromium, ferromolybdenum, silicomanganese, nickel, or other alloys. Following these additions, the bath is heated for a few minutes to ensure that the alloys have melted and that the desired tap temperature has been attained. The bath is quickly tapped out into a ladle and the refractory is preheated to minimize temperature loss and skulling from the refractories. Any alloy additions required to meet the final chemistry can be added to the ladle before tap out. The furnace should be tapped so as to hold back the furnace slag and to give a clean full metal stream. The furnace spout should be smooth, clean, and dry in order to obtain a full metal stream without a spray into the ladle. Deoxidation of the steel in the ladle can be accomplished in several ways. One method is to add the aluminum as ingots to the bottom of the ladle just before tap out with other ladle additions. In another method, a steel bar with stars or doughnut ingots attached to it is plunged into the ladle when the ladle is one-half to
Electric Arc Furnace Melting / 95 of rammable material. Silica materials or materials high in silica will cause melting and refining problems and will attack the basic lining materials. The slag additions for basic melting are lime and high-lime-containing minerals. The basic practice can remove phosphorus from the steel. Sulfur removal with the basic practice can be very successful if a very high lime-to-silica ratio (2.5 to 3.5) is maintained in the refining slag or if a synthetic desulfurizer is used at tap out in the ladle. Because of the ability of slag to remove phosphorus (by oxidation) and sulfur (by reduction), lower-quality scrap can be used in the charge. A great portion of the oxidizable elements can also be recovered from the lime slag during refining. For these reasons, the basic melting practice is the preferred method for the production of highly alloyed steels, stainless steels, high-manganese steels, and specialty alloys. However, the basic steelmaking practice will be slower than the acid melting practice when its abilities to reduce the phosphorus and sulfur from the steel and to reduce alloys from the slag are employed to capacity.
Basic Steelmaking
Fig. 12
Schematics of foundry charge buckets. (a) Orange-peel bottom. (b) Clam shell bottom
three-quarters full. Small additions of less than the weight of an ingot are made with aluminum shot or wire clippings weighed into a bag and tossed into the metal stream at tap out. The amount of aluminum used for deoxidation ranges from 1.0 to 1.5 kg/metric ton (2 to 3 lb/ ton); the higher-carbon heats require less aluminum, while the lower-carbon heats require more aluminum because of their higher oxygen levels in the steel. All alloy additions should be recorded on the furnace log with the wattmeter reading, along with any comments on the heat. Other deoxidizers are often used to provide a complex deoxidation of the steel. These include titanium and zirconium to tie up nitrogen plus calcium and barium to modify inclusions. The type and amounts of these supplemental deoxidizers are determined according to the need and choice of each foundry. After tap out, the furnace should be inspected for slag line and bottom erosion and should be patched with silica sand or a mixture of lowgrade alumina and silica. This prepares the furnace for the next heat. Loss of Alloys. In the acid melting process, alloying elements that are oxidizable are not
completely recoverable, and care must be taken to estimate their loss when the additions are made to the bath and the ladle, as well as the loss while the heat is being melted. This loss of alloying element can be compensated for based on the data given in Table 4. Alloy Additions. The typical analysis of alloy additions made to the electric furnace is shown in Table 5. Because this is an incomplete list, alloy manufacturers should be consulted for an expanded list and additional detail with regard to residual elements, material size, and packaging.
Basic Melting Practice The melting operation from charge to tap out will vary, depending on the type of charge materials and the alloy additions used to produce steel in the basic melting practice. The lining of the furnace bottom and sidewalls is made of basic materials, such as dolomite, magnesite, chrome magnesite, bricks, and/or rammable materials. The roof is generally made of highalumina brick, and the delta section is made
The basic steelmaking practice is divided into two parts: the oxidizing period and the reducing (refining) period. Both periods are important because of the reactions that occur in each portion. The charge for the foundry practice will consist of 30 to 50% shop returns, such as heads, gates, and casting scrap; the remainder will be 50 to 70% purchase scrap. Along with these materials, a carbon raiser and lime or limestone are added to the charge. The amount of carbon raiser added should be sufficient to increase the meltdown to 0.20 to 0.45% C over the finish carbon level. In addition to the charge to the furnace, a small amount of limestone or lime is normally added to the furnace bottom along with the recarburizer. (It is sometimes considered preferable to add lime to the middle or top of the melt.) Heavy scrap is first placed on the bottom, followed by light scrap and finally medium scrap on the top of the charge. Every effort is made to get the entire charge into the furnace at one time to reduce the time loss caused by delays in backcharging the required weight into the furnace. When stainless steels are being produced, a greater portion of the chromium addition is held out of the charge to conserve chromium, to help keep carbon levels down, and to cool the bath after the carbon boil when temperatures may reach 1870 to 1925 C (3400 to 3500 F) to achieve concentrations of 0.02 to 0.08% C. Because of the negative effect of silica in the basic practice, a conscious effort should be made to keep sand out of the charge. Oily and wet scrap should also be avoided in the charge materials to minimize gas pickup. The oxidizing period will require the addition of lime to form a slag on the bath during melting until approximately 1.5 to 2.0% CaO is
96 / Melting and Remelting A high lime ratio to negate the effect of the sil-
ica in the slag; because this slag must also be fluid, fluorspar additions may be required. High concentrations of manganese oxides also hinder the satisfactory removal of phosphorus. The desired reaction for the phosphorus is: 3CaO þ 5FeO þ 2P ! ðCaOÞ3 P2 O5 þ 5Fe
Fig. 13
Typical types of insulating ladles with cold disposable liners used in foundry operations. (a) Teapot. (b) Bottom pour. (c) Teapot Shank. (d) Lip-pour shank
Table 4
Alloy loss in acid melting practice Loss from charge after blow, %
Loss when making addition, %
Manganese Silicon Nickel Chromium
80–85 90–95 0.0–2 35–50
20–30 5 0.0–2 0.0–3
Molybdenum Niobium Copper
0.0–2 0.0–2 0.0–2
0.0–2 0.0–10 0.0–2
Element
Required element weight to be added as furnace or ladle addition to make up desired weight, %
125 95 100 For less than 1% Cr, and 150% to charge; over 1% Cr, add blow for a total of 120% for >4% Cr alloy. 100 100 100
present in the charge. Once the charge has been melted and temperatures of 1540 to 1595 C (2800 to 2900 F) have been reached, the oxygen lancing can begin. If phosphorus is to be removed, another approach requiring a doubleslag practice should be used. The following conditions must be met to remove phosphorus from the melt during the oxidation period:
after
C (2800 F), because the process is reversible with high temperature. High iron oxide content in the bath and the slag; the use of iron oxide in the early stages of meltdown is required as a source of oxygen to oxidize the phosphorus. Low silicon and silica in the bath and the slag Temperature below 1540
1=3
Tricalcium phosphate is formed in the slag. This slag must then be removed from the furnace to minimize the chance for phosphorus reversion which occurs upon de-oxidation of the melt. This is accomplished by raking off the slag. Limited success is achieved in reducing the phosphorus content, as manifested by the film visible around the sidewalls of the bath following slag removal. This film may contain a major portion of the tricalcium phosphate remaining in the bath, which, if left in the furnace, allows for only minor phosphorus reductions and the reversion of phosphorus to the bath with higher temperatures. The use of slag rakes or a slag flush have become the common methods for removing the slag. After the oxidizing slag has been removed, a new lime-based reducing slag must be added to the melt. Any oxidizable element lost to the first slag must also be added back to the bath during the reducing and final tap out. In the single-slag process, oxygen lancing will remove silicon, carbon, manganese, and chromium from the bath, with manganese and chromium partially reduced from the slag in the reducing period. To minimize chromium loss, a high temperature is required before lancing; therefore, a starting oxygen lancing temperature of 1595 C (2900 F) is suggested. The silicon and manganese contents after oxygen lancing will be between 0.01 to 0.05% and 0.03 to 0.20%, respectively, with carbons in the 0.10 to 0.20% range. An addition of ferrosilicon to bring the melt chemistry to 0.10% Si will stop the boil. Reducing or refining of the heat follows the oxygen lancing of the bath. At this time, a deoxidizer is added to the bath to lower the soluble contert of the melt, after which a reduction in the sulfur can be made. With the temperature of the bath increased after lancing, additions of lime or calcium oxide and fluorspar are made to the furnace to prepare for sulfur reduction and alloy recovery. Sulfur reduction depends on: A high lime-to-silica ratio in the slag, called
the V-ratio
High temperature Low oxygen content in the steel and slag
The reaction may be expressed as follows: ðCaOÞ þ S $ ðCaSÞ þ O
This reaction is reversible. An increase in silica will reduce the activity of CaO, thus retarding the reaction, while highly oxidizing conditions in the slag (high FeO in the slag) will limit the amount of sulfur that can be removed.
Electric Arc Furnace Melting / 97 Table 5
Analysis of alloy additions for the electric furnace Composition, %
Alloy(a)
50% ferrosilicon 75% ferrosilicon Standard ferromanganese 67% silicomanganese Pig iron Low-carbon ferrochromium Molybdenum oxide Low-carbon ferromolybdenum High-carbon charge chromium Standard ferromolybdenum Electrolytic nickel Electrolytic manganese Ferroniobium Copper Ferrovanadium Armco iron Manganese calcium-silicon Calcium silicon Calsibar alloy Hypercal alloy
C
Mn
Si
Cr
Ni
Mo
Cu
Nb
Al
Ca
P
S
V
Fe
Ba
... ... 47–51 ... ... ... ... ... 1.25 max 0.20 max 0.04 ... ... ... ... ... ... 74–79 ... ... ... ... ... 1.25 max 0.60 max ... ... ... ... ... 7.0 78 min 1.0 max ... ... ... ... ... ... ... 0.35 max ... ... ... ... 1.50–2.00 65–68 16–18.5 ... ... ... ... ... ... ... 0.20 max ... ... ... ... 4.0–4.5 0.30–0.50 1.00 max ... ... ... ... ... ... ... ... ... ... ... ... 0.025 max ... 2.00 max 67–72 . . . ... ... ... ... ... ... ... ... ... ... ... ... 4–7 ... . . . 55–60 0.70 max ... Fe + Al ... ... ... ... ... ... 0.12 max 0.25 max 1.5 max ... . . . 65 min 0.70 max ... ... ... ... ... ... ... ... 5–7 ... 2–5 65–70 . . . ... ... ... ... ... 0.03 max 0.07 max ... ... ... 2.5 max 0.25 max 1.5 max ... . . . 60 min 0.80 max ... ... ... ... ... ... ... ... ... ... ... ... 99.9 ... ... ... ... ... ... ... ... ... ... 0.005 max 99.9 0.001 ... ... ... ... ... 0.002 max ... ... ... ... ... ... 0.10 ... 1.3 ... ... ... ... 55.0 + 3.0Ta ... ... ... ... ... ... ... ... ... ... ... ... ... 99.7 ... ... ... ... ... ... ... ... 0.20 max ... 1.50 max ... ... ... ... ... ... ... ... ... 55–60 ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... 98.0 ... ... 14–20 49–55 ... ... ... ... ... 1.75 max 14–20 ... ... ... ... ... ... ... 60–65 ... ... ... ... ... ... 28–32 ... ... ... 5.0 max ... 1.00 max ... 57–62 ... ... ... ... ... 1.5 max 14–17 0.035 ... ... 7.0 max 14–18 ... ... 38–40 ... ... ... ... ... 19–21 10–13 ... ... ... ... 9–12
(a) Aluminum is not included in this list, but it is the primary deoxidizer used in steelmaking. The melter needs to understand that there are various grades of aluminum. For example, primary aluminum in the form of ingots and shot has a purity of 99%, while grades I, II, III, and IV can contain 86–96% Al plus other nonferrous alloy elements. Care should be taken to purchase and add this deoxidizer on the weight contained and not to buy on lowestcost basis.
The slag in the basic melting practice is the key to the production of quality steel. The reduction of iron oxide by the use of coal, aluminum, graphite, or low-sulfur coke is often done to deoxidize the slag and to reduce the content of the less oxides. With this deoxidation, much of the manganese and chromium in the slag are reverted (returned to the bath). After a check on the melt or bath chemistry, alloy additions are made to adjust the chemistry with ferrosilicon, silicomanganese, and ferromanganese, along with chromium, nickel, molybdenum, and other alloys if required. During tapping, final alloy additions are made to the ladle, and the heat is deoxidized with aluminum as the primary deoxidizer plus special deoxidizers as desired.
Melting Heats and Raw Materials The following discussion is a brief review of the basic melting practice. A heat and the raw materials used are also described. Historically, the charge was placed into the furnace with the heavy scrap carbon raiser and limestone on the bottom, the light scrap in the middle, and the medium scrap on the top. Today (2008), it is often common practice to use a cushion of busheling scrap on the bottom. Nonoxidizable materials such as nickel and molybdenum can also be charged into the furnace at this time. The roof is swung into place with an inspection of the top of the charge to ensure that the scrap does not contact the roof ring, and it is then locked down. Melting should proceed as soon as the charge is in the furnace on the No. 2 tap to burrow a hole into the scrap, followed by a change to the No. 1 tap until the charge is melted. The initial use of the No. 2 tap will reduce wear on the furnace delta section.
A pool of metal will form on the bottom of the furnace as the electrodes burrow through the charge. Once the electrodes have reached this pool, the charge will be melted by radiation from the pool and the arc splash off the electrodes. Rust on the scrap, along with oxidizing scrap from the melting charge, will react with the lime and eroded refractories to make a fluid slag. When the charge is completely melted, additional limestone or lime is added to the furnace to increase the volume of slag for an estimated 1½ to 3% of the charge weight (approximately 25 kg/metric ton, or 50 lb/per ton, of lime). At this point, the melt can be processed as a single-slag heat without phosphorus removal, or it can be processed as a double- slag heat where phosphorus removal is attempted. This procedure is described earlier in this section. If a double-slag practice is chosen, the oxygen content of the slag and the bath must be increased while holding down the bath temperature. The old slag must then be removed and a new slag built on the bath. Oxygen lancing can be done before or after this stage to lower the carbon content to the required level, or it could have been done to increase the oxygen content of the bath. Reducing the Heat. The slag must be reduced in iron oxide, and because only a small amount of slag has been added to the furnace up to this point, another 1½ to 3% of charge weight of lime is added to this slag, along with a measured amount of fluorspar. This additional slag volume will dilute the iron oxide content in the slag to reduce the oxygen content and to improve the ability of the slag to retain sulfur. A deoxidizer such as coke, coal, or aluminum can be added to the slag if desired to speed the reaction. Sulfur removal and alloy recovery are the primary concerns while the bath is being heated in preparation for finishing. If a specific sulfur level is needed, further laboratory testing and lime additions may be required
for achieving a higher lime-to-silica ratio, along with fluorspar to thin the slag. If the sulfur level is acceptable, the heat, blocked with ferrosilicon and alloy additions, is finished, and a temperature check is made prior to tap out. If desulfurization is required, it should be done after the heat is blocked. The melt is tapped into a ladle that has been preheated if it is refractory lined, or it is tapped into a cold ladle liner of insulating basic material or boards. Any further alloy additions are made to the ladle prior to tap out. The furnace should be tapped to hold back the slag and to provide a clean full stream. The furnace spout should be clean, smooth, and dry to obtain a full metal stream without a spray into the ladle. Additional sulfur removal may be achieved by adding a desulfurizing agent during tapping. Reductions in sulfur content by 20 to 50% are possible. Final deoxidation of the steel should be accomplished in the ladle; this can be done in several ways. One method is to add the aluminum to the bottom of the ladle and tap onto it. Wire feeding is a popular method that has become a common practice in many steelmaking shops. Another method (poling of aluminum) is to attach stars or doughnuts of aluminum to a steel rod and plunge the rod into the ladle when the ladle is onethird to three-quarters full. (Today, 2008, poling is usually done in the ladle. Historically, it was done in the furnace with even greater effectiveness.) Small additions of aluminum should be in the form of shot or wire clippings weighed and placed in a bag and thrown into the stream while tapping. Finally, ferroaluminum additions can also be used during tapping. The amount of aluminum used for deoxidation is approximately 1.0 to 1.5 kg/metric ton (2 to 3 lb/ton), although higher carbon heats require less, and lower carbon heats require more because of the oxygen levels in the steel.
98 / Melting and Remelting After tap out, the furnace should be inspected for slag line and bottom erosion and then patched with dolomite or other materials. Patching is done by shoveling in or gunning on the patch material. This prepares the furnace for the next heat. Loss of Alloys. A wide range of melting techniques are used in the basic melting practice by foundries—for example, single slag, double slag, deoxidation of slags, desulfurization, and alloy production. Therefore, the individual foundry and melter will need to adjust the recovery and additions to meet particular needs and final chemical analysis based on experience.
Newer Technologies Arc Melting of Iron. In addition to steel melting, electric arc melting of iron is practiced. The first continuous submerged direct current electric arc (Contiarc, SMS Demag AG) furnace for melting iron was installed by American Cast Iron Pipe Company in July 2001. The key to the EAF molten iron production is that it can tap, charge, and arc continuously and simultaneously. As a result, the power does not need to be turned off to tap or charge the furnace. The submerged arc design and the continuous nature of the furnace operation make it more energy efficient than the cupola. In addition, the furnace can take either low-grade steel scrap or direct reduced iron and/or hot briquette iron and combine it with coal and silica to melt steel or smelt iron in the same furnace system, producing 73 Mg/h (80 ton/h). EAF Steelmaking. Several technologies have become common practice in most high-volume
EAF steelmaking shops as well as some foundries. For more expansive treatment of these technologies, the reader is referred to either the article “Steel Melt Processing” in this Volume or to The Making, Shaping, and Treating of Steel, 11th ed. published by The AIST Foundation. The direct current arc furnace has been introduced into steel mills and should be changing the arc melting practices of the future from the three-phase arc furnace to the single-arc direct arc practice (Fig. 14). Use of the direct current arc furnace can depend on the stability of the electricity supply in a given local area. Steel mills using the direct current arc furnace have lower electrode consumption and slightly lower operating costs. The direct current arc furnace was developed for two main purposes: An alternative to the conventional three-
phase alternating current electric arc
For the final stage in practices for melting
prereduced material for the production of liquid iron Since its development, this furnace has been used in the production of ferroalloys as well as carbon, low-alloy, and stainless steels. An existing three-phase furnace can be modified into a direct current arc furnace by the addition of a thyristor converter to its controls and the incorporation of a water-cooling system. Modification of the existing furnace requires the removal of the two outboard electrodes and the positioning of the middle electrode to the center of the roof. It has also been shown that a larger electrode will require additional
hardware to conduct current and to provide support. The electrode is the cathode in this system. In the three-phase system, current passes from one electrode to another through the charge; in the direct current system, the current passes from the electrode through the charge or bath to the anode located at the bottom of the furnace. The current from the bottom of the furnace passes through graphite- or tar-impregnated refractories to a copper-base plate to the outside cables, or through pins of steel connected to outside current cables (Fig. 14). The heat distribution in the direct arc is 360 , while in the three-phase furnace, the heat pattern is off the edge of the electrode, causing localized refractory heating and erosion. The electrode usage per ton of steel is reported to be 50 to 60% of a conventional three-phase furnace. The shell bottom of the furnace needs to be modified to accept the current-carrying loads that will result and should be insulated to protect it from shorting out. Thermocouples and cooling air will be required for monitoring overheating or loss of current-carrying ability in the bottom. The furnace foundation may also require modification to accept power cables and air cooling ducts. Patching of the refractory is limited to the slag line, sidewalls, and roof. Noncurrent refractories on the bottom will cause poor conduction. Another consideration to be made by the melter is the heat sequence, because direct arc melting requires long start-ups or a heel of metal over the bottom to make electrical contact. Ladle metallurgy has become the processing method of choice in all large EAF steelmaking shops and several foundries. As their name implies, multipurpose ladle refiners have several important capabilities: precise temperature control, vacuum degassing, inert gas bubbling/ stirring for inclusion agglomeration/flotation and compositional uniformity, and alloy addition systems for precise compositional control. Argon stirring of heats of steel in ladles as a method to promote the removal of nonmetallic inclusions via agglomeration and flotation has become a primary tool for making clean steel. The need for dry gas is very important, because argon containing more than 5 ppm of moisture will induce gas pickup. Systems for the accurate, real-time assessment of hydrogen in the liquid metal have become commonplace and are an invaluable aid to the steelmaker. Calcium wire and other wire and powder injection systems have become important tools for specific deoxidation practices, compositional modification, and inclusion engineering. REFERENCES
Fig. 14
Cutaway view showing the basic components of a direct current arc furnace
1. J.F. Schifo and J.T. Radia, Theoretical/Best Practice Energy Use in Metalcasting Operations, U.S. Department of Energy, May 2004 2. Chap. 2, Steel Industry Technology Roadmap, AISI, Dec 2001
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 99-107 DOI: 10.1361/asmhba0005197
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Cupola Furnaces William L. Powell, ThyssenKrupp-Waupaca (Retired) Alan P. Druschitz, University of Alabama at Birmingham Jim Frost, ACIPCO Melting Department
THE CUPOLA is basically a simple cylindrical shaft furnace that burns coke for fuel. The heat generated is intensified by the blowing of air through the heated bed of coke. Alternate layers of scrap metal and coke are charged into the top of the cupola. In its slow descent, the scrap metal is heated to the melting temperature by direct contact with the upward flow of the hot gases from the coke combustion. The moltenmetal droplets collect in the inner portion of the cupola known as the well, where it is discharged into the foundry from the taphole by intermittent tapping or by continuous flow of molten metal and slag. Melting in cupola furnaces dates back several centuries to old cask-type units that were blown with hand-pumped bellows and hand charged with bits of metallic iron and scraps of cast iron. The cupola as we know it today (2008) was patented in 1794 by John Wilkinson. Prior to the 1950s, the cupola was the predominant furnace for iron melting because of its simplicity and the low cost of molten iron produced. Progressive improvements were made in electric induction furnace equipment since the 1950s. At their inception, induction furnaces required little or no pollution-control equipment because mostly clean, new scrap was charged. This mix of new scrap iron and pig iron yielded a low-sulfur iron that was satisfactory for the fast-growing ductile iron industry. Electric induction furnaces also began to replace many cupolas in small-tonnage operations where both gray and ductile iron or several other types of irons were being produced. The ability to make new chemical compositions in each melted batch provided more flexibility to the foundry operator. Electric induction melting was found to have a lower melting cost and better chemical control under many of these conditions. The number of cupolas in use in the United States has decreased from 4000 in the early 1950s to less than 1400 today (2008). However, the cupola remains the predominant tonnagemelting method, with over 60% of today’s iron tonnage still melted in cupolas. Through the years, the cupola has been progressively refined
and moved toward more efficient combustion, better chemical composition control, and more flexibility in the use of widely varying raw materials. However, during the melting process, the cupola emits considerable amounts of smoke and particulate matter. Known as emissions, this waste gas stream is now cleaned to meet the highest environmental standards, but this has required much engineering effort and considerable expense. With environmental laws requiring more elaborate pollution-control equipment, the cupola lost much of its advantage of low cost and simplicity, especially on small cupolas. In high-iron-tonnage operations, the cupola remains the most efficient source of continuous high volumes of iron needed to satisfy highproduction foundries or the multiple casting machines of centrifugal pipe producers. The continuous stream of cupola-melted iron is not interrupted by charging, slag removal, or chemical composition alteration. On a per ton basis, cupolas usually require less labor to operate and can use lower-cost forms of scrap. In such high-tonnage, continuous-demand situations, the cupola has remained the lowest-cost method of melting. Although the cupola is more difficult to control, many consider it to have favorable iron qualities, such as a low tendency toward chilling, good fluidity, better machinability, and improved shrinkage characteristics compared to electrically melted cast irons.
Progressive Refinements in Cupola Equipment Since the early days of the simple cupola, continuing technological refinements have been progressively applied to the cupola melting process to improve efficiency, flexibility, and control. Many of the refinements have stood the test of time and greatly added to the cupola melting technology. Some other ideas have not proved to be technically sound or have lost appeal because of ever-changing economic conditions, industry needs, state and federal law changes, and a host of other reasons to abandon
them by the melting community. Some of the more successful improvement technologies include the following. Preheated air blast was found to improve efficiency of the combustion reaction by allowing the input of British thermal units (Btus) from an external source, thus increasing melting rate, temperature, and carbon pickup with the same amount of coke, or alternatively, the amount of coke used to melt a ton of iron can be reduced for the same melt rate, temperature of the iron exiting the cupola, or greater carbon pickup. Early blast air heating systems could preheat the incoming air only 150 to 260 C (300 to 500 F), but even this limited additional temperature demonstrated the advantages of the preheated blast. The most successful hot blast systems have been externally fired with gas and are capable of up to approximately 650 C (1200 F) preheat. Higher blast temperatures would continue to improve performance, but the availability of heat-exchange tube materials has a practical economic limit of approximately 650 C (1200 F). Recuperative hot blast systems have begun to replace the externally fired hot blast system. The recuperative hot blast system is an integrated system, which combines both the blast air heating and gas cleaning system. Emission gases are usually first combusted to raise the heat content of the offgas, then the heat is extracted via a gasto-gas heat exchanger (the hot blast portion of the system), and finally, the cooler gases are cleaned of dirt particles. In some designs, the gases are cleaned and then burned. Both have distinct advantages and disadvantages but are effectively used to preheat and clean the offgas stream. Oxygen enrichment and injection, which adds 1 to 15% O2 to the blast air, also intensifies the combustion reaction by increasing temperature, which in turn increases the melting rate. The efficiency is derived from blowing in a higher number of oxygen molecules per volume of air and lower amounts of inert nitrogen molecules. Enrichment is limited to oxygen increases of approximately 2% and is accomplished by adding pure oxygen to the blast ductwork. Injection allows much higher additions of
100 / Melting and Remelting oxygen. Increases of 4 to 15% can be achieved. A pipe is inserted into the center of each tuyere for adding the oxygen. Often, a nozzle is added to the tip of the pipe to achieve supersonic flow rates for the oxygen. This increased velocity allows blast air and oxygen to penetrate deeper into the coke bed. Injection of fine coke and other materials, such as silicon carbide, through the tuyeres has proved to be effective in both increasing carbon and silicon content and shortening response time for chemical changes in the melt composition. The effectiveness of this technology is somewhat dependent on the operating parameters of the particular melting operation. Water cooling of the cupola shell has extended the length of the melting period beyond 24 h and has permitted a single cupola to be operated for weeks before requiring a refractory repair shutdown. Water cooling of the shell allowed the removal of the refractory lining and its effect on the slag chemistry. The same cupola then could operate with a choice of basic, neutral, or acid slag chemistry simply by regulating the amount of basic flux added. The other benefit is the reduction of sulfur in the molten iron, allowing those ductile iron treatment methods requiring low-sulfur-base irons to be used with no external desulfurizing step in the molten-metal processing. Duplex electric holders are large holding furnaces that receive the iron after leaving the cupola. They have ensured that a constant supply of iron is available at all times while also reducing variations of composition and temperature of the iron. Large electric channel induction furnaces hold between 30 and 90 min of iron. These electric channel furnaces can provide a modest boost in iron temperature or maintain iron temperature over cupola melting down-periods. Computer applications have been increasingly used to record, monitor, and control more cupola conditions. Today’s (2008) melting operations are as highly computer controlled as any manufacturing operation. Virtually all control of the motors, valves, flow devices, dampers, and so on are computer controlled. All the various inputs of temperature, flow, pressures, status of valving, and so on are fed back to a computer for data storage and analysis. Continuous analyses of top gases have made it possible to monitor combustion efficiency and to signal any water leaks to avoid explosive gas combinations. Basic slags have made possible the production of low-sulfur iron directly from the cupola and have increased the carbon pickup. By regulating slag basicity, a wide range of iron compositions can be obtained from a wide range of charge materials. Technology Failures. Some technologies have not passed the test of time in melting operations. These have failed the test of providing a true increase in efficiency, being technologically sound, or may have simply been a passing fad. Many are still reflected in
the technical literature of the time, and care must be taken by individuals to determine if they are truly an improvement to the cupola melting operation. Some of these technology failures include the following. Divided-blast cupolas divide the air between two rows of tuyeres in controlled proportions. This was reported to increase combustion efficiency and accomplish either coke reduction or increased melt rate and temperature at the same coke usage. Reversed taper shell was a cupola design meant to concentrate the air across the entire cupola cross section for more uniform melting area and lower coke usage. Ground rubber tires were tried as a coke substitute. Rubber tires are largely composed of carbon and have a high Btu content. Considered a waste product, their use in the cupola would be a way to reduce both a national waste item and coke cost. However, the volatile gases produced did not provide a carbon source and also proved difficult to scrub from the emission gases. Natural gas enrichment to either the blast or the offgases was proposed as a method to raise the Btu content of the blast air. However, the water vapor released during combustion proved to have a negative effect on the heat content within the cupola. Humidity control of the blast air has been employed on some cupolas to eliminate the effect of varying humidity when very close control was important. However, the cost of drying the absorption media proved to exceed the benefits of water removal.
melting needs and inputs: type and size of scrap available in the market place, desired melting rate, melting rate flexibility needed, carbon pickup desired, use or nonuse of additional oxygen, length of desired campaign, length of time for relining, chemistry control limits, and so on. Optimum cupola height has been well studied, and 8.5 to 10 m (28 to 34 ft) of effective melting height is universally accepted as the optimum height for a cupola. Cupola shell design (both diameter and height) as well as tuyere placement, offgas take-off location, and so on require detailed forethought and planning during the design and installation of a new cupola. The foundation of the cupola structure must be adequately engineered for the size of the cupola and the loads it will carry. Footing loadings will be for both the cupola itself and the burden contained within the cupola. A large cupola will often contain 136 to 181 Mg (150 to 200 tons) of scrap iron, coke, stone, and molten iron. An often overlooked component is the
Cupola Construction and Operation This section discusses the shell, intermittent or continuous tapping, tuyere and blower systems, refractory lining, water-cooled cupolas, emission-control systems, and storage and handling of the charge materials.
Shell Design The cupola shell is a vertical, cylindrical steel shell and is the container of the combustion and melting operation. Figure 1 shows a typical modern-day cupola. The size, or specifically, the diameter, of the cupola shell basically defines the limits of operation of the cupola. It limits the minimum and maximum melting rate per hour of operation, the carbon pickup ability, the length of the melting campaign, the maximum size of the scrap that can be charged, the size of the coke purchased for melting, and a host of other parameters. In the past, cupolas were purchased as somewhat standard pieces of equipment. For the past 30 years or more, cupolas have been very much a custom-designed piece of melting equipment. Cupola manufacturers can suggest an optimum diameter based on the user’s specific anticipated
Fig. 1
Typical modern-day cupola showing the construction and arrangements of the principal cupola components
Cupola Furnaces / 101 dynamic loading place on the cupola during charging. Charge drops of 4 to 7 Mg (4 to 8 tons) of scrap falling 1.5 to 6 m (5 to 20 ft) are not uncommon. Charge doors today (2008) are almost always not a door but are actually the opening at the top of the cupola. Traditionally, cupolas were side charged near the top of the cupola. Today’s emission systems demand controlled volumes of offgases. The side charge door allows uncontrolled volumes of air to enter the gas streams, thus causing erratic cleaning of the waste gases. The top-charged cupola employs a “plug” of scrap at the top of the cupola to effectively limit the amount of air being pulled into the cupola stack. A second benefit is the ability to soft-drop a charge onto the burden, thereby reducing the crushing of the coke charge during the charge drop. Bottom doors at the bottom of the cupola stack facilitate emptying of the stack of coke ash and residual metal at the end of an operational melting campaign in preparation for refractory repairs and the resumption of melting. The doors are usually constructed of heavy steel or ductile cast iron and are massive enough to support the entire burden of the charge material and coke as well as the dynamic forces encountered during charging. These doors cover the entire bottom cross section so as not to impose any restrictions during the dropping of the bottom. During the melting operation, the doors are held rigidly in place by heavy steel posts or similar mechanisms. There can be no flexing of the doors, either mechanically or thermally, during the melting operation, or cracks in the bottom refractory will occur that may result in an iron runout. A sand bottom is rammed over the closed bottom doors to contain the melted iron in the well. This sand bottom can either be all sand, all refractory, or a combination of both. In smaller cupolas with short-duration melting campaigns, the bottom can be as simple as foundry molding sand. For larger cupolas with longer melting campaigns, special sands and clays are preferred, with the addition of a refractory layer. There is no general agreement among melters on bottom sand materials and arrangements. What can be agreed upon is that once a system is determined, strict controls must be put in place to assure quality of installation every time a new bottom is put in service. A sand bottom is sloped toward the taphole to ensure that iron and slag continuously flow to the taphole and out of the cupola.
Intermittent or Continuous Tapping Cupolas Iron flow from a cupola can be made to be intermittent or continuous. The first cupolas were intermittently tapped by opening and closing the taphole as in a blast furnace operation. For intermittent tapping, a refractory plug is pushed into the open taphole after iron has
stopped flowing, and the plug is left in the taphole until a volume of iron has accumulated in the cupola well, then the plug is pulled out and iron begins to flow. This process is repeated for the duration of the melting campaign. Slag is accumulated within the cupola well area and flows out a back slag hole. Most modern-day cupolas are tapped continuously, with iron flowing under a refractory dam and over a lip in the cupola spout. The lighter slag floats on top of the stream of molten iron. A notch in the side of the cupola spout (also called the well) allows the slag to flow out of the front spout area. This slag flows continuously to a slag disposal area. Cupola slag has traditionally been quenched in a water basin or a trough of flowing water. However, due to wastewater rules prohibiting the discharge of process water, a variety of dry slagging methods have been developed. The various dams and notches in the front spout area are constructed of refractory materials. Due to refractory wear, the levels of metal dam and slag notch must be carefully calculated and maintained throughout the melt campaign. The dry-bottom cupola is a recent European development in which iron and slag are not collected in the well of the cupola but run into a special vessel outside and beside the cupola. In this outside vessel, the accumulated metal and slag are separated by gravity as they flow through the vessel. The advantages claimed for this method include more consistent iron analysis and easier repair of refractories. Cupolas often have two vessels attached to the front, allowing the repair of one while melting iron through the second vessel. This arrangement allows for longer melting campaigns.
Tuyere and Blower Systems Air for the combustion of the coke is introduced through holes in the shell that have water-cooled pipes, called tuyeres, protruding into the cupola. Tuyeres are constructed of either cast copper or fabricated copper coils twisted around a large piece of pipe. They are equally spaced around the cupola, with the number and size proportioned to cupola diameter. Tuyeres are sized to deliver a specific relationship of air volume, pressure, and velocity. On water-cooled hot blast cupolas, the desirable blast air velocity range is 4600 to 7600 m/min (15,000 to 25,000 ft/min). This velocity has been determined to give the blast air sufficient velocity to maximize the penetration of the coke bed. Blast air blowers generate the necessary blast air volume and pressure to melt iron. The air required to melt 900 kg (1 ton) of iron is 620 to 680 m3 (22,000 to 24,000 ft3) at standard temperature and pressure or 770 to 835 kg (1700 to 1840 lb). Air from the blast air blower is transmitted to the tuyeres through a large-diameter duct system to the cupola wind box. Blast pressure, usually called backpressure, normally ranges from 3.3 to 8.5 N (20 to 65 oz/in.2) or more, depending
on charge density. Blast blowers can be of several designs, such as positive displacement blowers, fans, or centrifugal-type blowers. The latest engineering design preference on large cupolas has been to use two or three blowers in parallel with each other and use automatic air weight controls to control the air volume produced by them. Blast blowers typically consume the second-largest amount of electrical energy in a melting department, second only to the emission system fans. The biggest waste of efficiency in the melting department is due to leakage in the blast air system. This is due both to missing insulation allowing heat loss before the heated blast air reaches the cupola, and leaks in the blast air ductwork system caused by thermal and mechanical movement.
Refractory Lining Conventional cupolas are usually lined with refractories to protect the steel shell against abrasion, heat, and oxidation. Figure 2 shows a cross section of a cupola with the various refractory types indicated. Lining thicknesses range from 114 to 305 mm (4.5 to 12 in.). Most linings are either fireclay brick or a sprayed or gunned alumina material. As the melting campaign progresses, the refractory lining in the melt zone is progressively fluxed away by the high temperature and oxidizing atmosphere, reducing the lining thickness. Early cupola melting practice was to stop melting, drop the bottom, and repair the lining at the end of the day’s melt. After cooling the cupola interior, repairs were made to the refractory lining, restoring it to its original diameter. Often, a pair of cupolas were alternated each day to facilitate the repair and preparation of the cupola lining after each day of use. Today’s (2008) improved refractories and shell water cooling have, for the most part, eliminated the need to have two operational cupolas. Many methods and materials are employed for cupola refractory repair. The most widely used method of lining repair uses a pneumatic gun to blow a mixture of clay and ganister through a nozzle containing water that becomes entrained in the stream as it exits the mixing nozzle. New methods such as shot creting have been satisfactorily used for cupola refractory patching and for extending the life of the cupola refractory. Shot creting creates a denser refractory layer and allows for a greater variety of refractory materials to be used. For repair of the taphole area, refractories containing silicon carbide particles are rammed in with pneumatic hammers. For the upper stack portion of the cupola, brick or blocks are layered with a thin mortar joint. Specialty refractories are being developed for every segment of the cupola to meet the exacting needs of the specific cupola operation. The advent and growth of ductile iron in the 1950s made it desirable to find a way to remove sulfur inside of the cupola. This was
102 / Melting and Remelting accomplished by using a sufficiently basic slag, as is done in basic steelmaking furnaces. However, a highly basic slag cannot be maintained in contact with acid refractories. After considerable experimentation, suitable basic refractories were successfully developed for use in cupolas. Originally, linings were either magnetite or dolomite brick. Patch materials composed of dolomite or magnetite aggregate were then
Fig. 2
developed that could be rammed or blown in using the pneumatic gun.
Water-Cooled Cupolas When the melting operations of foundries were extended to a continuous 16 to 20 h of operation, there was little or no refractory lining
remaining in the melting zone at the end of the melting campaign. When three-shift melting operations became a necessity, the water-cooled cupola was developed. The early water-cooled shells were simple tanks or jackets added to the outside of the cupola shell. Figure 3 shows an early water-cooled cupola design. Water circulated through the tank with little velocity, and cooling was basically just through heat conduction into the water. A limiting factor to the tank method of water cooling was that if a hot spot developed on the shell, it could not be observed nor could additional water cooling be added to the area during the melting campaign. Later water-cooling developments led to the cascading water shell, where a thin layer of water flows rapidly down a cupola shell, wiping heat from the shell in the process. At the top of the water-cooled area, a water curtain is applied to the shell by spray rings encircling the cupola. The water curtain cools the shell as it cascades down the shell, and the water is collected in a trough or moat at the well. One advantage of the open external cooling is that the outside of the cupola shell can be sand blasted at intervals to remove excess deposits of lime and other built-up materials. This sand blasting also restores the thermal conductivity of the steel shell. With increasing dependence on water cooling, it was found that the refractory lining could be eliminated from the melting zone and that the cupola could be operated continuously without a refractory lining. However, the advantages of water cooling are achieved at a loss in thermal efficiency. With demands on lower cost for coke and greater operating efficiencies,
Fig. 3 Cross section of a modern-day cupola showing typical refractories used in lining a cupola
Early water-jacketed cupola used to extend operating hours of a melting campaign
Cupola Furnaces / 103 today’s (2008) cupolas are a hybrid of water cooling and refractory linings. When iron is needed for only two shifts, the water-cooled cupola can be banked during the downshift and brought back on stream when needed. These water-cooled cupolas can even be banked over a weekend as well as operated without the need for refractory repairs for 2 to 4 weeks (or even longer under some operating conditions) before requiring the dropping of the cupola bottom for repairs. Some of the largest cupolas are equipped with double tapholes and double rear siphon slag spouts to enable the cupola to be operated continuously for even longer melting campaigns.
Emission-Control Systems Equipment has been successfully engineered to enable the present-day cupolas to meet the most rigid of environmental regulations. The two most common types of collectors currently in use today (2008) are the high-energy wet scrubber and the dry bag house. However, current environmental regulations have placed even higher demands on the cupola operator for cleaner air emissions, and only dry bag house systems are able to meet these new regulations. High-energy wet scrubbers use high-velocity water jets to impinge atomized water droplets on effluent dust particles passing through a venturi scrubber. The result is a dirty water sludge. This dirt and water sludge requires thickening, separation, and dewatering. Due to the various hazardous chemical elements (typically cadmium and lead) present in the cupola gas stream, the sludge and water are deemed toxic and require further chemical treatment to neutralize the hazardous elements. This is a complex science and a materials handling
challenge. Figure 4 illustrates a typical wet scrubber emission-control system connected to a recuperative hot blast system. The dry bag house environmental system is the best cleaning system for meeting today’s (2008) air emission standards. In this unit, the dust particles are collected on the surface of fabric bags, and then the bags are shaken to free them of the dust. Gases leave the cupola at temperatures ranging from 95 to 1100 C (200 to 2000 F), depending on what is happening in the cupola. The cloth bags can only withstand temperatures of approximately 230 C (450 F); thus, a system for cooling the gas stream before reaching the bag house is necessary to protect the bags from excessive heat. The two most popular heat-exchange systems are the air-to-air and air-to-liquid heat-exchange systems. The engineering for either is complex and a challenge for mass flow and heat-removal requirements. To better understand the operation of the recuperative, dry cleaning system, it can be broken into smaller subsystems. Figure 5 shows a modern dry bag house cupola emission system. The system design is basically simple in terms of what is being accomplished in each stage. Cupola offgas exits the cupola at a temperature between 95 and 540 C (200 and 1000 F) under normal operating conditions. This gas is rich in carbon monoxide (CO) but depleted in oxygen and has the potential for combustion. Air is added to the mixture in a controlled amount, and the mixture is ignited via pilot burners in a large combustion chamber. This combustion results in a stable temperature of approximately 820 C (1500 F). The entire gas stream is now passed through either an air-to-air or air-to-liquid heat exchanger to heat the blast air going to the cupola, or a secondary heat-removal step can be added to heat air for
building heat and/or for other processes in the foundry operation. The gas stream exits the cooling tower at a temperature of 135 to 165 C (275 to 325 F) and is now ready to enter the bag house at a temperature that is low enough not to harm the filter bags. While the pores of the cloth are very small, the particles in the gas stream are much smaller, so, in theory, the dust particles would pass through the bags and into the atmosphere and not be trapped. However, the dust layer that builds up on the bags traps the smaller dirt particles in a porous cake of dust. The bags are mechanically shaken or popped with a shot of air to remove some of the dust cake. As in the wet system, the dust may be a health hazard, depending on the concentration of hazardous elements present. Process chemicals may be added to the dust/gas stream before the bag house to neutralize these hazardous elements. Each lot of bag house dust is tested and then properly disposed of in a suitable landfill.
Charge Materials Storage and Handling of Charge Components Essential to efficient cupola operation is a good materials storage and handling system. Metal charge components are received via dump trailers, rail cars, or bulk containers. Unloading into storage bins can be by selfunloading dump trucks or by using an overhead charge crane with a magnet. The size of scrap storage bins varies widely with how much is to be melted each day, distance from the scrap-processing source, reliability of the transportation system, severe weather issues, and so on. Each charge material needs to be transferred to a weigh hopper, where a designated weight is measured out and transferred to a charge bucket for travel to the top of the cupola. The design possibilities for the charging area are endless. Figure 6 shows a typical charge yard design. Coke and limestone are usually stored in covered overhead storage bins and are discharged into a weigh hopper and on into the charge buckets. One factor in charge yard design that is often overlooked is planning for flexibility. Charge yards are expensive to construct and are longterm structures. The needs of the foundry and the availability of charge materials are likely to change over the lifetime of the facility, so it is necessary to build in some flexibility in design to accommodate for an unknown future need.
Coke
Fig. 4
Typical wet scrubber with heat recuperation hot blast used to clean emission gases
As the major fuel, coke is the most important charge material. High-quality, foundry-grade coke is essential for optimum cupola performance. The amount of coke in the charge mix usually ranges within 8 to 16% of the metal charge, depending on the iron temperature
104 / Melting and Remelting
Fig. 5
Fig. 6
Typical dry bag house emission and recuperative blast air heating system used to clean emission gases
Charge yard arrangement showing scrap delivery, storage bins, weigh hoppers, and charging bucket
required, the percentage of steel scrap used, and many other variables. Foundry coke must be strong to support the charge materials in the cupola stack and have sufficient impact strength to maintain its size during the impact of charging. These mechanical properties are important since they affect hot blast penetration and carbon pickup. Experience has shown that performance is optimized when coke size is consistent. Fine particles of coke should be screened out of the coke portion of the charge, because they fill the interstitial spaces between the coke and scrap iron pieces; this limits air distribution and overloads the emission system with excess carbon monoxide gas. On large cupolas, usually the largest available coke size is preferred, but on small cupolas, coke screened to any consistent size range provides good performance. Coke is manufactured to coke industry standards. Table 1 is a typical foundry specification for coke. Typical foundry coke should have a low ash content (<10%) and a chemical structure that provides for optimum carbon pickup and combustion efficiency. Low sulfur content (<0.70% S) is required for preventing excessive sulfur pickup by the molten iron. Volatiles should be less than 1.0%, and fixed carbon should be above 90% for maximum cupola performance. Physical properties of coke also must be controlled. Coke strength must be high to prevent crushing in the cupola stack. A dropshatter test best measures coke strength. Abrasion resistance, as measured by stability or hardness, must be as high as possible. Efforts have been directed toward the use of limited amounts of less-expensive substitutes
Cupola Furnaces / 105 Table 1 Typical specifications for foundry coke Fixed carbon Volatile Ash Sulfur Moisture Shatter strength 50 mm (2 in.) screen 75 mm (3 in.) screen 100 mm (4 in.) screen Stability index Coke size
91.5–94.5% 0.1–1.0% 5.0–7.5% 0.4–0.7% 0.1–5% 96.0% min 92% min 78% min 38.0% min As required by cupola size
for quality coke, such as briquetted coke fines, anthracite coal, injected fine coke, and various substitute fuels. However, all of these substitutes reduce cupola efficiency, and the effectiveness of each must be weighed against total fuel cost-savings.
Metallic Charge Materials The modern cupola can melt a wide range of metallic charge materials, most of which can be classified as scrap iron and steel, various cast irons, and other ferrous materials. Through its use of scrap materials, the casting industry contributes greatly to environmental cleanup and resource reutilization. Steel scrap was first melted in low percentages to lower the carbon content of gray iron in order to produce higher-strength irons. Today’s (2008) cupolas use steel scrap at a rate of 25 to 90% of the charge mix. Obsolete scrap iron is processed in a variety of ways. It can be cut with a cutting torch, sheared with a large blade shear, shredded in a hammer mill, or bailed into cubes. Each method produces a useable material for cupola melting, but each has its own melting characteristics. Cupolas are generally flexible enough in design to melt any of these prepared materials, but an optimized cupola design can make any scrap preparation method more efficient. Each preparation method has a unique cost associated with it and is local market dependent, so there is not one method or material that is preferred. Cast iron scrap is a common cupola charge material. It is broadly divided into four or five classifications of automotive, machinery, heating, and pipe. The grades useable in any particular foundry are largely dependent on the foundry products being produced. Scrap size and chemical composition are the basic source of the different grades of cast iron. Cast iron has a high carbon content and low melting point. It generally gives a consistent chemistry to the melt and lowers the chill tendency of the melt. Due to its good melting properties and limited availability, it is generally higher in price than scrap steels. Foundry returns, such as gates, runners, and internally generated scrap castings, usually constitute 30 to 50% of the pouring weight of a mold, and thus, they are returned to the charge
yard for melting in approximately this ratio. A melting operation will need to melt all of the foundry returns sent to it. As in cast iron scrap, carbon content is high and melting temperature low. Since this is internally generated scrap, it should have the most consistent chemical content of all metallic charge materials. Pig iron is considered a premium charge material because of its uniformity of composition and size, and it is free of detrimental tramp elements. In the early years, pig iron was the predominant charge component for ensuring quality iron with sufficient carbon content and low quantities of detrimental elements. Despite economics that encourage the use of more scrap and less pig iron, pig iron has remained a basic charge material because of its consistent size and quality. A wide range of scrap metals can be melted in the cupola. However, some forms of scrap metal are not practical to use in a cupola charge. Cast iron borings charged directly into the cupola are excessively oxidized if not blown out of the stack. With briquetting, limited proportions can be used. Steel turnings are more impractical because they are so thin that they readily oxidize high up in the cupola stack. Thin sheet steel has a very low density as received from the source and cannot be used without being compressed into bundles. Compact punchings tend to clog interstitial coke passages and hinder combustion efficiency. Very thick steel parts and large-diameter shafts are not completely melted within the melting zone as the charge descends into the cupola, and this results in low iron temperatures. Recently, a manufactured product consisting of cast iron borings, steel turnings, grinding dust, coke breeze (optional), and a binder has been successfully pressed into a brick. This allows the use of materials previously not usable in the cupola. Brick use must be limited, but use rates of up to 30% of the charge are possible.
Ferrous Alloys Any number of alloys can be added to the cupola for melting into the base iron. Most commonly, these alloys are purchased preblended with an iron-base material. Silicon, which is added to yield the desirable final silicon content, is usually added in any of the following forms: Lump ferrosilicon: 50 or 75 wt% Si content Silvery pig: 16 to 20 wt% Si content Silicon briquettes: composed of silicon fines
briquetted with a suitable binder, often portland cement, with the briquette containing approximately 40 wt% Si Silicon carbide briquettes: made of silicon carbide plus carbon briquetted together with a suitable binder Manganese briquettes of various compositions can be charged into the cupola in several forms to obtain the desired final manganese
content. The briquettes are generally composed of ferromanganese granular material and a cement binder. Contained manganese percentages are usually 50 to 75 wt%. Alloys of chromium, nickel, molybdenum, and other required alloys can be purchased and added to the cupola. Care must be taken in the selection of alloys to be sure that the size of the additive is sufficiently large to minimize oxidation and losses to the emission system. Briquettes must have sufficient strength to survive charging operations and the descent into the cupola melting zone. Binders must be evaluated to have sufficient hot strength to prevent premature breakdown of the binder system before entering the melting zone.
Fluxes Charge materials contain large amounts of nonmetallic materials, such as dirt on scrap, sand adhering to returns, slag attached to pig irons, and so on. Additionally, coke ash, oxidation products from melting, cement binders, and other noniron materials must be fluxed into low-melting-point compounds. These slags must be sufficiently fluid to be able to exit the cupola and float to the top of the iron so they can be skimmed off and not sent to the holding furnace. Limestone or dolomite additions of 2 to 8% of the metallic charge (or 30 to 45% of the coke portion of the charge) are generally needed in an acid cupola to contribute enough basic constituents of calcium oxide and magnesium oxide to impart sufficient fluidity for the slag to flow freely through the coke spaces, into the well, and down the front slagging trough. In the basic slag cupola, more basic fluxes are added to net an excess of basic slag constituents. Limestone is a natural calcium carbonate or dolomite stone. A lump size of 25 to 75 mm (1 by 3 in.) is preferred. Limestone composition varies widely across the country, so selection of stone for a melting operation is critical. Screening of limestone fines is mandatory for a good flux stone. For limestone (calcium carbonate limestone), CaO contents should be above 45%. For dolomite stones (those containing high levels of calcium-magnesium carbonate), CaO levels should be approximately 30% and MgO levels approximately 20%. Silica is an impurity and should be less than 2%. Several special or supplementary fluxes are occasionally used. When used, they are usually added very sparingly in addition rates of 0.2 to 2.0% of the metallic charge weight. Soda ash (Na2CO3) is a strong slag liquefier and is added as briquettes to the bed to supplement the limestone. Fluorspar (CaF2) is a powerful fluidizer that is used to clean up a cupola clogged with excessively viscous acid slag. In highly basic slags, fluorspar is needed to maintain fluidity. However, excessive fluorspar can cause environmental problems.
106 / Melting and Remelting Calcium carbide (CaC2) serves the multiple functions of deoxidizing, raising carbon, raising temperature, and providing a strong calcium oxide (CaO) flux contribution.
Cupola Control Principles Coke Bed. Preparation of the coke bed is vital to the initial start of the melt and to the maintenance of the melt at the specified level in the cupola combustion zones. Before metal is charged, a bed of good-quality, properly sized coke is ignited and blown to a red heat either internally in the cupola or externally and charging it hot. The height of the coke bed is adjusted to a specific height above the tuyeres that is determined by experience and based on blast volume, cupola diameter, and desired melting conditions. Bed flux and metal charges are then added, the blowers are turned on, and the melting begins within a few minutes. With a high coke bed, melting occurs higher in the cupola stack, producing slower melting rates and hotter iron with less oxidation and more carbon pickup. A low bed melts charge materials lower in the cupola, with faster melting rates and colder iron temperatures with more oxidation. Charge coke is the coke intended to replace the coke burned in the melting of each layer of charge materials. It is weighed and added to each charge to maintain the bed level within the cupola during the melting cycle. While the exact height of the bed during melting cannot be observed, a skilled operator knows its relative position by observing melting rates, slag temperatures, iron temperatures, and final iron chemistry. Melt Rate. A cupola is expected to maintain a melting rate sufficient to meet the needs of mold production. The cupola size must be chosen for the midrange of production demand. Several general rules can predict melting rates. One rule is that 4.5 kg (10 lb) of iron can be melted per hour for each square inch of crosssectional area. A second rule is that 3.2 to 3.3 Mg (3.5 to 3.7 tons) of gray iron can be melted with 0.5 m3/s (1000 ft3/min) blast air, and 2.9 to 3.0 Mg (3.2 to 3.3 tons) of ductile iron can be melted with 0.5 m3/s (1000 ft3/min) blast air. By increasing or decreasing the blast air rate, melt rate can be varied by approximately þ 20% with reasonably good efficiency and control. Beyond these limits, overblowing a cupola increasingly oxidizes the iron, causing lower carbon and silicon levels and decreasing the iron temperature. Underblowing will lower the iron temperature because the blast will not penetrate the cupola with sufficient air distribution for optimum combustion efficiency. Changes in melting rate can also be accomplished by changes in the charge mix ratio of cast iron components to steel and by varying the hot blast temperature. If oxygen injection is used, changes in oxygen volume can widen the melt rate control range.
Iron Temperature. The iron temperature at the cupola spout must be sufficient to overcome temperature losses due to metal handling from the cupola to the electric furnace or holding ladle, the holder to the pouring ladles, alloying additions, the treatment process (for ductile iron), and radiation losses from the pouring ladles. Iron temperature can be measured by optical pyrometers, immersion thermocouples, and/or continuously recording radiamatic instruments. Carbon and silicon concentrations are most important in the control of iron properties. Sufficient concentrations of carbon and silicon must be maintained to avoid chill and to encourage a microstructure suitable for the section size of the casting and the strength level required. In ductile iron, sufficient carbon and silicon concentrations must be maintained within controlled ranges for proper production of the nodular form of graphite and control of the matrix structure. Excessive carbon content can result in carbon flotation in ductile iron castings. Carbon content of the iron is influenced by charge mix, carbon content of the metallic scrap, the coke-to-metal ratio, and slag basicity. Carbon pickup is increased by increasing the blast temperature and by increasing the oxygen injection percentage. The final carbon content can be altered by changing any one or a combination of blast control parameters, which affect the combustion process. Silicon is controlled in the cupola with the addition of ferrosilicon alloy and/or SiC briquettes. Maintaining the desired silicon content in the final iron requires accounting for the oxidation losses of silicon in the charge. Sulfur in gray iron is a naturally occurring element of melting in a cupola. It primarily comes from the sulfur contained in the coke. Sulfur is not detrimental below approximately 0.150 wt% if the manganese content is maintained at a sufficient ratio in relation to the sulfur content. Sulfur inhibits the conversion of carbon to ductile iron. Ductile irons are treated with a magnesium-containing alloy, which preferentially reacts with sulfur, forming inert MgS compounds. The small amount of magnesium remaining in solution in the iron promotes the nodular graphite shape. Thus, sulfur in ductile base iron should be below 0.030 wt% and preferably below 0.015 wt% for efficient conversion costs. Manganese is needed in gray iron at an amount of 1.7 times sulfur plus an additional 0.20 to 0.30 wt% to ensure that the sulfur is precipitated as manganese sulfide rather than iron sulfide. In pearlitic ductile iron, manganese is added as an alloy to produce a pearlitic matrix. In ferritic ductile iron and austempered ductile iron, manganese should be kept as low as possible to minimize pearlite in the as-cast structure and to minimize segregation, respectively. Phosphorus is controlled by using a variety of charge materials that are low in phosphorus. In gray iron, phosphorus contents should be below 0.10 wt% and even below 0.05 wt%, because high phosphorus levels are detrimental
to machinability and casting soundness in complex shapes. In ductile iron, phosphorus content below 0.05 wt% and preferably 0.03 wt% is desirable for maximum impact resistance and ductility. For low phosphorus levels, careful selection of charge material is required. Steel scrap is generally the lowest-phosphorus-content charge material. Residual Alloys and Tramp Elements. The effects of various inadvertent alloys and tramp elements should be understood, and control ranges and maximum levels should be established. Rapid analysis by modern spectrometers makes it possible to monitor any number of trace element levels. Any scrap source contributing excessive concentrations of detrimental elements should be removed from the cupola charge. Certain tramp elements, such as lead and bismuth, are detrimental to graphite shape (particularly the nodular graphite shape) and must be carefully monitored and maintained below the maximum level allowed. Titanium in trace amounts is beneficial to gray iron but is detrimental to the nodular iron, since it inhibits the formation of the nodular graphite shape. Inadvertent alloys, such as chromium, copper, tin, and molybdenum, if maintained within controlled ranges and combinations, can contribute to the pearlitic matrix structure of high-strength gray irons and high-strength pearlitic ductile irons. However, for ferritic ductile iron requiring the highest ductility, acceptable limits for these residual alloys are much lower.
Control Tests and Analyses The chill test was one of the earliest control tests developed and is still a good indicator of the graphitizing balance of an iron. A chill specimen is cast either as an elongated wedge in a sand mold or as a bar cast against a metal chiller on one side. After a few minutes, the test specimen is broken, and the depth of carbidic white iron chill indicates the balance between carbon, silicon, and other graphitizing elements against carbide-forming elements. This chill depth signals the need for more or less silicon, more or less carbon, and/or an adjustment in the postinoculation. Eutectic arrest methods were developed to provide carbon equivalent and carbon content as well as calculated silicon content. These thermal analysis techniques are rapid and make possible immediate corrective action, that is, adjustment of carbon and silicon. Optical spectrometer equipment greatly facilitates the control of iron chemistry, since all of the important elements can be determined within a few minutes. In large plants, samples are conveyed to the laboratory in pneumatic tubes, and analysis reports are conveyed to the melting center by phone or computer within a few minutes of taking the sample. Mechanical tests of several types are conducted to ensure the uniformity of iron properties and compliance with specifications.
Cupola Furnaces / 107 Tension tests of gray iron ensure that the iron meets the strength level required for the specified class of iron. In ductile iron, ultimate tensile strength, yield strength, and percent elongation are usually measured. Brinell hardness testing of castings monitors part-to-part and batch-tobatch uniformity and is an indicator of the general matrix structure. For ductile irons, microscopic evaluation of each treatment batch is essential to ensure graphite nodularity before castings reach the shakeout operation and lose their identity. Nodular graphite structure in a casting can also be verified with ultrasonic velocity test equipment. Ultrasonic velocity testing of all castings is performed by some companies to ensure the nodularity of safety-critical parts.
Specialized Cupolas Cokeless Cupola. In the 1970s, a cokeless cupola fired with gas or oil was developed in Europe and put into operation in several countries. It is a refractory-lined shaft furnace with a water-cooled grate supporting a bed of refractory spheres, as illustrated in Fig. 7. Heat for melting was generated by either gas or oil burners located below the grate and refractory bed. The iron melted in a cokeless cupola has low sulfur content and low carbon content, both of which are due to the absence of coke. The iron is carburized to a specified carbon content by injecting a fine carburizer (either petroleum coke or graphite) into the iron in the well. Pollution from the cokeless cupola is minimal and reportedly would meet environmental standards in some countries without expensive emission-control equipment. In countries where coke must be imported at a high price and oil or gas is readily available, the cokeless cupola is economically attractive. Production experience has been limited to small cupolas in the 4.5 to 9.1 Mg/h (5 to 10 ton/h) range. No practical experience has been reported on large cokeless cupolas. Attractive fuel (oil or gas) cost is available in some countries, but the high cost of refractory sphere replacements and the restrictions on the percentage of steel scrap used limit the attractiveness of this equipment in many markets. Plasma-Fired Cupola. In the 1980s, a pilot plant plasma-fired cupola was built, and a number of experimental heats were run, with some encouraging results. In the 762 mm (30 in.) inside diameter cupola, a single 2 MW plasma torch was projected into the cupola at tuyere level, as shown in Fig. 8. In the water-cooled torch, an electric arc from a high-amperage direct current is maintained across a gap between two copper sleeves, and the arc is rotated by a magnetic field. Air (or gas) is passed through the electric arc, where it is ionized and heated to very high temperatures unattainable with current hot blast systems. The experimental heats conducted indicated several advantages of using a plasma-fired cupola:
Coke (fuel) could be considerably reduced
due to the supplementing electric power from the torch. Lower-quality coke could be used; blast furnace coke, fine waste coke, and anthracite coal were melted successfully. Direct melting of borings and turnings proved practical because of lower air velocity and lower oxidation level. Silicon could be produced by injecting silica sand plus coke breeze. Iron could be produced by injecting as much as 20% iron oxide through the tuyeres.
The many possibilities of the plasma-fired cupola merited further exploration. Economically, the most attractive immediate application was the direct melting of iron borings or steel turnings. In March of 1989, the first commercial plasma cupola started production. This cupola was 3963 mm (156 in.) in diameter and had six conventional tuyeres. A 1.5 MW plasma torch was mounted to the end of each tuyere. Compressed air at 0.7 MPa (100 psig) was used as a plasma source. A booster compressor increased the air supply to 0.9 MPa (130 psig). The compressed air was passed through driers and filters and then to a distribution manifold. After being collected in the manifold, the air was dispersed into the six torches. The operating consumption of air was approximately 0.08 standard m3/s (160 standard ft3/min) per torch. Deionized water was necessary for the cooling of the torches. A conventional 700 C (1300 F) gas heat exchanger was used to preheat the blast air. The air at the tip of the plasma torch was approximately 1650 C (3000 F). The iron temperature at the top of the coke bed was approximately 1540 C (2800 F). The air blast volume was 6.0 to 7.0 m3/s (13,000 to 15,000 ft3/min) at the tuyeres. The development of the plasma cupola for melt operations was to offer melters the opportunity to use low-cost charge materials such as borings and steel turnings. This production plasma cupola was able to use an average of 40% loose borings/turnings in the charge makeup. The highest level achieved was 70%. This offered a significant cost-savings. The increased electric energy cost was partially offset by reduced coke usage.
Fig. 7
Cokeless cupola showing scrap charging, watercooled grates, burner location, and so on
ACKNOWLEDGMENT S.F. Carter, Melting Furnaces: Cupolas, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 383–392. SELECTED REFERENCES Conference on Cupola Operation, American
Foundrymen’s Society, June 1980
Cupola Design, Operation, and Control,
BCIRA, 1979
Cupola Handbook, American Foundrymen’s
Society, 1975 and 1984
Fig. 8
Early version of an experimental plasma cupola used to confirm the concept of plasma heating as a melting technology
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 108-115 DOI: 10.1361/asmhba0005196
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Induction Furnaces Revised by Ralph Y. Perkul, Foundry Solutions & Design LLC
INDUCTION FURNACES have become the most widely used means for melting iron and, increasingly, nonferrous alloys as well. The key to the ready acceptance of induction furnaces has been its excellent metallurgical control coupled with its relatively low pollution operation compared to cupola furnaces. Induction furnaces are very efficient and are made in many sizes. Induction furnaces are relatively simple, and small quantities can be melted quickly. The melting time is relatively short, allowing metal to be delivered at small, regular intervals. A wide range of metals can be melted, although little refining of the metal is possible. Currently, induction furnaces are available in a wide variety of sizes. Coreless units range in capacity from a few pounds, favored by the precision cast metal producers, to 68 metric tons (MT) (75 tons) powered at 23,000 kW. Channel-type units have been built with a capacity of over 2200 MT (2400 tons) powered by 6 inductors rated at 6000 kW each for holding applications. Induction furnaces require much cleaner scrap than cupola furnaces and somewhat cleaner scrap than electric arc furnaces. The capital costs are higher than those of electric arc furnaces, but the operating costs are lower due to reduced refractory wear. In the past, large, high-powered induction furnaces operated at 60 Hz, the frequency of the incoming power. Thanks to major breakthroughs in electronic, solid-state frequency conversion techniques, it is now possible to build a highly efficient furnace operating at medium and high frequencies (70 to 10,000 Hz). Coreless units of 20,000 kW in power for ferrous metals and 7000 kW for nonferrous metals have been built in medium frequencies, and even channel-type units have recently been equipped with solid-state, medium-frequency power supplies. The advantage of medium frequency is that the power density of the furnace system can be increased substantially without increasing its size. Much more melting power can be applied, while maintaining the stirring action at desirable levels. Heat losses, which are a function of furnace size, are reduced, and
overall efficiency of the system is improved. The furnace is also easier to operate because a single potentiometer will typically provide efficient and stepless control of the power.
Types of Furnaces The basic principle of induction furnaces is that a high voltage in the primary coil induces a low-voltage, high current across the metal charge, which acts as a secondary coil. Because of electrical resistance in the metal, this electrical energy is converted into heat, which melts the charge. Once the metal is in its molten state, the magnetic field produces a stirring motion. The power and frequency applied determine the stirring rate. This is controlled to ensure complete melting of the charge and adequate mixing of alloy and fluxing materials, and to minimize temperature gradients in the charge. Excessive stirring, on the other hand, can increase lining damage, increase oxidation of the alloys, generate excess slag, and increase inclusions and gas pickup. There are two classifications of induction furnaces: coreless and channel. Cross sections of each are shown in Fig. 1 and 2. In a coreless furnace, the refractory-lined crucible is completely surrounded by a water-cooled copper coil, while in the channel furnace the coil surrounds only a small appendage of the unit, called an inductor. The term channel refers to the channel that the molten metal forms as a loop within the inductor. It is this metal loop that forms the secondary of the electrical circuit, with the surrounding copper coil being the primary. In a coreless furnace, the entire metal content of the crucible is the secondary. Although both channel and coreless induction furnaces employ similar electrical principles, the two types of furnaces are quite different in their capabilities and operation. The coreless furnace is more widely used for melting and superheating, whereas a channel furnace is better suited for superheating, duplexing, and holding. The channel furnace has been used by nonferrous foundries for many years. Recently,
however, this type of furnace has been increasingly replaced by the coreless and resistance furnace in aluminum-melting applications. It is still commonly used for melting copper alloys. Technological innovations in the channel furnace have been slower in developing than advances in the coreless furnace. Today (2008), medium-frequency power supplies are being installed on some new channel furnaces. Many older furnaces still rely on 60 Hz power supplies and use power-regulating autotransformers for control. Inherent in the channel concept is the fact that only the small amount of metal in the inductor channel receives energy through the surrounding coil. This necessitates superheating the metal in the loop and pumping it out of the channel into the main vessel, where it mixes with the cold metal and heats it by conduction. The ratio of heat transfer is limited, as is the stirring action in the main vessel. This is the primary reason this type of furnace has seen only limited service as a melter. The power that can be connected to a vessel is much smaller than for a similarly sized coreless vessel. The mild stirring action also makes it difficult to get alloy and carbon additions dissolved and mixed rapidly, particularly important in the production of synthetic irons. In addition, the channel furnace must be started with a supply of molten metal, or a melt-out form and a metal heel must be continuously maintained in the furnace. The furnace must not be shut down even for holidays or for other extended periods, because the metal will freeze, and considerable refractory damage will be incurred. The channel furnace is more economical to manufacture in large sizes and uses electric power more efficiently than a coreless furnace. If more power is desired, additional inductors can be connected to the main vessel. Several geometric shapes of vessels have been used. These include configurations that are described as vertical (upright cylindrical), drum (horizontal cylindrical), and low profile (spherical). Typically, all of these units are hydraulically tilted to permit the emptying of the contents.
Induction Furnaces / 109
Fig. 1
Components of a coreless-type induction furnace. (a) Operational elements. (b) Cross section showing water-cooled copper induction coil and key structural components. The entire molten metal bath (which serves as the secondary) is surrounded by the coil (the primary) that encircles the working lining.
Because the magnetic field is transmitted in all directions, special vertical laminations of transformer iron, which form the magnetic yokes (shunts), are evenly spaced around the circumference of the coil to provide additional strength and to pick up the stray flux that would otherwise heat up the furnace frame. Coreless furnaces are typically tilted about the spout through a 95 to 100 angle to empty their contents. The medium-frequency units can be completely emptied for alloy changes or factory shutdowns and quickly started up again with a cold charge. It is not necessary to maintain a molten heel or to use precast starter blocks, as is the case with the 60 Hz units.
Electromagnetic Stirring Fig. 2
A cross section of a channel-type induction furnace showing the water-cooled copper induction coil, which is located inside of a 360 loop formed by the throat and channel portion of the molten metal vessel. It is the channel portion of the loop that serves as the secondary of the electrical circuit, in which the copper coil is the primary.
The coreless furnace is completely surrounded by a spiral of copper tubing of a special cross section that features a water-cooling channel in its center. It must also provide good electrical coupling and strength to withstand substantial electromagnetic forces, yet with provision for thermal expansion.
When alternating current is applied to an induction coil, it produces a magnetic field, which in turn generates a current flow through the charge material, heating and finally melting it. The amount of energy absorbed by the charge depends on the magnetic field intensity, electrical resistivity of the charge, and the operating frequency. A second magnetic field is created by the induced current in the charge. Because these two fields are always in opposite directions, they create a mechanical force that is perpendicular to the lines of flux and cause metal movement, or stirring, when the charge is liquified. The mechanical force stays perpendicular to the
field only in the center of the coil; on both ends of the coil it changes direction. The metal is pushed away from the coil, moves upward and downward, and flows back. Figure 3 shows the four-quadrant stirring. It is this stirring that allows excellent alloy and charge absorption and aids in producing a melt that is both chemically and thermally homogeneous. The stirring is directly determined by the amount of power induced and is inversely proportional to the square root of the frequency. Therefore, the higher the power and the lower the frequency, the more intense the stirring for a given size furnace. The reasons for an increased interest in induction melting for aluminum applications are many, but perhaps most significant is the question of metal loss. Because of the inherent electromagnetic stirring action in the coreless furnace, charge materials quickly become immersed in the bath, minimizing oxidation losses. This effect is particularly pronounced when melting light-gage scrap. Thus, the growing emphasis on recycling favors induction melting as a cost-effective alternative to other metal melting processes.
Power Supplies 60 Hz Line Frequency. In order to meet the range of melting needs, the proper power supply must be selected. Because the vast majority of existing furnaces use line frequency, this arrangement is discussed first (Fig. 4).
110 / Melting and Remelting
Fig. 3
A cross-sectional view of a coreless-type induction furnace illustrating four-quadrant stirring action, which aids in producing homogeneous melt
Fig. 4
Modifying 60 Hz line frequency supplied by utility company to serve as power supply for induction furnaces
Protection of the furnace and the power system, phase balancing, and power correction are required in order to operate an effective installation. The typical 60 Hz (line frequency) power supply basically consists of:
Primary switchgear Furnace transformer Mainline contactor Starting resistor and bypass contactor (for installations with off-load tap changers) Phase balancing system Power factor correction capacitor bank Instrumentation Controls and supervision
A primary transformation from the highvoltage grid of the local utility is usually
required because the voltage applied to the coil is between 500 and 3000 V. Because of its high inductance, an induction furnace coil has a low power factor. In order to bring the power factor to near unity, capacitors are installed and connected in parallel to the furnace coil. When starting a furnace, in order to reduce the inrush, a soft start is provided. When the transformer with an on-load tap changer is installed, the tap changer travels down to a lower tap whenever the furnace is switched off. This provides a start with a low voltage, resulting in low inrush current. When an off-load tap changer is applied, a starting resistor bank with bypass contactors is provided. Because the induction furnace is a singlephase furnace but its feed is commonly a three-phase line, a phase balancing system consisting of capacitors and reactors must be provided. In ferrous applications, refractory lining wear decreases the coupling distance between the coil and charge, resulting in a power increase. Because power input is limited, the voltage must be decreased to keep power from exceeding its maximum. This is typically done by using steps on a transformer tap changer. As the lining wears and the power increases, the voltage must be reduced by going to a lower tap. Additionally, in order to compensate for the amount of metal in the furnace and the condition of the refractory lining, a portion of the capacitor bank that is installed in parallel with the coil needs to be switched by adjusting the coarse and fine settings on the electrical panel. Medium Frequency. In the modern medium-frequency power supplies, (Fig. 5), many of the functions described are now performed automatically or have been eliminated altogether. The typical conversion is in an ac-dc-ac sequence. While frequency conversion entails electrical losses, modern power supplies incorporate solid-state frequency converters, and the increase in thermal efficiency usually outweighs the conversion losses. Modern converters exceed an efficiency of 97%. In the power conversion commonly used, the incoming 60 Hz ac power is rectified to dc and then chopped by an inverter to a higher frequency. Series and Parallel Inverters. Two types of inverters are typically used, namely, series and parallel. In each case, the furnace and compensating capacitor bank (fixed, in this case) form part of a resonance circuit. In the voltage series converter, the rectifier produces dc voltage, and the energy is stored in a filter capacitor. In a current parallel converter, the rectifier produces dc current, and the energy is stored in a smoothing reactor. The melting power for each of the two types of converters is also controlled differently. In the voltage series converter, the output is controlled entirely by variation of the inverter commutation frequency. The voltage is not controlled and is defined only by the line voltage.
In the current converter, the melting power is controlled by the variations of this inverter frequency as well as the magnitude of the dc voltage. The latter is regulated by variation of the conduction angle or phase control in the rectifier section. In both series and parallel circuits, it is possible to get power control that allows full power to be drawn regardless of the level of the charge in the furnace or the condition of the refractory lining. The midfrequency converters operate on a variable frequency. The load-commutated converter increases frequency as the charge contents of a furnace increase, while it simultaneously compensates for lining wear. No capacitor switching is required as in the line frequency units. In addition, a single potentiometer offers stepless and infinite control, eliminating the need for tap changer systems for power control. Modern induction furnaces take advantage of the latest technology, and many new installations are equipped with programmable controllers in lieu of relay logic. An even more recent advance has been the use of industrial computers for controlling the furnace and auxiliaries. With the variety of software that is currently available, the operator can choose screens such as systems data, automatic lining sintering, automatic cold starting of the furnace after a prolonged shutdown, automatic melting programs for various alloys, continuous calculated-temperature monitoring, fault detection, demand limiting, tabulation of energy consumption, melting cost data, metal inventory control, and chemistry adjustment. Also, by tying the computer to support systems such as charge make-up and hot-metal delivery, the utilization of the furnace can be increased. Packaging of Systems. Until the early 1970s, the electrical components of coreless induction furnaces were typically assembled on site in a concrete vault. Field construction costs made this installation method quite expensive. Modern component assembly techniques have allowed for the power supply to be packaged into steel modules at the equipment manufacturing plants. For furnaces up to 3.6 MT (4 tons) in capacity, current practice allows the manufacturers to completely assemble, wire, and pipe all components on a steel frame, to test, and subsequently to ship the unit as a single module, thus minimizing installation and start-up costs. For furnaces with power above approximately 1250 kW, a step-down isolation transformer from the main distribution voltage is normally used. Below 1250 kW, a 480 V distribution line backed by a disconnecting device can be used. The step-down transformers, when required, can be of the oil-filled outdoor type, which is the least expensive, or the indoor dry or silicon-filled type. Some manufacturers use the 480 V directly in their power supply; others choose to provide a step-up transformer and operate at the 575 V level.
Induction Furnaces / 111 In addition, with channel furnaces, standby generators are often employed to maintain cooling to the inductor.
Lining Material
Fig. 5
Typical medium-frequency induction power supply incorporating (a) a parallel inverter and (b) a series inverter
Water Cooling Systems Because the furnace coils and power supplies of the induction furnaces need to be cooled, a closed-loop water system is generally called for. In a medium-frequency power supply, the electronic components of the system contain a dc leg; it is therefore important to have clean, demineralized water circulating, typically with conductivity set at 50 mS (50 mmho) in order to prevent electrolysis from taking place. The heat exchanger to be used depends entirely on local circumstances such as climate, availability of raw water, and so on. The most widely used are the evaporative, the water-towater, and the air-to-water types. The latter has been frequently adapted to heat recovery systems. In those applications, the hot water from the furnace power supply and coil is piped to a location in the plant where heat is required. The water is circulated through a tube bundle, and a fan dispenses the warm air to the
surroundings. Considering that up to 10% of the power from a channel furnace and 20% of the power from a coreless furnace is lost to the water system, heat recovery systems for induction furnaces are often practical. In all coreless- and channel-type furnaces, the water is typically circulated through the power supply and then routed through the coil and on to the heat exchanger. Sometimes the water system is separated into two — one for the power supply and another for the coil. Usually, two pumps are provided, one as a standby with an automatic switchover, because it is extremely important to circulate water through the coil continuously to prevent it from being damaged. To protect induction furnace equipment during momentary power interruptions or prolonged power failures, manufacturers provide a special pressure-operated valve that, when activated, allows emergency deployment of city water to circulate through the coil. Such a system is shown in Fig. 6.
The selection of lining material is determined by the metallurgical requirements, the operating temperatures, and the type of operation. Refractories used in induction furnaces are oxides of minerals. Typically used are silica (SiO2), alumina (Al2O3), or magnesia (MgO) linings. From a chemical point of view, the silica is classified as an acid, alumina as a neutral, and magnesia as a basic material. Silica is the clear choice in iron melting, because it does not readily react with the acid slag typically produced in high-silica iron, has an extremely forgiving expansion curve, and is easy and economical to use. Alumina is the usual choice for aluminummelting furnaces, and steel-melting furnaces usually employ alumina or magnesia lining. Each of the refractory groups mentioned previously has a different thermal expansion characteristic, as shown in Fig. 7. As can be seen, alumina and magnesia have a nearly linear expansion. Silica, however, completes its expansion at approximately 815 C (1500 F) and remains constant at higher temperatures. With the selection of silica lining, the furnace can be shut down, cooled, and restarted without running the risk of metal penetrating thermal cracks. Magnesia is very sensitive to thermal shock and has the greatest expansion. Once cracked, this type of lining is not likely to seal itself, thus resulting in metal penetration. Alumina has approximately one-half the expansion of magnesia and is therefore less crack sensitive. It is important, however, to heat the furnace lining slowly when undergoing a cold start. This programmed start-up allows the cracks to seal before the metal adhering to the side walls melts and penetrates. This applies particularly to aluminum-melting furnaces because the metal melting temperature is relatively low. In the case of aluminum and some copper alloys, the oxide and metal buildup on the crucible not only poses a potential refractory penetration problem but also effectively increases the lining thickness and reduces the power that goes into the furnace. If not removed, this layer, which contains a small amount of metal, absorbs energy, and becomes superheated. The superheated metal can penetrate the lining and cause premature failure. For these reasons, the furnace crucible walls must be cleaned periodically. Crucible Wall Scrapers. The task of scraping the crucible wall has been made easier with the introduction of a mechanical scraper (Fig. 8), which cleans the sides in a full, hot furnace. The device essentially consists of a number of oscillating blades assembled on a steel frame. To start the operation, the scraper is placed on top of the furnace, and the
112 / Melting and Remelting
Fig. 6
A layout of a closed-loop cooling system showing routing of furnace-cooling water under both normal (black arrows) and emergency conditions (white arrows). Standby emergency pump is activated if the primary pump fails.
Temperature, ⬚F 32 2.0
392
752
1112
1472
1832
2192
2552 A
1.8 1.6
Linear thermal expansion, %
B 1.4
C D E F
1.2 1.0
Fig. 8
A cross section of a patented mechanical scraper used to clean the refractory lining of induction furnaces that melt nonferrous metals
G 0.8
H I K
0.6 0.4 0.2 0 0
100
200
300
400
500
600
700
800
900
1000 1100
1200
1300
1400
Temperature, ⬚C
Fig. 7
Thermal expansion curves of various refractory brick oxide materials used for linings in induction furnaces: A, magnesia; B, chrome magnesia; C, chromite; D, silica; E, zirconia; F, corundum 99; G, corundum 90; H, fireclay; I, sillimanite; J, zircon; K, silicon carbide
oscillating tools travel down the depth of the crucible. Radial pressure on the furnace is adjustable. The loose materials float to the top of the bath and are skimmed off. Ramming Mixes. Most induction furnaces operating in iron foundries are, as stated
previously, lined with a silica ramming mix. Alumina, magnesia, or zircon are typically used for melting steel and alloys. Alumina is typically used for aluminum. The current practice is to use granular materials. For example, the silica ramming mix is granular material combined with the sintering agent boron oxide (B2O3). The amount of boric
acid depends on the operating temperature and the wall thickness of the crucible. The flux or sintering agent fuses the silica particles to each other, forming a hard, glazed, ceramic surface. Ideally, the lining, when sintered, consists of one-third hard, sintered ceramic; one-third fritted material, in which fine particles are sintered and larger particles can still be identified; and one-third loose particles. This loose material has its purpose. When starting up a furnace from a cold start, the lining naturally expands. The loose material provides a cushion to handle the pressure. All granular materials are rammed on the furnace bottom using an electric vibrator equipped with a tamping plate. For the side walls, the lining material is spread behind a lining form that is vibrated with a pneumatic vibrator attached to it. For ferrous applications, the lining form is consumed in the first heat. For nonferrous applications, the lining form can be removed and reused. It may be tapered or collapsible.
Induction Furnaces / 113
Fig. 9
The principal components of a patented push-out device used to extract deteriorated refractory lining from coreless induction furnaces
The sintering process itself takes some time because it is necessary to heat the lining at the rate of approximately 110 C/h (200 F/h) until a temperature of about 30 C (50 F) above operating temperature is achieved. This temperature is held for 1 h before the furnace is ready for production. Lining Push-Out Device. When the linings are to be replaced, they can be chipped out with pneumatic hammers, and the lining procedure is repeated. A push-out device is available that eliminates the need to chip out the spent linings manually, saving both labor and time. The key components of this device, a schematic of which is provided in Fig. 9, are a locking plate that is bolted to the bottom of the furnace and a hydraulic ram assembly that is brought into position and latched to the plate. Once activated, the hydraulic ram moves forward, exerting pressure on a specially designed refractory cone installed in the bottom of the furnace. This distributes the force over the cross section of the lining. The push-out is aided by a slip joint between the coil grout and the working lining. The whole operation takes but a few minutes, with additional time being saved because the cool-down time is cut in half.
Melting Operations There are two distinct ways of operating a coreless induction furnace. One method is a batch operation, in which the entire contents of the furnace are emptied and the unit is
recharged with solids, typically using a conveyor or bucket charger. The other method involves a tap-and-charge operation, in which a portion of the furnace contents, typically one-third to one-half, is tapped and the identical weight in solids is recharged using a charge bucket or charge conveyor. Batch operation has now become feasible because medium-frequency power supplies allow the furnace to be started with ordinary scrap, rather than with starting blocks, which are required of 60 Hz (line frequency) furnaces. In terms of electrical consumption, the batchmelting method provides a 7% higher efficiency in iron-melting furnaces and a 4% higher efficiency in nonferrous-melting units over the tap-and-charge method. This is because the cold charge material couples with the magnetic flux more effectively than does molten metal. In addition to having greater electrical efficiency, a medium-frequency batch operation allows the drying of wet or oily charge material in the furnace, minimizing the need for a scrapdrying system. The tap-and-charge operation is still the most popular approach, however, since the chemistry is much easier to control. Less than one-half of the furnace contents are replenished at any one time, making alloying simpler because most of the furnace contents are already at the current specification. Finally, in order to use a batch operation effectively, the contents of the unit must be quickly and completely discharged, usually into a holding furnace or large ladles. This option may not be possible if small amounts of molten metal need
to be taken out at defined times and the capital and operational costs of a holding furnace cannot be economically justified. In either method of operation, it is particularly important to use the proper power supply efficiently and keep the cover closed to minimize potentially high radiation losses. Utilization is calculated in terms of the time that the power is on. It is obvious that power cannot be on constantly, because certain operations, such as charging, slagging, sampling, and tapping, are best and most safely done while the power is off. The objective is to keep the power on at least 75% of the time. Sometimes, a second furnace added to the power supply improves utilization, because the power can be switched from one furnace to the other while the above-mentioned operations take place. Utilization can also be improved by quickly charging and tapping to minimize waiting time and heat losses as well as by increasing the charge size. Today (2008), induction power supplies can operate two or more furnaces at the same time without the use of mechanical switches. These designs allow the system utilization to go to over 90%. Although medium frequency has made the need for heavy charge materials obsolete, it is still good practice to choose materials that have a reasonably high bulk density, are relatively clean and dry, and are not too long. In order to avoid bridging, the length of the pieces should be no more than two-thirds the diameter of the crucible. If longer pieces such as casting gating systems and structural steel sections are available, they should be broken or cut first. A further aid to utilization efficiency is the recent method of conveyor charging. Made possible by furnace covers acting as close-capture hoods, the furnace lid is kept up for much of the charging cycle. The charging is done continuously, adding more solid charge material as the contents of the furnace liquefy and collapse, allowing room for additional material. This adds safety to the often hazardous charging operation by removing the operator from the proximity of the furnace and allows control to take place from the melt shop control room. There are a variety of conveyor charging concepts. They are usually vibrating steel decks that can swivel between furnaces or extend-retract and move laterally between furnaces. Dry charge material is best suited for melting operations. Oils on the charge surface generate fumes from the furnace, which may need to be exhausted. Moisture on the charge also can be a problem, because a steam explosion can result. If dry materials are not readily available, a charge dryer is often employed to drive off any moisture remaining on the charge. The most common dryer consists of a refractorylined hood equipped with gas burners, which heat the solids on a vibrating conveyor by means of forced convection of the hot gases through the material. Figure 10 shows the various major components of the system. It is
114 / Melting and Remelting
Fig. 10
The major components of a dryer used to preheat induction furnace charge material. Integrated scrap preheat process combines (1) weigh hoppers, (2) preheat hood, (3) material transfer mechanism, and (4 and 5) furnace-charging apparatus into a single automated process.
Fig. 11
A mechanical clamshell-type power skimmer used to remove slag before each tap of an induction furnace
important to get the preheated charge into the furnace quickly to take advantage of the thermal energy in the material and to reduce the electrical energy output required for melting. The average temperature used for drying is typically 315 C (600 F); for preheating, it is 540 C (1000 F). When applied properly, the preheating of the charge material can increase the melt rate by 10 to 15%. Direct impingement
Fig. 12
Robotic attending of furnace. This robot can slag the furnace, take the bath temperature, as well as sample the metal.
Induction Furnaces / 115 of the flames on the charge material, particularly if it is thin, should be avoided because oxidation losses can occur. When dirty material is charged into the furnace, slag forms from the dirt and rust in the scrap as well as from the foundry return sand. Slag is undesirable in induction furnaces because it absorbs energy and erodes the lining. Therefore, it is removed each cycle with a skimmer to prevent slag accumulation. Mechanical Skimmers. Furnace operators can increase the power supply utilization by the use of mechanical skimmers (Fig. 11). This is particularly important in large furnaces because mechanical slagging units can minimize the time that power is either turned off or reduced for skimming during each charge-to-tap cycle.
A power skimmer can easily boost the output and productivity of the melting furnace. A skimmer is typically mounted on a jib. This tool has a long handle, which keeps the operator farther away from the hot furnace. The best and most effective type of slagger is the clamshell type. The skimmer is positioned over the furnace and lowered into the slag. With an automatic raking motion, the jaws of the unit close around the slag, which is then lifted clear of the bath and deposited into a container. The blades are typically dipped into a slurry to reduce slag adhesion during the next charge-to-tap cycle. Today (2008), most larger coreless units use back-slagging, where the furnace is back-tilted to allow the slag to be more easily removed.
A further development has been introduced recently that makes use of an industrial six-axis robot for the first time. In this innovation, the operator, from the safety of the melt deck control room, opens the cover and deploys the robot to skim the furnace, using preset programs (Fig. 12). As an added benefit, this system allows the operator to remotely obtain bath temperature, take samples, and add alloys. This further adds safety to the melting operation. ACKNOWLEDGMENT Updated from: R.Y. Perkul, Melting Furnaces: Induction Furnaces, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 368–374
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 116-123 DOI: 10.1361/asmhba0005200
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Vacuum Induction Melting MELTING UNDER VACUUM in an induction-heated crucible is a tried and tested process in the production of liquid metal. It has its origins in the middle of the 19th century, but the actual technical breakthrough occurred in the second half of the 20th century. Commercial vacuum induction melting (VIM) was developed in the early 1950s, having been stimulated by the need to produce superalloys containing reactive elements within an evacuated atmosphere. The process is relatively flexible, featuring the independent control of time, temperature, pressure, and mass transport through melt stirring. As such, VIM offers more control over alloy composition and homogeneity than other vacuum melting processes. Vacuum induction melting can be used to advantage in many applications, particularly in the case of the complex alloys employed in aerospace engineering. The following advantages have a decisive influence on the rapid increase of metal production by VIM: Flexibility due to small batch sizes Fast change of program for different types of
the melting process. Accordingly, the vacuummelted superalloys (compared to EAF/AODmelted alloys) are improved in fatigue and stress-rupture properties. Control of alloying elements also may be achieved to much tighter levels than in EAF/ AOD products. However, problems can arise in the case of alloying elements with high vapor pressures, such as manganese. Vacuum melting also is more costly than EAF/AOD melting. The EAF/AOD process allows compositional modification (reduction of carbon, titanium, sulfur, silicon, aluminum, etc.). In vacuum melting, the charge remains very close in composition to the nominal chemistry of the initial charge made to the furnace. Minor reductions in carbon content may occur, and most VIM operations now include a deliberate desulfurization step. However, the composition is substantially fixed by choice of the initial charge materials, and these materials are inevitably higher-priced than those that are used in arc-AOD.
The charge generally consists of three portions: A virgin portion, which consists of material
that has never been vacuum melted
A refractory portion, which consists of those
virgin elements that are strong oxide formers and have the tendency to increase the Crucible Power supply
Induction coil
To vacuum pumps
steels and alloys
Easy operation Low losses of alloying elements by oxidation Achievement of very close compositional
tolerances
Precise temperature control Low level of environmental pollution from
dust output
Removal of undesired trace elements with
high vapor pressures
Removal of dissolved gases, for example,
hydrogen and nitrogen Vacuum induction melting is indispensable in the manufacture of superalloys. Compared to air-melting processes such electric arc furnaces (EAF) with argon oxygen decarburization (AOD) converters, VIM of superalloys provides a considerable reduction in oxygen and nitrogen contents. Accordingly, with fewer oxides and nitrides formed, the microcleanliness of vacuum-melted superalloys is greatly improved compared to air (EAF/AOD)-melted superalloys. Additionally, high-vapor-pressure elements (specifically lead and bismuth) that may enter the scrap circuit during the manufacture of superalloy components are reduced during
Process Description A VIM furnace is simply a melting crucible inside a steel shell that is connected to a highspeed vacuum system (Fig. 1). The heart of the furnace is the crucible (Fig. 2) with heating and cooling coils and refractory lining. Heating is done by electric current that passes through a set of induction coils. The coils are made from copper tubing that is cooled by water flowing through the tubing. The passage of current through the coils creates a magnetic field that induces a current in the charge inside the refractory. When the heating of the charge material is sufficient that the charge has become all molten, these magnetic fields cause stirring of the liquid charge. The optimal induction coil frequency for heating the charge varies with the charge shape, size, and melt status (liquid or solid). Older equipment used a single frequency, but newer power supplies are able to be operated at variable frequencies and are adjusted throughout the melt to obtain the most rapid heating/melting conditions.
Fig. 1
Basic elements of a vacuum induction melting furnace
Shunt Heating coil
Brick crucible
Cooling coil Ground detection
Fig. 2
Schematic of vacuum induction melting crucible (shell, coil stack, backup lining, and working lining)
Vacuum Induction Melting / 117 Bulk charger Crucible (pouring)
Melting chamber
Cover (removed) Launder Tundish car
Operating platform
Mold chamber
Electric room Power supply
Molds
Shop floor
Mold car
Fig. 3
Vacuum system
Schematic of a top-opening, double-chamber vacuum induction melting furnace
Table 1
Typical refractories used to line vacuum induction melting crucibles Maximum melt temperature
Refractory
MgO Al2O3 MgO-spinel Al2O3-spinel ZrO2 Graphite
C
1600 1900 1900 1900 2300 2300
F
2910 3450 3450 3450 4170 4170
Refractory density g/cm3
lb/in.3
Resistance to thermal shock
2.8 3.7 3.8 3.7 5.4 1.5
0.101 0.134 0.138 0.134 0.195 0.054
Good Good Poor Relatively good Poor Excellent
solubility of oxides and nitrides in the virgin charge A revert (or scrap) portion, which consists of both internal and external scrap that previously has been vacuum melted Vacuum-melted scrap has already had its gas content reduced to levels consistent with vacuum production. Scrap, however, has the possibility of having become contaminated during the production process, and care (expense) must be taken in the segregation and preparation of scrap materials for vacuum melting. In most VIM furnaces there is a vacuum lock bulk charger located directly over the crucible (Fig. 3). Charge material may be added to the heat through the bulk charger while melting is in process in the crucible. The material to be added is placed in bottom-opening buckets, placed in the bulk charger, and the charger is evacuated. The valve isolating the charger from the melt chamber is opened, and the bucket is lowered to a point close to the crucible top and the bottom opened so as to drop the charge material into the crucible. In constant operation, if a furnace is not to be opened to the atmosphere, all charge material for a heat will be added by this process.
Applications
Superalloys, high-quality Superalloys, high-quality Superalloys, high-quality Superalloys, high-quality Superalloys, high-quality Copper, copper alloys
steels steels steels steels steels
Older VIM furnaces may have been designed as single-chamber systems with the mold put inside the furnace before the beginning of the melt. The molten charge is then poured into the mold inside the furnace. Single-chamber furnaces thus must be opened after each heat to extract the molds and put in the new molds. Most furnaces have some system of large vacuum locks for transferring the prepared molds into the melt chamber. In double-chamber furnaces (Fig. 3), there is a separate chamber for the molds. The molten metal is transferred via launders (refractory-lined steel troughs) to a refractory tub (tundish). Some systems are designed to pour directly from the crucible into the tundish. The tundish contains a considerable volume of metal and allows residence time for entrained slag to float to the top of the tundish and be removed from the pour stream. The pour stream exits the bottom of the tundish. The pour time is regulated by pour temperature and the nozzle diameter of the tundish. The typical tundish is designed to provide a low-velocity flow path from the point of entry of the metal to the bottom nozzle. Refractories. Crucibles with a capacity of approximately 4500 to 22,500 kg (10,000 to 50,000 lb) are generally built up from refractory brick. Smaller furnaces, used for production of
master melt, may use single-piece crucibles. Refractory brick linings are usually two layers. The backup lining protects the induction coil in the event of a failure of the outer or working lining. The working lining is the primary interface with the metal and is replaced when erosion of the lining becomes excessive. Refractory life is also affected by the expansion of the refractory during the repeated melting cycles. Refractory brick is chosen with regard to resistance to erosion and expansion. Commercially available refractory brick may be incompletely sintered and expands during use, causing loss of crucible integrity. The refractory material used for the crucible lining is based on oxides such as Al2O3, MgO, CaO, or ZrO2 (Table 1). The lining is almost always rammed and sintered; prefabricated brick is used in larger furnaces. Dried silicate, combined with small oxide additions, appears to be very suitable for crucible lining because of its thermal characteristics. Because of an irreversible thermal expansion of 8% above 1000 C (1830 F), a high densification of the lining takes place during sintering. For this reason, this active lining is suitable for foundries. The behavior of the lining refractory with regard to stability at high temperature under vacuum must also be considered. The melting crucible material is not inert and is actually another source of oxygen and other impurities, depending on refractory type and condition. Therefore, both melt refining temperature and refining duration are carefully scrutinized. Proper melt stirring is integral to the deoxidation process and must be optimized through proper furnace power frequency and application procedure to prevent refractory lining erosion, a potential problem particularly during the controlled but more vigorous CO boiling portion of the process. Process Sequence. Figure 4 shows a typical process profile for the VIM of nickel- and cobalt-base superalloys. Before operation or at the completion the preceding heat, the melt chamber is isolated from the mold chamber and the vacuum integrity of the furnace is verified. Specific practices with regard to vacuum measurements will differ in detail. The deterioration rate of the vacuum is a measurement of the inherent vacuum integrity of the furnace. This integrity cannot be measured by vacuum alone, because the large pumping capacities of the pumps used in VIM are able to achieve low vacuum pressures even with significant leakage into the furnace. It is considered bad practice to continuously draw air into the vacuum furnace and across the melt. Immediately after ensuring that the furnace is vacuum-tight, the virgin portion of the charge is placed into the VIM furnace first. This may be done by opening the furnace or, more commonly, by charging the furnace through hoppers lowered through a large vacuum lock (bulk charger) located over the crucible. The furnace is capable of quickly pumping down to or maintaining vacuum levels below 100 mm (and often into the
118 / Melting and Remelting 105
1000
Installed power 1200
104
100 Pyrometer temperature
900
1200 1100
10
1.0
1000 900 800
2732 10 Melting power Pressure rise test
800
1.0
700 Casting end Casting start
600 500
All liquid
400
Additions/Sampling Pressure rise test
300 0.1
200 100
0.01
0
0.1
2552
2192 2012 1832 1652
0.01
1472
10−3
392
Additions Sampling
200
212
100 0
2372
Pyrometer temperature, ⬚F
1300 Melting power, kW
100
1400
1000
Pyrometer temperature, ⬚C
Melt chamber pressure, Pa
103
1100
Melt chamber pressure, mbar
1500
1. Charge 0
3. Charge
2. Charge 1
Melting
2
3
4 Refining
5
6 Superheating
10−4
7
32
Casting
Time, h
Fig. 4
Typical vacuum induction melting protocol for nickel- and cobalt-base superalloys
<10 mm) range. The virgin material is melted by application of current to the induction coils surrounding the refractory crucible. When the virgin material has been completely melted (all molten), it outgasses. The outgassing is monitored until it is complete. The outgassing is a response of the gas in solution in the melt to the low partial pressures of gases in the vacuum chamber. Some degassing is accomplished because carbon in the charge will form CO with the oxygen and also be evolved from the melt. The progress of degassing is followed by measuring the leak-up rate at set time intervals. When a constant leak-up rate is obtained, this indicates that degassing (refining) of the melt is complete. After degassing is complete, the reactive and revert charges are added to bring the charge to its planned weight. A ladle sample is taken after all additions are complete. Based on the analysis of this sample, trim additions are made to bring the melt into a very precise compositional range. Because there are no ongoing chemical changes in the melting, as there are in EAF/AOD, the compositional requirements of a melt may be met as closely as allowed by the reproducibility of chemical analysis. After the trim additions have been made, the temperature of the heat is brought precisely to the desired point, and the heat is poured. The heat, although produced in vacuum, will still have generated significant amounts of slag from the products of deoxidation, desulfurization, and the deterioration of the refractory crucible lining.
Process Controls. Process computerization and automation continue to enhance better reproducibility of the melts and computer modeling. Using charge and alloy calculations, it is easy to achieve the required chemical analysis with minimal cost. For this application, in the ideal case, the computer is linked to the analysis system computer so that the additives can be calculated and, if necessary, weighed and added directly after the analysis. If this alloy calculation is laid out as a charge calculation, the charges and the amounts of scrap necessary can be optimized. The complete calculation system enables a calculation of the alloying elements, starting with the amount of scrap needed and proceeding to the necessary alloy addition at the end of the refining period. A computer-controlled mass spectrometer system, specially developed for the VIM process, can make a significant contribution to the optimization of the melting process and its economics. Using a computer evaluation of the gas composition in the furnace chamber, information about the state of the degassing process and of the chemical reactions in the melt can be continuously obtained. The melting process can thus be controlled more exactly. Process parameters that can be closely controlled include: Determination of the correct time for the
addition of chemically active elements and the sequence of addition of such elements
Determination of the process steps, such as
completion of the refining period and the suitable time for tapping Advance detection of leaks in the furnace chamber, cooling water system, or hydraulic system Monitoring of the state of degassing of the crucible refractory lining Remelting by VIM for Shape Casting. Vacuum furnaces for precision casting are used to remelt alloys that have already been treated in vacuum. The remelted alloy is then precision cast into preheated ceramic molds. Heating and melting in such furnaces almost always takes place by induction, although some electron beam melters and vacuum arc furnaces are used. Electron beam furnaces offer the advantage of ceramic-free melting, but they do not permit a sufficiently accurate and reproducible temperature control for charges of more than approximately 5 to 7 kg (11 to 15 lb). Vacuum arc furnaces are used for titanium precision casting. Remelting with VIM furnaces for precision casting is done with two-chamber furnaces (Fig. 5a), where the melting crucible and mold are in separate chambers. The mold chamber is preferably placed under the melting chamber so that, in connection with another chamber for the charge or melting material, semicontinuous operation with minimal cycle times is possible. For cobalt- or nickel-base superalloys, the melting chamber pressure is of the order of 0.01 Pa
Vacuum Induction Melting / 119 (104 mbar, or 7.5 105 torr), while the mold lock would be evacuated to approximately 5 Pa (0.05 mbar, or 0.038 torr) within 1 min. In addition to the casting of equiaxed, directional solidification, and single-crystal blades for gas turbine engines, vacuum precision casting can be used for large parts (melt weights up to 1000 kg, or 2200 lb) or for the large-scale production of small parts. Examples of large parts produced using vacuum precision casting include compressor sections of jet engines and structural parts for stationary turbines.
A typical small casting produced by this method is a turbocharger wheel for automotive engines (Fig. 5b). Automotive turbocharger wheels weighing 400 to 500 g (0.9 to 1.1 lb) are produced in automatic precision casting machines. Total cycle time, including 40 s for melting, is typically 90 to 100 s. In the production of such small parts, the crucible is often integrated into the mold. Casting yields of 80 to 90% are obtained. Filters. One of the most critical stages with respect to cleanliness is the pouring of the melt.
Ceramic foam filters are used in some master metal operations to remove relatively large melt inclusions by means of entrapment. Foam filters are most effective where extremely high pour rate conditions and gross cleanliness problems prevail. Filter performance often varies because of the occasional use of filters with poor mechanical strength and/or thermal shock resistance. Foam cell particle breakage often results from handling during shipment or tundish installation and, if undetected, results in filter particulate in the alloy bar stock and subsequently cast components. Optimized VIM technology and practice without filters provide a clean alloy, without the inherent risks associated with filter use when applied to master metal production.
VIM Metallurgy
(a)
Fig. 5
(b)
Shape casting with vacuum induction melting, (a) Computer-controlled vacuum furnace with mold chamber. (b) Precision-cast turbocharger wheels for automotive engines. From left: mold with integrated crucible, bar stick, cast part, machined turbocharger wheel
Fig. 6
Vacuum induction melting is often done as the primary melting operation followed by secondary melting (remelting) operations such as electroslag remelting (ESR) and/or vacuum arc remelting (VAR).Various processing routes following VIM are illustrated in Fig. 6. Some superalloys are produced by a triple-melt sequence (VIM/ESR/VAR), where the VIM ingot is typically referred to as an electrode for subsequent ESR and/or VAR operation. The casting weight can vary from a few kilograms to 30 Mg (33 tons), depending on whether the VIM furnace is being used for
Potential processing routes for products cast from vacuum induction melting (VIM) ingots or electrodes. VAR, vacuum are remelting; ESR, electroslag remelting; EB, electron beam; HIP, hot isostatic pressing. Source: Ref 1
120 / Melting and Remelting precision casting or for the production of ingots or electrodes for further remelting. There are many different metallurgical factors that arise with induction furnace melting. The crucible material has an extraordinary effect on the metal/slag reaction because the ceramic outer wall reacts with the liquid metal and with the slag. In a VIM furnace, slag would be transported to the crucible wall by the characteristic bath movement. The result is that the slag solidifies at the wall and therefore has an insufficient reaction with the melt. It is therefore beneficial to use a slag that is more or less saturated with the oxides of the crucible lining in order to minimize heavy slag attack of the lining. In contrast to the ladle metallurgy processes with EAF, the crucible wall lining is susceptible to significantly higher erosion than the brick wall of a ladle. For this reason, metallurgical operations such as dephosphorization and desulfurization are limited. The metallurgy of VIM is primarily limited to the pressure-dependent reactions, such as carbon, oxygen, nitrogen, and hydrogen, and the evaporation of undesired elements with high vapor pressures, such as copper, lead, bismuth, tellurium, and antimony. Cleanliness can be significantly improved if a reactive liquid slag capable of absorbing oxides and sulfides is in contact with the melt. Vacuum induction furnaces are generally not operated with active slags. Therefore, reaction products can precipitate only on the crucible walls, and the melt may not be as clean as with other processing methods. Because the various alloys produced in vacuum induction furnaces must meet the highest quality requirements and because the vacuum induction furnace is primarily a consolidation unit and only secondarily a refining unit, the following methods are used to produce clean melts: Selection of a more stable refractory mate-
rial for the crucible lining
Rinsing of the melt with inert gas Minimizing the contact time of the melt in
the crucible
Exact temperature control to minimize cru-
cible reactions with the melt
Fig. 7
Vacuum induction refining process
Suitable deslagging and filtering techniques
during pouring Conception of a suitable tundish and launder system for good oxide removal For particular applications, however, the quality of the material produced by VIM will not be sufficient to satisfy the highest quality requirements with respect to cleanliness and primary structure. In this case, the VIM-produced material must undergo secondary melting processes (Fig. 6). VIM Refinement Process. The primary process of melt refinement is the removal of meltcontained oxygen by means of a reaction with carbon to form carbon monoxide (CO). The reaction occurs most readily at or near the melt surface, with the reaction kinetics being affected by crucible geometry and melt stirring. The CO bubbles form along the walls and, sometimes, bottom of the melt/lining-refractory interface (Fig. 7). This occurs preferentially at small crevices existing in the lining, with the bubbles growing during movement toward the molten metal/vacuum interface (Ref 2). Actual bubble formation is dependent on the number of gas molecules present, the pressure in the liquid at the level of the bubble, the temperature of the gas, and, for very small bubbles, the interfacial tension between the gas and the liquid metal. Following formation, bubble growth and mass transport within the liquid toward the liquid/vacuum interface is dependent on: The quantity of the dissolved gas The decreased pressure exerted on the bub-
ble as it rises in the melt
The bath temperature The time it takes for the bubble to rise
through the melt to the surface, which, in turn, is a function of melt stirring The pressure above the melt The interfacial tension between the bubble and the liquid metal The relatively vigorous, but controlled, portion of the boiling process results in the greatest CO removal. Concurrently, a slight nitrogen
loss is realized because of scavenging associated with the CO bubbles, and a slight sulfur reduction may occur during the CO supersaturation stage via sulfur dioxide (SO2) evolution. Minor tramp elements such as lead, silver, bismuth, selenium, and tellurium are partially evaporated during this period as well as throughout the entire refining process (Ref 3). Some undesirable elements, however, such as arsenic and tin, must be controlled through raw materials selection because they are not removed by vacuum refining. Bath refining from a somewhat vigorous CO boil is undertaken at a temperature and duration long enough to reach the so-called system equilibrium conditions, the assurance of which is provided by the attainment of consistent furnace leak-up rates. At this point, there is the addition of the refractory charge, which includes elements with reactivity toward oxygen (for example, aluminum, titanium, zirconium, and hafnium) that were withheld from the virgin charge because of their reactivity. Revert charge is also added. Once the boiling subsides, surface desorption of additional CO occurs, and it is during this nonboiling period that nitrogen removal (desorption) is most effective (Ref 4). Trace Element Removal. The removal of undesired volatile trace elements, such as arsenic, antimony, tellurium, selenium, bismuth, and copper, from the vacuum induction furnace is of considerable practical importance. These elements must be held to very low concentrations to avoid the risk of premature part failure, particularly for the production of superalloys for critical applications, such as jet engine parts. Figure 8 shows the influence of some trace elements on the stress-rupture properties of alloy 718. Because of the high vapor pressures of most of the undesirable trace elements, they can be kept to very low levels by distillation during melting under vacuum. Figure 9 shows how the different trace elements behave under vacuum. In a nickel-chromium melt, arsenic, antimony, and tin cannot be reduced over the gas phase, while copper, lead, selenium, tellurium, and bismuth can be reduced to a level far below 10 ppm. Nitrogen Degassing. Nitrogen gas can be decreased because its solubility at constant temperature is directly proportional to its partial pressures. As long as no strong nitride-forming elements are present in the melt, there is no difficulty in reducing the nitrogen content to 20 ppm (Ref 7–9). When nitride-forming elements such as chromium, vanadium, aluminum, and titanium are present, the activity of the nitrogen is very low; therefore, removal of nitrogen under high vacuum is difficult. If low nitrogen contents are desired in alloys containing nitride-forming elements, raw materials low in nitrogen must be used. Figure 10 shows the reduction in nitrogen levels obtained in the VIM of a die steel. From an average level of approximately 400 ppm at the beginning of treatment, the nitrogen has
Vacuum Induction Melting / 121 200
99.9 99.5 99
Silver Concentration, %
Rupture life, h
Selenium 120 Lead 80
Frequency, %
160
97.5 95 90
10
40
Fig. 0
10
20
30
50
40
60
70
80
Effect of trace elements on the stress-rupture properties of alloy 718. Test conditions: 650 C (1200 F), 690 MPa (100 ksi). Source: Ref 5
(0.5)
99.9 Arsenic Tin Frequency, %
0.1
Concentration, %
Antimony
Selenium 0.01
99.5 99 97.5 95
After treatment With argon bubbling
90
Without argon bubbling
80 70 60 50 40 30
Before treatment
20
Copper
10 0
100
200
300
400
500
Nitrogen content, ppm Bismuth
Lead
10−3 Tellurium (3⫻10−3)
0
20
40
60
80
100
120
Time, min
Fig. 9
Evaporation of trace elements from Ni-20Cr melts under vacuum. Source: Ref 6
been reduced to a level of 50 ppm. Additional argon purging during melting of the steel did not improve nitrogen removal. In the case of iron-nickel melts, however, the additional argon purging reduces nitrogen to much lower values. Hydrogen Degassing. Like nitrogen, hydrogen can be decreased via the gas phase because its solubility is directly proportional to its partial pressure. The solubility of hydrogen at atmospheric pressure and at 1600 C (2910 F) in iron or nickel melts is in the range of 30 ppm. During VIM, the hydrogen can be removed to very low concentrations. Figure 11
Before treatment
0
1.0
2.0 3.0 Hydrogen, ppm
4.0
5.0
11
Reduction in hydrogen content of X38CrMoV51 die steel after vacuum induction degassing
90
Impurity content, ppm
Fig. 8
80 70 60 50 40 30 20
Bismuth
0
After treatment
Fig. 10
Reduction in nitrogen content of X38CrMoV51 (Fe-0.38C-5.2Cr-1.3Mo-0.4V-1Si-0.4Mn) die steel after vacuum induction melting processing
shows the hydrogen contents of a die steel before and after vacuum treatment. It is evident from Fig. 11 that final hydrogen contents below 1 ppm can be routinely achieved. The reduction in hydrogen during the degassing treatment amounts to approximately 80%. Argon purging had no perceptible influence. The bath agitation in the case of this furnace, which is operating at normal frequency, is sufficient for hydrogen removal. Similar results were also obtained for the VIM processing of superalloys. Deoxidation of the melt in the vacuum induction furnace can be done via the generation of CO gas (C + O ! CO). The removal of oxygen from the melt as CO is favored by decreased melt chamber pressure, elevated bath temperature, and increased carbon activity (Ref 10). Proper melt stirring is integral to the deoxidation process and must be optimized through proper furnace power frequency and application procedure to
prevent refractory lining erosion, a potential problem particularly during the controlled but more vigorous CO boiling portion of the process. The final attainable oxygen content for a given carbon content is directly proportional to the CO partial pressure. The theoretical equilibrium values for oxygen are below 1 ppm in carbon-containing iron or nickel melts at 1600 C (2910 F) and at a pressure of 0.1 Pa (103 mbar, or 7.5 104 torr). However, the actual values are up to 1 order of magnitude higher. These higher oxygen contents result from impurities in the crucible lining, crucible outgassing and leakage, and, above all, the fact that the reaction does not reach equilibrium with decreasing oxygen content because of the difficulty of CO nucleation. In addition, the hydrostatic pressure of the liquid metal in the melt must be taken into account because only a relatively small bath surface area is in direct contact with the prevailing vacuum pressure. However, the influence of the liquid-metal hydrostatic pressure can be minimized through the use of additional agitation effects. The CO reaction occurs in two stages (Ref 1). The first stage is boiling, that is, the formation of CO bubbles within the melt along with a strong bath agitation as a result of this gas formation. The second stage is desorption, in which no more CO bubbles form inside the melt, and CO formation takes place only at the bath surface. Deoxidation with carbon is also a decarburization reaction, which is used for the production of low-carbon high-chromium steels. By using the CO reaction under vacuum, very low carbon contents can be achieved without a noticeable loss of chromium. The oxygen content in such alloys is higher because of the decrease in oxygen activity due to chromium. In this case, precipitation deoxidation using aluminum, silicon, or titanium must be applied. The CO partial pressures at which the lining oxides are reduced by carbon in the melt are:
122 / Melting and Remelting Carbon monoxide partial pressure Lining material
CaO ZrO2 MgO Al2O3 SiO2
kPa
0.04 0.13 0.53 0.53 81.1
torr
0.3 1.0 4.0 4.0 610.0
Source: Ref 1
It is evident that CaO represents the stable crucible material, while SiO2 is reduced at relatively high pressure. Apart from that, SiO2 can be reduced by alloying elements such as manganese, chromium, aluminum, titanium, or zirconium. This can cause a heavy chemical erosion of the lining, accompanied by an undesired silicon pickup in the melt. Because of these undesired reactions with a silica lining, a spinel-forming basic refractory material is used for the melting of high-grade steels and superalloys. The spinel formation during sintering leads to volume growth; therefore, as with the oxidic material, a densification of the rammed lining takes place. The vacuum induction degassing and pouring (VIDP) furnace (Fig. 12) is a variation of the conventional VIM furnace. The VIDP furnace design employs a modular concept that allows for connecting it to casting units for ingot casting, horizontal and vertical continuous casting, or powder production. Because of the smaller volume of the VIDP furnace compared to the VIM furnace and significantly lower desorption and leakage rates, it is possible to obtain very low pressures with lower pumping capacity. The lower part of the furnace can be decoupled and replaced rapidly; therefore, the VIDP furnace enables faster replacement of different furnace vessels with changes of alloy.
Melting Bath Agitation. As noted, melt agitation is an important process factor. Agitation is caused by the induction coil itself, depending on the power input and the installed frequency. Stirring is directly proportional to the amount of induced power and inversely proportional to the square root of the frequency. Therefore, a more intense stirring for a given size furnace occurs at a higher power and a lower frequency. Bath agitation from the induction coil is basically limited to the meltdown period. It can be done with the induction coil alone only if there is a high-power input and therefore strong bath heating. This high temperature increase is not always metallurgically desirable. For this reason, two agitation systems are offered in stateof-the-art vacuum induction furnaces:
vacuum. This ensures that the admitted purging gas flows through the melt and does not take the path of least resistance through the lining of the crucible. Both agitation processes offer advantages with regard to higher reaction rates, correct temperature adjustment, homogenization with simultaneously lower erosion of the crucible lining, and better cleaning of the melt. Figure 14 shows the three different agitation methods operating in a typical VIM process sequence. The agitation effect at 200 to 500 Hz in the melting phase assists in shortening the melting time. In the refining and superheating periods, it is better to use electromagnetic agitation at 50 to 60 Hz or argon purging through a porous plug.
Electromagnetic agitation with an additional
coil
Production of Nonferrous Materials
tom of the crucible
Apart from melting high-grade steels and superalloys, VIM is being increasingly used for the production of nonferrous metals and alloys. Table 2 shows some examples for possible use in nonferrous metallurgy. Aluminum alloys with additives such as zirconium, titanium, beryllium, cerium, tellurium, and cadmium must be melted under vacuum or under inert gas atmosphere because of their high reactivity with air and, in some cases, their toxicity. Aluminum-lithium alloys are also candidates for VIM processing. Copper Alloys. The production of highpurity copper having less than 2 ppm O can be accomplished only in a vacuum induction furnace. Oxygen content influences the electrical conductivity of copper alloys; the lower the oxygen content, the higher the electrical conductivity (Ref 11). For the production of
Agitation by argon purging through the bot-
The difference between the electromagnetic agitation by a separate coil and the agitation caused by the electromagnetic forces of the main coil lies in the fact that the bath can be more vigorously stirred without a temperature increase. A second transformer would be needed. Another method of bath agitation is derived from the principle of argon purging usually applied in ladle metallurgy. Argon purging has a long history of use in atmospheric melting furnaces. A characteristic of this purging is that the porous plug is not in direct contact with the liquid melt. Instead, the plug is covered with refractory material identical to that used for the crucible lining (Fig. 13). The basic difference is that the porous plug is in direct contact with the melt and is suitable for operating under
Crucible lining
Exchangeable porous plug
Basic set
Fig. 12
Schematic of the vacuum induction degassing and pouring furnace
Fig. 13
Argon purging system for vacuum induction melting furnaces
Vacuum Induction Melting / 123 oxygen-free copper, melting and casting must be carried out under vacuum. Selective Evaporation of Alloying Elements. The use of vacuum metallurgy is primarily linked with degassing and decarburization. A side effect of these treatments is the evaporation of
elements with high vapor pressures. In nonferrous metallurgy, this effect is used for the distillation of metals—for example, for the separation of lead and zinc in lead refining, in zinc production, and for the reduction of magnesium and nonalkali metals. Similarly, copper can be refined
from copper scrap by using the vacuum to evaporate volatile elements such as lead and zinc (Ref 12). ACKNOWLEDGMENTS Portions of this article were adapted from: A. Choudhury and H. Kemmer, Vacuum
Melting and Remelting Processes, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 393–425 M. Donachie and S. Donachie, Superalloys: A Technical Guide, ASM International, 2002 G.L. Erickson, Polycrystalline Cast Superalloys, Properties and Selection: Irons, Steels, and High-Performance Alloys, Vol 1, ASM Handbook, ASM International, 1990 REFERENCES
Fig. 14
(a) Melting and (b) stirring modes of the vacuum induction melting process
Table 2 Possible applications for vacuum induction melting in the processing of nonferrous metals and alloys Metal or alloy
Metallurgical results
Aluminum, aluminum alloys High-purity metals
Low oxygen content, effective removal of oxides Oxygen-free copper (2 ppm O)
“New” alloys
Al-Mg, Al-Ce, Al-Zr, Al-Hf alloys; copper alloys with beryllium, cobalt, titanium, zirconium, and lithium; avoidance of lithium losses in production of Al-Li alloys Production of calcium, barium, lithium, strontium, and magnesium; recycling of hard metals; purification of alloys or melts; removal of zinc and lead
Metal production and recycling
Achieved by
Vacuum tank degassing without additions of polluting gases or materials Melting and pouring under steady, controlled atmosphere or vacuum; gas stirring; carbon as reducing agent Furnace and casting system under inert gas atmosphere; vacuum-tight, highly automated furnace
Thermic/alumino-thermic vacuum distillation; leaching with zinc and subsequent vacuum distillation melting; vacuum distillation
1. J.W. Pridgeon et al., in Superalloys Source Book, American Society for Metals, 1984, p 201–217 2. D. Winkler, Thermodynamics and Kinetics in Vacuum Metallurgy, Vacuum Metallurgy, O. Winkler and R. Bakish, Ed., Elsevier, 1971, p 42–54 3. Evaporation of Elements from 80/20 Nickel-Chromium During Vacuum Induction Melting, Transactions of the Vacuum Metallurgy Conference, American Vacuum Society, 1963 4. V.M. Antipov, Refining of High-Temperature Nickel Alloy in Vacuum Induction Furnaces, Stal0 , Feb 1968, p 117–120 5. W.B. Kent, Int. Voc. Sci. Technol., Vol 11 (No. 6), 1974, p 1038–1046 6. P.P. Turillon, in Transactions of the Sixth International Vacuum Metallurgy Conference (Boston), American Vacuum Society, 1983, p 88 7. G.A. Simkovich, Int. Met., Vol 253 (No. 4), 1966, p 504–512 8. H. Katayam et al., in Proceedings of the Seventh International Conference on Vacuum Metallurgy (Tokyo), The Iron and Steel Institute of Japan, 1982, p 933–940 9. A. Choudhury et al., World Steel and Metalworking Manual, Vol 9, 1987–1988, p 1–6 10. D.R. Gaskell, Introduction to Metallurgical Thermodynamics, Scripta Publishing, 1973, p 268–273 11. O. Kamado et al., Method of Producing Electrical Conductor, European Patent 0121152, 1986 12. J.G. Kru¨ger, Proceedings of the Fifth International Vacuum Metallurgy Conference (Munich), 1976, p 75–80
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 124-131 DOI: 10.1361/asmhba0005201
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Electroslag Remelting ELECTROSLAG REMELTING (ESR) is a secondary refining process (Fig. 1), where a consumable electrode is melted by immersion in a molten flux (slag) consisting primarily of fluorspar (CaF2). The heating is generated by an electrical current (usually ac) that flows from the base plate through the liquid slag and the electrode. Oxides may be added to the slag to raise the resistivity (so as to make melting more efficient electrically) or to modify the slag chemistry with regard to its reactivity with the metal being melted. Electroslag remelting is similar to vacuum arc remelting (VAR), except that ESR is carried out at normal atmospheric pressure and has a greater melting rate than VAR. When the tip of the electrode is melted by the molten slag, metal droplets fall through the liquid slag and are collected in the water-cooled mold (Fig. 1). The slag not only shields the metal pool from contaminating gases but can also be so chosen that it acts to segregate and absorb impurities from the molten metal in its passage through the slag from the electrode to the molten pool. During the formation of the liquid film, the metal is refined and cleaned of contaminants, such as oxide particles. The high degree of superheat of the slag and of the metal favors the metal/slag reaction. Melting in the form of metal droplets greatly increases the metal/slag interface surface area. The intensive reactions between metal and slag refine the melt. The remaining inclusions are very small and are evenly distributed in the remelted ingot. The absence of macrosegregation or heterogeneity in the distribution of nonmetallic inclusions in ESR ingot improves yield to justify the additional cost of remelting. In addition, the very clean and smooth surface of the ESR ingot helps to reduce production costs, because surface conditioning before hot working is not necessary. Another special feature of the ESR process, as in VAR, is the directional solidification of the ingot from bottom to top. The macrostructure is marked by an extraordinarily high density and homogeneity as well as by the absence of segregations and shrinkage cavities. Electroslag remelting is commonly used to produce the highest levels of quality in plate steels, particularly in thick plates. The homogeneity of the ingot results in uniform mechanical properties in the longitudinal and transverse directions after hot working. For steels, the slag
composition can be adjusted to serve as a desulfurizer or dephosphorizer as well as a reservoir for floating inclusions. Sulfur can be lowered to below 0.002%. The cleanliness and homogeneity of ESR ingots results in steels with excellent mechanical properties and superior toughness (Fig. 2) compared to conventionally manufactured steel. In superalloys, ESR is
Fig. 1
often used as an intermediate remelt in a socalled triple-melt process (VIM-ESR-VAR), where the ESR electrode is produced by vacuum induction melting (VIM) and then remelted a third time by VAR after ESR (Ref 4). The theory of ESR was developed in the 1930s and was the subject of a U.S. patent in
Schematic of an ESR furnace with multiple electrodes for large ingot production
Electroslag Remelting / 125 1945 (Ref 5). However, it took more than 30 years for a general breakthrough on a production scale (Ref 6–12) in primary (mill) products. Currently, ESR metallurgy is a well-established production technology (Ref 13). Global production capacity of ESR steel is over a million (106) tons, which includes low-alloy highstrength steel, bearing steel, tool and die steel, and stainless and heat-resistant steels (Ref 14). Ingots from high-quality heavy ESR are over 200 tons, and production cost of heavy ESR manufacturing is significantly reduced by electroslag hot topping. Related technologies include electroslag casting, electroslag turning casting, Boher electroslag tapping, and electroslag centrifugal casting. Process variations of ESR technology include vacuum ESR and pressurized ESR. Electroslag remelting under vacuum (VACESR) is a process variation that has received some attention, because the traditional ESR process is essentially atmospheric and thus ineffective in reducing dissolved gases such as hydrogen and nitrogen. Vacuum ESR is not a superalloy production technology, given the predominance of triple-melt (VIM-ESR-VAR) for superalloys. Pressure electroslag remelting Temperature, ⬚F −4
Impact strength, J
80 60
32
68 104 140 176 212 73.8 ESR (tangential rim) 59.0
ESR (axial core)
44.3
40
29.5
20 Conventional (tangential rim)
0 −40 −20
20
0
40
Conventional (axial core)
60
Impact strength, ft-Ib
−40 100
is a process variation that increases dissolved nitrogen in the melt of austenitic steels for strengthening. Another ESR process variation is electroslag rapid remelt (ESRR). This technique allows production of smaller cross-sectional shapes (rounds and squares), which can be directly processed into smaller shapes on conventional open-die presses and rolling mills, saving conversion time and costs as compared to starting with typical larger round ingots. The ESRR technology is also being developed for nucleated casting (Ref 15), which is a type of spray casting.
ESR Furnaces Major components of modern singleelectrode ESR plants are shown in Fig. 3 (with fixed mold) and Fig. 4 (with a retractable bottom plate). In the case of a retractable bottom plate furnace, current feedback does not take place through a closed copper pipe; instead, it is accomplished with four symmetrically arranged current feedback tubes. The ESR process, like VAR, is a fully automated operation. Unlike a VAR crucible, the ESR crucible is normally self-contained in that the inside and outside shells through which the cooling water runs are assembled as a single piece. A watercooled stool is assembled to the shell to form the bottom of the crucible. A starter or striker plate is usually enclosed in the junction between the crucible shell and the crucible stool.
For crucibles to be used with hot slag starts, the crucible may contain an opening at the bottom (a “mousehole”) through which the molten flux would be introduced. The top of the crucible does not mate to the rest of the furnace, so it operates exposed to air. The electrode is attached to the ram through a stinger (Fig. 5). The ESR stinger is welded to the electrode outside the furnace. The stinger is often tubular, because only the periphery of the electrode is involved in the welding process. The top of the ESR furnace contains the ram drive and load cells. It is connected electrically to the crucible stool through, typically, four vertical bus bars located at 90 intervals around the crucible. The ESR head commonly is able to be translated in the horizontal direction by an XY drive, which allows centering of the electrode in the crucible prior to the start of the melt. (Note that XY capability is inherently different in ESR compared to VAR.) The power supply is most often ac. Most ESR furnaces have two melting stations (crucible setups). While an electrode is melting in one station, the next station is prepared for melting. When one melt is finished, the ESR head is rotated in the horizontal plane to be located over the
14.8
0 80 100
Temperature, ⬚C (a)
167
212
257
3822
in.
122
3500 ESR (0.002% S)
3185
2800
2548
Fracture toughness, ksi
Fracture toughness, MPa
m
Temperature, ⬚F 77
4200
1911
2100 Conventional (0.001% S)
1400
1274 637
700 25
50
75
100
125
Temperature, ⬚C (b)
Fig. 2
Comparison of properties of steel rotor forgings made from ESR and conventionally melted ingots. (a) Impact strength of grade X22CrMoV121. (b) Fracture toughness of grade 30CrMoNiV511. Specimen orientation and location are indicated next to curves. Source: Ref 1–3
Fig. 4 Fig. 3
Basic design of a modern ESR furnace with a fixed mold. 1, ram drive system; 2, electrode ram; 3, XY adjustment; 4, load cell system; 5, sliding contact; 6, four bus tubes; 7, pivoting drive; 8, electrode; 9, mold assembly; 10, coaxial bus tube; 11, base plate; 12, multicontacts
Schematic of an ESR furnace with retractable bottom plate. 1, electrode drive system; 2, load cell system; 3, stinger crosshead; 4, two ball screws; 5, three bus tubes; 6, mold assembly; 7, ingot; 8, ingot withdrawal table; 9, three bus tubes; 10, three ball screws; 11, three sliding contacts; 12, base plate; 13, furnace head drive; 14, electrode; 15, three sliding contacts; 16, XY adjustment
126 / Melting and Remelting second crucible/electrode stinger setup. The ESR process is also capable of using multiple electrodes, which can be melted into a single mold. Electrode preparation is not as critical a part of quality ESR operations as for VAR. Oxides on an electrode surface are incorporated into the slag. Thus, electrodes for ESR are generally not ground. An exception to this is for production of premium-quality material where the end-user specifications require grinding to ensure the removal of potential iron-mold pullout on the electrode surface. Also, the shrinkage cavity at the top of the electrode is not a problem, if the electrode is melted top down. Thus, superalloy ESR electrodes are not commonly hot topped in VIM or cropped prior to melting. There are two possible start-up scenarios: cold start and hot start. In cold start, the slag and small particles of the alloy to be melted (usually machining chips) are placed on the starter plate. The electrode is touched into the slag-alloy mix and backed out to establish an arc. High melt power is used in this start-up phase. The arc melts down both the slag and the metal particles, at which point the electrode becomes immersed in the slag, and the melt process shifts to melting of the electrode surface by the slag. It is necessary that the start-up power be sufficiently high that actual welding (melt-in) occurs between the embryonic ingot and the starter plate. This is required so that the electrical conduction path will be predominantly through the electrode, then the slag to the ingot, from the ingot through the starter plate, and then to the stool. In hot slag starts, the slag is melted externally by electric arc in graphite crucibles. The molten slag is introduced, generally through a bottom mousehole, into the crucible. The electrode is lowered into the slag, and melting
Fig. 5
ESR electrode with stinger
commences. Although high initial melt power is not required to melt slag and metal starter material, a high-power profile start-up is generally used to ensure the melt-in of the ingot to the starter plate (to ensure good electrical conduction paths). ESR Control Parameters. Control of the ESR process is considered from two different aspects: control of pool depth and shape (control of solidification structure) and development of a satisfactory ingot skin. Pool depth and shape are controlled primarily by melt rate (power input). Electroslag remelting generally is run using melt rate control, with continuous amperage adjustments to maintain a uniform melt rate. Due to the high thermal inertia of the ESR process, the melt rate of ESR is more variable than in VAR. Secondary control of pool depth and shape is achieved by maintaining a uniform, shallow immersion of the electrode in the slag. The passage of molten metal droplets through the slag affects slag resistivity and thus causes variation in the operating voltage of the system. The operating voltage appears as a band of values due to this variation. The width of that band is called the volt swing. The electrode immersion is proportional to the volt swing, and the ram drive of the ESR advances the electrode to maintain the desired swing value. Production of satisfactory ingot skin is not actively controlled during the melt but is due to the choice of variables. The most important of these are melt rate, depth of the slag pool, electrode immersion depth in the slag (analogous to arc gap in VAR), ingot and electrode diameter (as in VAR), and choice of slag composition. Higher melt rates produce better ingot surfaces. Because higher melt rates also produce deeper molten pools and move the process closer to the conditions for freckle formation, selection of the correct balance of parameters is critical. The choice of ingot and electrode diameter defines the clearance (annulus) between the crucible and the electrode. Insufficient clearance is generally regarded as a problem in processes using cold slag starts, in that the slag-chip mix may get hung up in the annulus. However, a tight annulus may result in more uniform heat distribution across the molten pool. There is no published information to support either viewpoint. Thus, annulus becomes, much more than in VAR, a parameter dictated by the individual producer’s philosophy. Electrode immersion is generally thought of as the depth of penetration of the electrode into the slag. There is no definitive method for measuring the immersion. In superalloy melts, observation of interrupted melts suggests that the meniscus of slag reaching up the side of the electrode will generally not exceed 6.4 mm (0.25 in.) (Ref 4). Thus, schematics of the ESR process are misleading when they imply significant immersion into the molten slag. However, the degree of immersion, as affected by the ram drive, can produce drastically
different electrical signals. The quality of the melt has been found to be responsive to these signals. Thus, while there is no realistic understanding of what depth of immersion means physically, the parameter that controls it can be demonstrated to have a significant effect on melt quality. The electrode immersion plus the thickness of the slag cap may be considered as the resistance (ac impedance) in the circuit. Changing the weight of slag in the system or the slag composition changes the impedance of the circuit and thus changes the amperage and voltage required to maintain a given power input (for constant melt rate). The slag cap thickness is the primary driver, while electrode immersion may account for 10 to 20% of the total impedance. Measurement of the changes in this resistance or of the corresponding changes in voltage may be used to control the immersion of the electrode in the slag. This suggests that the voltage and impedance changes associated with electrode immersion are actually the impedance of the circuit associated with the interfacial resistance between the electrode melt surface and the slag surface. It thus should be realized that in ESR of superalloys, the electrode immersion is always extremely shallow, and, in fact, the most common electrode positioning is that of “skittering” on the top of the molten slag cap. The degree of this skittering is commonly controlled by the repetitive change in voltage in the process (volt swing). The electrode melt rate is the other important parameter to be controlled. The melt rate is dependent on the melt current. Although some VAR processes are run in constant current rather than in melt rate control, ESR processes are exclusively melt rate control. Melt rates are generally calculated using rolling averages. Load cells are subject to the generation of false signals and to errors in the absolute value. Thus, load cell design and maintenance are important factors in maintaining the quality of ESR. In ESR, the melt current responds to changes in the melt rate and attempts to hold the melt rate constant at the set point. Because of the presence of the slag, which must be heated or cooled, in the system, the response time of ESR to the melt current change is much slower than that of VAR. A fourth factor controlling melt quality and unique to ESR is the choice of slag and the volume of slag used. The resistance of the slag cap depends on both the resistivity of the slag and the thickness of the slag cap through which the melt current must travel. An additional consideration is the degree of chemical reactivity of the slag with elemental components of the electrode. Figure 6 is a nominal plot of process parameters for the remelting of a nickel-base alloy. Figure 7 shows voltage swings and the major control parameters for a hypothetical ESR melt. The melt begins with a high-amp start-up in amperage control and, at the completion of the start-up, is placed into melt rate control.
Electroslag Remelting / 127 6
Melt rate, kg/min
Ingot diam: 460 mm Ingot weight: 2070 kg
Set rate 3.5 kg/min
4
6.6
3 2
4.4
Start
1 Power control 0 5 Power, kW ⫻ 10−2 Voltage, V ⫻ 10−1 Current, kA ⫻ 10−1
8.8
2.2
Hot topping
Melting rate control
Melt rate, Ib/min
11.0
5
0 Power control
4
Voltage
3 Power 2 Current
1 0
Ensuring consistent remelting over a period 0
1
2
3
4
5
6
7
8
9
10
Time, h
Fig. 6
Electroslag Remelting of Heavy Ingots. At the end of the 1960s, the concept of using ESR plants to manufacture large ingots weighing up to 350 Mg (385 tons) gained acceptance. Increasing demands for energy and the trend toward larger electrical power-generating units required cast ingots weighing 100 Mg (110 tons) and more for the manufacture of generator and turbine shafts. With ESR, it is possible to achieve higher yields and to avoid such internal defects as porosity, macrosegregation, and the accumulation of nonmetallic inclusions. An early example of a large ESR furnace in the 1970s had a retractable bottom plate and manufactured ingots 2300 mm (90.5 in.) in diameter, 5000 mm (197 in.) in length, and weighing up to 160 Mg (176 tons) (Ref 2) (Fig. 8). Remelting of large ingots of this size in retractable bottom plate furnaces demanded the solution of specific technical and metallurgical problems, such as: of several days without interruption
Prevention of steel and slag breakout when
remelting with a retractable bottom plate
Achieving a good ingot surface Ensuring directional solidification to avoid
ESR process parameters for the melting of a nickel-base alloy
macrosegregation and shrinkage cavities
Controlling the steel and slag compositions
Volt swing - Ram travel (in.) - Resistance Voltage swing and resistance
over the entire ingot height
Ram
Adjusting for low hydrogen content Adjusting for a low aluminum content
iion Posit
Volt Swing Resistance
Voltage
Melt rate - Voltage - Amps 40 35 30 25 20 15 10 5 0
Voltage
Melt Rate Amperage
Elapsed time during melt
Fig. 7
Melt trace for ESR showing variation of melt current, melt rate, ram travel, slag resistance, voltage, and volt swing
Note that the change in voltage is the result of the deliberate change in amperage. Formation of a shrinkage cavity several hundred pounds deep into the ingot can occur from solidification of the large molten pool after the end of the melt, if the electrode is completely consumed or if the power abruptly shuts off. This cavity must be removed during subsequent processing. To minimize the amount of material that must be cropped because of the end-of-
melt shrinkage, it is common to step down the melt current as the end of melt approaches and leave a small amount of electrode. Unlike VAR, this biscuit is generally not used to put electrical power into the melt. Rather, it simply prevents radiation loss from the slag. Because of the presence of the slag as a heat source at the top of the ingot, ESR hot topping practices are not as extended as are those for VAR.
(<0.010%) for rotor steels; this has a decisive influence on the creep-rupture properties of the rotors After these problems had been solved, heavy rotors for electrical generators could be manufactured from ESR ingots. Interior defects such as macrosegregation, shrinkage cavities, and nonuniform distribution of inclusions can be avoided in large ingots only if directional solidification is ensured over the entire ingot section. By maintaining the correct melting rate and temperature of the slag, directional solidification can be achieved for large ingot diameters. Accordingly, the ESR ingot is free of macrosegregation in spite of the large diameter. Several electrode changes are necessary to produce such large ingots. In the case of a multi-electrode operation in a single mold, the formation of the ingot can be concluded by a hot topping procedure performed by withdrawing all but one of the electrodes and continuing the process in single-phase mode with the remaining electrode centralized in the mold, for example, by appropriate sideways displacement of the latter on a wheeled support. In such a case, loss of bottom contact to the ingot during the multiphase operation could prevent the subsequent single-phase hot topping step from proceeding. Moreover, loss of contact is more likely to occur with such a multiphase operation than with a single-phase operation, because in the latter the flow of current between the ingot bottom and the mold base has some tendency to keep a thin slag film
128 / Melting and Remelting
Fig. 8
160 Mg (176 ton) ESR ingot before forging into a generator rotor
Fig. 9
Schematic of solidification relationships in VAR (left) and ESR (right) melting processes. Source: Ref 16
between the ingot and mold base in an electrically conductive state. The refining process for producing steel ingots of large sizes (e.g., 200 to 400 ton class) is possible by controlling the rate of solidification in electroslag remelting or with electroslag hot topping techniques.
Ingot Solidification The solidification structure of an ingot of a given composition is a function of the solidification rate and temperature gradient at the
solid/liquid interface. When remelting a consumable electrode in a water-cooled copper mold, the solidification rate is constant because of the constant heat transfer through the cooling water; therefore, a relatively high temperature gradient at the solidification front must be maintained during the entire remelting period to achieve a directional dendritic structure. The structures in the solidified ingot are governed by the same principles as those governing a solidifying VAR ingot. However, the nature of the heat transfer (through an oxide skin) and the presence of a heat reservoir (the molten
slag) at the top of the ingot/molten pool make ESR structure inherently different from that of VAR. Figure 9 schematically illustrates the relationship between the ESR electrode, the solidifying ingot, and the shape and depth of the molten pool. A major feature of ESR solidification is that a “skin” of oxide is incorporated between the ingot surface and the crucible wall. Similar to VAR, the distance between the side of the electrode and the crucible wall is the annulus. The depth of the slag pool is an important operating parameter that controls the electrical resistance of the system. The depth of immersion of the electrode into the slag is an important parameter in controlling the molten pool shape. As shown in Fig. 9, the heat extraction in ESR is of a much greater magnitude than for VAR. Due to the insulating characteristic of the ingot slag skin, heat extraction in ESR is inherently less efficient than in VAR. Thus, ESR is a process with high thermal inertia and slow responses (change in pool shape) that occur in response to changes in power input. The presence of an ingot slag skin not only improves the ingot surface quality but also, more importantly, eliminates the presence of a solute-lean shelf layer on ESR ingots. (As oxide particles from the electrode are dissolved into the slag, they do not agglomerate on the ingot surface.) Because ESR ingots do not have an ingot shelf, they are inherently free of the formation of drop-in (discrete) white spots from this source, but drop-in electrode may still be a source of white spots in ESR. Solidification white spots have not been reported in ESR product. However, one consequence of the ingot slag skin is that heat extraction in ESR is less efficient than for VAR. This is shown schematically in Fig. 9 by the difference in pool shape between the two processes. ESR pools tend to be both deeper and more V-shaped (compared to a U-shaped VAR pool). Deep pools favor freckle formation, which can only be counteracted by reducing ingot size. Consequently, the maximum size of ESR ingot produced (for any given alloy system) is smaller than one that can be produced by VAR. Figure 10 schematically illustrates the structure from an ESR molten pool (the relative sizes of annulus and immersion are not to scale). Again, the most important feature shown is the depth of the molten pool (where the depth and angle of the liquid-solid zone also controls freckle formation in superalloys). Compared to VAR, few studies have been done on the nature of the molten pool in ESR. The pool differs fundamentally from a VAR pool in that the molten slag cap provides a source of heat to the system, and the slag skin on the solidifying ingot provides insulation, reducing heat extraction. Consequently, for comparable melt rates and crucible diameters, an ESR ingot will have a deeper, steeper-sided pool. Compared to VAR, ESR ingots have traditionally been regarded as being more prone to the formation of positive segregation effects. However, even though the ESR pool may inherently be more
Electroslag Remelting / 129 prone to positive segregation, that shape may be altered by judicious choice of melting parameters. Even in the case of directional dendritic solidification, the microsegregations increase with increasing dendrite arm spacing. For optimal results, the objective should be a solidification structure with dendrites oriented as close to parallel to the ingot axis as possible. However, this is not always possible. To achieve a good ingot surface, there is a minimum energy requirement, and therefore a minimum melting rate, which is a function of ingot diameter. This means that the melting rate for large-diameter ingots cannot be maintained for axis-parallel dendrite growth. Like VAR, the melting rates for different steels and alloys are a function of ingot diameter. Ingot Defects. In spite of directional dendritic solidification, remelted ingot may include various defects, such as the formation of: Tree ring patterns Freckles White spots
These types of defects also occur in VAR ingots and are discussed in more detail in the article “Vacuum Arc Remelting” in this volume. In Fig. 10, the schematic for the edge of a solidifying ingot indicates a processing advantage for ESR that has made it desirable to develop this process for use in critical components. The insulating nature of the slag
creates conditions at the edge of the ingot that do not produce a shelf. Thus, with no arc instabilities of concern and no shelf to be undercut, the formation of discrete white spots in ESR product could only occur by some infrequent occurrence, such as drop-in of pieces from an unsound electrode. Positive Segregation. Alloys that have highdensity interdendritic liquids may form unique structures in ESR. The high-density liquid tends to flow downhill toward the center of the ingot. Although channel defects may not be formed, large pools of high-solute alloy may be formed. These regions may form in widely separated parts of an ingot and be separated by “normal” material. Because of the greater number of variables inherent in ESR, the underlying causes for such segregation have not been determined. The nature of these areas and their effect on properties have not been reported. Electroslag Remelting Ingot Surface. The interaction of the molten metal with the slag cap should be understood with regard to obtaining good surface on ESR ingots. As shown in Fig. 10, the molten slag solidifies against the water-cooled crucible side just as the molten metal would. The molten metal pool, in addition to being V-shaped (containing the mushy zone), also contains a fully molten straightsided component or metal head. The metal head remelts some of the solidified slag on the crucible wall, so a thinner slag layer is then contained between the ingot and the crucible wall. Should this slag skin be locally penetrated by the metal head, a thin stream of metal will be driven through the hole. Such a metal fin or bleedout is illustrated in Fig. 11. Note that the penetration of the slag is through a circular region that is essentially the size of a dime.
Metal head
Liquid metal penetration of slag
Ni balls Metal fin
Fig. 11
Fig. 10
Longitudinal schematic of the structure developed in an ESR ingot during melting
Longitudinal section through a metal fin that was formed by a bleedout on the surface of an ESR nickel superalloy ingot. Note: The circular, darketching features are nickel balls introduced in an experiment to delineate the shape of the molten pool. They are not related to the bleedout mechanism.
The molten metal not only moves down the ingot surface in response to gravity but is forced upward from the penetration due to the combined hydrostatic pressure of the molten metal head and the molten slag cap. Bleedout formation is sensitive to melting conditions and slag composition. While there is no published association of bleedouts with detrimental structure, there is an economic factor involved, in that an ingot with bleedouts will require grinding prior to forging or other hot working.
Process Variations Several variations of ESR have been developed: remelting under high pressure, known as pressure electroslag remelting; remelting under reduced pressure; and electroslag rapid remelting. Pressure Electroslag Remelting. The strength of austenitic steels can be increased by alloying with nitrogen as long as the nitrogen is interstitially dissolved in a stable condition. Nitrogen solubility, however, is limited in the case of melting under atmospheric pressure for a given steel composition. Because the solubility of nitrogen in iron is proportional to the nitrogen partial pressure in the furnace atmosphere, there is the possibility of alloying nitrogen into the melt under higher pressure and allowing it to solidify. The ESR process is suitable for this technique. The first trials in this area were carried out in 1971 (Ref 17). For some years, a high-manganese steel has been successfully remelted under a pressure of 4.2 MPa (610 psi). Up to 0.5% interstitial solubility of nitrogen in steel has been achieved. The nitrogen-alloyed austenitic steels are used for the manufacture of retaining rings in generators. Electroslag remelting under vacuum (VAC-ESR) is another interesting process. Alloy 718 has been remelted in a VAC-ESR pilot plant with a fluoride-free slag based on lime and alumina (Ref 18). Because the remelting is carried out under vacuum, problems of oxidation of the melt do not arise. In addition, dissolved gases, such as hydrogen and nitrogen, can be removed. Thus, the advantages of both ESR and VAR are combined in one process (Ref 18). Nonetheless, VAC-ESR is not a production practice for superalloys. Electroslag rapid remelt is a process technology where electrodes are ESRed into smaller cross-sectional shapes (rounds and squares) that can be directly processed into smaller shapes on conventional open-die presses and rolling mills, saving conversion time and costs as compared to starting with typical larger round ingots.
Steel ESR There are a number of publications that deal primarily with improvements in the physical properties of remelted steel (Ref. 6–12).
130 / Melting and Remelting However, they discuss not only the directional solidification of the metal but also the influence of liquid, superheated, and therefore reactive slags. The extent and direction of the metallurgical reactions in the ESR process are determined by the steel composition, the slag used, and the atmosphere used (inert gas, air, etc.). The principal feature of the ESR process is the slag bath. A continuous transport of liquid metal through the slag takes place. During this transport, the slag and the metal compositions change, according to the kinetic and thermodynamic conditions. To perform its intended function, the slag must have some well-defined properties, for example: Its melting point must be lower than that of
the metal.
It must be electrically efficient. Its composition should be such that the
desired reactions, such as removal of sulfur and oxides, are ensured. It must have suitable viscosity at remelting temperature.
of easily oxidizable elements, such as aluminum and silicon, during remelting. To counteract this oxidation, the slag should be continuously deoxidized, preferably with aluminum. With proper slag composition and
Table 1 Compositions of slags commonly used in ESR Composition, % Slag designation(a)
CaF2
CaO
MgO
Al2O3
SIO2
Comments
100F 70F/30
100 70
... 30
... ...
... ...
... ...
70 70 50 70
20 15 20 ...
... ... ... ...
10 15 30 30
... ... ... ...
Electrically inefficient; use when oxides are not permissible. Difficult starting; high conductivity; use when aluminum is not allowed; risk of hydrogen pickup
40 60 80 60 ...
30 20 10 10 50
... ... ... 10 ...
30 20 10 10 50
... ... ... 10 ...
70F/20/0/10 70F/15/0/15 50F/20/0/30 70F/0/0/30 40F/30/0/30 60F/20/0/20 80F/0/10/10 60F/10/10/10/10 0F/50/0/50
Good general-purpose slags; medium resistivity Some risk of aluminum pickup; good for avoiding hydrogen pickup; higher resistivity
g
Good general-purpose slags Moderate resistivity; relatively inert Low-melting, “long” slag Difficult starting; electrically efficient
60 ⬉10% Al2O3
50 40 30 20
(Eq 1)
(Eq 2)
It is evident that a saturation of the slag with sulfur does not take place; therefore, the desulfurization capacity of the slag remains intact throughout the entire remelting process. With a highly basic slag (CaO/SiO2 > 3), more than 80% of the sulfur can be removed. Behavior of Oxygen. As mentioned previously, the ESR process is usually carried out under a normal air atmosphere. Oxidation of the metal is unavoidable. Oxygen can be transferred into the metal in several ways: Oxidation of the electrode surface above the
slag bath Oxidation on the slag surface of elements with variable valences, such as iron and manganese Oxides attached to the electrode surface Transfer of oxygen into the slag also takes place, and oxygen is transferred into the metal according to Eq 1. This results in losses
for C16 10 Average oxygen content, ppm
The second reaction is the slag/gas phase reaction. In this case, the sulfur absorbed by the slag is removed by the oxygen of the gas phase in the form of gaseous sulfur dioxide: 3 CaS þ O2 ðgasÞ ! CaO þ SO2 ðgasÞ 2
g
(a) The universal designation system for slags is the percentage of constituents listed in the order CaF2/MgF2/CaO/MgO/Al2O3. Thus, a common base slag such as 60%CaF2/20%CaO/20%Al2O3 would be referred to as 60/0/20/0/20. However, industrial designations often omit the 0s. Thus, the previously mentioned slag is generally referred to as 60/20/20. Source: Ref 7
Slags for ESR are usually based on calcium fluoride (CaF2), lime (CaO), and alumina (Al2O3). Magnesia (MgO), titania (TiO2), and silica (SiO2) are also added, depending on the alloy to be remelted. Table 1 shows the compositions of some common ESR slags. Behavior of Sulfur. One of the primary advantages of the ESR process for steel is the good desulfurization of the metal. The final desulfurization is determined by two reactions. The first is the metal/slag reaction, in which sulfur is transferred from the metal to the slag: FeS þ CaO ! CaS þ FeO
remelting techniques, oxygen contents of less than 20 ppm in unalloyed steel are possible. Figure 12 shows the influence of slag composition on the final oxygen content of remelted ingots.
0 40 15–25% Al2O3 30 20 10 0 40 30–45% Al2O3
30 20 10 0
1
2
3
4
5
6
7
8
9
Slag basicity
Fig. 12
Effect of slag composition on the oxygen content of the remelted ingot. Slag basicity is the ratio of CaO to SiO2.
Electroslag Remelting / 131
Superalloy ESR Generally, superalloy ESR uses a cast VIM electrode as a consumable electrode. Unlike VAR, high-vapor-pressure tramp elements are not removed by ESR. Thus, magnesium may be retained at higher levels than is possible in VAR (with beneficial effects on hot workability). A beneficial chemical reaction related to workability is the removal of sulfur. Detrimental compositional changes may occur in reactive elements (notably titanium and aluminum), because these elements change to attain chemical equilibrium with components of the slag. Oxides incorporated in the electrode being melted are effectively removed by the ESR process, producing an ingot that is generally cleaner than a VAR ingot. All ESR slags contain CaF2 as the primary constituent. Oxides may be added to the slag to raise the resistivity (so as to make melting more efficient electrically) or to modify the slag chemistry with regard to its reactivity with the metal being melted. Additions of CaO, MgO, and Al2O3 are made to modify both the resistance of the slag and its melting point. All commercial-purity CaF2 contains SiO2. During melting of titanium-bearing alloys, an exchange will take place between the titanium in the alloy and the silicon in the slag. The slag becomes enriched with TiO2, and the silicon content of the alloy increases until an equilibrium is established, usually within the first few hundred pounds of melting. To minimize this exchange, some commercial slags are buffered with intentional additions of TiO2. Titanium and zirconium are superalloy elements that often interact with the slag during ESR. Many proprietary slags contain additions of TiO2 and/or ZrO2 to prevent excessive loss of these elements into the slag during melting. The loss is due to exchange with any oxide in the slag that is less stable than the element in the electrode. A common problem with titanium-bearing superalloys is that commercially available CaF2 contains SiO2. During ESR, the SiO2 is converted to silicon, which is picked up by the metal. Thus, the slag is enriched in TiO2, which reduces the titanium content of the metal. This reaction generally results in a titanium gradient in the ingot, because the SiO2 content is consumed early in the melt and no longer plays a part in the reaction after the first few hours of melting. When
ESR slag is recycled, it is already depleted in SiO2, and this reaction no longer occurs. The slag always tries to move to chemical equilibrium with the molten metal at the operating temperature of the slag. Thus, repetitive recycling of slag with a given alloy composition produces a slag that is naturally buffered against causing elemental loss from the electrode during ESR. Unfortunately, slag loss (left as a skin on the electrode during the process) generally keeps recovery levels of the slag below 100%, so some fresh slag is always needed in the ESR process.
ACKNOWLEDGMENTS Reviewed by Scott A. Balliett, Latrobe Specialty Steel Company, with portions adapted from Ref 4 and from: A. Choudhury and F. Knell, Electroslag Remelting (ESR), Metals Handbook, 9th Edition, Volume 15, Casting, 1988, pp 401–405.
REFERENCES 1. R. Jauch, A. Choudhury, F. Tince, and H. Steil, Herstellung und Verarbeitung von ESU-Blo¨cken bis zu 160 t, Stahl Eisen, Vol 101, 1981, p 41–44 2. A. Choudhury, R. Jauch, and F. Tince, Low Frequency Electroslag Remelting of Heavy Ingots with Low Aluminium Content, Proceedings of the Sixth International Conference on Vacuum Metallurgy (San Diego), 1979, p 785–794 3. A. Choudhury, R. Jauch, H. Lo¨wenkamp, and F. Tince, Application of the Electroslag Remelting Process for the Production of Heavy Turbine Rotors from 12 %Cr-Steel, Stahl Eisen, Vol 97, 1977, p 857–868 4. M.J. Donachie and S.J. Donachie, Superalloys: A Technical Guide, ASM International, 2002 5. R.K. Hopkins, Method and Apparatus for Producing Cast Metal Bodies, U.S. Patent 2,380,238, 1945 6. B.E. Paton, B.I. Medovar, and W.E. Raton, Ein neues Verfahren—Das Elektroabgiessen von Blo¨cken, Bull. Tech. Informazil NTO, No. 1, Maschprom., Kiew, 1956
7. G. Hoyle, Electroslag Processes, Applied Science, 1983 8. W. Holzgruber and E. Plo¨ckinger, Metallurgische und verfahrenstechnische Grundlagen des Elektroschlacke-Umschmelzens von Stahl, Stahl Eisen, Vol 88, 1968, p 638–648 9. M. Wahlster, A. Choudhury, and K. Forch, Einfluss des Umschmelzens nach Sonderverfahren auf Gefu¨ge und einige Eigenschaften von Sta¨hlen, Stahl Eisen, Vol 88, 1968, p 1193–1202 10. M. Wahlster and A. Choudhury, Beitrag zum Elektroschlacke-Umschmelzen von Sta¨hlen, Rheinstahl Technik, Vol 5, 1967, p 31–37 11. M. Wahlster, G.H. Klingelho¨fer, and A. Choudhury, Neue metallurgusche und technologische Ergebnisse einer 10 t ESUAnlage, Radex-Rundsch., Vol 2, 1970, p 99–111 12. G.H. Klingelho¨fer, A. Choudhury, and E. Ko¨niger, Etude compare´e des caracte´ristique des aciers elabore´s a` l’air et des aciers refondus selon le proce´de´ sous laiter electroconducteur (ESR) pour les cylindres de laminage a` froid et les cylindres calandreurs, Rev. Metall., Vol 67, 1970, p 512–522 13. B.I. Medovar, Arc-Slag Remelting Steel and Alloys, Cambridge International Science Publishing, England, 1995 14. T. Li, Spec. Steel (China), Vol 25 (No. 5), Sept-Oct 2004, p 1–5 15. W.T. Carter, J.J. Jackson, R.M. ForbesJones, and R. Minisandram, Clean Metal Nucleated Casting of Superalloys, Sixth International Symposium on Superalloys 718, 625, 706, and Derivatives, Oct 2–5 2005 (Pittsburgh, PA) 16. G.E. Maurer, Primary and Secondary Melt Processing—Superalloys, Superalloys, Supercomposites and Superceramics, Academic Press, 1989, p 49–97 17. Ch. Kubisch, Druckstickstoffsta¨hle—eine neue Gruppe von Edelsta¨hlen Berg.—und Hu¨ttenma¨nnische Monatshefte, 1971, p 84–88 18. A. Choudhury, “Vacuum Electroslag Remelting—A New Process for Better Products,” Paper presented at the Vacuum Metallurgy Conference, June 1986 (Pittsburgh, PA)
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 132-138 DOI: 10.1361/asmhba0005202
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Vacuum Arc Remelting THERE ARE TWO COMMONLY USED remelting processes for metal refinement: electroslag remelting (ESR) and vacuum arc remelting (VAR). In both processes, an electrode is melted as it advances into the melting region of the furnace. As the working face of the electrode is heated to the melting point, drops of liquid metal that fall from the electrode face are collected in a lower crucible and rapidly solidified. However, ESR and VAR have very different melting methods, which have implications regarding the magnitude of cooling rates obtained and the nature of defects created by the remelting process itself. The VAR process also is carried out in a vacuum, while the ESR process typically occurs under atmospheric conditions—although electroslag remelting under vacuum (VACESR) or under pressure is a process variation of ESR (see the article “Electroslag Remelting” in this Volume for a brief description of VAC-ESR). The VAR process is widely used to improve the cleanliness and refine the structure of standard airmelted or vacuum induction melted (VIM) ingots (typically referred to as consumable electrodes
Fig. 1
Key components of a vacuum arc remelting furnace. Source: Ref 1
for remelting). The VAR process is also commonly used in the triplex production (VIMESR-VAR) of superalloys. The VAR process was the first commercial remelting process for superalloys. It was used in the late 1950s to manufacture materials for the aircraft industry. Current applications of VAR processing include the production of steels and superalloys as well as the melting of reactive metals such as titanium and zirconium alloys. The melt cleanliness and homogeneity from VAR provides benefits in both primary ingot and foundry shape casting for high-integrity applications where improved fatigue and fracture toughness of the final product are essential. This includes aerospace applications, high-strength steels, ball-bearing steels, die steels, tool steels (cold and hot work steels), power generation, defense, and medical and nuclear industries.
Process Description The basic process of VAR (Fig. 1) is the continuous melting of a consumable electrode by means of a direct current arc (electrode negative, melt pool positive) in a vacuum on the order of 0.1 to 1 Pa (7.5 104 to 0.0075 torr). In some cases, the melting is carried out under inert gas with a pressure up to 1000 Pa (7.5 torr). Evaporation losses of volatile alloying
Fig. 2
elements are minimized at this pressure. Remelting currents are up to 40 kA. Melting rates with VAR furnaces are as high as 1150 kg/h (2500 lb/h), and ingot diameters are as large as 1.5 m (5 ft). Figure 2 shows melting rates for various steels and alloys as a function of ingot diameter. These are empirical values that were obtained from experience in operation. These melting rates gave low microsegregation while achieving acceptable surface quality. The diameter of the electrode and its relationship to the crucible is critical and must be matched for the melting rate. The VAR crucible is immersed in a tank of cooling water. Often, a tube is used to further enclose the crucible and limit the thickness of the water stream moving past the crucible surface. This increases the water velocity and thus increases heat removal from the crucible surface. Such tubes are called water guides. The crucible base (stool) has provision for the inlet of a gas (usually helium). The top of the crucible mates with the VAR head. The head contains the vacuum ports and the ram, which drives the electrode into the crucible as it is consumed. The electrode position is controlled to maintain an arc gap between the melting electrode tip and the molten pool on the top of the solidified ingot in the water-cooled copper mold. When an ingot is solidifying in the crucible, it pulls away from the crucible wall as it
Typical vacuum arc remelting rates for various steels and nickel- and cobalt-base superalloys
Vacuum Arc Remelting / 133 cools. This reduces heat conduction to the crucible wall. To maintain the highest possible heat removal from the solidifying ingot, a high-heat-capacity gas such as helium is introduced into the gap that is formed by the ingot shrinkage. The ram in newer designs is capable of X-Y movement. In VAR, this is not movement in the horizontal plane, which would allow centering of the stinger/electrode assembly in the crucible. It is actually angulation of the ram from the vertical, so that if the stinger is not perfectly parallel with the electrode, then the ram may be angled to make the electrode parallel to the crucible sides. The VAR head mates to the crucible flange through an O-ring seal. Most VAR furnaces have two melting stations (crucible setups). While an electrode is melting in one station, the next station is prepared for melting. When one melt is finished, the VAR head is rotated in the horizontal plane to be located over the second crucible/electrode setup. The electrode is attached to the ram through a stinger, which is welded to the electrode. Although the basic design of a VAR furnace has remained largely unchanged over the years, significant advances have been made in the fields of control and regulation of the process with the object of achieving a fully automatic melting procedure. Modern VAR furnaces (Fig. 3) make use of computer-controlled process automation with continuous weighing of the consumable electrode and closed-loop process control of melting parameters such as remelting rate and arc gap based on arc voltage or drop-short pulse rate. This in turn has had a
decisive positive influence on the metallurgical properties of the products. Close control of all remelting parameters is required for reproducible production of homogeneous ingots with a controlled solidification structure free of macrosegregation. Process Capabilities. Because of the high arc temperature and the small pool of liquid metal, sound ingots with dense crystal structure, low hydrogen and oxygen contents, and minimal chemical and nonmetallic segregation are produced. The primary benefits of remelting a consumable electrode under vacuum are: Removal of dissolved gases, such as hydro-
gen, nitrogen, oxygen, and CO
Reduction of undesired trace elements with
high vapor pressure
Improvement of cleanliness by removal of
oxides
Achievement of directional solidification of the
ingot from bottom to top, thus avoiding macrosegregation and reducing microsegregation In VAR, center porosity and segregation in the ingot resulting from ingot casting are also eliminated because of progressive melting and solidification. Directional solidification of the ingot from bottom to top avoids macrosegregation and minimizes microsegregation. Oxide inclusion removal is optimized because of the relatively short reaction paths during melting of the hot electrode end and because of good drop dispersion in the plasma arc. Oxide removal is achieved by chemical and physical processes. Less stable oxides or
nitrides are thermally dissociated or are reduced by carbon present in the alloy and are removed via the gas phase. However, in special alloys and in high-alloyed steels, the nonmetallic inclusions (e.g., alumina and titanium-carbonitrides) are very stable. Some removal of these inclusions takes place by flotation during remelting. The remaining inclusions are broken up and evenly distributed in the cross section of the solidified ingot. Because the molten metal drops immediately by gravity into the molten pool, it is inherently impossible to superheat the metal in the VAR process. This, coupled with the very high heat-extractive capability of the process, makes VAR the choice for economic manufacture of the largest-diameter ingots of segregation-prone superalloys. Additionally, the exposure of small volumes of molten metal to high vacuum is capable of removing detrimental high-vaporpressure elements. Unfortunately, beneficial high-vapor-pressure elements such as magnesium are also reduced greatly in concentration. Melt Pool in the VAR Process. Figure 4 illustrates the features of a molten pool in a VAR ingot. The most important feature shown is the depth of the molten pool. It generally is assumed that the depth of the liquid zone is representative of the depth of the liquid-solid zone. For a given crucible diameter, the size (depth) and shape of the liquid-solid zone has a direct relationship with heat input (melt rate) and heat extraction (water cooling). The depth and angle of the liquid-solid zone also controls freckle formation for any given alloy. The melt rate is directly proportional to the applied melt current (Fig. 5). However, as seen in Fig. 6, the relationship between pool depth and melt current is direct but is divided into
Fig. 4
Fig. 3
Modern vacuum arc remelting (VAR) furnace. (a) 30 ton VAR. (b) Operational components: 1, electrode feed drive; 2, furnace chamber; 3, melting power supply; 4, busbars/cables; 5, electrode ram; 6, water jacket with crucible; 7, vacuum suction port; 8, X-Y adjustment; 9, load cell system. Courtesy of ALD Vacuum Technologies GmbH
Molten pool region in a vacuum arc remelted ingot. The relative sizes of annulus and arc gap are not to scale. Competing currents are illustrated for stirring by thermal buoyancy (rising of lower-density hot metal along the ingot centerline) versus the stirring by electromagnetic fields (Lorentz stirring) that drive liquid metal from the edge of the crucible to the center of the crucible and down the centerline. Source: Ref 1
134 / Melting and Remelting
Fig. 5
Melt rate versus vacuum arc remelting current (50 cm, or 20 in., ingot; a composite graph from several independent sources). Source: Ref 1
liquid metal from the edge of the crucible to the center of the crucible and down the centerline. Figure 6 thus has two regimes, where the inflection point (nominally 6600 A for a 50 cm, or 20 in., IN-718 ingot) marks the point when Lorentz stirring becomes stronger than thermal buoyancy stirring. The apparent depth of the molten pool is altered not because there is a change in the volume of molten metal, but because the centerline stirring at higher amperages causes the pool to depart from a nominal U-shape and thus become deeper in the center. Due to Lorentz stirring, high VAR amperages thus produce deeper pools and also alter the angle of the liquid-solid zone, with respect to the ingot axis. In addition to deepening the molten pool and increasing the angle of the liquid-solid zone, there is another ramification when Lorentz stirring is dominant. At low currents, oxides and nitrides in the electrode are melted out, drop onto the molten pool surface, and are swept to the sides of the electrode. When melting at high currents, these particles may not migrate to the edge of the ingot but may be incorporated to some extent in the general structure of the ingot, perhaps with a concentration along the centerline. It is thus considered inadvisable to melt at VAR currents in the Lorentz-dominated region. As the volume of metal on the stool (the ingot) builds up, an equilibrium state is reached in which there is a solid ingot, a mushy zone of both liquid and solid above that, and then a zone that is totally liquid. Because the heat extraction in the steady-state condition of melting is fastest through the sidewalls and slower down the ingot and through the stool, the mushy zones and liquid zones are shallower near the sidewalls and deeper in the ingot center. The thickness and growth angles of the mushy zone determine if the interdendritic liquid regions will coalesce to form positive segregation channel defects. The thickness of the mushy zone at the center of a VAR ingot is controlled by the efficiency of heat extraction, the size (diameter) of the crucible, and the melt rate of the electrode face. The melt rate is controlled by the magnitude of current passed through the electrode. Other factors that are controlled so as to positively affect the solidifying structure are the distance of the melting face above the molten pool (arc gap) and the clearance of the sides of the electrode from the crucible (annulus).
Solidification Fig. 6
Vacuum arc remelting pool depth versus melt current. Source: Ref 1
two separate regions. The reason for this is indicated by the two lines in Fig. 4, which indicate the direction of liquid metal flow in the VAR pool. Competing currents are generated
by the tendency of lower-density hot metal to rise along the ingot centerline (thermal buoyancy stirring) versus the tendency of electromagnetic fields (Lorentz stirring) to drive
The solidification structure of a VAR ingot is a function of the local solidification rate and the temperature gradient at the liquid-solid interface. To achieve a directed dendritic primary structure, a relatively high temperature gradient at the solidification front must be maintained during the entire remelting process. The growth
Vacuum Arc Remelting / 135 direction of the cellular dendrites conforms to the direction of the temperature gradient, that is, the direction of the heat flow at the moment of solidification at the solidification front. The direction of the heat flow (and thus the growth direction of the dendrites) is always perpendicular to the solidification front (Fig. 4). The edge of a VAR ingot also is the first metal to solidify. It is thus alloy-lean, as dictated by the phase rule. It is present as a skin around VAR ingots and is known as shelf. The shelf, in addition to being alloy-lean, contains both the oxides and nitrides collected from the melt pool surface as well as portions of the vapor deposits and splash BBs that have been deposited on the crucible wall above the advancing molten pool. If not adequately removed, this shelf material carries over from the ingot. As pool depth increases with the remelting rate, the growth angle of the dendrites, with respect to the ingot axis, also increases. In extreme cases, the growth of the directed dendrites can come to a stop. The ingot core then solidifies nondirectionally, for example, in equiaxed grains, leading to segregation and microshrinkage. Even in the case of directional solidification, microsegregation increases with dendrite arm spacing. A solidification structure with dendrites parallel to the ingot axis yields optimal results. However, a good ingot surface requires a certain level of energy input, resulting in respective remelting rates. Optimal melt rates and energy inputs depend on ingot diameter and material grade, which means that the necessary low remelting rates for large-diameter ingots cannot always be maintained to achieve axis-parallel crystallization. In spite of directional solidification, defects do occur. This can lead to rejection of the ingot, particularly in the case of superalloys. For the most part, these defects also occur in discrete regions, as opposed to continuous defects. Solute-lean (negative segregation) defects are not inherently as detrimental to the properties of a component as are defects resulting from positive segregation. Types of ingot defects include: Tree ring patterns Freckles White spots
Tree ring patterns can be identified in a macroetched transverse section as light-etching rings. They usually represent a negative crystal segregation. Tree ring patterns seem to have little effect on material properties. Freckles and white spots have a much greater effect on material properties as compared to tree ring patterns. Both defects can represent a significant cause for premature failure of turbine disks in aircraft engines. Tree ring patterns are the result of a wide fluctuation in the remelting rate. In modern remelting plants, however, the remelting rate is maintained at the desired value by means of an exact control of the melting rate during
operation, so that the melting rate exhibits no significant fluctuations unless caused by electrode defects. Freckles are dark-etching circular or nearly circular spots that are generally rich in carbides or carbide-forming elements. The formation of freckles is usually a result of a high pool depth or movement of the rotating liquid pool. The liquid pool can be set into motion (rotation) by stray magnetic fields during remelting. Freckles can be avoided by maintaining a low pool depth and by avoiding disturbing magnetic fields through the use of a coaxial current supply. White spots are typical defects in VAR ingots. They are recognizable as light-etched spots on a macroetched surface. There are several mechanisms that could account for the formation of white spots: Residues of unmelted dendrites of the con-
sumable electrode
Pieces of the ingot crown that fall into the
pool and are not dissolved or melted
Pieces of the shelf region transported into
the solidifying interface All three of the aforementioned mechanisms, individually or combined, can be considered as sources of white spots. This indicates that white spots cannot be completely avoided during VAR. To minimize the frequency of occurrence of these defects, the following conditions should be observed during remelting: Use the maximum acceptable metal rate
permitted by the ingot macrostructure.
Use a short arc gap to minimize crown for-
mation and to maximize arc stability.
Use a homogeneous electrode substantially
free of cavities and cracks.
Use a proper melting power supply.
Under ideal conditions, the arc is distributed uniformly across the surface of the molten pool. (This is called a diffuse arc.) However, the introduction of conductive ionic species into the melt or increases in the resistance (arc gap) in the system may cause the arc to become locally concentrated (constricted arc). Constricted arcs do not dwell in one place on the molten pool surface but move about the surface. When constricted arcs undercut the crown/ shelf, small pieces of the shelf may fall into the molten pool. These pieces move by gravity down the pool profile. They do not readily remelt into the pool, because they have a higher melting point than the pool in general. Ultimately, they are not dissolved and are incorporated into the ingot as regions of low solute content, containing stringers of oxide and nitride. When detected in the final product, they are seen as light-etching defects containing stringers of oxide and nitride and are known popularly as dirty white spots. More correctly, these regions are called discrete white spots, in keeping with their nature as a distinct but isolated structure within the ingot matrix.
Theoretically, a piece of shelf, free of oxides and nitrides, may be undercut and form a discrete white spot that is not dirty. Often, because of the lack of solute, the grain size in discrete white spots will be larger than that of the matrix. When present, dirt stringers act as stress raisers. Thus, discrete white spots, which have a high probability of being largegrained and dirty, are acknowledged to be detrimental to properties in superalloys. Unfortunately, although best practice for electrode preparation and for VAR melting may minimize the frequency of occurrence of discrete white spots, there is no way to guarantee their elimination in VAR products. The probability of their occurrence in the final component must be considered in the design of the component. Currently, there is no method of detecting a subsurface defect except to the extent that, combined with forging deformation, a crack may be generated that will be detectable by ultrasonic inspection. Macroetching of the surface of critical components is a common method for detection of discrete white spots and any other melt-related structure.
Process Variables The manufacture of homogeneous ingots with minimal segregation requires careful control of remelting parameters. The three major parameters defining the melt process are ingot and electrode diameter, arc gap, and melt rate. The choice of ingot and electrode diameter defines the clearance (annulus) between the crucible and the electrode. Insufficient clearance, either through a poor choice of annulus or through an off-center set-up, will lead to excessive current loss from the electrode to the crucible wall, rather than through the arc to the ingot. The choice of ingot diameter controls the amount of possible heat extraction and dictates the choice of melt rate so as to avoid positive segregation defects in the ingot. These choices, of course, are long-term choices dictated in the purchase of VIM molds and VAR crucibles. Control of the VAR process also includes: Pool depth and shape (control of solidifica-
tion structure)
Arc stability (control of drop-in defect
formation) As noted, the depth and angle of the solidification front (pool shape) needs to be maintained so that unacceptable solute-rich solidification structures do not occur. Pool shape and depth are controlled by regulating heat input and heat extraction. Of these, the melting current density has the largest influence on the melting bath geometry and conditions of solidification. Melt Rate. The melt rate is dependent on the melt current. Many producers choose to select a given melt amperage and hold it constant. Other producers use load cells built into the
136 / Melting and Remelting VAR furnace to measure electrode weight and thus the change in weight. The change in weight is used to calculate a melt rate, and the applied current is varied to attempt to maintain uniform melt rate. Individual melt rate measurements will vary widely depending on the time frame chosen for the measurement. Melt rates are generally calculated using rolling averages. A common averaging time is 20 min. Modern VAR furnaces are equipped with a load cell system to measure the weight of the electrode at a particular interval of time. Load cell design and maintenance are important factors in maintaining the quality of material melted in melt rate control, because load cells are subject to the generation of false signals. The actual values of the melt rate are compared by computer with the desired set values. Any difference between the measured melt rate and the desired value is eliminated by the proper accommodation of the power input. Figure 7 shows the melt rate and the melting current at startup, during steady-state melting, and during hot topping. Startup and hot topping are usually controlled based on time. The melting phase is controlled based on weight. Hot topping begins when a preselected residual electrode weight is reached. A computer controls the melting parameters, and melting rates can be controlled with a precision of better than þ2%. Heat input is determined by the melt rate of the electrode. Melt rate is often monitored and corrected over very short time periods by changes in melt current (melt rate control). Alternatively, control of the median melt rate over the duration of a melt may be maintained by operating at constant melt amperage (amperage control). Heat extraction is a function of the contact surface/ingot volume (ingot size) being melted. Larger ingots have poorer heat extraction (lower surface/volume) and must be melted at proportionately lower melt rates to prevent the formation of large liquid + solid zones and thus the formation of unacceptable positive segregation. While, for a given ingot size, high
Fig. 7
melt rate may promote positive segregation, a melt rate that is too low may cause the formation of visible solute-lean structures. The mechanism of formation of these solidification white spots is not completely understood. They differ from discrete white spots in that they are a solidification effect, not a shelf or electrode drop-in, and therefore do not contain oxide/ nitride stringers. Nevertheless, such solute-lean regions are potentially detrimental because they may make it difficult to control grain growth in subsequent processing. Arc Stability. Active control of arc stability is primarily by control of the arc gap throughout the melting of an electrode. The arc gap is the nominal distance of the melting electrode surface above the molten pool at the top of the solidifying ingot. Commercially useful gaps range from a minimum of approximately 2.5 mm (0.1 in.) to a maximum of 12 mm (0.5 in.). The arc gap may be considered as the resistance in this circuit. Because the melt current is held constant, any increase or decrease in the resistance of the circuit (change in arc gap) is seen as a change in voltage. Early VAR controls actually did measure these changes in voltage as a means for maintaining a uniform arc gap throughout the melt. Modern systems control the arc gap by measurement of drip short frequency (DSF). When a drip (drop) of molten metal is formed on the melting electrode surface, there is a point in time at which it is in contact with both the electrode and the molten pool. This causes a sudden decrease in voltage (an electrical short) across the arc gap. The duration of the short is measured in milliseconds. Control systems measure the number of drips of a given duration. It has been demonstrated that the frequency of drip shorts (usually in drips/min) is related to the width of the arc gap. The relationship is not linear over the whole range of melt conditions that may be experienced but does allow useful measurement and control of the arc gap in the ranges that are currently used
Process control parameters and set point functions in vacuum arc remelting
commercially. The higher the DSF, the smaller the arc gap. Making the gap smaller increases the duration of a short and thus increases the number of shorts counted. The relationship between arc gap and DSF also changes with changes in the melt current. At higher currents (higher melt rate), the drips are larger but less frequent. Thus, as in Fig. 8, raising the current requires a reduction in DSF to maintain equivalent arc gap. Drip short frequency is measured over a period of time, usually fractions of a minute. Individual DSF values vary widely. Thus, control signals for the ram drive (electrode feed speed) fluctuate. Normally, the ram is in continuous motion. The response to changes in DSF is to increase or to decrease the ram speed. Monitoring of the consistency of the arc gap may be accomplished by computing a rolling average or by simple observation of the width of the band generated in recording each individual DSF measurement. For melting of superalloys, arc gaps are generally controlled in the range of 6.5 to 13 mm (¼ to ½ in.). To maintain constant arc gap, the electrode is fed into the melt either faster or slower (ram drive) according to the fluctuation in voltage that occurs when the resistance of the system changes in response to changes in the arc gap. Modern control systems for superalloys generally do not use voltage for arc gap control but rather use a related voltage measurement: the DSF. A drip short is the transient downward voltage spike that occurs when a molten metal droplet is in contact with both the electrode and the molten pool. The number of shorts occurring in a given time period is proportional to the arc gap. The relationship for arc gap versus DSF is not melt rate independent. Passive control of arc stability is maintained by controlling the annulus and controlling the quality of the electrode being remelted. Electrodes for VIM generally contain porosity and low levels of residual refractory, both of which may act to destabilize the arc. Vacuum induction melting and pouring practice
Fig. 8
Drip short frequency as a function of arc gap and melt current. Source: Ref 1
Vacuum Arc Remelting / 137 is thus a significant factor in ensuring VAR arc stability and in minimizing the formation of drop-in-type defects. Vacuum Arc Remelting Control Anomalies. Variations in the traces of VAR parameters may indicate changes in melt conditions that will be detrimental to the quality of the solidification structure. Two of the more easily recognized problems are dead shorts and melt rate excursions (MREs). Dead shorts occur when the electrode is driven into the top of the molten pool, causing direct transfer of the melt current to the solidifying ingot. Such a problem may be caused by an unusually shaped primary shrinkage cavity in the electrode. If the normal ram drive programming cannot accommodate the unusual shape, the electrode may be driven into the pool. This will show on the recorded parameters as a sudden reduction of voltage to zero and an immediate ram back-out from the dead short situation. Drip short frequency, of course, would also go to zero while the electrode face was in contact with the pool. (This would not show as zero DSF, because DSF is measured over a time period.) Dead shorts will cause a major change in solidification characteristics in the molten pool in that region. Another cause of multiple dead shorts is the presence of high volumes of oxides and nitrides in the VAR electrode. If the oxide/nitride volume being melted onto the molten pool surface exceeds the capability of the process to sweep the “dirt” to the side and incorporate it into the ingot surface, then the ability to sustain an arc will be deteriorated. Melt rate will drop, and the ram travel, unable to react rapidly to sudden changes in arc gap, will drive the electrode into the pool, causing a short. This may happen repetitively as the ram backs out of the pool and then advances into it again. Transverse cracks in the electrode cause MREs, the most readily definable VAR chart anomaly. Electrodes may develop transverse internal cracks due to thermal stresses. This may happen in the cooling of the electrode, or, often, due to the thermal stresses generated as the VAR melt front moves up the electrode. When the melt front approaches a transverse crack, heat transfer across the crack is diminished, and the material below the crack becomes hotter than normal. Thus, the melt rate begins to increase. The ram speed, which responds to a decreasing DSF, does not respond fast enough to maintain arc gap, and thus, the arc gap increases.
Superalloy VAR The VAR process typically uses a cast (generally VIM or VIM-ESR) electrode. Elements such as bismuth and lead (which can be present in low concentrations even after VIM) are highly undesirable but are reduced to negligible levels by the VAR process. Magnesium, which is considered desirable for improved workability (and
is an addition element in VIM), is reduced in concentration but not completely removed. With the exception of high-vapor-pressure elements, the VIM chemistry is representative of the VAR chemistry. The flotation of oxides and nitrides to the melt surface in VAR improves the cleanliness and reduces the gas content of material processed through VAR. A common factor affecting consumable remelt quality for both ESR and VAR is the quality of the electrode. Electrode quality is affected by composition, cleanliness (low oxide and nitride content), and soundness (freedom from porosity and cracks). The most important characteristic of an electrode is composition. With the exception of reductions in volatile elements, VAR does not change the composition of the electrode. For high-nitrogen electrodes, some reduction of nitrogen is accomplished by flotation of nitrides. As the solubility levels of TiN in the alloy are approached, this mechanism is no longer effective. Electroslag remelting reduces sulfur content but may also cause minor changes in composition through reaction of titanium, aluminum, zirconium, and silicon with components of the slag. These changes are predictable in a mature process. Thus, there is no practical way to achieve composition modification once the master heat has been cast, but departure from normal practice in ESR may change titanium, aluminum, zirconium, or silicon content. Cleanliness. The cleanliness of the cast electrode with regard to entrained oxide and nitride is an important characteristic for VAR electrodes. The introduction of a high volume of oxide or nitride onto the molten pool in VAR may degrade the efficiency of melting of the arc as well as disturb the electrical characteristics measured to maintain control of the arc gap. This is not a problem in electrodes to be melted by ESR. Porosity. The porosity (secondary shrinkage) of an electrode may be of some concern in VAR. Vacuum arc remelting of electrodes with centerline porosity results in the preferred melting of the face in the porosity area. Thus, the melt face is no longer flat. The effect of this with regard to gap control during melting has not been quantified. A secondary problem connected with centerline porosity is the concern that projecting dendrites in this region may become detached from the electrode and fall, unmelted, into the molten pool. Then, they may be incorporated into the structure without being remelted. The composition of such a region thus will be that of a primary dendrite (alloy-lean) rather than representative of the bulk composition. Theoretically, these concerns may also apply to ESR processes, but the necessity for melted or unmelted material to pass through the slag makes ESR products inherently less sensitive to electrode porosity. Cracking. The final electrode characteristic affecting remelt quality is the freedom of the electrode from transverse cracking. In highly alloyed superalloys, the thermal stresses
generated upon cooling of electrodes or upon heating in the remelt process may be sufficient to form transverse cracks. When the melt front of either VAR or ESR approaches such a crack, the heat transfer up the electrode is diminished, because of the thermal barrier presented by the crack. Thus, the material in front of the crack becomes hotter than would occur under equilibrium conditions and will tend to melt at a faster rate. After the melt front has gone through the crack, it encounters cold material, and the melt rate will tend to drop. In both VAR and ESR, the result is that the controls respond with changes in applied power to try to keep the melt rate constant. This, however, is seldom accomplished without some discernable disruption to the continuous nature of the growth of the solidification front. The only defense against these MREs (also called events) is to produce an electrode free of cracks. In the future, it is hoped that more sophisticated melting controls will be able to detect the melt rate changes earlier and thus compensate for them more efficiently.
Titanium VAR Since the commercial introduction of titanium alloys in the 1950s, VAR has been the principal method for the production of titanium ingots. The process is capable of high melt rates with titanium, but steady-state solidification is not possible, and the molten metal pool is relatively deep. Vacuum arc remelting has many advantages, including high purity, good control, and reproducibility in the product. However, if high-density components of an electrode fall into the melt as solids, they can survive in the melt by rapidly dropping to the bottom of the molten pool and solidifying before they can be melted and homogenized in the solidifying ingot. An alternate to VAR is cold hearth melting. The electrodes for making titanium ingots are compacted aggregates (compact or briquettes) of sponge and alloy elements, including both master melt and elemental materials. For commercially pure titanium, the sponge is compacted, and the electrode is melted with appropriate oxygen control. When alloys are to be produced, granules of titanium sponge are mixed and then blended with appropriate alloy-containing material. The resultant alloy mixture is pressed into a compact or briquette. Most titanium and titanium alloy ingot is melted twice using the VAR process. The procedure is known as the double consumable-electrode VAR process. Double melting is considered necessary for all applications to ensure an acceptable degree of homogeneity in the resulting product. For certain critical applications, a third, or triple, melting step was specified at times, especially after the middle of the 1960s when defects were found in rotor-grade titanium alloys produced by double melting. Triple melting achieves even better
138 / Melting and Remelting uniformity of chemistry and structure. Triple melting also reduces oxygen-rich or nitrogen-rich inclusions in the microstructure to a very low level by providing an additional melting operation to dissolve them. In the two-stage process, titanium sponge, revert, and alloy additions are initially mechanically consolidated and then are melted together to form an ingot. Ingots from the first melt are used as the consumable electrodes for secondstage melting. Ingots from the second melt are used as the consumable electrodes for thirdstage melting when it is done. Processes other than consumable-electrode arc melting may be permitted in some instances for first-stage melting of ingot for noncritical applications.
Usually, all melting is done under vacuum, but in any event, the final stage of melting for noncritical applications must be done by the consumable-electrode vacuum-arc process.
“Vacuum Arc Remelting,” ALD Vacuum
Technologies, GmbH
REFERENCE ACKNOWLEDGMENT
1. M. Donachie and S. Donachie, Superalloys: A Technical Guide, ASM International, 2002
Portions of this article were adapted from: A. Choudhury and E. Weinga¨rtner, Vacuum
SELECTED REFERENCES
Arc Remelting (VAR), Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 406 M. Donachie and S. Donachie, Superalloys: A Technical Guide, ASM International, 2002
M. Blum, Vacuum Metallurgy, Vulkan, 2005 M. Donachie and S. Donachie, Superalloys:
A Technical Guide, ASM International, 2002
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 139-141 DOI: 10.1361/asmhba0005203
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Skull Melting Reviewed and revised by G. Keough, Consarc
SKULL MELTING refers to the use of furnaces with water-cooled crucibles that freeze a solid “skull” of material on the crucible wall. This technique is useful when melting highly reactive metals (e.g., titanium and zirconium) that would otherwise attack the linings of a crucible with refractories. The solidified metal creates a corrosion-resistant protective layer on the containment vessel surface. Two methods of skull melt casting have evolved. The first, vacuum arc skull melting, was originally developed in the early 1950s by the U.S. Bureau of Mines in Albany, Oregon (Ref 1, 2). Both methods are used to melt significant quantities of contamination free titanium for ingots or near net shape castings. Because the processes are used to melt reactive metals, they are normally carried out in a vacuum, but, in some cases, may also be done in an inert atmosphere. Vacuum arc skull melting was once the dominant method of melting titanium. Currently, other methods for melting titanium have become prevalent. Vacuum arc refining (VAR) furnaces with static crucibles are used in titanium primary melting. Recently, much of the melting of titanium sponge has been converted to plasma cold hearth (PCH) melting. Also, some of the largest rectangular ingots produced for rolling into sheet are now made on large multimegawatt electron beam (EB) furnaces. The total tonnage of material melted by VAR, PCH, and EB is significantly higher than that melted in vacuum arc skull caster furnaces. Vacuum arc skull melting works reasonably well for medium-to-large pour weights, typically in the range of 110 kg (250 lb) and up, but it suffers from a relatively large skull (typically 20% of metal melted). This lowers the thermal efficiency of the system and produces little superheat, so that precision and fine detail are relatively difficult to reproduce. The consumable reactive metal electrode also is expensive to produce, because the alloy shape and size must be preformed. An alternative technique is induction skull melting. Induction skull melting furnaces can melt a wide range of material forms including sponge (compacted), scrap, revert, or master alloy billet. The induction field also promotes vigorous stirring, which ensures uniform alloy composition throughout the melt
and increases superheat (compared to vacuum arc skull casting). Induction skull melting is also used in atomization of melts for powder production of reactive metals such as titanium.
Vacuum Arc Skull Melting Although the possibility of titanium skull melting was considered as early as 1948 and 1949, the first castings were made in 1953. In the late 1950s, this technology was applied by research institutes, which were looking for a practical means of liquefying and pouring uranium into graphite molds, for example, to produce uranium carbide. An early industrial vacuum arc skull melter was built in 1963 for the continuous production of uranium carbide. This furnace had a crucible volume of approximately 0.01 m3 (0.35 ft3) and used a nonconsumable graphite electrode to liquefy the uranium pellets fed into the crucible. The crucible tilting system was hydraulically driven. The molds were stationary. It was not until 1973 that one of the first skull melters for titanium went into operation; this furnace started production in 1974 in West Germany. State-of-the-art titanium vacuum arc skull melting furnaces are often equipped with turntable systems for centrifugal casting (up to 350 rpm). Casting weights of more than 1000 kg (2200 lb) are possible. Vacuum arc skull melting and casting is used for many titanium investment castings for aircraft, aerospace, medical, and chemical applications. Electron beam skull melting is also used for titanium alloys (see the article “Electron Beam Melting” in this Volume).
Furnaces Vacuum arc skull casting furnaces basically consist of a vacuumtight chamber in which an electrode is driven down into a water-cooled copper crucible. The direct current power supply provides the fusing current needed to strike an electric arc between the consumable electrode and the crucible. Because the crucible is water cooled, a solidified skull forms at the
crucible surface, thus avoiding direct contact between melt and crucible. The mechanism of skull formation is similar to the slag splashing on a refractory lining in the steel industry or the formation of a “frozen ledge” on refractories in the aluminum industry. Once the predetermined amount of liquid titanium is contained in the crucible, the electrode is retracted, and the crucible is tilted to pour the melt into the investment casting mold positioned below. For optimal mold filling, the mold can be preheated and/or rotated on a centrifugal turntable. Figure 1 shows the basic components of a vacuum arc skull melting furnace. Operating pressure is in the range of 1 Pa (0.0075 torr), and the specific working current ranges from approximately 1 kA/kg for small furnaces to approximately 0.2 kA/kg for large pouring weights. This batch-type skull melting furnace allows for cycle times of approximately 1 h for a full 50 kg (110 lb) pumping/melting/casting cycle. In principle, three consecutive pours can be obtained from one electrode. This furnace basically consists of a vacuum chamber, an arc voltage-controlled electrode drive system, a skull crucible, a centrifugal casting system with stepless adjustable turntable speed, an automatically sequenced vacuum pump system, a power supply, and an electrical control system with control desk. The cylindrical vacuum chamber is equipped with two large dished doors that support the crucible with the tilting mechanism, the mold platform with the centrifugal casting system, and the casting tundish with its cover. The crucible support system with an additional detachable device is also used for electrode loading. The chamber is jacketed for water cooling in regions that are subject to heat radiation. A top flange with a throat carries the electrode chamber and the electrode feeding system. Viewing ports allow for video monitoring of the melting and pouring. A vacuum pumping port is located in the cylindrical portion of the chamber. Figure 2 shows a semicontinuously operated vacuum arc skull melter for charge weights of up to 1000 kg (2200 lb). The principal difference between this furnace and the smaller model (Fig. 1)—apart from capacity-related
140 / Melting and Remelting
1
2 3 4
5
layout features—is the rectangular vacuum chamber. Again, the crucible and the tilting mechanism are carried by a dished door. In this furnace design, the centrifugal casting system is introduced from the bottom of the chamber to allow the horizontal connection of a separate cooling chamber, if desired. Molds are loaded from the back side into the chamber and can be discharged through a front door, which also allows the use of a cooling and charging chamber with lock valves for continuous mold transport flow. Modern vacuum arc skull melting furnaces are usually equipped with: Coaxial power feed directly to the skull
6 7
8 9 10
Fig. 1
Schematic of a vacuum arc skull melting and casting furnace. 1, fast retraction system; 2, power cables; 3, electrode feeder ram; 4, power supplies; 5, consumable electrode; 6, skull crucible (50 kg, or 110 lb); 7, tundish shield; 8, mold arrangement; 9, centrifugal casting system; 10, chamber lid carriage
crucible to avoid electromagnetic fields that can disturb the melt bath Programmable control systems for crucible tilting to allow repeatable pouring profiles for consistent parameters Highly accurate electrode weighing system for precise determination of pouring weights X-Y adjustment system for coaxial positioning of the electrode in the skull crucible Compact air-cooled power-supply modules of high capacity for achieving high melt rates, thinner skulls, and correspondingly increased yields above 80% Proven vacuum pumping and measuring systems Forced argon cooling systems for faster mold cooling
Induction Skull Melting 1 2
3 4
5
6
7
8
9 10
11
Fig. 2
Schematic of a semicontinuously operating vacuum arc skull melter for charge weights of up to 1000 kg (2200 lb). 1, fast retraction system; 2, power cables; 3, power supplies; 4, electrode feeder ram; 5, consumable electrode; 6, skull crucible; 7, crucible carriage; 8, tundish shield; 9, mold arrangement; 10, vacuum pumping system; 11, centrifugal casting system
In the 1970s, the Bureau of Mines developed an induction melting process for reactive metals known as the induction slag melting process. However, the drawback of that process was that a slag based on calcium fluoride was used to inhibit electrical shorting between segments of the crucible. Later, the Duriron Company further developed the process and, in particular, the crucible design to allow the slag-free melting of highly reactive metals and alloys in what is now referred to as the induction skull melting (ISM) furnace. Induction skull melting has evolved into a process with unique applications. As noted, ISM furnaces have some important advantages compared to vacuum arc skull furnaces. Induction skull furnaces can be charged with various material forms (scrap, revert, billet), and custom alloys can be alloyed in the ISM crucible because the induction field promotes vigorous stirring of the melt. A combination of stirring and levitation ensures uniform alloy composition throughout the melt and more uniform heating and increased superheat (compared to vacuum arc). As an indication of the flexibility of the ISM process, it is claimed that over 3000 alloys have been melted in this process. Part of this flexibility comes from the fact that the process is able to operate at pressure levels from vacuum through 101 kPa (1 atm). The process
Skull Melting / 141
Fig. 3
A 10 kg (22 lb) induction skull melting/casting unit in a vacuum/inert gas chamber
ac coil ISM crucible
Top dc coil
Water in Water out Bottom dc coil Pole piece (a)
ac coil
ISM crucible
Top dc coil
Bottom dc coil Water in Water out Pole piece (b)
Fig. 4
Cross section of induction skull melting (ISM) crucible with two different configurations of alternating current (ac) and direct current (dc) coils used in trials. Source: Ref 3
is also very rapid. A 10 kg (22 lb) charge, for example, can be loaded, melted, cast, and prepared for the next charge on a 15 min cycle. The key design elements of a reliable ISM crucible include effective water cooling and a segmented crucible that allows the field from the induction coil to couple to the load. Modern ISM crucibles use narrow gap technology to minimize the risk of liquid metal penetrating in between the segments and causing a short circuit. Each of the segments is intensively water cooled. A well-designed crucible will typically last 15,000 to 18,000 heats before extensive repair or replacement is required. Figure 3 shows a 10 kg (22 lb) induction skull furnace mounted in a vacuum/inert gas chamber. The chamber and vacuum or inert gas backfill is essential to avoid excessive oxidation of the reactive alloy elements, particularly titanium. Directly below the furnace is a centrifuge table that the molds are mounted on so that the molds can be centrifuged during filling. Skull weights in the range of 5 to 10% of the charge weight can be achieved in a welldesigned ISM furnace. Subsequent charges of the same alloy can be charged into the skull of a prior melt, thus minimizing the material requirements. At the end of a run for a particular alloy, the skull can be removed, identified, and saved for use on a future run. To optimize the efficiency of the process, a relatively high power together with a corresponding frequency are used to achieve both rapid melting and effective stirring of the liquid metal. The individual segments are intensively cooled to protect the copper from erosion and to minimize the risk of contamination of the melt. The combination of high power and high frequency also achieve higher superheat in the melt. This allows the process to be used for smaller pour weights (down to approximately 1 kg, or 2 lb) and for high-precision castings. Typically, the ISM process is recommended for pour weights in the range of 1 to 100 kg
(2 to 220 lb), although further extension of the range for higher pour weights is being developed. Of particular interest is a relatively new development using a combination of direct current electromagnetic fields (Fig. 4) with the traditional alternating current electromagnetic fields to achieve a significant further increase in available superheat of the molten charge (Ref 3). The ISM process also offers other advantages in casting. It can, of course, be used for conventional tilt-pour casting into both static molds and molds mounted on a centrifuge. However, unlike other methods, it can also be used for rollover casting as well as countergravity casting. Rollover casting, in particular, has been shown to be one of the least turbulent methods of skull casting, with promising results for casting complex and thin-walled near-net shape parts. ACKNOWLEDGMENTS Updated from: F. Mu¨ller and E. Weinga¨rtner, Vacuum Arc Skull Melting and Casting, Volume 15, Casting, Metals Handbook 9th Edition, 1988. REFERENCES 1. R.A. Beahl, F.W. Wood, J.O. Borg, and H.L. Gilbert, “Production of Titanium Castings,” Report 5265, U.S. Bureau of Mines, Aug 1956, p 42 2. W.J. Kroll, C.T. Anderson, and H.L. Gilbert, A New Graphite Resistor Vacuum Furnace and Its Application in Melting Zirconium, Trans. AIME, Vol 175, 1948, p 766–773 3. R.A. Harding, M.Wickins, G. Keough, R.J. Roberts, K. Pericleous, and V. Bojarevics, “The Use of Combined DC and AC Fields to Increase Superheat in an Induction Skull Melting Furnace,” LMPC Conference 2005 (Santa Fe, NM), ASM International
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 142-148 DOI: 10.1361/asmhba0005204
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Electron Beam Melting Revised by David W. Tripp, TIMET, USA
ELECTRON BEAM MELTING AND CASTING TECHNOLOGY is accepted worldwide for the production of niobium and tantalum ingots weighing up to 2500 kg (5500 lb) in furnaces with electron beams of 200 to 1500 kW. Other applications in Germany and former Soviet bloc countries are the production of steel ingots weighing 3.3 to 18 Mg (3.6 to 20 tons) using electron beams of up to 1200 kW. Furnaces of up to 4200 kW in electron beam power have been used since 1982 for recycling titanium scrap to produce 16 Mg (17.5 ton) slabs 1420 mm (56 in.) wide. Furnaces of 200 to 1200 kW are used to refine nickel-base superalloys. Other metals, such as vanadium and hafnium, are melted and refined in furnaces between 60 and 260 kW. Many furnaces with melting powers ranging from 20 to 300 kW are in operation in research facilities. These furnaces are used in the development of new grades and purities of conventional and exotic metals and alloys, for example, uranium, copper, precious metals, rare earth alloys, intermetallic materials, and ceramics. The total power
of industrial installed electron beam melting and casting furnaces worldwide was approximately 35,000 kW at the end of 2007. Electron beam melting and casting includes melting, refining, and conversion processes for metals and alloys. In electron beam melting, the feedstock is melted by impinging highenergy electrons. Electron beam refining takes place in vacuum in the pool of a water-cooled copper crucible, ladle, trough, or hearth. In electron beam refining, the material solidifies in a water-cooled continuous casting copper crucible or in an investment ceramic or graphite mold. This technology can be used for all materials that do not sublimate in vacuum. Competing processes include sintering (for example, for refractory metals), vacuum arc melting and remelting (for reactive metals and superalloys), and electroslag melting and vacuum induction melting (for superalloys, specialty steels, and nonferrous metals). Some advantages and limitations of the competing vacuum processes are given in Table 1. Additional information on some of these processes is available in
the articles “Vacuum Arc Remelting,” “Electroslag Remelting,” and “Vacuum Induction Melting” in this Volume.
Process Description The characteristics of electron beam melting and casting technology are: The flexibility and controllability of the pro-
cess temperature, speed, and reaction
The use of a wide variety of feedstock materi-
als in terms of material quality, size, and shape
The different methods of material processing
available
Product quality, size, and quantity
Contamination of the product is avoided by melting in a controlled vacuum and in watercooled copper crucibles (Fig. 1). The energy efficiency of electron beam processing exceeds that of competing processes because of the control of the beam spot dwell
Table 1 Comparison of characteristics of electron beam melting and competing processes Sintering Metal
Advantages
Vacuum arc melting Limitations
Tungsten, molybdenum
Small grain size; most often used
Refining limited; small batches; high energy consumption
Tantalum, niobium
Small grain size; good workability
Same as above; rarely applied
Hafnium, vanadium
...
Same as above
Zirconium, titanium
...
Not used
Advantages
Limitations
Refining limited; costly Moderate grain size; electrode preparation; acceptable melting dangerous workability; large ingots; low energy consumption Alloying; moderate grain Refining limited; expensive size; large ingots; low electrode; melting energy consumption dangerous Alloying during Almost no refining; costly remelting electrode preparation; melting dangerous
Electron beam melting Advantages
Limitations
Highest possible purity; economical feedstock preparation; large ingots; low energy consumption
Large grain size; brittle product; very rarely applied
Same as above; most frequently used
Alloying limited
Good refining; economical feedstock preparation and ingot production; most often used Economical feedstock Very low contamination; Limited refining; expensive feedstock preparation; only preparation; refining of wide range of alloying round ingots high-density inclusions; possible; large ingots melting of slabs, ingots, low energy and rods; high production consumption; rate; low energy economical melting consumption
High melting costs
Alloying limited; material losses from splatter; high furnace investment
Electron Beam Melting / 143 Large power range of 0 to 1200 kW Long free beam path of 250 to 1500 mm (10
to 60 in.) and the adjustable beam power distribution Beam deflection angle of þ45 and spot frequency up to 500 Hz Usable vacuum pressure range between 1 and 0.0001 Pa (102 and 106 mbar, or 7.5 103 and 7.5 107 torr) Reliability of the gun and cathode system The power control and distribution system allows for very accurate distribution of beam power and energy for achieving the required heating for material melting, superheating, refining, and electrothermal effects.
Drip Melting
Fig. 1
Schematic of the electron beam melting process
Fig. 2
Examples of electron beam melting and casting processes. (a) Button melting with controlled solidification for quantitative determination of low-density inclusions. (b) Consolidation of raw material, chips, and solid scrap to consumable electrodes for vacuum arc or electron beam remelting. (c) Drip melting of horizontally or vertically fed feedstocks. (d) Cold hearth refining/melting. (e) Floating zone melting. (f) Investment casting. (g) Pelletizing (manufacture of pellets from scrap and other materials for scrap recycling). (h) Atomization and granulation of refractory and reactive metals
time and distribution at the areas to be melted or maintained as liquid. In addition, unnecessary heating of the ingot pool, as occurs in vacuum arc remelting, for example, is avoided. Power losses of the electron beam inside the gun and between the gun nozzle and the target are very small, but losses up to 40% of the beam power can be incurred because of beam reflection, radiation of the liquid metal, and heat conductivity of the water-cooled trough and crucible walls. From the large variety of electron beam melting and casting processes shown in Fig. 2, only the processes illustrated in (c) and (d) are related to processes used in foundry technology:
Drip melting process for the preparation of
refractory and reactive metal feedstock material for electron beam and vacuum arc remelting skull melting and casting Cold hearth melting process for the feedstock refining of superalloys for vaccum induction melting and electron beam casting Electron Beam Heat Source Specifications. For all electron beam melting and casting processes, except for the crucible-free floating zone melting process, Pierce-type electron beam guns with separately evacuated beam generating and prefocusing chambers are used. The essential features of these guns are:
The drip melting processes (Fig. 3c and d) are primarily used for the production of clean, mostly ductile ingots of refractory and reactive metals. The feedstock for the first melt (Fig. 3c) can be compacted sponge, granular, powder, or scrap, which may be presintered in a vacuum-heating furnace. In some cases, loose raw materials can be consolidated in a watercooled copper trough (Fig. 3a). The consolidated ingot can then be fed horizontally for drip melting. Raw material that is continuously consolidated in a water-cooled copper crucible with a retractable bottom plate (Fig. 3b) can be fed horizontally or vertically for drip melting. In both consolidation processes, only 20 to 80% of the material is melted. Refining and losses of material by splattering are negligible. Drip melting of horizontally fed compacts is the most frequently used process for the production of ingots from refractory or reactive metals. The resulting ingot is of sufficient purity but has an area of inhomogeneity caused by the shadow of the horizontally fed bar. Two or more electron guns are used in drip melting to make use of reflected electron beams and to reduce evaporation and splattering. The end of a compact is welded to the front of the following one to avoid dropping semisolid material into the pool. Table 2 lists processing parameters that have been successfully used to electron beam melt various reactive and refractory metals. Vertical Feeding of Ingots. Refractory and reactive metal ingots of high purity, homogeneity, and smooth surface are remelted by vertical feeding (Fig. 3d). The molten metal droplets run down the conical, rotating electrode tip, are refined, and then drop into the pool center. The crucible pool is normally of the same diameter as the electrode, but it is not necessary. The crucible pool is kept in the liquid state to allow final refining and to guarantee ingot homogeneity. Because two or more electron guns are used, the entire pool can be equally bombarded; thus, shadow effects of the electrode can be eliminated. Simultaneous melting of horizontally and vertically fed electrodes (Fig. 3e) can be used
144 / Melting and Remelting
Fig. 3
Schematics of electron beam consolidation and drip melting processes. (a) Consolidation of coarse and solid scrap. (b) Continuous consolidation of raw material, chips, and solid scrap by direct feeding into a continuous casting crucible. (c) Drip melting of horizontally fed compacts, sintered bars, or consolidates for initial melting of reactive and refractory metals. (d) Drip melting of vertically fed vacuum induction melted or conventionally melted electrodes. (e) Drip melting of horizontally and vertically fed materials for the production of alloys from feedstocks with very different melting points
Table 2 Melting and refining data of refractory and reactive metals and alloy steels gained in laboratory and pilot production furnaces
Metal
Feedstock size, mm (in.)
Tungsten
40 (1.6) diam
Tantalum
55 (2.2) diam, 100 (4) long 60 (2.4) diam
60 (2.4) square, 160 (6.3) long Molybdenum 100 (4) diam 100 (4) diam 120 (4.7) square, 180 (7.1) long Niobium 80 (3.2) square, 150 (6) long 120 (4.7) square, 180 (7.1) long Hafnium 60 (2.4) square 80 (3.2) diam 100 (4) diam 100 (4) diam Zirconium 60 (2.4) square, 100 (4) long 60 (2.4) square, 150 (6) long Vanadium 60 (2.4) diam, 80 (3.2) long Titanium 100 (4) diam Ti-6Al-4V ... Ti-8Al100 (4) diam 1Mo-1V 4340 steel 80 (3.2) diam
Ingot diameter, mm (in.)
Ingot weight, kg (lb)
Integral melting rate, kg/h (lb/h)
Electron beam power of second melt, kW
... 60 (2.4) ... 115 (4.5) ... 80 (3.2) ... 160 (6.3) ... 100 (4) ... 180 (7.1) ... 150 (6) ... 180 (7.1) ... 80 (3.2) ... 130 (5.1) ... 100 (4) ... 150 (6) ... 80 (3.2) ... 100 (4) ... 100 (4) ... 150 (6) 150 (6)
... 37 (80) ... 200 (440) ... 65 (145) ... 523 (1150) ... 64 (140) ... 408 (900) ... 227 (500) ... 326 (720) ... 40 (90) ... 173 (380) ... 19.5 (43) ... 179 (395) ... 7.7 (17) ... 28.2 (62) ... 28.2 (62) ... 31 (70) 31 (70)
... 10.5 (23) ... 20 (44) ... 16.7 (37) ... 38.4 (85) ... 12.5 (27.5) ... 50.2 (111) ... 17.6 (39) ... 13.2 (29) ... 1.7 (3.7) ... 7.5 (16.5) ... 14.3 (31.5) ... 73 (161) ... 2.8 (6.2) ... 45.2 (100) ... 22.5 (50) ... 10 (22) 80 (176)
... 119 ... 300 ... 130 ... 371 ... 130 ... 290 ... 240 ... 218 ... 80 ... 110 ... 80 ... 250 ... 80 ... 87 ... 60 ... 52 80
Operating vacuum pressure of last melt, Pa (torr)
... 8 103 (6 105) ... 2 102 (1.5 10 4) ... 8 103 (6 10 5) ... 3 103 (2 10 5) ... 8 104 (6 10 6) ... 103 (7.5 10 6) ... 102 (7.5 10 5) ... 5 103 (3.8 10 5) ... 5 103 (3.8 10 5) ... 4 103 (3 10 5) ... 8 103 (6 10 5) ... 2 102 (1.5 10 4) ... 8 103 (6 10 5) ... 0.8 (6 10 3) ... 0.4 (3 10 3) ... 4 103 (3 10 5) 2.0 (0.015)
Interstitial elements in feedstock and final ingot, ppm
Total specific melting energy, kW 7 h/kg
Material yield, %
C
O
N
... 10.3 ... 9.9 ... 6.0 ... 8.30 ... 10.4 ... 5.2 ... 12.3 ... 15.9 ... 38 ... 14.7 ... 4.6 ... 2.6 ... 3.1 ... 1.91 ... 2.66 ... 1.67 1.0
... 93.1 ... 90 ... 92 ... 91.8 ... 95 ... 96.8 ... 87 ... 96.8 ... 93.1 ... 93.5 ... 88.2 ... ... ... 91 ... 99 ... 98 ... 93 99
70 10 200 10 100 10 ... 6 170 12 200 10 160 12 80 6 ... ... 500 100 ... ... ... ... ... ... ... ... ... ... 3900 4300 3622
4100 8 5000 8 1200 75 650 15 810 10 750 12 5220 106 4500 111 1870 170 900 200 950 545 950 735 1045(a) 235(a) 2180 1850 890 730 63 2.4 10
30 5 50 7 140 30 13 13 50 10 60 11 554 60 330 52 95 25 100 50 95 30 95 48 210(a) 95(a) 100 80 70 50 100 26 78
H
10 1 20 1 10 1 10 2 10 2 10 2 35 5 40 8 2 1 2 1 30 3 30 9 16(a) 13(a) 30 16 ... ... 0.3 0.08 0.10
Hardness, HB
... 200 ... 210 ... 65 ... 69 ... 140 ... 140 ... 77 ... 66 ... 160 ... 170 ... 120 ... 125 ... 30–100 ... ... ... ... ... ... ...
(a) The reproducibility of the refining data could not be confirmed.
for the production of critical alloys. In this case, the feedstock should be of the desired purity. Horizontally Fed Ingots. In this process, the feedstock size is smaller than the pool diameter to minimize the shadow effect of the horizontally fed bar. In production units, feeding can be carried out from two opposite sides. Other Process Considerations. To ensure the production of clean, homogeneous metals and alloys in electron beam drip melting furnaces, various aspects of material processing and handling must be controlled. Key considerations include:
Dimensions and quality of the feedstock,
and the feeding system used Ingot cooling and unloading during melting of another ingot Passivation and removal of condensates from the melt chamber Planning of melt sequences to minimize the number of furnace cleanings required Routine preventive furnace maintenance to ensure reliability Material yield and energy consumption
Equipment for Drip Melting. The essential equipment groups required for drip melting—
melting furnaces, control systems, and power supply units—are all important for achieving optimum productivity. The melting furnace (Fig. 4) includes the electron beam gun as the heat source, material feeding and ingot withdrawal systems, a crucible for material solidification, and a vacuum system to maintain the low pressure. Process observation, both visually and with video systems, is possible through viewports. The melt chamber flanges are equipped with x-ray-absorbing steel boards, and interlocking systems prevent operation failures and accidents.
Electron Beam Melting / 145
Fig. 4
Single-gun furnace for horizontal drip melting
Table 3 Principal applications for vacuum-arc-remelted (VAR), electron-beam-melted (EB), and powder metallurgy (PM) reactive and refractory metal ingots Metal
Reactive metals, VAR and EB melting Hafnium Vanadium
Zirconium Titanium
Refractory metals, EB melting and PM Tungsten
Tantalum
Molybdenum
Niobium
Applications
Glow discharge tubes for the electronics industry; control rods and breakoff elements in submarine nuclear reactors Targets for high-deposition-rate sputtering processes in the electronics industry; breakoff elements, fixtures, and fasteners in nuclear reactors; standards for basic research; alloying element for certain high-purity alloys Getter material in tubes in the electronics industry; fuel claddings, fasteners, and fixtures for nuclear reactors Components for bleaching equipment and desalination plants in the chemical industry; superconductive wires; turbine engine disks, blades and housings, rain erosion boards, landing legs, wing frames, missile cladding, and fuel containers in the aircraft and aerospace industries; shape memory alloys; biomedical fixtures and implants; corrosion-resistant claddings Heating elements, punches and dies, and nonconsummable electrodes for arc melting and gas tungsten arc welding for metal processing equipment; targets for x-ray equipment and high-sputtering-rate devices such as very large-scale integrated circuits, cathodes and anodes for electronic vacuum tubes in the electronics industry; radiation shields in the nuclear industry; cladding and fasteners for missile and reentry vehicles Condensers, autoclaves, heat exchangers, armatures, and fittings for the chemical industry; electrolytic capacitors for the electronics industry; surgical implants; fasteners for aerospace applications Dies for conventional and isothermal forging equipment; electrodes for glass melting; targets for x-ray equipment; cladding and fasteners for missile and reentry vehicles Superconductive wire for energy transmission and large magnets for the electrical and electronics industries; heavy ion accelerators and radio frequency cavities for nuclear applications; components for aircraft and aerospace applications
The control system allows the adjustment and control of such operating process parameters as electron beam power, operating vacuum level, material feed rate, and ingot withdraw speed. The control system also records and logs the process data.
Power Supply Units. One or more high-voltage power supply units are needed to supply the electron beam guns with the required continuous voltage (30 to 50 kV). The beam power of each gun can be adjusted between zero and maximum power with an accuracy of þ2%.
Other Equipment. Large-production furnaces are equipped with lock-valve systems to allow simultaneous melting and unloading of ingots without breaking the vacuum in the melt chamber. Production is thus limited only when the condensate remaining in the melt chamber requires cleaning or when a different alloy is to be melted. Characteristics of Electron Beam DripMelted Metals. Electron beam melted and refined material is of the highest quality. The amount of interstitials present is very low, and trace elements of specific high vapor pressure can be reduced to very low values (Ref 1, 2). Tantalum and niobium ingots have smooth surfaces and are of sufficient ductility that they can be cold worked, and sheets and wires can be produced. Tungsten and molybdenum ingots are also of the highest possible purity, but the ingots are brittle because of the very large grain size and the concentration of impurities at grain boundaries. Hafnium. Electron-beam-melted hafnium is of higher ductility than the vacuum-arcremelted metal (Ref 3). The main application of electron-beam-melted hafnium is as control elements for submarine nuclear reactors. Vanadium is refined by electron beam drip melting. The alumino-thermically produced feedstock is drip melted in several steps. During this procedure, the ingot diameter is reduced at each step by approximately 30 to 40 mm (1.2 to 1.6 in.) to obtain an ingot 30 to 40 mm (1.2 to 1.6 in.) in diameter, regardless of the initial ingot diameter. The clean vanadium ingots are primarily used in nuclear reactor applications (Ref 4). Applications for electron-beam-melted refractory and reactive metals are listed in Table 3.
Cold Hearth Melting The cold hearth melting process (Fig. 5) was developed approximately 10 years after drip melting (Ref 5). Cold hearth melting is mainly used for refining and recycling reactive metal scrap, especially titanium turnings machined with high-density tungsten carbide tool tips (Ref 6) and for refining specialty steels and superalloys. Principles of Cold Hearth Melting. Cold hearth melting (Fig. 6) is the most flexible vacuum metallurgical melting process. It is a two-stage process. The first step (material feeding, melting, and refining) takes place in a water-cooled copper trough or hearth. In the second step, solidification occurs in one of several round, rectangular, or specially shaped water-cooled continuous copper crucibles. Both process steps are nearly independent from each other; they are linked only by the flow of the liquid metal stream. The major refining actions are carried out in the hearth, but some
146 / Melting and Remelting
Fig. 5
Schematic of the cold hearth melting process
postrefining takes place in the pool of the continuous casting crucible, similar to the drip melting of horizontally fed billets. Refinement in cold hearth melting occurs by vacuum distillation in the hearth pool, superheating, and stirring of the molten metal pool. Removal of Impurities. Most impurities with densities lower than that of the melt (for example, oxides in superalloys) can be segregated by flotation and formed into a slag raft. The raft is then held in place by either mechanical or electrothermal means. Impurities denser than the melt, such as tungsten carbide tool tips in titanium, are removed by sedimentation. Inclusions with densities such that efficient flotation or sedimentation does not occur can be partially removed by adhesion to the slag raft or by dissolution in the liquid pool. Hearth dimensions are based on the type and amount of refining required. For example, hearths for vacuum distillation should be nearly square and relatively deep to allow sufficient melt stirring. For flotation refining, the hearth should be
long and narrow. Hearths for titanium alloy scrap recycling can be relatively short if all the materials can be transported to the pool of the hearth rather than to the ingot pool. Longer hearths provide for longer residence times, improving the removal of impurities by dissolution. Feeding. Material feeding criteria include 100% homogeneous material transportation to avoid uncontrolled evaporation of alloying elements and correct feeding into or above the hearth pool. Horizontal feeding of compacted, premelted, or cast material is most often used. Loose scrap and raw material can be used. Feeding of liquid metal was used in one of the first cold hearth melting furnaces to produce ferritic steel in a vacuum induction furnace (Ref 7). Postrefining was carried out in a cascade of five hearths 1.5 m (60 in.) long and 1 m (40 in.) wide. Casting and Solidification. The criteria for material casting and solidification include the shape of the final product and the solidification rate required to avoid ingot tears or other
defects and to ensure a homogeneous ingot structure. The multiple casting of small ingots is sometimes used, especially when forging is impossible because of the brittleness of the solidified material (for example, MCrAly wearresistant coating alloys). The casting of round and rectangular ingots and slabs is common practice. Cold Hearth versus Drip Melting. Table 4 compares the essential features of drip melting and cold hearth melting. Generally, cold hearth melting is used for titanium and superalloys especially when flotation or sedimentation of inclusions is required. Drip melting is used for other refractory metals because of the high melting points and the resulting high heat losses to the water-cooled copper crucible. Depending on production quantity, double or triple drip melting may require less energy than a single cold hearth melt of some materials, such as niobium. Refining and Production Data. Data on cold hearth electron beam melting and refining in laboratory and pilot production furnaces are given in Table 5. The data demonstrate the effectiveness of the process in reducing impurities and interstitial elements. It can also be seen that the selective evaporation of chromium from superalloys can be controlled by the distribution of beam power at the trough pool and by controlling trough pool area and melt rate. The selective evaporation of aluminum from Ti6Al-4V alloy is much more difficult to control; additional aluminum must be used to compensate for the aluminum evaporated. Equipment for Cold Hearth Electron Beam Melting. The equipment required for cold hearth melting is different from that used in drip melting, mainly because of the hearth and the larger melting chamber. In addition, because of the materials often melted in the cold hearth process (superalloys and titanium alloys), additional instrumentation is often provided. This may include an ingot pool level control system, metal vapor and partial pressure analyzers, a two-color temperature control system, and a data logging system. Accurate beam power distribution is achieved in multiple-gun furnaces by microprocessor control, which allows the splitting of a single beam to many locations and the adjustment of dwell time at each location. The beam spot at each of the locations can be scanned to create spot patterns as required by the application. With such systems, the required heating can be achieved without unnecessary power consumption and evaporation of alloying elements. Process observation is accomplished with a video monitoring system. Samples can be obtained from both the trough pool and the ingot pool for nearly continuous control of material quality. Feeding systems for cold hearth furnaces must maintain homogeneity along the length of the feed material. The trough and crucible should be easily accessible for convenient maintenance, especially when different alloys are to be melted in the same furnace.
Electron Beam Melting / 147 Characteristics of Cold-Hearth-Melted Materials. Titanium ingots and slabs can be produced from titanium scrap contaminated with tungsten carbide tool tips. The electronbeam-melted product will not contain any tungsten carbide particles. Oxygen content can be controlled by mixing titanium sponge with titanium revert material in the appropriate ratio. Cold-hearth-melted titanium ingots can be remelted in a vacuum arc remelting furnace. In some cases, depending on alloy and application, the product of the cold hearth furnace can be used without additional melt processing. Superalloys. Cold-hearth electron-beammelted superalloy ingots 150 to 200 mm (6 to 8 in.) in diameter are often used in vacuum induction melting investment casting furnaces. Such ingots are nearly free of nonmetallic inclusions and trace elements (Ref 9).
ACKNOWLEDGMENT Revised from article by W. Dietrich and H. Stephan, Electron Beam Melting and Casting, Metals Handbook, 9th Edition, Volume 15, 1988 pp 410–418
REFERENCES
Fig. 6
Four-gun 1200 kW combined electron beam drip melting and cold hearth melting furnace
Table 4
Comparison of the characteristics of drip melting and cold hearth melting
Characteristic
Power density Inclusions Ingot shape and structure Mass production Competitive economical processes Preferred method
Refractory metals
High Irrelevant Round; coarse grain Low Vacuum arc remelting Drip melting
Reactive metals, superalloys, and specialty steels
Soft; smoothly distributed Must be removed Round or flat; fine grain, segregation-free High Vacuum arc remelting; electroslag remelting Cold hearth melting
1. R.E. Lu¨ders, Tantalum Melting in a 800 kW EB Furnace, Proceedings of the Bakish Conference on Electron Beam Melting and Refining (Reno, NV), R. Bakish, Ed., 1983, p 230–244 2. J.A. Pierret and J.B. Lambert, Operation of Electron Beam Furnace for Melting Refractory Metals, Proceedings of the Bakish Conference on Electron Beam Melting and Refining (Reno, NV), R. Bakish, Ed., 1984, p 208–218 3. H. Sperner, Hafnium, Metallwis. Technik., Vol 7 (No. 16), 1962, p 679–682 4. R. Ha¨hn and J. Kru¨ger, Refining of Vanadium Aluminium Alloys to Vanadium 99.9% by EB Melting, Proceedings of the
148 / Melting and Remelting Table 5 Refining and production data for the melting of reactive and refractory metals and stainless steels in laboratory and pilot production furnaces
Metal
Feedstock size, mm (in.)
Hafnium
60 (2.4) square
Zirconium
100 (4) square
Zirconium
80 (3.2) square
Vanadium
50 (2) square
Ti-6Al-4V
Swarf
Ti-6Al-4V
Solid scrap
Ti-6Al-4V
125 (5) diam
Commercially 160 (6.3) pure titanium Commercially Sponge pure titanium Stainless steel 150 (6) diam Alloy 718
133 (5.2) diam
AISI type 316 stainless steel
150 (6) diam
Trough size, mm (in.)
Ingot size, mm (in.)
Ingot weight, kg (lb)
Melt rate, kg/h (lb/h)
Electron beam power, kW
120 250 (5 10) 120 300 (5 12) 120 300 (5 12) 120 300 (5 12) 120 300 (5 12) 120 300 (5 12) 150 400 (6 16) 150 250 (6 10)
100 (4) diam
83.0 (183)
40 (88)
180
150 (6)
90.5 (200)
42 (92.5)
185
100 (4)
40.2 (89)
80 (176)
140
Operating pressure, Pa (torr)
Specific melting energy, kW 7 h/kg
Composition of feedstock and product
C, ppm
4 10 2 (3 10 4) 3.5 10 2 (2.6 10 4) 3.5 10 2 (2.6 10 4) 1.5 10 2 (1.1 10 4) 2 10 2 (1.5 10 4) 7 10 2 (5.3 10 4) 6 10 2 (4.5 10 4) 6 10 2 (4.5 10 4)
4.5 4.4 1.75
O, ppm
N, ppm
H, ppm
Al, %
V, %
Cr, %
... ... ... ... ... ... ... ... 400 200 1520 ... ... ... ... ...
900 . . . 600 . . . 950 95 540 30 4000 800 1520 210 1045 210 277 50 2600 110 2700 110 1520 75 1320 76 ... ... ... ... ... ... ... ...
... ... 30 3 10 3 10 3 84 22 15 8 ... ... ... ...
... ... ... ... ... ... ... ... 6.0 4.4 6.0 4.8 6.0 3.6 ... ...
... ... ... ... ... ... 99 99 4.0 4.2 4.0 4.1 4.0 4.3 ... ...
... ... ... ... ... ... ... ... ... ... ... ... ... ... ... ...
100 (4)
...
20 (44)
130
150 (6)
62.6 (138)
40 (88)
122
150 (6)
62.6 (138)
70 (154)
140
2 75 (3) diam 100 400 (4 16)
2 32 (70.5)
91 (200)
147
96.4 (213)
86.3 (190)
148
150 500 (6 20)
100 400 (4 16)
103.0 (227)
41.2 (91)
226
8 10 2 (6 104)
5.5
...
...
...
...
...
...
...
150 400 (6 16) 150 400 (6 16) 150 400 (6 16)
2 75 (3) diam 2 75 (3) diam 3 65 (2.6) diam
2 55 (121)
136 (300)
144
1.06
2 57 (126)
136 (300)
156
3 41.5 (91.5)
136 (300)
156
6 10 2 (4.5 10 4) 6 10 2 (4.5 10 4) 6 10 2 (4.5 104)
701 536 417 363 ...
97 33 14 17 ...
155 68 52 34 ...
... ... ... ... ...
... ... ... 0.72 ...
... ... ... ... ...
18.25 18.11 19.11 18.73 ...
6.5 3.0 2.0 1.61 1.71
1.15 1.15
Source: Ref 8
Bakish Conference on Electron Beam Melting and Refining (Reno, NV), R. Bakish, Ed., 1986, p 53–67 5. C.d’A. Hunt and H.R. Smith, Electron Beam Processing of Molten Steel in Cold Hearth Furnace, J. Met., Vol 18, 1966, p 570–577 6. H.R. Harker, Electron Beam Melting of Titanium Scrap, Proceedings of the Bakish Conference on Electron Beam Melting and
Refining (Reno, NV), R. Bakish, Ed., 1983, p 187–190 7. C.d’A. Hunt, H.R. Smith, and B.C. Coad, The Combined Induction and Electron Beam Furnace for Steel Refining and Casting, Proceedings of the Vacuum Metallurgy Conference (Pittsburgh, PA), June 1969, p 1–22 8. H. Stephan, Production of Ingots and Cast Parts from Reactive Metals by Electron
Beam Melting and Casting, Proceedings of the Third Electron Beam Processing Seminar (Stratford, U.K.), 1974, p 1b1–1b69 9. M. Krehl and J.C. Lowe, Electron Beam Cold Hearth Refining for Superalloy Revert for Use in Foundry Production, Proceedings of the Bakish Conference on Electron Beam Melting and Refining (Reno, NV), R. Bakish, Ed., 1986, p 286–296
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 149-154 DOI: 10.1361/asmhba0005205
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Plasma Melting and Heating Reviewed and revised by Robin Lampson and Todd Telfer, Retech Systems
PLASMA MELTING is a material-processing technique in which the heat of a thermal plasma is used to melt a material. Plasma torches produce heat from an electrical current that passes through a partially ionized gas (or plasma). The plasma is generated by the phenomenon known as gaseous discharge, whereby a gas is ionized by electrical breakdown. The electrical breakdown of the gas can be accomplished by an applied voltage between two electrodes (as in the case of fluorescent lights). Plasmas also can be generated by various “electrodeless” methods that involve radio frequency (RF) discharges, microwaves, high-frequency induction coils, shock waves, lasers, or highenergy particle beams. Plasma arc torches have several advantages. Unlike heating by consumable electrodes (such as tungsten arc or graphite electrodes), plasma arc heating involves a gas medium instead of a consumable electrode. The gas environment and flow rates can be controlled, and only a small percentage of ionized gas (less than ½% of the total gas flow) is sufficient to create heating temperatures. Without a consumable electrode, plasma arc can generate higher temperatures than the electric arc methods. Products of combustion are eliminated or reduced. In addition, high inert gas pressure (5 kPa, or 50 mbar, or 37.5 torr, and higher) can be obtained with plasma torches. This is an important prerequisite for the avoidance of selective evaporation of alloying elements from the melt, an aspect that is particularly important in the case of titanium and superalloys. Plasma arc torches also supply electric heat to a ladle of molten metal to either maintain the temperature or raise the temperature of the melt. For numerous steel grades, especially those with extremely low carbon contents, plasma arc ladle heating is an alternative to heating with graphite electrodes and the risk of carbon pickup. Highpower plasma torches can be used with minimum modification to the existing ladle design. Prior to their integration into steel ladle metallurgy stations, plasma torches had been used primarily for scrap melting and heat-loss compensation before and during continuous casting. Only thermal (hot) plasmas are considered in this article. The term thermal differentiates
these plasmas from cold (nonequilibrium) plasmas, which are characterized by high electron temperatures and low sensible temperatures of the heavy particles. Cold plasmas are produced in various types of glow discharges, low-pressure RF discharges, and corona discharges. Cold plasmas are not suitable for melting processes. Thermal plasmas are suitable for melting in the production of ingots, slabs, castings, or powders.
Plasma Torches Breakdown (that is, the creation of ionic charge carriers) establishes a conducting path for resistance heating within a plasma torch. Most plasma generators (plasma torches) for melting processes use an electric arc to produce gaseous discharges. The characteristics of an electric arc include relatively high current densities, low cathode fall, and high luminosity of the column. A typical potential distribution along an arc is shown in Fig. 1. The column cross section at the cathode is smaller than at the anode, with contrary current and power densities. This fact influences the design principle of plasma torches. Most plasma heaters operate on direct current (dc) and are designed as either nontransferred or transferred arcs (Fig. 2). However, threephase alternating current plasma torches are employed in plasma heating and degassing applications. These units do not need a bottom electrode but do require more sophisticated power supplies than dc plasma torches do in order to limit the torch current and to ensure a quick, uninterrupted current pass through zero. In a nontransferred plasma torch, the gas is the only resistive element that generates heat. Nontransferred arcs are desirable for hot gas contact in shaft heaters, melters (cupolas), and reactors such as blast furnaces. This mode is suitable for the melting of conductive and nonconductive materials, but it has the disadvantage of lower efficiency compared to the transferred mode. Examples of nontransferred plasma torches are illustrated in Fig. 3. Magnetic coils are used to direct the flow of the plasma in the torch. For the electrodeless RF
torch, the plasma does not depend on the physical connection of electrodes. Transferred Arc Plasma Torches. In the transferred mode (Fig. 2b), one electrode is inside the torch, and the counterelectrode is the material to be melted. With the workpiece being one of electrodes, it provides the medium for resistive heating of both conductive and nonconductive materials. Most nonconductive materials become conductive when heated. Transferred arc heaters are useful for bulk heating of conducting solids and melts, where the gas phase is not a reactant but can be used as an inert shield. Figure 4 shows two typical design principles for torches in the transferred mode: the tungsten tip design and the hollow copper electrode design. In the tungsten tip design (Fig. 4a), the torch electrode is connected as the cathode because high current densities are required. The cathodes are usually made of tungsten, along with small additions of thoria to lower the thermionic work function of tungsten. Still, electron emission requires high electrode temperatures (3500 to 6000 K, or 5800 to 10,300 F) at the attachment of the arc. Therefore, the cathode material is liquid at the arc attachment point even with water cooling at the rear of the
Fig. 1
Typical potential distribution along a plasma arc. Va, anode voltage; Vc, cathode voltage; da, anode current density; dc, cathode current density
150 / Melting and Remelting
Fig. 2
Configurations of (a) Nontransferred arc plasma torch and (b) transferred arc plasma torch
Fig. 3
Examples of some typical nontransferred plasma torch configurations. (a) Electrode-based torches. (b) Radio frequency “electrodeless” plasma torch
cathode. For this reason, oxidizing gases generally cannot be used in direct contact with the tungsten tip. Argon shielding is one solution. Further, it should be kept in mind that portions of the small liquid pool at the cathode tip are ejected into the material being melted. The high melting point of tungsten can result in inclusions in the product; such defects are not tolerable in high-strength metals such as titanium alloys and superalloys for aircraft applications. With argon as the plasma gas, the service life of the cathode in the tungsten tip design is typically between 30 and 150 h. The inert gas consumption and plasma column stabilization of such torches are low. The hollow copper electrode design is shown in Fig. 4(b). The electrode material, usually copper alloyed with chromium or zirconium, is intensively water cooled. Copper is a field emitter of electrons (as opposed to tungsten, a thermionic emitter), because its boiling point is substantially below that required for thermionic electron emission. Current densities at the attachment points on cold cathodes are double those for thermionic emission on hot cathodes. Cathode spots are of high intensity, rapidly moving on the surface of the cold cathode and causing erosion by locally vaporizing copper. For this reason, the polarity is usually reversed in plasma torches with cold electrodes. The rear electrode acts as the anode, and the workpiece or melt is the cathode. The anodic arc attachment is smoother, and erosion is reduced. The plasma-forming gas is blown in tangentially between the rear electrode and the nozzle to create a swirl or vortex to stabilize the arc and to rotate the arc attachment in the rear electrode. With the movement of the arc attachment, a substantially lower power density can be achieved in the rear electrode. Electrode life of 1000 h has been demonstrated in torches using argon or helium as the plasma gas. All inert and reactive gases, or mixtures of these, can generally be used in such torches, but the influence of the gases on electrode erosion varies. The gas consumption of vortex-stabilized torches is higher than that of other designs. The efficiency of plasma torches is less affected by design principles than by process parameters and by the type of plasma gas used. Efficiency is defined as Pw/Pt, where Pt is the electrical power input of the torch, and Pw is the thermal power applied to the workpiece or melt. The efficiency of a 650 kW torch measured in a cold furnace ranges from approximately 25 to 50%, depending on the plasma gas used, furnace pressure, and the distance between torch and workpiece. The highest efficiency is obtained with helium as the plasma gas and a short distance between material and torch. The power loss in torches with hollow water-cooled electrodes is approximately 25%. Remaining additional 25% losses are due to radiation and convection of the plasma column between torch and workpiece. These losses are primarily a function of
Plasma Melting and Heating / 151
Melting Processes
Fig. 4
Design concepts for plasma arc torches in the transferred mode. (a) Torch with tungsten tip and concentric gas flow. (b) Torch with hollow copper electrode and vortex generator
the distance between torch and workpiece, but radiative losses of the plasma column also increase with pressure.
Furnace Equipment Plasma melting furnaces usually consist of double-wall, vacuum-tight, water-cooled constructions that can withstand the radiative and convective heat transfer from the plasma column. Furnaces also usually have material feeding systems. These can consist of vibratory or rotary feeders with bins (for bulk materials) or bar feeding systems. If the furnace is to operate continuously, the feed material must be “locked in” to prevent disturbance of the furnace atmosphere. The furnace must also have a vacuum pump so that it can be evacuated before backfilling. If the furnace is operated under reduced pressure, an offgas pump with an on-line gas control valve is required.
Atmosphere Control Plasma melting furnaces are usually operated under slightly positive pressure to prevent the potential atmospheric contamination by oxygen and nitrogen. However, state-of-the-art furnaces are vacuum-tight and can be operated at pressures between 5 and 200 kPa (50 and 2000
mbar, or 38 and 1500 torr). Vacuum-tightness is essential because the back diffusion of oxygen, nitrogen, and moisture through small leaks can be easily demonstrated even with positive pressure in the furnace. As noted, plasma torches are normally operated with argon or helium plasma gas. A typical gas purity of 99.999% indicates 10 ppm gaseous impurities and corresponds to an absolute pressure of 1 Pa (0.01 mbar, or 0.0075 torr) in a vacuum process. Additional sources of atmospheric contamination are the moisture and gases in the feed material, desorption of the furnace wall, and leaks in the furnace construction. In the case of metals with high affinities for oxygen or nitrogen, these gases are totally absorbed by the melt and will be found as oxide or nitride inclusions in the product. To obtain a clean initial furnace atmosphere, the vacuum-tight plasma furnace is evacuated with a mechanical pump to final pressure (typically approximately 1 Pa, or 0.01 mbar, or 0.0075 torr). The furnace walls can be degassed with the help of warm water running between the double walls of the furnace to enhance moisture desorption. The furnace is then backfilled with high-purity inert gas up to 1 kPa (10 mbar, or 7.5 torr), evacuated again to 1 Pa (0.01 mbar, or 0.0075 torr), and backfilled to operating pressure. This procedure dilutes residual gas impurities by a factor of 1000.
Plasma furnace equipment is of great interest when selecting a suitable heat source for melting, remelting, casting, and atomizing of reactive metals, titanium, and superalloys. Plasma torches, however, are the only nonconsumable heat sources for melting under high inert gas pressures in skull crucibles. High pressure is an essential requirement for preventing the selective evaporation of alloying elements that are characterized by high activity coefficients and/or vapor pressures, such as chromium and manganese in superalloys and aluminum in titanium alloys. Various plasma melting processes have been developed, including plasma consolidation, plasma arc remelting, plasma cold hearth melting, and plasma casting. Plasma cold hearth melting and plasma cold crucible casting are thought to have the most potential for industrial application. Table 1 lists some typical data on feed materials, operating parameters, and products of furnaces for some of these processes. With plasma furnaces, several companies are producing 3 to 5m (120 to 200 in) slabs with sections 50 to 89 cm (20 to 35 in.) thick and 100 to 150 cm (40 to 60 in.) wide. Plasma Consolidation. Consolidation implies that low-density feed material is converted into a high-density product. Plasma consolidation is used for titanium scrap and sponge to produce an electrode for further remelting in vacuum arc remelting (VAR) furnaces. The material is usually fed directly into a watercooled copper crucible and melted by plasma heat from one or several plasma torches (Fig. 5). The water-cooled copper baseplate of the crucible is continuously withdrawn. The shape and depth of the liquid pool in the crucible depend on feed rate and power input. Under these conditions, 100% melting cannot be ensured; some material can be trapped at the liquid/solid interface before melting. Figure 6 shows ingot density as a function of the raw material feed rate and power consumption in the plasma consolidation of titanium. Plasma consolidation processes can reduce manufacturing costs for the preparation of titanium VAR electrodes, which are usually plasma melted from pressed sponge or chip compacts. With plasma consolidation, metal chlorides can be removed from titanium sponge to sufficiently low levels. Tungsten carbide tool tips in titanium chips cannot be removed, and they eventually form high-density inclusions in the product. Therefore, the conventional x-ray inspection of scrap is still necessary. Low-density inclusions such as titanium nitride particles have not been found after plasma consolidation of contaminated titanium scrap and subsequent VAR remelting. Plasma consolidation has also been successfully applied to the processing of titanium aluminide (TiAl), a low-density alloy with excellent high-temperature service properties.
152 / Melting and Remelting Table 1 Characteristics and operating parameters of furnaces for plasma melting processes Type
Feed material
Plasma consolidation Titanium scrap and furnaces sponge Plasma consolidation Titanium scrap furnaces and sponge Plasma arc remelting Titanium, titanium furnaces alloys Plasma arc remelting Specialty steels, furnaces heat-resistant alloys
Plasma cold hearth melting furnaces Plasma cold hearth melting furnaces
Specialty steels, heat-resistant alloys Titanium scrap and sponge Titanium scrap and sponge, superalloys
Plasma cold crucible Titanium scrap melting and and sponge, casting furnaces superalloys
Fig. 5
Product form and size, mm (in.)
Ingot 355–430 (14–17) diam, 3000 (120) long; density >90% Ingot 700 (27.5) diam, 3800 (150) long; density 98% Ingot 125 (5) diam, 500 (20) long Ingot 150–200 (6–8) diam, 1200 (47) long; slabs 70 300 1200 (2.8 12 47) Ingots 650 (25.5) diam, 1500–2300 (60–90) long Slab 245 1125 2450 (9.6 44.3 96.5)
Specific power consumption, kW h/kg
Plasma gas
250–260 (550–570)
1.4–1.6
Argon
Atmospheric
Tungsten tip with external ignition bar; 6 90 kW
250–670 (550–1480)
0.89–2.4
Argon
Atmospheric
50–70 (110–155) 90–135 (200–300)
2.0–2.5
...
Atmospheric
1.5–2
...
Atmospheric
Hollow copper electrode with graphite liner; 600 kW Probably tungsten tip; 500 kW Probably tungsten tip; 500–1000 kW
400–900 (880–1985)
1.2–1.8
...
Atmospheric
Probably tungsten tip; 300–3600 kW
1–13 Pa (0.01–0.13 mbar)
Six hollow cathode torches; 700 kW on hearth; 1170 kW on crucible Hollow copper electrode, vortex stabilized; 200 kW on hearth, 100 kW on crucible Hollow copper electrode, vortex stabilized; 650 kW
Typical melt rate, kg/h (lb/h)
Typical furnace pressure
330 (730)
9
...
Ingot 195 (7.7) diam
90–227 (200–500)
1.3–3
Helium; gas recycling
Atmospheric
Casting ingots
380 (840)
1.3
Argon; gas recycling
50 kPa (500 mbar)
Plasma consolidation furnaces. (a) Multiple electrode. (b) Single electrode
Plasma Arc Remelting. An electrode that has already been melted by a primary melting process such as electric arc or induction melting can be plasma arc remelted into a water-cooled withdrawal crucible. The major objectives for plasma arc remelting are: To obtain directional solidification without
changing the chemical composition of the feed material To improve cleanliness by removal, size reduction, shape control, and redistribution of inclusions To lower gas impurity content To alloy with nitrogen Plasma arc remelting is used most frequently in the former Soviet Union countries to produce high-temperature alloys, bearing steels,
superalloys, and titanium alloys. Plasma arc remelting can be advantageous for the nitrogen alloying, deoxidation, or desulfurization of nonreactive alloys. An example of successful reactive melting is the nitrogen alloying of steel during plasma arc remelting with argon-nitrogen mixtures as the plasma gas. Figure 7 shows nitrogen concentrations in various metals as a function of nitrogen partial pressure in the gaseous phase after plasma arc remelting. The data indicate that excited molecules of nitrogen in the plasma react with the metal. In plasma cold hearth melting, bars, sponge, or scrap can be continuously fed into a water-cooled copper trough, usually called a hearth, and melted by means of plasma heating (Fig. 8). The liquid metal flows over the lip of the hearth into a withdrawal mold. The plasma cold hearth melting process is adapted from
Torch design principle and power
the electron-beam cold hearth refining process, which is a well-developed refining and casting method for superalloys. Plasma heat, instead of an electron beam, is employed to minimize selective evaporative losses that occur during electron-beam melting in the pressure range of 0.01 to 1 Pa (10 4 to 0.01 mbar, or 7.5 10 5 to 0.0075 torr). In the hearth, high-density inclusions from tungsten carbide tool tips will sink down to the skull. Low-density inclusions can be dissociated or dissolved in the base metal during interaction with the superheated melt and the plasma jet. Process optimization is necessary to prevent short-circuiting of the hearth by unmelted particles and to establish sufficient dwell time at high melting rates. The power distribution of one or more plasma torches required for smooth melting and fluid flow can be ensured by automated torch-moving devices (Fig. 9). In unconventional configurations, there is no mechanism provided for torch movement. Figure 8(b) shows a modern plasma cold hearth melting furnace concept with a doublewall, water-cooled, vacuum-tight furnace chamber and three plasma torches for the production of ingots 710 mm (28 in.) in diameter. Two torches serve the hearth, and the third the ingot mold. The torches use hollow water-cooled copper electrodes. The high gas consumption of this torch design means that an effective gas recycling system must be provided, especially if helium is used as the plasma gas. A key technology in the plasma melter is gas recycling systems that include reactive filters, oil-free high-capacity pump sets, and compressors. If necessary, the recycling system can be supplemented with cleaning systems for oxygen, carbon dioxide, nitrogen, hydrogen, and moisture. Helium gas recycling systems are
Plasma Melting and Heating / 153 now operating at 99.5% recycle efficiency (less than 0.5% of full gas flow). Plasma cold crucible casting is an alternative process to vacuum arc skull melting and casting or electron-beam skull melting and casting, and it can overcome some of the disadvantages of these processes. The disadvantages of vacuum arc skull melting in the melting and casting of titanium and superalloys are: Control of superheating is not possible;
increased power results in increased melting rate instead of superheating.
Expensive electrodes must be used instead
of the high-quality unconsolidated revert material that can be used for plasma cold crucible casting. For small batches of superalloys (up to 8 kg, or 18 lb), electron-beam skull casting is a well-established technique. Electron-beam melting and casting is rarely used for large castings of either superalloys or titanium alloys. This is because of the shallow melt pool in the crucible due to lack of stirring action and the potential for selective evaporation of alloying elements during melting.
The concept of plasma cold crucible melting is shown in Fig. 10(a). Solid material can be charged continuously into the water-cooled copper crucible, and additional material can be fed into the crucible during melting. The melt is poured into stationary molds or centrifugal casting molds situated in the mold chamber. The pouring stream can be superheated during casting or for production of atomized powder. An example of a cold crucible casting furnace (Fig. 10b) shows a horizontally installed mold chamber with the melt chamber on top of the mold chamber. In this example, a programmable torch-moving device is installed on top of the melt chamber. The pressure in the furnace during the melting operation can be preselected; pressures typically range from 30 to 50 kPa (300 to 500 mbar, or 225 to 375 torr). The melting operation can be started at every pressure in this range. The typical casting weight of this furnace is 27 kg (60 lb) of titanium with a 650 kW plasma torch. To achieve a deep molten pool, a high heat flow from the plasma torch to the melt must be established in combination with high stirring velocity in the melt to enhance the effective heat-transfer coefficient. The stirring action in the melt can be regulated by controlling current, gas flow, and torch-to-melt distance. Continuous feeding of the metal to be melted is required in order to achieve a deep pool. In contrast to vacuum arc skull melting, the crucible need not be inspected after each melt. The power densities in the plasma attachment area are adjusted to avoid damaging the crucible if it is inadvertently exposed to the plasma for short periods of time.
Fig. 6
Relationship among raw material feeding rate, ingot bulk density, and specific power consumption in the plasma consolidation of titanium. Source: Ref 1
Fig. 7
Nitrogen concentrations in various metals as a function of the partial pressure of nitrogen after plasma arc remelting. (a) Iron. (b) 16–25–6 stainless steel. (c) Austenitic stainless steel. Curves A and B, induction heating; minimum and maximum nitrogen concentrations, respectively. Curve C, plasma arc remelting. Source: Ref 2
154 / Melting and Remelting
Fig. 8
Plasma cold hearth melting. (a) Concept. (b) Configuration of system for recycling scrap
Fig. 9
Plasma torch-moving device. Courtesy of Leybold AG
REFERENCES 1. T. Yagima et al., Development of the Plasma Progressive Casting Process and Its Application for Titanium Melting, Titanium 1986: Products and Applications, Vol II, Proceedings of the Technical Program from the 1986 International Conference, Titanium Development Association, 1987, p 985–993 2. G.K. Bhat, Plasma Arc Remelting, Plasma Technology in Metallurgical Processing, J. Feinman, Ed., Iron and Steel Society, 1987, p 163–174
Fig. 10
Cold crucible plasma melting. (a) Melting, pouring, and superheating for powder atomization. (b) Industrialsized plasma cold crucible casting furnace with 650 kW torch power
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 155-159 DOI: 10.1361/asmhba0005198
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Crucible Furnaces George Giles, Morgan Molten Metal Systems
CRUCIBLES, in the form of well-fired clay pots, were present at the dawn of the metals industry, some 6000 years ago, and gave early metal workers the ability to melt metal in a durable container that could then be used to pour the molten material into a mold. Few, if any, metallurgical or industrial processes have a more distinguished history than the crucible melting of metals. Crucibles or earthenware smelting pots that were used to refine copper from King Solomon’s mines can still be found in that area of ancient Israel. Some pots had capacities of up to 0.40 m3 (14 ft3). The Wootz steels of India, the famous Damascus sword blades, and the equally famous steels of Toledo, Spain, were produced using crucible melting processes. At their most basic level of form and function, crucibles today (2008) are little changed from their early prototypes. However, in terms of materials, construction, thermal efficiency, and service life, modern 21st century crucibles (Fig. 1) are superior to those in use just a short time ago. Designed for a wide range of traditional applications and new processes, crucibles are formed using a wide variety of materials and processes that impart specific characteristics, capabilities, and limitations to each of the many types of crucibles available from crucible manufacturers. Crucibles may be as small as teacups or may hold several tons of metal.
Fig. 1
Large electric bale-out crucible furnace. Courtesy of Morgan Crucible
They may be fixed in place within a furnace structure or may be designed to be removed from the furnace for pouring at the end of each melt. Crucibles are used in gas- and oil-fired furnaces, electric resistance furnaces, and induction furnaces. They come with or without pouring spouts and in a wide variety of traditional and specialized shapes. There is no direct impingement of the flame on the metal, and the heat loss to the outside is limited by the refractory walls. However, the energy efficiency of the crucible furnace can be low (7 to 19%), with over 60% of the heat loss attributed to radiation (Ref 1). Without proper care, the life of the crucible furnace may be short. Temperature control may also be difficult.
Applications In the metals industry, most crucibles are used to melt, hold, and/or transfer nonferrous metals—commonly aluminum, zinc, and their alloys and copper and copper-base alloys—as well as precious metals. A crucible furnace is the least expensive melting method preferred for melting small volumes of nonferrous metals. It is also used for higher-temperature alloys, such as nickel bronze and cupronickel, and, to a lesser extent, for melting ferrous metals, such as gray iron. Crucible furnaces are popular in jobbing foundries and die casting shops because they are easy to tap and charge with another alloy. Foundry melting and handling operations with crucible equipment include furnace melting, pouring ladle equipment (either separately heated or as a ladle liner), and liners for electric induction melting furnaces. Crucible melting can also be used as a component of qualityadjustment equipment between melter and casting furnaces to alter the chemistry and temperature of the melt for control of such factors as porosity content prior to casting. Metal-processing operations include in-line filtration, fluxing, and hydrogen gas content control. Crucible Advantages. Crucible melting is a simple and flexible process. Crucible furnaces can generally be started or shut down at a moment’s notice. With their unique ability to
melt, hold, and transfer metal using a single vessel and to allow even incompatible alloy changes to be made simply by switching vessels, crucibles provide significant operational flexibility for various foundry applications. Even when fixed within the furnace structure, crucibles offer important advantages when compared to directly heated fuel-fired furnaces and to electric resistance or induction furnaces with rammed refractory linings. Advantages include: Small, compact size and relatively lower
acquisition cost are economical.
Smaller (and round) furnaces provide and
maintain good, even temperature profiles.
Open crucible furnaces, or even those with
pivoting lid covers (Fig. 2), allow easy access for ladling. Lower Metal Loss/Cleaner Metal. In fuelfired furnaces where the metal is heated directly by a flame produced by gas or oil burners, there are both high levels of metal loss and increased oxidation due to the combustion products generated. In fuel-fired crucible furnaces, however, the flame is isolated from the metal, and the metal is heated indirectly with heat conducted through the crucible. This indirect heating greatly reduces metal loss and metal oxidation, resulting in lower metal purchase costs, lower energy costs per pound of metal poured, and less slag or dross formation. Alloy Flexibility. Even when fixed in place in a furnace, the high-density interior surface of a crucible reduces the likelihood of cross contamination when alloy changes are made. Quick Replacement. Removing a worn or damaged crucible from a furnace and installing a new crucible is a faster, simpler, and less technically demanding process than removing and replacing a worn or damaged rammed refractory furnace lining. Also, because crucibles are preformed to exact material and dimensional specifications, there are few concerns about improperly compacted refractory or incomplete sintering. Matching Crucible and Application. Determining the required properties of an application is the most important step in selecting the proper crucible. Knowing the metal to be
156 / Melting and Remelting Glass: Glass is used to provide a semiperme-
able barrier to gases, controlling the ingress of oxygen and egress of combustion products. Typically, 10 to 20% of a crucible is glass. Aluminosilicates: Aluminosilicates are used to modify the refractory properties of the crucible or internal glass composition. Unfortunately, there is not a single crucible material recipe or manufacturing process that offers the highest level of every desirable characteristic for every application. Therefore, the best course is to prioritize the list of crucible properties most important for the application and review those requirements with the crucible provider to find the crucible product best suited for the specific needs.
Furnace Types
Fig. 2
Crucible furnace in die casting foundry. Courtesy of Morgan Crucible
melted, held, and/or transferred, however, is just the start in selecting the right crucible type for the application. The specific crucible properties that are important for a particular application must be considered. These properties include: Maximum temperature limits: It is important
that the highest temperature reached by the application does not exceed the maximum temperature limit for the crucible; otherwise, the crucible will fail. Tolerance of thermal shock: If the application results in rapid heating or cooling of the crucible, a crucible type must be selected that resists thermal shock. Impact resistance and mechanical strength: If the crucible in the application is used to transfer metal or if it is subject to the impact of dropped charge materials, a crucible with the strength and durability to survive these activities must be selected. Erosion/corrosion resistance: If the application produces slag, dross, or other by-products or uses fluxes that mechanically or chemically attack the inside surface of the crucible, then a crucible type must be selected that resists such attacks. Oxidation resistance: When exposed to air at high temperatures, graphite in a crucible will oxidize. If the application is likely to cause melt oxidation or excessively high temperatures (which exist in many foundry casting applications), a crucible should be selected with best oxidation resistance. Nonwettability: It is important in aluminum foundry application that the inside surface of the crucible resist the intrusion of molten metal into pores in the crucible material. Consequently, a crucible should be selected that possesses specific nonwetting agents.
Electrical properties: If the application
requires a crucible that heats inductively, select a crucible designed for use in an induction furnace and with a specific electrical resistivity that optimizes the coupling power of the crucible. Thermal conductivity: Not all crucibles offer the same level of heat transfer through the crucible and into the metal. To achieve the lowest energy costs, select the crucible with the highest level of thermal conductivity that meets all of the other requirements of the application. Crucibles have been called the first composites, and modern crucibles use an eclectic mix of materials to achieve the important characteristics listed here: Graphite: A foundry crucible contains 30 to
50% by weight of flake graphite, depending on the properties required. The carbon content of graphite imparts high thermal conductivity and nonwettability and, coupled with its platelike layered structure, provides high thermal shock resistance. This is critical to foundry applications where temperatures can change by several hundred degrees in seconds. Silicon carbide: A foundry crucible typically will contain 10 to 50% of silicon carbide, based on the desired characteristics. Silicon carbide provides excellent resistance to high-temperature erosion and corrosion and imparts a high level of thermal conductivity to the crucible. Silicon metal: Fine powders of silicon metal or alloys are used to increase strength and erosion resistance, to protect the graphite and carbon bonding matrix from oxygen, and to help ensure an even glaze coating.
There are many types of crucibles and shapes produced by crucible manufacturers for the foundry market. Most crucibles used in the metals industry for melting, holding, and/or transferring operations are preformed carbonbonded or ceramic-bonded vessels. Silicon carbide crucibles are a widely used type of carbonbonded crucibles, and clay-graphite crucibles are a good example of ceramic-bonded vessels. It is important to discern which type is appropriate for a given application. Table 1 lists the several shapes and the application that is generally appropriate for their use. The basic types of foundry crucible furnaces can be classified as stationary, tilting, or movable. Stationary furnaces use a crucible set inside a refractory-lined shell that has heating equipment attached. A lift-out crucible is shown in Fig. 3. Normally, molten metal is dipped out with hand ladles, or the crucible containing the molten metal is removed from the shell for the casting operation. By using tong and shank tools to lift and carry the bilge-shaped crucible to the mold, pouring of the casting is done directly from the crucible without the need for potentially damaging metal transfers. The maximum crucible size normally used in this manner holds 140 kg (300 lb) of aluminum or copper. The weight and the strength of the crucible limit this type of operation. Another variation is a pushout furnace (Fig. 4) with induction coils and mechanical lifts to raise and lower the crucibles within the coils. Some stationary crucible furnaces are equipped for hand-ladle dip-out pours. For most metals, it is imperative that the flame and flue gases do not contact the charge. Two methods of controlling this critical factor are illustrated in Fig. 5. The first is the side flue configuration used in an aluminum melter installation (Fig. 5a), and the second is the cover plate configuration used for a brass foundry furnace (Fig. 5b). The cover plate setup incorporates holes on the outside diameter of the crucible top. Also shown is the extensive use of highly insulating fiber refractory material. This material has a low heat
Crucible Furnaces / 157 Table 1
Application of specific crucible shapes
Shape
Advantages
Bilge
Reduced top diameter provides less metal surface area. This helps to conserve heat and cuts metal loss and oxidation. Provides large metal surface area to facilitate manual or automated baling out of metal. Wider lower profile provides more efficient heat transfer in electric resistance furnaces. This shape is also suitable for low-pressure furnaces where die casting stalks are used to feed a die above the metal bath. Provides large surface area to facilitate manual or automated baling out of metal. Narrow lower profile is advantageous in fuel-fired furnaces to provide space for burner flame to circulate. Large-capacity crucibles for bulk melting of metal. Metal is usually then transferred to a holding bale-out furnace. Parallel profile means crucible is equidistant from the induction coil, allowing even heating.
Basin
Bowl
Spouted basin
Cylinder
Furnace type
Application
Typical user
Lift-out (crucible is lifted out of furnace chamber for casting)
General melting work for small castings
“Jobbing” foundries
Bale-out (crucible remains fixed in furnace)
Holding and melt/holding
Die caster
Bale-out (crucible remains fixed in furnace)
Holding and melt/holding
Die caster
Fig. 3
Typical lift-out version of a stationary crucible furnace specifically adapted to the foundry melting of small quantities (<140 kg, or 300 lb) of copper alloys
Melting Basin tilting furnaces (crucible remains fixed in furnace; furnace is tilted to pour metal)
Die caster, ingot producer
Induction tilting furnaces Melting (crucible remains fixed in furnace)
Master alloy producer
Fig. 5
Stationary crucible furnaces equipped for handladle dip-out pours. (a) Side flue configuration. (b) Configuration with cover plate openings on outside diameter of crucible top
Fig. 4
Cross section of a double pushout stationary crucible furnace
content and serves to conserve fuel and to minimize the temperature override that often occurs with the heavier refractory furnace linings. The added refractory flame face is required to accommodate the higher melt temperature. Steel crucibles can also be used in stationary crucible furnaces, as shown in Fig. 6. The high
heat conductivity of the steel offers fuel and meltdown time savings. Steel and ductile iron crucibles are extensively used for magnesium and zinc melting. In these molten alloys, there is little attack of the crucible walls. On the other hand, aluminum alloys will attack the steel crucible walls by forming sludge
(combination of iron, manganese, and chromium), which eventually causes hot spots in the castings produced from this melt. However, there have been successful aluminum foundry uses of steel crucibles. This requires a thorough refractory coating on the crucible surfaces in contact with molten metal and careful cleaning of the rust that tends to accumulate in the firing chamber (for example, flame contact with the
158 / Melting and Remelting
Fig. 6
Cross section of a stationary fuel-fired furnace used for the open crucible melting of magnesium alloys
crucible causes rust and scaling). This can collect in the firing chamber. A special clean-out door is used to remove this residue. Both aluminum and magnesium can react explosively with iron oxide. The use of electrical resistance furnaces and steel crucibles reduces the seriousness of this problem, and long-term use of steel crucibles and resistance furnaces has been successful in the casting of aluminum alloys. Tilting furnaces incorporate a crucible that is supported in the shell such that both are free to rotate on an axis. This allows the molten metal to be poured into ladles or molds. The axis of the pour can be either at the center of gravity or at the lip of the furnace. This design simplifies the handling of molten metal in many foundries. Tilting crucible furnaces are designed to simplify metal removal and have been extensively used in numerous foundries for many years. Figure 7(a) illustrates a center-axis furnace; Fig. 7(b) shows a cross section of a lip-axis furnace. Tilting crucible furnaces are useful where castings require more metal than can be easily handled by stationary crucible furnaces. These units are more flexible and can easily handle production requirements that are too small for the efficient operation of larger hearth furnaces. Figure 8 illustrates the use of a crucible liner for an induction melting tilt furnace. Use of the crucible liner overcomes the problems that would occur as a result of the strong erosive currents existing at the crucible sidewalls. The low electrical conductivity of clay-graphite crucibles makes them ideal for use in such applications. A gradual buildup of molten metal in the heel and voids in the backup refractory, both caused by repeated tilting of the furnace, are detrimental to crucible life. Crucible liners, which provide a durable wear-resistant surface for molten metal, are often used in the brass, steel, and ductile iron foundry industries. This application for the foundry use of crucibles is similar to that for the induction furnace liners shown in Fig. 8.
Fig. 7
Two variations of a tilting crucible furnace. (a) Center axis. (b) Lip axis
Movable crucible furnaces are used in conjunction with automated pouring system methods. Methods include multiple crucible melters arranged on either a shuttle or a turntable, such that one furnace is moved into position for pouring while others are positioned for recharging, stabilizing, and metal treatment.
Design and Operation Proper Furnace Atmosphere. With few exceptions, a slightly oxidizing flame should be used. With aluminum alloys, this is a pale yellow color. For copper-base alloys, this will be tinged with green (due to the oxidation of the copper). A traditional test for the correct furnace atmosphere is to place a cold piece of zinc in the flame. A yellowish color indicates an oxidizing flame, and a black color indicates a reducing flame. The objective is to obtain a flame having a tinge of yellow, which indicates the presence of the desirable, slightly oxidizing atmosphere. Oxidizing conditions were maintained in old coke-fired crucible furnaces at the crucible walls by burning the coke in a 75 mm (3 in.) space around the crucible sidewall. The burning coke maintained a slightly oxidizing atmosphere at the walls and ensured good heat transfer. Burner Control. For efficient melting and optimal crucible life, it is necessary to have even heating of the crucible and to avoid hot and cold areas around the sides of the crucible bowl. Some furnace designers use two or four burners to minimize this problem. This will contribute to longer crucible life, better fuel economy, faster melting, and lower noise levels. Precise periodic adjustment of the burner(s) is essential to efficient crucible furnace operation.
Fig. 8
Cross section of a tilt furnace for the highfrequency induction melting of brass and bronze alloys. Crucible is of clay-graphite composition. Also shown are the locations of the molten metal buildup and the voids in the backup refractory, which shorten crucible life.
Special care is required to avoid overheating of the bath. Approximately 80 to 95% of the heat required to prepare a heat for the foundry is used to heat the charge to the melting temperature and then to convert the solid to a liquid. Very little heat is required to raise the temperature of the newly melted liquid to the required pouring temperature. This can cause problems. For rapid melting of the crucible contents, high heat inputs are normally used. As soon as the charge becomes entirely liquid, little or no additional heat is required. If the equipment or the operator does not immediately turn the heat down or off, the bath temperature can become far higher than required. For most metals, this causes quality-control problems. Currently, most foundries depend on their furnace operators to control this harmful variable. Burner Design. Because of the wide diversity in crucible applications using many different fuels, a summary of these data is well
Crucible Furnaces / 159 beyond the scope of this discussion. It should be noted that developments in the design and operation of burners have significantly improved this furnace component. Any new installations of significance probably should include recent developments in both burner designs and burner control units. Charging Techniques. The most efficient melting (in terms of heat transfer or efficient use of fuel) is achieved with faster melt times, and the highest metal quality is obtained when a heel of metal is used to start the melting. Here, the heat travels easily from the flame through the crucible wall and into the metal. A crucible that contains only a solid charge provides little contact between the crucible wall and the metal to be melted. As a result, there is a considerable time lapse between charging and the initial melting of the charge. For efficient melting, a good heel size is roughly 25% of the crucible capacity. A second point concerning charging is the need to keep a loose charge. When a cold charge is tightly jammed into the crucible, the expansion of this charge upon heating will often break or crack the crucible. The third point concerns the need for gentle charging of all heavy or sharp-cornered material. Crucibles are fragile and easily broken. Many crucibles have been broken by rough and careless charging with heavy and/or sharp ingot or scrap. Handling of Crucibles During Melting and Pouring. It is important to keep in mind that hot crucibles filled with molten metal are fragile and therefore deserve careful handling. Care should be taken to set them on sand or other nonconductive surfaces. Placing a hot crucible full of molten metal on a steel plate or damp floor can easily break or seriously weaken the bottom of the crucible. Crucible Care. With proper care, crucibles will provide long service life. However, like every system used in the processing of molten metal, crucibles must be handled properly and
must be operated within their design parameters. All crucibles may be damaged by: Excessive temperatures: Crucibles must not
be heated beyond their specified maximum temperatures. Excessive temperatures will cause crucible failure. Thermal shock: Structural damage may be caused by too rapid heating or cooling of the crucible. Uneven heating or frequent thermal cycling between cold and hot also can result in crucible damage, especially with frequent metal-level changes. Physical shock: Damage may be caused to the crucible by the impact of charge materials and/or, in the case of removable crucibles, by impacts during movement. Damage to crucible glazes: Crucible glazes serve as barriers against oxygen attack of the carbon material making up the crucible. These glazes may be damaged by any of the previously noted hazards, and damage to the glaze will lead to crucible failure. Chemical attack: Use of inappropriate or excessive metallurgical additives may lead to corrosion of crucible components, particularly at higher temperatures. Fluxing salts that contain fluorides are especially harmful.
Crucibles must be thoroughly inspected for shipping damage when received. If cracked, chipped, or otherwise damaged, they must not be installed. Crucibles should be stored inside in a warm and dry environment and must be completely dry before use. All clay-graphite crucibles must be properly annealed before being subjected to the first full heat. This will help the crucibles withstand thermal shock throughout their service life. All crucibles should be preheated gradually whenever entering use from a cold, that is, room-temperature, condition, following the crucible manufacturer’s specific recommendations. During operations, the crucible should be inspected
for damage after each heat, and slag or dross buildup inside the crucible must be carefully removed, avoiding damage to the crucible wall. Damaged or worn crucibles must be replaced immediately. Selecting the right crucible for the application; installing it properly; operating it within its design parameters; protecting it from physical, chemical, and thermal damage; and replacing it immediately when damaged or worn are best practices. Following these rules will enhance service life, operational efficiencies, and significant economies in crucible furnace operations. Cleaning of crucibles is impor tant in preventing premature failure. In many, if not most, metal-melting operations, slag or dross buildups develop at the metal surface line, and heavy, unmelted slush residue collects on the bottom. Both of these residues shorten crucible life by causing hot and weak wall areas. The poor and uneven heat transfer around the crucible walls causes hot spots and uneven stresses within the crucible walls. The metallic residues that will expand faster than the crucible upon heating can easily rupture the wall. At times, these residues expand and fracture the crucibles during the normal heating and cooling cycles. Cleanliness and careful workmanship on the melt room floor should be encouraged. ACKNOWLEDGMENT Portions of this article were from Casting, Volume 15, Metals Handbook, 9th ed., 1988. REFERENCE 1. Advanced Melting Technologies: Energy Saving Concepts and Opportunities for the Metal Casting Industry, BCS, Inc., Nov 2005
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 160-169 DOI: 10.1361/asmhba0005350
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Reverberatory and Stack Furnaces David White, Schaefer Furance Martin Reeves, Striko-Dynarad Paul Campbell and David Neff, Metaullics
REVERBERATORY FURNACES have a shallow hearth for melting metal, which is heated indirectly by burners mounted in the roof or in a sidewall of the furnace (Fig. 1). The flame used for melting the metal does not impinge on the metal surface itself but is reradiated off the walls or the roof of the furnace. The term reverberatory furnace is a throwback to the World War II era, when the most common fuels for hearth aluminum melting were coal or coke (Fig. 2). This necessitated locating the heat source or fire box at one end of the bath and the flue at the other end. The flue (stack) was designed to be strong enough to draw the flame over the bath. The roof was slanted to bounce or reverberate the flame (similar to sound echoes) off the ceiling, down across the metal to be melted in the molten bath, and then onto the flue. A primary objective was to heat the refractory walls and ceiling and then use these surfaces to radiate heat to melt the charge. This same flame pattern and the method of using a flame to heat the bath were implemented when
the fuel used was changed from coke to oil and gas. This older concept of aluminum metal melting is now uncommon but may still be used in some foundries or smelters. The modern method is to use roof burners or side burners directed toward the metal surface. In effect, the objective is to heat the metal, not the furnace, for efficient operation. Gas- or electric-heated burners are used in reverberatory furnaces. Aluminum foundry reverberatory furnaces range from small electric or gas-fired vessels of just a few tons volume to upward of 40 to 50 tons. Secondary smelters that produce ingot or supply liquid metal to foundries operate even larger furnaces. Gas reverberatory furnaces are commonly used in large nonferrous casting foundries because they provide a large reservoir of molten metal, ensuring a steady and reliable supply of molten metal to the shop. Small gas or electric everbretatory furnaces (Fig. 3) are also used for melting. Units are also used for holding furnaces, although typically the term reverbretatory furnace is used for melting
units—not holding furnaces. Practical sizes range from 1 to 45 Mt (2000 to 100,000 lb). The advantages of reverberatory furnaces include high-volume processing rates and low maintenance costs. Many foundries and die casters have used reverberatory furnaces for many years. Large- and medium-sized shops use them as central melters to supply multiple operations located within the shop. By using double-end clean-out and preheat hearths, these furnaces are some of the easiest to clean and are much more efficient than just a few years ago. Advances in insulating back-up linings have made these furnaces cooler to work around and more energy efficient than ever before. Many of the newer furnaces can reach an energy efficiency capability of 1250 Btu/lb of metal melted. Typical capacity of the furnace can range widely, from 1 to 75,000 tons. The energy efficiency of reverberatory furnaces is in the range of approximately 20 to 30% (Ref 2). Energy is mainly lost through the hot flue gases. In addition, as the molten
Fig. 2
Fig. 1
Basic components of a reverberatory furnace. Source: Ref 1
Early version of a coal-fired melter hearth furnace used to produce aluminum castings. (a) Plan view. (b) Vertical cross section
Reverberatory and Stack Furnaces / 161
Fig. 5
Fig. 3
Aluminum reverberotory furnace (a) 180 kg (400 lb) per hour electric reverberotory furnace, (b) small gas reverberotory furnace, Courtesy of Schaefer Furnace
Table 1 Melt loss and themal efficiency of electric and gas reverberatory furnaces Melt
Aluminum Aluminum Zinc Zinc
Type
Melt loss, %
Thermal efficiency, %
Electric Gas Electric Gas
1–2 3–5 2–3 4–7
59–76(a) 30–45 59–76(a) 32–40
Major components of a stack melter
metal comes in contact with the furnace gases, it forms slag/dross. The melt loss rate is 3 to 5% in aluminum gas reverberatory operations (Table 1). Improvements in burner technology, fuel/air ratio control, insulating refractories, and temperature control have contributed to a slow but steady improvement in the efficiency of reverberatory furnaces. Efficiency is also obtained in furnaces with preheating hearths (Fig. 4). Fuel savings of approximately 15% per year are possible by allowing a sow or ingot to sweat on the hearth as the internal temperature is raised to approximately 480 C (900 F). There is no moisture left in the sow or ingot, thus making it safer to charge. The stored thermal energy (Btus) provides further energy savings in maintaining the melt.
(a) The primary energy efficiencies of these furnaces are much lower (approximately one-third) than the listed efficiencies due to the use of electricity, which involves sizeable energy losses during generation and transmission. Source: Ref 2
Furnace Types
Fig. 4
In most large facilities, the conventional means for melting and holding aluminum is a gas-fired reverberatory furnace. In this arrangement, radiation from the burners, roof, and sidewalls heats the bath and solid charge. Forced convection through stirring or pumping is also needed to overcome a number of process limitations. Without forced circulation, the concentration of heat at the bath surface readily produces undesirably high temperatures. The high heat flux generated by direct radiation at the bath surface is typically much greater than conductive and natural convective forces, which distribute the heat through the bath. In turn, surface oxidation and dross formation ensue. There are two types of gas reverberatory furnace, based on whether the sows, ingots, and revert materials are preheated before melting or not. A dry hearth reverberatory charges and
Preheat hearth
preheats sows, ingots, and revert materials on a slope between the charging door and the molten metal bath, removing any contamination or water residue prior to melting, which prevents splashing or explosions. A wet bath reverberatory directly dumps the sows, ingots, and reverts into the molten metal bath. The wet bath system is beneficial for melting scrap types with large surface-to-volume ratios, which may oxidize excessively unless melted quickly. Another type is a side-well reverberatory furnace that consists of a number of burners firing inside the hearth, against the furnace hot wall or door. A charging well and a pump well are attached to the furnace hot wall on the outside of the furnace. Both wells are connected to each other and with the furnace hearth by arches, which permit aluminum circulation between the furnace chambers. During the charging period of the furnace cycle, molten aluminum circulates through the charge well where it melts the scrap, then circulates back into the hearth for reheating. Charging continues until the molten bath reaches the desired level. Subsequent steps in the smelting process may include degassing and alloying. Improvements in energy efficiency have also been made with the use of stack furnaces (Fig. 5), which can be considered as a modified reverberatory furnace with improved efficiency by better sealing of the furnace and the use of the flue gases to preheat the charge materials. Stack furnaces are more common than reverberatory furnaces for melting aluminum in Europe. In the United States, 95% of aluminum is melted in gas reverberatory furnaces, which operate with lower energy efficiency but have lower capital cost and provide easier operation and maintenance than stack melters (Ref 2). Dry Hearth Reverberatory Furnace. In a dry hearth furnace, the charge of solid metal
162 / Melting and Remelting is positioned on a sloping hearth above the level of the molten metal so that the charge is completely enveloped by the hot gases. Heat is rapidly absorbed by the solid charge, which melts and subsequently drains from the sloping hearth into the wet holding basin or chamber. Dry hearth furnaces are typically used where castings and/or returns contain nonaluminum inserts or in sand casting operations where casting scrap containing sand must be reclaimed. These castings can be melted on the hearth, scraping the sand or insert off the hearth and keeping them out of the molten furnace bath. This keeps the furnace cleaner and the aluminum from being contaminated with iron (in the case of cast iron inserts). The advantages of the dry hearth furnace include:
the metal is transferred to a holding furnace(s) that feeds the casting line(s). These furnaces are designed to use radiant heat (a highly efficient method of transferring heat into aluminum), where the heat source (burners) is located 46 to 61 cm (18 to 24 in.) off the bath for best practices. The closer the heat source is to the aluminum, the faster the transfer of Btus. The temperature differential between the thermal head and the metal bath provides the driving force for heat transfer. The advantages of the wet bath furnace include: The up-front cost is lower than high-headroom
Safety: Everything melts on the hearth in the
atmosphere of the furnace and not under the bath. Such preheating avoids the potential for moisture entrapment and subsequent explosive force in and beneath the bath surface. The dry hearth allows the saving of inserts and avoids quality problems with undesirable impurities being incorporated into the melt. Of course, there are other potential sources of impurities that must be considered. The dry hearth holding chamber is typically smaller, approximately a 4 to 1 hold-to-melt ratio. The furnace holds four times what it melts or less, so the furnace is smaller. If the flue is maintained correctly, very good temperature control can be achieved in the hold zone. The disadvantages of the dry hearth furnace include: Typically, these furnaces have two flues: one
reverberatory furnaces and shaft melters. Efficiencies can be achieved on average of 1500 Btu/lb of metal melted. Everything is melted under the surface of the bath for greatly reduced metal melt loss (2 to 5% in gas-fired furnaces and less than 1% with electric furnaces). Many options are available: preheat hearths for sows and ingots, circulation wells for scrap, vortexing for chips, flue pressurization, well space for continuous degassing and filtration, and so on. With circulation and preheating, efficiency can be increased to 1250 Btu/lb of metal melted. By adding radiant roof regenerative burners, energy efficiencies can be further improved to 1000 Btu/lb of metal melted.
The disadvantages of the wet bath furnace include: These furnaces have a higher hold-to-metal
ratio and therefore are larger (higher capital cost). With a typical hold-to-melt ratio of 8 to 1, a minimum of eight times the solid charge weight must be maintained to help
dilute the cold metal charge and not cause excessive bath temperature reduction. Larger furnaces require extra floor space. Sows and ingots must be preheated before loading in the furnace, whether on the charge well wear plates (i.e., furnace sill) or on a preheat hearth. Electric wet bath reverberatories must be charged evenly or they will lose temperature and take time to recover. Larger electricpowered melting furnaces are extremely expensive and not as efficient as large gasfired furnaces. Low headroom requires care when cleaning. Sows or ingot bundles cannot be double stacked on hearths because of lower headroom.
Stack Melters. For aluminum metal casting, there are substantial opportunities for energy savings in replacing reverberatory furnaces with stack melters (also referred to as shaft melters). Stack melters improve efficiency by better sealing of the furnace and use of the flue gases to preheat the charge materials. The charge materials slide down the shaft and reach the melting zone, where they are melted by the burners, and the molten metal flows down to the holding area (Fig. 5). The hot exhaust gases from the melting zone flow through the shaft to preheat the incoming charge, improving the energy efficiency of the stack furnace by 40 to 50% (Ref 2). A typical example of a shaft melter furnace in a foundry is shown in Fig. 6. One of the central advantages of the stack melter is the physical separation of the charging, preheating, melting, and holding areas of the furnace. Advanced burner controls can augment the advantages of this kind of furnace. The stack furnace also can be coupled with
over the hearth and one over the hold chamber. Two separately controlled burner systems are required for optimal operation. High metal melt loss (possibly 7 to 12%) can result due to melting in the atmosphere, which creates a very high oxidation rate at these temperatures (700 to 760 C, or 1300 to 1400 F). When melting light-gage scrap and small foundry ingot, the metal melt loss is high and the thermal efficiency is low (1800 Btu/lb or higher). When melting larger sows (i.e., typically 540 kg, or 1200 lb), the efficiency is better and approaches the 1500 Btu/lb range. Metal melt loss is lower as the volume-to-surface area of the charge material increases. Wet Bath Reverberatory Furnaces. In a wet bath furnace, the products of combustion are in direct contact with the top of the molten bath, and heat transfer is achieved by a combination of convection and radiation. Wet bath reverberatory furnaces generally are used where there is just one alloy being melted or minimal alloy changes. Melting can be performed in a central furnace, sometimes called a break-down furnace, and then
Fig. 6
Stack melter employed in foundry practice. Courtesy of Striko-Dynarad Corporation
Reverberatory and Stack Furnaces / 163 other melting systems to improve efficiency. For example, it can be used with a crucible furnace, reducing its melt losses since the crucible melts by indirect heat to the metal through the wall of the furnace. The disadvantage of the stack furnace is that, due to the charging mechanism, the height of the stack furnace must be more than 6 m (20 ft) (Ref 3). In addition, as the furnace is charged from the top of the shaft, the refractories at the bottom of the charging door are subjected to repeated impact and wear. This requires additional maintenance. The shaft construction and melting design are at the heart of the advantages offered. The choked shaft design allows the consistent preheating of the charge material prior to melting. The hot exhaust gases from the holding bath pass through the melt zone and from there into the preheating zone before exhausting from the furnace. The additional energy required for melting comes from burners located in the melt zone and focused into the area at the base of the preheating area. Exhaust gases from these burners also exit through the shaft and therefore ensure a greater degree of preheating as well as greater speed and efficiency of melting. Energy Efficiency. Compared to other furnace designs, stack furnaces can result in substantial savings from lower fuel consumption and consequently lower melting costs. In the practical foundry operation, for example, with the use of 50% ingot and 50% bulky return material and a bath temperature of 720 C (1330 F), a power requirement is only 600 kW per hour per ton. This corresponds to a thermal efficiency of over 50%. Improved Melt Cleanliness, Quality, and Yield. By the separation of melting and preheating zones, the production of clean, high-quality metal can be achieved with stack melters. Moisture and other contaminants are evaporated or burnt off in the shaft by preheating with exhaust gases, ensuring that only dry material reaches the melting chamber. The primary source of contamination is moisture from storage of the scrap or ingots; the result of this is an increase in the hydrogen content of the metal, with its consequent quality problems in the cast parts. A further advantage is the avoidance of excessive contact between the melt and atmospheric oxygen, thereby reducing the tendency to form oxides. In order to minimize this danger, the burners in the melt chamber are arranged to avoid direct metal impingement. This allows melting on the hearth only and allows the metal to flow quickly but without turbulence into the holding bath without superheat. This ensures an optimal value for metal yield; typical values lie between 0.8 and 1.5%. A third measure for the maintenance of metal quality is the avoidance of contamination by impurities entering the metal holding bath. The separation of melting and holding ensures that contaminated charge material is not charged directly into the bath. In a crucible furnace or wet bath melter, what goes in stays in and must be removed by additional
metal-cleaning operations. Cleanliness is also assured by allowing the metal to run laterally into the holding bath. In this way, dross and other impurities remain on the melt hearth, to be removed during cleaning. An additional advantage from the separation of melting and holding is an almost continuous availability of molten metal, on demand, with a very consistent temperature (to +/5 C, or 9 F), which itself represents an important feature of quality. Melting continues while the holding bath level remains below maximum. The holding bath also offers the foundryman a source of molten metal at a consistent and correct temperature without fluctuations due to melting. Automated charging (Fig. 7) also contributes to this factor by ensuring that material is always available to melt when required. This is done using an automated lift-and-tilt system. Stack melters can also be combined with more conventional sow-melting technology to improve efficiencies and provide more options for charge materials. Various levels and forms of automation are available to cater to the different needs of the foundry and material supply. This can be by conveyor feed and automatic bin loading, allowing up to several hours, depending on melting rate and metal demand.
Reverberatory Furnace Practice Radiant Heating Factors. Radiant roof reverberatory furnaces rely on three major factors to make the process more efficient: Distance from the heat source (the burners)
to the load (the aluminum)
Hold-to-melt ratio Even charging of the load
One major design criterion for radiant roof melting is based on the Stephan-Boltzmann law of radiant heat, which states that the rate of heat transfer increases the closer an object is to a heat source and the greater the temperature differential. By using low-velocity, flat-flame, highthermal-release burners that heat the refractory block construction to 1400 C (2600 F), a huge temperature differential is obtained between the aluminum and the heat source. In this type of furnace, efficient heat transfer into the aluminum is principally by radiation. The greatest efficiency is obtained when the temperature differential is the greatest between the radiant source (the burner blocks, not the furnace upper chamber) and the receiving medium (the molten aluminum bath). Heat transfer is also more efficient when the radiant source is closer to its receiving medium. When the radiant heat sources are closer, more burners are needed to obtain coverage of the medium. Most reverberatory, radiantly heated melting furnaces will employ practices in aluminum melting wherein the metal charged is either all solid (ingot, foundry scrap, gates, runners, etc.) or a combination of solids plus liquid
Fig. 7
Charge for stack melter
(metal delivered in over-the-road ladles). Consequently, a portion of the charge will be cold (solid) metal. The hold-to-melt ratio is critical as a dilution factor for the cold metal being charged into the furnace. The higher that ratio, the less temperature differential is obtained when the furnace is evenly charged. Either radiant roof or highheadroom side-fired melters should be charged at ¼ of their hourly rated capacity every 15 min. This will ensure that an even temperature is maintained in the bath at +/10 during the melting process. Charging. Overcharging of these furnaces can create sludge in the furnace (if the metal temperature drops below the sludge point of the alloy), which will result in poor-quality metal for casting. Such overcharging and causing greater temperature swing in the melting and already molten bath can also result in additional dross formation at the melt surface, due to increased exposure time (i.e., the overall melt cycle) and the absolute temperature at the bath surface as it continues to absorb heat during recovery of the bath temperature. Typically, the radiant roof approach of melting everything under the surface of the bath greatly reduces metal melt loss (2 to 3% in gas furnaces and less than 1% in electric furnaces) and increases melt efficiencies for melting small, lightweight scrap. Recuperation. While there is little doubt that there is a great deal of savings possible in recouping heat lost out of a flue, some care is needed on how far this can be taken. For example, keeping the recuperation temperature below 370 C (700 F) can substantially reduce the system cost. Below 370 C (700 F), insulated mild steel manifolds can be employed as opposed to more costly stainless steel. Even at 370 C (700 F) preheated combustion air, a 25% savings in fuel usage is possible. An example of the cost trade-off is based on numbers produced by a major burner manufacturer at $10.00/106 Btus on an 1800 kg/h (4000 lb/h) radiant roof melter. One can expect to pay approximately $85,000.00 more for a burner system using an installed recuperator and save approximately $54,750.00 per year, or have a
164 / Melting and Remelting 17.4 month payback on this investment differential. Actual costs may vary due to the age of the furnace and components or the type of furnace and the number of burners. Other factors to consider in evaluating recuperation are: The flue gases should be dampered to be
able to bypass the exchanger when the furnaces are fluxed. Otherwise, excessive wear and maintenance on the heat exchanger will result. The nozzles in the burner should be changed to stainless steel to handle the additional temperature of the preheated combustion air. Older furnaces should consider changing to an ultralow-NOx burner. As the combustion air temperature increases, the NOx emissions go up. Regenerative burners enhance melting efficiency by providing a steady supply of hightemperature preheated combustion area. These burners contain dual firing ports, reversing firing as the air is drawn over heat-storage media in each port. In most cases, initial first-cost of recuperation provides a faster payback than does regeneration on smaller (2700 kg/h, or 6000 lb/h, or less) furnaces. This assumes a new furnace. For larger furnaces, the regenerative burners really begin to pay off. Payback varies depending on the type and age of the furnace in which regenerative burners are installed. Certainly, if the furnace is melting metal at a high end (1800 to 2000 Btus/lb of metal melted), then the return on investment becomes more lucrative. Regenerative burners become cost-efficient with high natural gas prices. As a further example, buying a 3200 kg/h (7000 lb/h) melter with regenerative burners to replace a less energyefficient high-headroom furnace that melts at 2000 Btu/lb of metal melted would result in a savings of over $452,000.00/year in fuel cost. For a melter of this size, the return on investment payback can be two years or less. (These figures are based on 3200 kg/h, or 7000 lb/h, for 20 h of melting/day, 340 days/year, and $10.00/1000 ft3 of gas.) Reducing Fixed-Heat Losses. Highly insulated furnaces with substantial linings help provide for the lowest fixed-heat loss. Fixed-heat loss is the amount of energy being absorbed into and passing through the lining of the furnace and out the doors and open wells on the furnace. An example of fixed-heat loss is: If the sidewall casing temperature of a given
furnace is 96 C (205 F), then 300 Btus/ft2/h are lost from that casing. By simply reducing the temperature to 64 C (148 F), 152 Btus/ft2/h are saved. The older refractory configuration concept taught that the freeze plane (the point where the molten metal would solidify if it leaked through the lining) had to be in the hot face or first layer of refractory material. With the
availability of newer, nonwetting insulating materials, the freeze plane can now be designed to be further back in the thickness of the lining. This gives much more flexibility to highly engineer this lining to not only contain the metal but actually save energy. Melting furnace casing temperatures are being designed to 50 to 60 C (125 to 135 F) instead of the old 75 to 85 C (165 to 185 F) of just a few years ago. New superinsulating materials can be used in the roof and floor of the furnaces to further reduce energy consumption. Some furnaces need to have individual burner controls, just to keep the furnaces from overheating over long weekends, because of this thermos-bottle approach to highly insulated linings. The fixed-heat loss has been reduced to the point that the burners on the furnace at low fire will overcome the heat losses and actually raise the temperature if they are not turned off. Covering up open wells will reduce heat loss substantially. An open well of 700 C (1300 F) aluminum with a 0.5 emissivity of the bath surface in the well will lose 8214 Btus/ft2/h. If the average charge well is 2.4 m long by 1 m wide (8 ft long by 3 ft wide) (2.2 m2, or 24 ft2), then 197,136 Btus/h are being lost off of that well. That corresponds to 5.6 m3 (197 ft3 of gas per hour, or 134 m3 (4731 ft3) a day by 340 days a year is 45,560 m3 (1,608,540 ft3) of gas lost. That is $16,085.40 worth of natural gas at $10.00/therm lost a year. Improving and Maintaining Furnace Efficiencies. Burners are not stoichiometric. The combustion system must be set up properly, or more dross will be generated. Typically, most natural gas-fired burners are set up on a 10 to 1 air-to-gas ratio. At this ratio, 95% of the oxygen in the burner system is consumed. However, if a furnace is set up with ratios in excess of this, dross increases in the furnace. On the other hand, a furnace that is set up fuel-rich will produce a colder flame and make the furnace work harder to get the job done. Keep the burner in ratio. A qualified technician should check the combustion system for the proper set-up. The burner system should be firing stoichiometric, or in ratio, air to gas. Typically in reverberatory furnaces, this is a 10 to 1 air-to-gas mixture. By analyzing the flue gases, a technician can determine if indeed the system is running at peak efficiency. This can also be done with a manometer to test pressure (inches of water column) of gas and air, if the technician knows the burner manufacturer’s recommended set-up pressures. Fuel-rich burners produce a colder flame, and some of the gas will ignite when it reaches the atmosphere at the flue level. This results in a wall of flame coming out of the flue. If the burner is set up too lean or has excess air, a colder flame results, together with an increase in oxide formation and dross in the furnace. This may cause the burners to work harder to produce the same Btus, resulting in lower energy efficiency and therefore higher costs of production.
In addition to burner ratio, improving and maintaining furnace efficiencies can be achieved in several ways, as summarized subsequently. Remove all fans that blow on the furnaces. This will result in less heat loss from the casing where the fan is hitting the furnace. Instead, use individual cooling stations that bring the air in from overhead to a specific spot where the operator can stand to cool off. Install well covers (Fig. 8) whenever the furnace is idle for longer than a half-hour. Install a half-flue cover over the flue during long idle conditions. Warning: The flue cover must be removed before turning the furnace to high fire. At low fire, the flue opening is still sized for 100% output of the burners. Without automatic flue pressurization, the holding efficiencies will drop considerably. The furnace may not need to be on high fire as much. Charging. Make sure the furnace is used at full capacity and charged ¼ of its hourly rated capacity every 15 min. Overcharging the furnaces causes dramatic swings in temperature and increased sludge buildup on the floor of the furnace. The more sludge on the floor, the less likely the furnace will melt at its rated capacity. The furnace may not melt anywhere close to the Btus/lb of metal melted that it was designed to do. Reduce molten metal temperature over weekends by 11 C (20 F) or more if feasible. Clean furnaces during idle shifts or idle times, but clean them daily! This will cut down on the amount of oxides growing in the furnace. Oxide is dense and absorbs additional Btus from the metal. Do not overflux the melters. Overfluxing of the walls can break down the binder in the refractory and cause premature erosion of the belly-band area. Clean or change combustion air filters regularly. Hold a dollar bill 50 mm (2 in.) away from the filter. If the bill is not drawn to the filter rapidly, it is time to clean or change the filter on the blower. Double-check and monitor combustion set-up for maximum burner performance (requires a manometer). If a wall of flame is coming out of the flue, a fuel-rich condition probably exists inside the furnace. The raw gas that does not have enough air in the furnace for combustion
Fig. 8
Pneumatic well cover
Reverberatory and Stack Furnaces / 165 ignites when it reaches the atmosphere just outside the flue, because of all the free air available to it at the flue and hood junction. Overpressurization of the furnace can also cause this condition. Preheat Ingots. In reverberatory (radiant roof low-headroom or high-headroom) furnaces, preheat all ingots on a charge well (wear plates) or on a hearth before putting them into the bath. Preheat all sows on a hearth before pushing them into the bath. When a sow or ingot begins to sweat, the internal temperature is close to 480 C (900 F). There is no moisture left in the aluminum, and it can be safely charged under the bath of aluminum, where the stored Btus in the metal help finish the melting process. Ingots left on a well wear plate to preheat will reach 120 C (250 F) after 20 min. This is sufficient to load into the bath. Some bubbling may occur from moisture left in the very center of the ingot. Metal delivery to the holding furnaces should be every 20 to 30 min; less frequently will mean overcharging the melters and drawing the metal down too low in the holders. The closer the metal is to the heat source, the faster the transfer of Btus. A loss of 10% in efficiencies could take place by drawing the holders that low. Ensuring that the metal tenders put back what they take out in the melters results in a difference in metal temperature stability. Remove dross buildup whenever it becomes thicker than 13 mm (½ in.), or the dross will begin to act like an insulator and make the furnace work harder to do the same job. A completely clean bath surface will reflect the heat, so it should not be cleaned any more frequently than once per shift. If a furnace creates a dross layer over 13 mm (½ in.) in less than 6 to 8 h, then the furnace is creating more dross than normal. Electric Pump Motors. To further conserve energy, the pump should have an electric motor rather than an air motor. Plant air is the most expensive energy source in the plant. A small blower keeps the pump cool. Treat the cause of dross buildup, not the symptom. Investigate why the furnace is generating dross. Typical reasons are: Air leaking in and around the doors Drawing the furnace down below
the submerged arches or door blades Burners that have a flame that touches the bath or scrubs the solid metal on a hearth Too large of a flue opening for the Btus connected, causing the furnace to draw in more air Keep the burn in ratio!
Molten Metal Circulation In the melting of aluminum and its alloys, forced convectional heat transfer within the melt often plays an important role. There is a considerable latent heat of fusion (95 cal/g) associated with taking aluminum at its melting
point from a solid to a liquid. Forced convection assists this transition, particularly for larger volumes of metal, by better mixing of the solid metal with already molten metal, thus reducing time and energy input into the melting process. Other nonferrous alloys, such as zinc, magnesium, and copper, could also benefit from molten metal circulation, but pumping systems for circulation are not usually employed. Mechanical stirrers are effective, particularly for zinc, due in part to the lower heat of fusion of zinc and the smaller melt sizes usually used. In reverberatory furnaces, forced convection enables the radiant heat energy absorbed by the bath surface to move, that is, to be transmitted throughout the bath depth, so that a steady heat flux is developed and continuing heat transfer to the surface occurs by radiation. Without such forced convection, the surface layer of the molten metal becomes heat saturated, and the surface temperature rises. Ultimately, excessive surface temperature will result in oxidation and dross buildup, which reduces energy efficiency. Forced convection is produced by the circulation of the molten metal bath. Current methods to achieve circulation or metal movement include manual stirring, mechanical pumps, electromagnetic pumps, induction stirrers, porous plugs, jet pumps, and towmotor/boom agitation (see the section “Molten Aluminum Circulation Methods” in this article). Manual stirring is only practical in smaller furnaces (in small pot furnaces, crucibles, and so on). Forced circulation is usually best achieved in large furnaces (2000 to 110,000 kg, or 5000 to 250,000 lb, capacity) by using a submerged discharge molten metal pump. Several types of molten metals pumps have appeared over the years (such as electromagnetic and pressure pumps), but the centrifugal pump has been a common type. These pumps are usually installed in the pump wells and discharge into the charge well of reverberatory furnaces for remelting in large aluminum foundries, die casting operations, secondary smelters producing foundry ingot, and in mill product casting houses. Typical installations are illustrated in Fig. 9 (a and b). A summary of the benefits of forced circulation in large furnaces employed in aluminum foundry practice, as well as a comparison of the different methods of circulation, is in the following sections.
much greater than conductive and natural convective forces, which distribute the heat through the bath. In turn, surface oxidation and dross formation ensue. Figure 10 (Ref 4) provides the oxidation rate of molten aluminum as a function of temperature. Because aluminum oxide is thermally insulative, the variation between surface and bath temperatures increases with the growth of the oxide layer. Temperature differentials in a 915 mm (3 ft) deep bath with a clean surface and without forced convection can typically reach 50 to 85 C (90 to 150 F). Significant melt loss is incurred due to ever-increasing dross formation under these conditions. Without effective heat transfer to the metal, other quality and furnace performance factors may be less than optimal.
Fig. 9
Molten metal circulation pumps. (a) Typical pump installation in a sidewell furnace. (b) Circulation well and pump
Benefits of Forced Circulation The benefits of circulating molten metal in reverberatory furnaces are well documented, especially for aluminum. Forced convection through stirring or pumping overcomes a number of process limitations of reverberatory furnaces for melting and holding. Without forced circulation, the concentration of heat at the bath surface readily produces undesirably high temperatures. The high heat flux generated by direct radiation at the bath surface is typically
Fig. 10
Oxidation of aluminum versus temperature
166 / Melting and Remelting The result of forced convection is to reduce metal temperature stratification in the bath, improve melting rates, enhance metal quality through alloy homogeneity, reduce energy consumption as a result of improved heat transfer, extend refractory life, reduce sludge (dross) formation and therefore lower furnace-tending requirements, and accelerate alloy dissolution rates. Operational benefits of more homogeneous bath temperature of the entire bath of metal (top to bottom, inside and out in the well) include: Quicker melting with reduction of melt
cycle times by 10 to 50% Reduce energy consumption by 15 to 25% Extend refractory life at least 25% Improve recovery rates by a minimum of 0.25% With these major factors at stake, it is no wonder that a wide variety of techniques have been developed to provide forced circulation. While many production parameters may be difficult to fully quantify, forced circulation clearly improves many key aspects that affect total quality. These benefits of circulation are influenced by a variety of production/operation factors, as briefly described here. The next section, “Molten Aluminum Circulation Methods,” in this article also summarizes the performance of currently applied circulation technologies in large-furnace foundry practice. The review covers six methods that have evolved to some degree of acceptance: mechanical pumps, electromagnetic pumps, induction stirrers, Alcan jet stirrer, porous plugs, and towmotor/boom. With each technique, the production manager must assess the accompanying benefits, characteristics, and drawbacks. Reduced Temperature Stratification. Bath temperature is a key factor that impacts overall product quality. Forced circulation improves bath product quality by ensuring the bath temperature will be consistent throughout the furnace. Typically, reductions in temperature differentials for a 915 mm (3 ft) bath are from 50 to 85 C (90 to 150 F) in an uncirculated furnace to 3 to 7 C (5 to 13 F) for stirred furnaces. Circulation ensures that the temperature of the metal exiting the furnace is virtually the same as that indicated by the furnace thermocouple. Production parameters can be set and more easily controlled when the metal delivered is at a constant temperature. Additionally, there will be less superheat of the metal required because there will be less variance in the furnace temperature. Increased Melt Rate. Constant circulation of the furnace will increase the melt rate of the furnace as well as improve the thermal profile. The high-velocity stream from a circulation pump will increase the melt rate of large submerged solids in the furnace. Insulating boundary layers that are formed around large solids as they melt are swept away by the hot metal stream. Such a device equipped with gas injection can provide bath fluxing and degassing during the melting cycle, eliminating
or greatly reducing the time required at the end of the cycle for metal treatment. The same holds true for temperature control. Continuous circulation provides excellent temperature control, so temperature adjustments are usually not required prior to transfer. Cunard, using a direct-charge furnace in a billet casting process, reported in a study over a five month time period that furnace cycle times increased between 8.1 and 20.6% when the pump was not in operation (Ref 5). He estimated that using a pump produced a conservative value of 10% cycle-time savings. Charging time reductions of 25 to 50% have often been observed in sidewell furnaces when a pump is used. Savings may be obtained by increasing output with existing equipment or by maintaining production levels with fewer furnaces. Manufacturers of foundry reverberatory furnaces estimate that the melting capacity of a given furnace will increase by 10 to 15% when forced circulation is added to an uncirculated furnace. Alloy Homogeneity. One method for stirring a furnace was developed by Alcan during the 1980s. They attached a large pressure vessel to the side of the furnace. The pressure vessel communicates with the furnace through an opening below the metal surface. When a vacuum is applied to the vessel, metal is drawn in from the furnace through the opening. The vessel is then pressurized, forcing metal back into the furnace through the opening. The vacuum is reapplied to the vessel, and the cycle repeats. The results of an Alcan jet stirrer (AJS) installed in a top-charge 70 ton round melter have been documented (Ref 6). Figure 11 schematically represents the main elements of the AJS, while Fig. 12 shows a typical adaptation to a top-charge round melter. The results were obtained during the preparation of an AA4000series alloy with a high silicon content. The silicon specification for this alloy is 7.15 to 7.45%. The silicon content was monitored in the trough during the cast and compared with samples taken from the furnace at the end of the batch casting process. A significant difference was observed between the two samples for the same batch. For eight runs without stirring, the average silicon composition in the furnace at the end of the cast was 7.28%. In contrast, the average for trough samples during the cast was 7.22%, with a standard deviation of 0.08%. In fact, two ingot batches were outside of specification. In five runs where stirring was used, no significant difference between furnace and trough samples was observed. All ingots were within specification, and the standard deviation for the cast samples dropped to 0.0%. In foundry applications, similar results can be expected when the furnace is circulated. Without forced circulation, alloy additions must be stirred in mechanically to ensure uniformity. Since mechanical stirring is done over a relatively short period of time, it is difficult to ensure that the alloy is mixed uniformly
Fig. 11
Schematic of the main elements of the Alcan jet stirrer
Fig. 12
Typical Alcan jet stirrer installation in a top-charge melter
throughout the furnace. Some additions will also begin to sink to the bottom or float to the surface if not continually stirred. When a stirring device such as a mechanical pump is used, alloy elements are constantly mixed with the rest of the metal in the furnace. In a typical foundry application, the volume of metal in the furnace passes through the pump every 15 min. Since the pump is operated 24 h per day, samples taken from any point in the furnace will show uniform chemistry. One other consideration with foundry alloys is the formation of sludge. This is an Fe-CrMn compound that precipitates out of the metal if the temperature drops below a critical point. This is discussed in detail by Jorstad (Ref 7). With a temperature differential of 50 to 85 C (90 to 150 F) between the top and bottom of the bath, it is easy to see how the metal at the bottom of the bath can cool below the critical point if the furnace temperature is monitored and controlled by a thermocouple located near the top of the bath. As previously discussed, temperature is uniform throughout the bath when the furnace is constantly stirred. This allows for accurate temperature control and the avoidance of sludge formation. Reduced Energy Consumption. The rate at which energy is transferred from the burners
Reverberatory and Stack Furnaces / 167 and refractories to the metal is referred to as the heat flux (q). Heat flux in a reverberatory furnace is determined by the following equation (Ref 8): q ¼ Ah½ðt1 Þ4 ðt2 Þ4 where A is the surface area of the bath inside the furnace, h is the heat-transfer coefficient of the area inside the furnace above the bath, t1 is the temperature of the heat source or refractory walls and ceiling, and t2 is the temperature of the bath surface. From this equation, it can be seen that the rate at which energy is transferred into the metal is dependent on the difference between the surface temperature of the metal and the temperature of the refractories, to the fourth power. Constantly stirring the furnace allows heat to be removed from the surface to the rest of the bath by convection. As heat flows more efficiently away from the surface of the bath, the bath surface temperature is reduced. As the surface temperature drops, the temperature differential between the refractories and the metal increases, which translates into much higher heat-transfer rates. This subject is covered in more detail by Campbell (Ref 9). Since stirring occurs during the melting process and with the doors closed, energy that would have been lost out the door when the bath is stirred manually is saved. Additional savings are possible from the elimination of throughthe-door fluxing operations if the circulation pump is equipped with gas injection capability. Actual furnace efficiency improvements vary from operation to operation. Experience has shown that a minimum of a 15% improvement can be expected. A 25 to 30% improvement is possible, depending on furnace and operationspecific variables. In the study done at VAW in Ellenville, New York (Ref 5), savings ranged from 10.3 to 20.2% over a five-month period. Longer Refractory Life. By reducing the metal surface temperature, more heat is transferred to the metal and less is absorbed by the refractories. Roof and stack temperatures are usually reduced by 80 to 115 C (145 to 205 F). These lower operating temperatures translate into longer refractory life. Lower temperatures will also result in less corundum formation at the metal line. Less buildup means less need for mechanical cleaning and therefore less opportunity for damage from cleaning tools. While specific information on savings from improved refractory life is not widely shared, dramatic improvements have been reported by some companies using circulation pumps (Ref 10). Decreased Dross Formation. Another aspect that should not be ignored is melt loss. Oxidation increases with temperature (Fig. 10); at temperatures above 775 C (1425 F), the rate at which aluminum oxidizes increases quickly. Reducing the bath surface temperature can significantly reduce the rate of dross formation. Minimizing surface agitation with submerged circulation removes another source of dross formation.
Many foundries do not pay enough attention to melt loss that results from dross formation. Every pound of aluminum that is lost to oxidation in the form of dross is lost to the manufacturing process and must be replaced. While not widely documented, the typical dross reduction in the hearth when using a centrifugal pump for forced circulation is approximately 0.5%. Using a conservative value of 0.25%, a foundry that melts and casts 454 MT (1 million 1b) of aluminum a month would save the replacement cost for 1.1 MT (2500 1b) of metal every month. In another study, a large die casting company augmented its production facility with centrifugal pumps (Ref 11). They estimated a 98% recovery of charged material during the first six months of pump use versus a 96% recovery prior to its use. Easier Alloy Adjustments. The benefits associated with alloy additions can manifest themselves in three areas: alloy addition dissolution, treatment during melting, and improved melt quality before downstream treatments. By continually providing forced circulation, alloy additions and adjustments are easier to control, since the bath will remain homogeneous. Silicon and other alloy additions go into solution much faster when exposed to the discharge stream of a pump. When heavy chlorination is required due to scrap charge composition and desired alloy chemistry, a gas injection pump is an extremely effective tool (Ref 12). The dispersion and subsequent reaction of chlorine gas as it exits the pump in the molten metal stream creates a very effective kinetic environment for removal of magnesium from the melt. Figure 13 provides a typical representation of demagging results in the production of foundry ingot in a 68 MT (150,000 lb) furnace with a 34 MT (75,000 lb) heel from the previous heat. The magnesium content was decreased consistently and effectively during the normal charge cycle.
Fig. 13
Demagging results using a gas injection pump
Results from tests conducted at Golden Aluminum in 1995 show that with a mixture of 95% Ar and 5% Cl, a gas injection pump can significantly reduce sodium and calcium levels without affecting magnesium content (Ref 13). The tests also showed an improvement in porous disc filtration analysis results and a reduction in gas levels while the gas injection pump was in operation. By removing inclusions and hydrogen in the melting furnace, the load is reduced on downstream operations such as degassers and filters.
Molten Aluminum Circulation Methods As noted, methods of molten metal circulation include manual stirring, mechanical pumps, electromagnetic pumps, induction stirrers, porous plugs, jet pumps, and towmotor/ boom agitation. These technologies have evolved to facilitate efficient and consistent molten aluminum processing in melting and holding furnaces as the commercial importance of aluminum continues to grow. The features, limitations, and costs of each device must be considered with the requirements of the operation to determine the best solution. While all of these devices impart motion to the molten aluminum bath, the effectiveness of heat and mass transfer, capital investment, running expense, maintenance issues, ease of installation, and capabilities for alloy modification vary significantly. One of the most important factors in today’s (2008) competitive environment is product quality. While the economic benefits of quality can be difficult to quantify for a specific operation, the benefits of improved productivity, reduced cycle time, and better thermodynamics can be readily translated into economic savings. Many papers have been published on the application and performance of mechanical circulation and gas injection pumps, as provided
168 / Melting and Remelting in the references. These are typically centrifugal pumps with a rotating impeller in a volute casing. Materials of construction consist of graphite and/or ceramics so that they can survive in contact with the molten metal. They require an external sidewell (Fig. 9a) on the furnace. Mechanical pumps offer the most flexibility, low capital cost, and moderate maintenance requirements. On the other hand, they are not suitable for every application. The operating principle of an induction stirrer and an electromagnetic pump is the same as that of a linear motor (Ref 14). An arrangement of a bottom-mounted induction stirrer with a furnace is provided in Fig. 14. The melt is stirred by the force created by the inductor attached to the bottom of the furnace. A variation of this mounts the stirrer on the side of the furnace. Electromagnetic stirrers typically have the highest capital and energy costs, but with no metal contact, they tend to be the most reliable of all the systems considered here. In contrast, the electromagnetic pump moves molten metal through a tube within a coil by the generation of an electromagnetic force created by an applied magnetic field and induced eddy currents (Ref 15). Figure 15 depicts the plan view of an electromagnetic pump in a furnace with charge and dross wells. Electromagnetic pumps offer many of the advantages of a mechanical pump, generally require less maintenance, and can be used in applications not suitable for mechanical pumps. Capital costs
Fig. 14
Fig. 15
Schematic side view of an induction stirrer
Plan view of an electromagnetic pump, charging well, drossing well, and furnace
and energy costs are higher than mechanical pumps. The operating principle of the AJS is based on the repetition of two phases: a vacuum/accumulation phase followed by the pressurization/ expulsion phase. Metal is drawn into the unit through an opening with the furnace during the vacuum phase, then forced out through the same opening during the pressurization phase. For several years, Reynolds Metals used porous plugs in melting and holding furnaces at the alloys plant in Muscle Shoals, Alabama (Ref 16). This technology uses porous refractory plugs that are permanently installed in the furnace floor. Gas, usually nitrogen or argon, passes through porous refractory and produces a plume of fine bubbles over each plug. This gas plume creates localized circulation over each plug. Finally, a towmotor/boom system is used in some foundries and is considered to provide circulation within the furnace hearth. With this method, the door is opened, and the towmotor operator works a boom-and-rake arrangement back and forth to stir and skim the molten melt. Maintenance Comparisons. As a general rule, mechanical pumps will require more maintenance than electromagnetic pumps and stirrers. Mechanical pumps have seen significant improvements in designs and materials of construction in recent years, and reliability has improved to a point where at least 1 manufacturer now offers a 1 year warranty on the major components. Electromagnetic pumps and stirrers have no moving parts and therefore require little maintenance (Ref 17). This is somewhat offset by higher capital and energy costs. Porous plugs require no maintenance. Life and reliability are determined by installation and operation techniques. Major refractory repairs are required when they need to be replaced. Flow Rate/Effectiveness Comparison. As described in the first section of this article, there are many benefits of strong, forced circulation. Mechanical pumps, electromagnetic pumps, and induction stirrers receive high marks for delivering these benefits to the production process. Porous plugs, which are probably better categorized as a treatment technique rather than forced circulation, rate as moderate in overall performance. The towmotor/boom system is a “get-what-you-pay-for” arrangement. Not only is the circulation poor and intermittent, there are quite a number of negative factors that are associated with this technique. Since the furnace doors must be opened for access, there is a high heat loss. With this type of thermal cycling and exposure, the refractory life is significantly reduced. While this method is inexpensive from a capital standpoint, the increased dross formation and associated metal losses can be a large operating loss. Finally, safety is an issue. Certainly, the towmotor operator is at risk in this environment; it is often a topic of lively discussion at aluminum
meetings. Additionally, the back-and-forth action that is required presents a hazard in an aisle that is common to both plant personnel and equipment traffic. Alloy Modification. In today’s (2008) demanding production environment, concurrent processes are often added to classical charging, melting, and homogenization schemes to reduce overall production time. These processes include chlorine injection to remove magnesium from the melt, gas injection to remove hydrogen and increase inclusion flotation, and solids addition for alloy modification. A mechanical pump structured as a gas injection model has documented performance for all three key needs. An electromagnetic pump can be used for alloy modifications by the addition of solid materials, but it has not been shown to be effective for gas injection. In contrast, the induction stirrer does not provide any of these benefits as it is currently configured. Porous plugs are excellent for hydrogen and inclusion reduction. They are not employed for demagging, and their effectiveness for alloy additions is not documented. Towmotor/booms are often used to stir in alloy additions, but they cannot effectively improve operations over any of the other stirring methods. Ease of Installation/Retrofit. The application of a mechanical pump requires an open well and a residual molten metal heel during normal furnace cycles. A floor-mounted induction stirrer system is usually only considered for new construction, since the installation requires a basement area to house the inductor, and the furnace floor must be stainless steel. Side-mounted induction stirrers can be added to a furnace fairly easily. Electromagnetic pumps can be adapted to most furnaces. The furnace refractory and steel shell must be modified so that the equipment can be bolted to the side of the furnace. In order to install porous plugs, a major furnace reline is required. The towmotor with boom may be used in all instances. REFERENCES 1. “Refractories for Industrial Processing: Energy Reduction Opportunities,” Oak Ridge National Laboratory, Dec 2004 2. Advanced Melting Technologies: Energy Saving Concepts and Opportunities for the Metal Casting Industry, BCS, Inc., Nov 2005 3. J.F. Schifo and J.T. Radia, Theoretical/Best Practice Energy Use in Metalcasting Operations, U.S. Department of Energy Industrial Technologies Program and Keramida Environmental, Inc., May 2004, p 36 4. D.V. Neff, Molten Aluminum in Motion, Light, Metals, TMS, 1993, p 197–211 5. T. Cunard, Application of Graphite Pump to Direct Charged Aluminum Melting Furnace, Aluminum Industry Energy Conservation Workshop X, Oct 20–21, 1987, Aluminum Association, p 159–185
Reverberatory and Stack Furnaces / 169 6. M.A. Thibault, F. Tremblay, and J.C. Pomerleau, Molten Metal Stirring: The Alcan Jet Stirrer, Light Metals, TMS, 1991, p 1005–1011 7. J. Jorstad, Understanding Sludge, Die Casting Engineer, NADCA, 1986 8. M.C. Mangalick, Reduced Fuel Consumption Through Mechanized Molten Metal Circulation, Soc. Die Cast. Eng., Vol 17 (No. 5), Sept–Oct 1973, p 22–31 9. P.S. Campbell, Use of Molten Metal Pumps to Reduce Energy Consumption in Reverberatory Furnaces, Aluminum Industry Energy Conservation Workshop XI, Nov 1–2, 1990, Aluminum Association, p 537–546
10. T. Thornton, C. Hammond, J. van Linden, P. Campbell, and C. Vild, “Improved UBC Melting Through Advanced Processing,” 136th Annual Meeting and Exhibition, TMS, 2007 11. L-35 gas injection customer, private communication 12. D.V. Neff, Use of Gas Injection Pumps in Secondary Aluminum Metal Refining, Recycle and Secondary Recovery of Metals, Dec 1–4, 1985, Metallurgical Society of AIME, p 73–95 13. R.S. Henderson et al., “Quantification of Molten Metal Quality Improvements Using an L-Series Gas Injection Pump,” Third Int. Symposium on Recycling of Metals
14. 15.
16.
17.
and Engineered Materials, Nov 12–16, 1995, TMS C. Clark et al., A Molten Metal Pump Installation at New Zealand Aluminum Smelters Limited, TMS, 1988 R.M. Jones, A New Electromagnetic Metal Pumping and Stirring Technology, Fourth Australian-Asian-Pacific Course and Conference, TMS, 1995, p 275–289 J.R. Guttery and W. Evans, A New Generation of Fluxing in Aluminum Melting and Holding Furnaces, Light Metals, TMS, 1994, p 921–927 A. Katyal, Metal Stirring Using an Electromagnetic Pump, Optimization ’95 Conference, Proc. BCIRA, Vol 8, 1995, p 1–20
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 173-184 DOI: 10.1361/asmhba0005352
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Transfer and Treatment of Molten Metal—An Introduction MOLTEN METALS are processed and handled in various ways, depending on the alloy and melting practice. The very nature of melting is a way of combining elements for the purpose of alloying based on the laws of physical chemistry (see the article “Chemical Thermodynamics and Kinetics” in this Volume). However, the reactivity of metals and the high temperatures of melting also introduce the deleterious effects of dissolved gases and inclusions in molten metals, as described in the articles “Gases in Metals” and “Inclusion-Forming Reactions” in this Volume. Molten-Metal Treatments. Of primary importance during melting and the handling of molten metal is the mitigation and reduction of dissolved gases and inclusions in the melt. This objective may be accomplished in different ways, depending on the base metal, alloying constituents, and residual impurities. Fluxing of molten metal is often the first step in obtaining a clean melt. Fluxing is a term commonly used in foundries and melting departments that refers to molten-metal treatments from the addition of reactive chemical compounds or inert gases. The various functions of molten-metal fluxing additives, either singly or in combination, include:
To render the high-melting refractory oxide
impurities or gangue fusible and separable from the molten metal To provide a medium with which the impurity elements or compounds would combine in preference to the metal In steelmaking, the primary basic fluxes are either limestone (CaCO3) or dolomite ((Ca, Mg)CO3), depending on the other slag constituents present and the required ability of the slag to remove sulfur. Limestone is preferred for large sulfur removal. The slag-forming compounds in steel melting can be classified as either basic or acidic (see the article “Electric Arc Furnace Melting” in this Volume). Likewise, secondary refinement of steel melts may include various fluxing techniques in converter and ladle metallurgy (see the article “Steel Melt Processing” in this Volume). Fluxing is also commonly done to some extent with virtually all nonferrous mill and foundry melting operations. In particular, fluxing is used extensively in controlling dross formation and removing oxides from aluminum melts and melting furances. The basic types of aluminum fluxes are: Cover fluxes designed to barrier “blanket”
on the surface of molten aluminum
the article “Casting of Copper and Copper Alloys” in this Volume) Fluxing of zinc alloys with chloride-containing fluxes that form fluid slag covers, which can be used to minimize melt loss if they are carefully skimmed from the melt before pouring (see the article “Casting of Zinc Alloys” in this Volume) Degassing. Removal and reduction of dissolved gases in molten metal is typically accomplished with injection of inert gases in the melt (see the article “Degassing” in this Volume). In addition to dissolved gases, other sources of gas porosity include various mold-metal reactions (see the article “Gases in Metals” in this Volume). Molten-Metal Filtration. In addition to fluxing, inclusions in molten metal can also be controlled with various types of filters, such as screens or foam. Filters are commonly used in both aluminum melting furnaces and in-mold systems (see the article “Molten-Metal Filtration” in this Volume). Filters are also used in steel foundries. Molten-metal handling systems are also important components in the effective transfer, holding, and pouring of melts. After the melting process, molten metal may be handled and processed in various systems that may include:
Removing inclusions, impurities, or dissolved
Drossing fluxes with wetting action that
Launders, which are channels for transport-
gases from the melt Preventing excessive formation of solidphase inclusions Preventing and/or removing oxide buildup on furnace walls Injecting or distributing alloying constituents
promotes coalescence of aluminum Melt-cleaning fluxes designed to remove aluminum oxides from the melt Wall-cleaning fluxes specifically designed for the softening and removal of excessive aluminum oxide buildup that occurs on melting furnace walls
Thus, fluxes are commonly used to some extent with virtually all molten-metal operations in both the foundry and in the production of mill products. Fluxing can occur throughout the various stages of melting and handling of molten metal—from the melting or holding furnaces to the transfer ladles. Molten-metal fluxing agents may be chemically active solids, inert gases, or a chemically active gas. The type of flux depends, of course, on the composition of the melt and refractories. In iron and steelmaking, for example, the function of fluxes is twofold:
These fluxing treatments of molten aluminum are described in more detail in the article “Aluminum Fluxes and Fluxing Practice” in this Volume. Fluxing of other nonferrous alloys includes:
ing molten metal from melting units to tundishes, ladles, or molds Tundishes, which are intermediate holding vessels that have the ability to regulate molten-metal flow using slide gates, sliding nozzles, or a stopper Holding furnaces or transport crucibles that maintain a melt (or superheat) over a period of time Molten-metal transfer pumps Teeming ladles Dosing furnaces for delivery of controlled quantities (doses) of molten metal to the mold or die
Fluxing of magnesium alloys using salt
fluxes or inert gas as a cover (see the article “Magnesium and Magnesium Alloy Castings” in this Volume). Fluxing of copper alloys to remove gas or prevent its absorption into the melt, to reduce metal loss, and to remove specific impurities and nonmetallic inclusions (see
This article describes these typical systems for transporting, holding, or delivering molten metal to the mold/die system. Specific moltenmetal treatments and refractories are covered in other articles, as noted in the preceding overview on fluxing.
174 / Molten Metal Processing and Handling
Launders Launders are refractory-lined steel troughs or channels for transporting molten metal. The refractory may be brick or, if the launder size is not excessive, cast and sintered as single pieces of refractory. Launders often incorporate projections (dams and weirs) into the line of flow of the molten metal to help prevent the transfer of slag into the molds during pouring. Dams are the projections rising from the bottom of a launder. Weirs are projections coming down into the top of the pour stream. Launders are used in the tapping of blast furnaces (several hundred in 30 min) and cupolas (1=10 that of blast furnace tapping). Launders are also used in smaller furnaces, where melt is transferred from one chamber to another (see, for example, the article “Vacuum Induction Melting” in this Volume). With systems opertaing under atmospheric conditions, the launder is usually covered so the metal is not
Fig. 1
Fig. 2
Electric low-energy holding furnace with laundered inlet. Courtesy of Schaefer Furnace Inc.
directly exposed to the atmosphere. The covers are heavily insulated and are hinged to permit inspection of any portion of the launder. The trough is lined with heavy insulation having low thermal conductivity. This insulation is also nonwettable by the molten metal. Metal conveyance free from turbulence throughout the launder system is important, so the metal moves smoothly and continuously usually under a protective layer of stationary metal oxides. Most importantly, metal temperature fluctuations are thus significantly reduced, and manual work is lessened. The size of launders varies, of course, depending on the scale of the production operation. Small holding furnaces may have an integrated launder inlet (Fig. 1). A much larger launder unit is shown in Fig. 2 with its connection to a tilting aluminum furnace. In this case, the tilting furnace is connected to the launder with two cast iron parts. These cast iron parts need maintenance after each cast in order to keep them working properly. If they start to leak, it takes quite some time to prepare them or replace the cast iron pieces before the next cast. Periodic skimming of the launder is also another maintenance task. Flexible launder joints (Fig. 3) are also used for the transfer of molten aluminum. They are very lightweight and easy to install. One-layer launder joints are used only once because of possible leakage. Flexible launder joints also include multiple-layer designs that can extend the load-carrying capacity and be used multiple times. The thickness and density of the blanket and the hot-face cloth can be adjusted according to the operating temperature, possible chemical reaction, and whether or not there is any flame contact.
Transfer area from a tilting aluminum furnace to a launder. Courtesy of AGN Nachrodt
Tundishes The tundish is an intermediate holding vessel used in steel continuous casting systems or in the handling of other metals such as superalloys. The metallurgical role of a tundish is to facilitate separation of inclusions and slag, which floats with sufficient residence time of the metals in the tundish. Tundishes also perform the function of a metal reservoir that allows unabated casting during ladle changeovers for sequence casting of heats. In this case, it is imperative that the ladle change be done in the shortest possible time. The tundish is a ceramic bricked or cast tub into which the metal flows after leaving a teeming ladle. Some systems are designed to pour directly from the melting crucible into the tundish. A typical tundish design has significant depth and provides a low-velocity flow path from the point of entry of the metal to the bottom installed nozzle through which the metal exits into the molds or feeding system. Tundishes are preheated prior to metal pouring and are often covered to minimize radiation heat losses. Tundishes may also be heated. For highly deoxidized steels, they are protected with an argon gas cover to reduce reoxidation and inclusion formation. An artificial slag is usually added to absorb inclusions that are floating out. Metal is poured away from the nozzle location directly on a wear resistant pad. A constant metal height above the tundish nozzle is maintained to discharge the metal at a constant rate and, in turn, to maintain constant casting speeds. The height of metal in the tundish is kept constant by regulating the metal pour rate from the ladle. The height of the liquid head in the tundish makes the flow (pour rate) out of the tundish nozzle a fixed rate. For specialquality steels or other melts with stringent surface and cleanliness requirements, shrouding is commonly employed to protect the steel from air, particularly for aluminum-killed grades, which have a potential of forming alumina inclusions. The shrouding systems include inert gas shrouding and refractory-tube shrouding. The concept of a refractory-tube shrouding
Fig. 3
Flexible launder joint in aluminum unit. Courtesy of Alunorf and Ceranex
Transfer and Treatment of Molten Metal—An Introduction / 175
Fig. 4
Schematic of a refractory-tube shrouding system for minimizing oxidation during pouring. Courtesy of Metaullics Systems Division, Pyrotek, Inc.
system is shown in Fig. 4 for ladle-to-tundish and tundish-to-mold pouring. Refractory-tube shrouds for steel are made of fused silica or alumina graphite. Nominal residence time of the metal in the tundish is important, because this controls the extent to which entrained slag will float to the top of the tundish and be removed from the pour. This has led to a trend toward tundishes with larger volumes and thus longer residence times for a given throughput. Flowcontrol devices, such as refractory dams and weirs, are attached to the tundish to distribute the metal flow, to minimize turbulence, and to eliminate dead flow zones (Ref 1). Refractory baffle walls in tundishes are also used to help direct the flow of liquid, so inclusions have an easier chance to float out at a later point. The holes in the baffle walls are typically approximately 50 mm (2 in.) and do not catch or trap inclusions (as in the case of molten-metal filters for inclusions in aluminum melts). Dams and weirs also are used in the tundish to collect slag particles. Bottom-projecting dams direct metal flow to the top of the tundish as a flotation aid for the slag particles. Flow modifiers have no influence on the very small inclusions collected by ternary refining, but they can help clean the steel of midsized inclusions in the 50 to 200 mm (2 to 8 mils) range. Flow controls are not needed for larger inclusions (which have higher Stokes rising velocities). The refractory lining in a tundish includes a service lining and a back-up safety lining. Refractories in the tundish are subject to corrosive wear, slag penetration, and thermal shock. Until recently, refractory wear in a tundish was usually of secondary importance to the ease of installation and tearout, as well as metallurgical considerations of the steel. During deskulling, the hot-face service lining is usually lined with a thin, 2.5 to 7.5 cm (1 to 3 in.), coating layer of olivine (magnesium iron silicate), MgO, or a mixture of these. This coating is
usually a slurry spray material or a dry vibratable powder that is installed behind a steel form that is heated to set a phenolic resin binder. With the drive for longer and longer cast sequences, the emphasis of tundish linings is changing toward more corrosion-resistant coatings that last longer and are cost-effective but that also maintain good deskulling properties. In steel continuous casting, the tundish continues to play a critical role in postfurnace steelmaking practice (along with steel ladle metallurgy treatments, as described in the article “Steel Melt Processing” in this Volume). Properly designed tundish flow modifiers provide substantial benefits in terms of both quality and productivity, and extensive work in physical and mathematical modeling has been done to understand and improve the fluid flow pattern in the tundish (Ref 2–4). Mathematical prediction of flow pattern has become more widely used in analyzing turbulence. Molten-Steel Filters. For critical applications of steel, the presence of inclusions with a diameter greater than approximately 50 mm (2 mils) needs to be prevented. There was a move in the industry in the 1990s to use filters (foams or screens) for such applications to help eliminate inclusions under approximately 50 mm (2 mils) in size, which are not susceptible to flow modifiers. However, this use of filters in tundishes for steel continuous casting did not prove effective, because the large number of inclusions would quickly clog the filter.
Molten-Metal Pumps Molten metals have traditionally been transferred from the melting vessel to a mold or casting unit by tapout or lip pouring (essentially gravity flow systems) or by using an intermediate ladle. Molten-metal transfer/recirculation pumps are now commonly used with aluminum, magnesium, zinc, and lead alloys. The pumps may be a mechanical device or a device operating on the principle of electromagnetic pumping. Mechanical pumps are typically centrifugal pumps with a rotating impeller in a volute casing. They can be used as either transfer pumps or as recirculation pumps in melting or holding units. Mechanical pumps are also used for gas injection. Electromagnetic pumps are also widely used in industry, because they have no moving parts, so routine maintenance is minimized.
Mechanical Pumps Mechanical pumps for molten-metal recirculation and transfer (Fig. 5) traditionally have graphite posts, which must be rebuilt approximately three times per year when they become oxidized or broken. Newer techonolgy includes molten-metal pumps with ceramic posts (Fig. 5) that are held in compression from a tension rod made of high-temperature alloy steel.
By replacing the posts with a stronger material that will not oxidize, the life of the pump has been extended, with fewer rebuilds. In this design, fragile ceramic sleeves are no longer needed to protect graphite posts from oxidation and abrasion. Ceramic sleeves can also hide post oxidation, which could result in unexpected failures. The ceramic posts held in compression are stronger than graphite; if they do break, the base is still supported by the steel rods. Mechanical recirculation pumps are used to improve heat transfer by forced convection. This is particularly important in aluminum reverberatory furnaces (see the article “Reverberatory and Stack Furnaces” in this Volume). In reverberatory furnaces, forced convection keeps the entire melt at a more uniform temperature, so that the burners above the melt transfer more radiant heat to the surface of the melt. If the surface temperature becomes hotter than the lower parts of the bath, then the heat transfer for melting become less efficient in reverberatory furnaces. Other nonferrous alloys, such as zinc, magnesium, and copper, also benefit from molten-metal circulation, but pumping systems for circulation are not usually employed. Mechanical stirrers are effective, particularly for zinc, due in part to the lower heat of fusion of zinc and the smaller melt sizes usually used. Mechanical transfer pumps are now commonly used with aluminum, magnesium, zinc, and lead alloys. A mechanical pumping system has several advantages. Pressurized flows permit larger volumes of metal to be transferred per unit time, with resultant productivity gains. The ability to move greater volumes of metal in shorter times also means that less superheat must be put into the metal to accommodate temperature losses during transfer. Thus, shorter overall furnace melting cycles are possible. Because most metal-transfer pumping systems have submerged intakes, dross entrainment from the surface is avoided, providing better metal quality. Transfer pumping per se does not produce additional oxidation or dross losses, because molten nonferrous metals are always subjected to turbulent flows (high Reynolds numbers), even with gravity flow tapout systems. Transfer pumps can often be used with closed piping systems to provide transfers surmounting physical obstacles (over other equipment, for example) or overcoming head and height limitations when gravity flow is limited or not possible. Melting and casting facilities can therefore be mixed and matched and need not be directly in line with one another. Transfer pumps are now commonly used in many segments of the aluminum casting industry to move molten metal from melting furnace to holding furnace, to feed or empty auxiliary refining vessels (used for alloying, degassing, and filtration), to empty and fill transfer ladles, to directly pour castings and fill molds, or to feed casting machines in the wrought industry. In the magnesium, zinc, and lead metal casting industries, mechanical pumps transfer metal
176 / Molten Metal Processing and Handling
Fig. 5
Fig. 6
Molten-metal pumps with ceramic posts and steel tension rods. (a) Recirculation pump. (b) Transfer pump. (c) Pump elements. Courtesy of Metaullics Inc.
Schematic of metallic centrifugal pump used to transfer magnesium
between furnaces, fill ingot and pig molds, feed die casting machines, cast electrogalvanizing anodes, and empty galvanizing kettles when repairs or desludging is required. Molten copper alloys are also transferred by pressure pump systems. Of the various kinds of mechanical pumping systems, a positive-displacement pressure pump is often quite useful for small furnaces or small flows (<45 kg, or 100 lb). When actuated, external air pressure displaces the volume of air in the cylinder, which forces an equivalent volume of liquid metal through the delivery tube. Another type of pump is the siphon pump. Both of these pumps are quite capable of filling small molds where the accuracy of fill demanded is moderate (several percent variation) and where pressure head requirements are minimal (0.3 to 0.6 m, or 1 to 2 ft, maximum). When larger volumes of metal, more sustained flow, or higher head is required, a centrifugal pump can be used. Flow is directed by an impeller up through a riser tube. Operated either by air or electric drive, these pumps are
typically constructed of graphite and are easily operated in molten aluminum, magnesium, and zinc alloys. It is important to keep the riser tube open, using insulation and careful external heating if necessary, to prevent freezeups, especially if use is intermittent. Another type of centrifugal mechanical pump often used with magnesium alloys is shown in Fig. 6. Pump construction is usually cast iron, low-carbon steel, or 400-series stainless steel. Nickel impurities in metal pump components should be specified to be below 0.20% to avoid the dissolution of nickel in the molten magnesium and the subsequent reduced corrosion resistance of magnesium castings. Centrifugal pumping systems for zinc are usually constructed of graphite, but metallic centrifugal pumps are often used as well. Pumping systems for lead are usually constructed with cast iron because little corrosion is experienced. Cast iron can also be used to pump zinc as long as the temperature is kept below 480 C (900 F), above which molten zinc begins to dissolve iron. Austenitic stainless
Transfer and Treatment of Molten Metal—An Introduction / 177 steels can be used as pump materials with aluminum-containing zinc alloys (0.2 to 3.5% A1). Stainless steel must not be used, however, with pure zinc, because liquid metal embrittlement and intergranular attack occur quite rapidly. With zinc-aluminum galvanizing alloys (40 to 60% Al), the high operating temperature (595 to 650 C, or 1100 to 1200 F) is in the region where austenitic stainless steels become sensitized to greater intergranular carbide precipitation and subsequent accelerated corrosion. Stainless steel pumps should therefore not be used for this application. A useful but expensive alternative for the latter two applications is an Mo-30W or Mo-40W pump shaft, which provides excellent corrosion resistance but is fairly brittle at high temperatures.
Electromagnetic Pumps Electromagnetic pumps have been developed for circulation of molten metal and controlled mold-filling capabilities. This type of pump is constructed from ceramic materials and has a linear motor that induces pumping on the basis of the left-hand rule of electromagnetic field theory, using the molten metal as the secondary field. It is particularly useful for metals with good electrical conductivity. One specific advantage of this type of pump is that there are no moving parts. However, the moltenmetal channel opening must be kept very clean and free of oxide buildup or dross; otherwise, the field strength and resulting pumping capacity can be greatly reduced. Electromagnetic Pumps for Furnace Circulation and Scrap Recycling Technology. Electromagnetic (EM) pumping systems are used for application of forced circulation in both foundry and casthouse aluminum melting. The EM pumps are supplied to suit the furnace size and application, with mass flow capability from 1 metric tonne per minute to greater than 20 metric tonnes per minute. All pumps are operated at main voltage and frequency and are fully controllable for both speed and direction. The consumable pump tube is replaced approximately every 6 months, especially with mixed-charge scrap melting, although with certain applications, up to 18 months of continuous operation are possible. An EM pump is especially suitable for those melting furnaces that may not have sidewells, or furnaces that are not stationary but rather tilt-pour. With the consolidation in the foundry industry, larger, volume-oriented castings producers are increasing their central melting furnace sizes well above 10 tons, making the application of an EM pumping system to circulate the furnace a viable option. In terms of operation, an EM pump is connected via insulated tubes to the furnace charge chamber and an external refractory-lined charge well, where a unique internal bowl shape develops a vortex with, low-turbulence scrap submergence (LOTUSSE Metaullics
Systems) in the molten-metal stream. The vortex creates an ideal entry point for high-surface-area, low-density scrap and alloying materials that need rapid submergence to avoid oxidation with direct flame impingement. These include machining chips, borings, turnings, flashings, used beverage containers, and so on. While many foundries would only melt a small portion of such materials, cast wheel producers as well as producers of castings with inhouse extensive machining produce a sufficient quantity of such light-gage materials (machining chips) to make the EM pump/LOTUSS scrap recycling technology quite feasible and economical. Indeed, the same LOTUSS scrap recycling technology can be employed with mechanical pumps as well. The choice of which system to employ—EM pumping or mechanical pumping—will depend on many specific user’s economical and process factors. Metal flow capability and electrical requirements of typical EM pumping parameters are: Metal flow rate, t/min
Kilovolt amperes (VA) rating
8–10 8–10
110 110
15
130
8–10
110
15
130
20 12–14
150 110
Features
100 mm (4 in.) pump 100 mm (4 in.) pump and redesign tie bars 100 mm (4 in.) pump and laminations 150 mm (6 in.) pump and redesign tie bars 150 mm (6 in.) pump and laminations 150 mm (6 in.) pump, 9 poles 200 mm (8 in.) standard pump
Use of a either a centrifugal or electromagnetic metal pumping system has the following benefits: Increased metal yield (typically for dry
Alloying materials in foundry melting are prone to problems with dissolution, incomplete melting, and less-than-optimum recovery. The combination of rapid-submergence charging capability, coupled with the forced convectional circulation provided by the pumping action, helps to ensure quick dissolution and higher recovery than would otherwise be possible. Pure copper, silicon, or aluminum master alloys can be processed by this technology. Smaller lump forms, granules, and small ingot bars are preferable to powdered additions. Recent advances (Ref 5) in EM pumping technology have integrated several other functions into the heretofore single application of forced circulation. The EM pumps are now capable of providing transfer of metal out of the closed system into a transfer ladle, launder, or a holding furnace, and with suitable gas injection systems, degassing is also possible. Additionally, such systems can be outfitted with noncontact sensor devices (video camera, level sensor, etc). Mold-Filling EM Pumps. Electromagnetic pumps are easily used for controlled mold filling in such applications as permanent mold casting, large die casting, low-pressure die casting, and perhaps small ingot casting. They offer exceptional control of metal flow throughout mold filling, and turbulence is minimal since there are no moving parts. This puts the highest-quality metal into the mold. Other benefits are: Metal filtration is easily performed using a
chips, 98%)
compared to a static bath)
yield; silicon, +95% yield)
Increased melt rate (typically 10 to 20% High-efficiency alloying (magnesium, +95% Reduced energy use (equivalent to increased
output)
Chemical and thermal homogeneity within
the furnace (þ /5 C, or 9 F, within the furnace) Reduced dross formation (depending on scrap types)
With the LOTUSSE scrap recycling system, many different materials are possible to be charged into the special-shaped bowl. However, it is not possible to charge large sows or large scrap castings through the system. The following are typical smaller, lighter-gage aluminum materials that are suitable for processing in this vortex scrap charging system: Machining chips, turnings, borings (which are
most prevalent in integrated foundry operations)
Shredded sheet scrap Extrusions
Large fragmented castings Small castings Scalper chips Used beverage containers Aluminum chips
filter over the pump inlet, filtering all metal to the casting. The response to a desired change in flow rate is very fast. Plant safety is increased due to automation— no need for hand ladling. The alloy in the furnace can be tested at any time, and additions made without interfering with production. In low-pressure casting, the furnace is not a closed pressure vessel. The temperature of the metal delivered by a pump is very close to that of the moltenmetal bath, so superheating of the furnace is not required.
Types of EM Pumps. The types of EM molten-metal pumps include:
Single-phase conduction pumps Jumping ring pumps Multiphase linear pumps Pinch-effect pumps Direct current conduction pumps
Single-phase EM pumps (Fig. 7) can be applied to almost any metal casting process. They are particularly suited to low-pressure casting,
178 / Molten Metal Processing and Handling
Fig. 7 Example of two sizes of single-phase conduction pumps with metal filters over the bottom inlet where they fill the mold from below while exhibiting very good control over the rate of metal delivery, along with the ability to almost instantly change the flow rate. Repeatability is generally excellent. It is also very easy to hold the metal just below the mold between castings to minimize oxide generation and to reduce overall cycle time. Other uses include squeeze casting through the side of a vertical shot sleeve, just above the top of the shot tip, to eliminate dross from top-filling; high-pressure die casting; bottom-fill lost foam; and all kinds of gravity casting. Single-Phase Conduction Pump. The singlephase conduction pump uses alternating current applied to a coil of many turns of wire that is wound around a laminated ferrous core, which constitutes the primary winding of a transformer. Around the same laminated core is a ceramic material that is compatible with aluminum, and this ceramic has an internal passage into which molten aluminum flows from the metal bath. This passage forms a complete circuit (loop) around the laminated core; hence, the molten aluminum therein forms a one-turn shorted secondary winding on the transformer. By virtue of this single turn, a high-amperage flow is induced into the molten aluminum. A portion of the loop is flattened to allow for close spacing of the magnet pole pieces and to make the current density in that portion of the loop relatively uniform throughout the cross section. This is the area (pumping vein) in which the pumping force will be produced. Two pole pieces are placed as closely as possible to the broad face of the pumping vein, and additional electrical coils are around the laminated core that forms the magnet pole pieces. This creates an alternating magnetic field across the pumping vein. Since the transformer primary and the magnetic field coils are powered from the same alternating current source, both the current flow in the metal and the applied magnetic field alternate polarities together. If either the current in the metal or
the magnetic field were static while the other one alternated, the metal would experience an alternating force, first in one direction, then in the other. Since the magnetic field and the current in the metal are both alternating together, the force on the metal simply pulsates from zero to maximum at twice the applied line frequency. These pressure pulsations may prove beneficial when pumping thixotropic materials such as metal-matrix composites, because the pulsations introduce shear into the moving metal, thus greatly increasing the fluidity, which helps the metal to flow into mold details. Since the electrical current is induced directly into the molten metal, there is no need for electrodes in contact with the metal, also eliminating the need for heavy cables and very large direct current power supplies. This is the type of pump most commonly used for mold-filling applications in aluminum foundries, because of its ease of use, reliability, and repeatability. Jumping Ring Pumps. If a coil of wire is wound around a straight ferromagnetic bar or rod, and a ring of any conductive metal is placed near it over the same bar, with alternating current applied to the first (primary) coil, the ring is very strongly repelled from the coil. This is called the jumping ring pump. These pumps do not generate much force in the molten metal and are most often used for stirring metal in a furnace. Multiphase Linear Pump. If a multiphase electric motor stator (the stationary part with the electrical windings) were cut open and rolled out so the windings are all in a line, when power is applied, there would be a moving magnetic field. If a conductive material is placed above this stator, electrical current would be induced into that material. This would also produce corresponding magnetic fields, which would interact with the moving magnetic field, and the interaction between these magnetic fields would produce a force on the conductive material, tending to move it in the direction of the moving magnetic field. As described, this would be a very inefficient pump due to poor control over the magnetic field. There are two common ways to improve this. One is to place a plain pole piece opposite the stator while keeping the molten metal between the stator and the coil assembly. The second improvement is that the stator could be rolled to form a long tube, with the moving magnetic field traveling down the length of the tube. This type of pump is typically used for high-volume metal transfer and stirring of metal in large furnaces. Pinch-Effect Pumps. Another form of pump induces current into a cylinder of molten metal that has a tapered ceramic mandrel in the center. Since the metal is trying to collapse on itself, it is forced toward the narrower end of the taper, which is also the metal discharge area. These pinch pumps typically use a highfrequency coil to improve the efficiency of the induction, and they also typically use water cooling to keep the coil from overheating.
Direct current conduction pumps use electrodes to pass a large direct current into the fluid to be pumped, and a permanent magnet supplies the magnetic field. These pumps are not suitable for use with aluminum, since they would require electrodes to be in contact with the molten metal, and there is no suitable material available for this purpose. Additionally, the amperage required is very high, which would need very heavy cables and a direct current power supply capable of supplying as much as 10,000 A.
Pouring and Dosing As noted, molten metals traditionally have been transferred to a mold or casting unit by tapout or lip pouring by using an intermediate ladle. Traditional metal transfer from the molten-metal supply to the mold is done with open-top ladles ranging from hand held to multiton, crane-supported, mechanized tilt ladles. Underpouring from a ladle, either through the incorporation of a teapot spout or bottom stopper rod and nozzle, reduces slag inclusions in castings (Fig. 8). In addition to gravity-feed pouring, a number of other methods are also used to fill molds. For example, molten-metal pumps (as briefly described in the preceding section of this article) are used. Other methods include the use of dosing furnaces in die casting operations (e.g., see the article “Automation in High-Pressure Die Casting” in this Volume). Dosing furnaces are replacing ladle furnaces and are described in the next section of this article. Another method is low-pressure filling systems for sand casting of heavy metals. In lowpressure pouring (Fig. 9), the molten metal is pushed into the open mold from below. Compared with gravity casting, the low-pressure method offers a number of important advantages: Reliable, simple, and tight mold-to-nozzle
docking action
Adaptable pressure buildup for active control
of mold-filling characteristics
High dosing accuracy based on pressure and
(where applicable) level control
Essentially laminar mold-filling flow Cleaner melt supply
Finally, automated methods encompass a wide range of mechanized methods and specialized equipment. This includes mechanized ladles and pouring robots that are progarmmed to transfer measured amounts of metal from a storage vessel or a controlled pouring unit to the mold. Pumps and dosing furnaces are also suited to integrated automation. Benefits of automated pouring include: Increased productivity due to the speed and
consistency of pours from one mold to the next
Transfer and Treatment of Molten Metal—An Introduction / 179
Fig. 8
Four types of pouring ladles. (a) Open lip-pour ladle. (b) Teapot ladle. (c) and (d) Bottom-pour ladles
Improved process control with tracking of
production and inventory information. In some cases, automatic pouring equipment also allows a reduction in pouring temperatures with significant savings in energy and refractory costs. Improved working conditions in terms of health and safety, environmental, and operational factors
Pouring Methods
Fig. 9
Low-pressure mold filling. (a) Bottom fill. (b) Fill from top of the mold
Increased yield due to fewer overpours,
spills, and allowance for smaller pour basins. Consistent, controllable pour rates also may allow for the design of less complex gating and riser systems as well as for simultaneous multiple-pour streams in the same mold. Reduced scrap costs due to reduction of inclusions in castings and the reduction of pouring failures such as underpours, interrupted pours, and irregular pour rates Improved quality control with better temperature control within the pouring vessel, more consistent metallurgy, and automatic feeding of alloy and inoculant additions into the pouring ladle or metal pour stream
The two key variables in pouring are pouring rate and duration. The optimum pour rate is the maximum rate at which a mold will accept metal, and the optimum duration is the time from the start of the pour until the mold will no longer accept metal. Other important considerations in automatic pouring systems include the relative positioning of the mold and pour spout or nozzle and the sequencing of the pour. In manual pouring, the worker compensates for variables in the mold, ladle, and metal while pouring to produce the best results for each pour. Duplication of this effort is the challenge of automatic pouring. Optimum pouring time is based on the weight and shape of the casting, the temperature of the molten metal and its solidification characteristics, and the heat-transfer characteristics and thermal stability of the mold. Most foundries may also pour many different castings from one heat, or even from one ladle. Therefore, these foundries do not attempt to control pouring time, but they do control the speed at which molten steel enters the mold cavity. This control is achieved through the design of the gating system. Pouring Ladles. The three basic types of pouring ladles are: Bottom pour Teapot Lip pour
Ladle capacity normally ranges from 45 kg to 36 Mg (100 lb to 40 tons), although ladles having much larger capacities are available.
Ladle Linings. Ladles with capacities above approximately 360 kg (800 lb) are lined with layers of firebrick covered with a thick refractory coating. Smaller ladles are protected by monolithic linings prepared by ramming or by pouring and vibrating refractory mixtures around forms. For molten steel, the optimum refractory for this purpose is high-purity alumina, but because of the high cost of alumina, alternative materials such as olivine and zircon are more widely used. The service lives of these monolithic linings, which are most commonly used for small ladles, vary considerably, depending on the refractory mixture and the type of service. Thermal expansion and contraction can cause hairline cracks, into which hot metal enters. This results in erosion of the lining. Erosion not only shortens the life of the lining but also contributes to excessive dirt in the steel. Disposable fiber linings that do not require preheating are gaining acceptance. Preheating the ladle helps to prolong the life of the ladle lining. Heating the lining to a temperature close to that of molten steel serves to lessen the thermal shock to the lining when the molten steel contacts it. Common practice is to preheat ladle linings to 815 to 1095 C (1500 to 2000 F) with oil-fired or gas-fired torches. Ladles lined with high-purity alumina can be preheated to temperatures as high as 1370 C (2500 F). Bottom-pour ladles have an opening in the bottom that is fitted with a refractory nozzle. Stopper rods (Fig. 10a) were commonly used to control flow from the unit, but more typically a valve-type mechanism (Fig. 10b) is often used. Bottom pouring is preferred for pouring large castings from large ladles, because it is difficult to tip a large ladle and still control the stream of molten steel. Also, the bottom-pour ladle delivers cleaner metal to the mold. Inclusions, pieces of ladle lining, and slag float to the top of the ladle; thus, bottom pouring greatly reduces the risk of passing nonmetallic particles into the mold cavity. Pouring ladles may also include options for melt treatments or shrouding. For example, Fig. 11 shows a schematic for a bottom-pour
180 / Molten Metal Processing and Handling
Fig. 13 Fig. 10
Bottom-pour mechanisms (a) stopper rod (b) sliding value
Fig. 12
Fig. 11
Bottom-pour teeming with argon insertion ring. The term teeming typically applies to the pouring of molten metal from a ladle into ingot molds.
mold setup with an argon insertion ring that shrouds the pour from oxygen and nitrogen pickup from the atmosphere. Mold powders or fluxes are added to the advancing metal front to improve surface quality by providing thermal insulation from the mold surface. At the end of the pour, exothermic compounds may be added to the top of the mold to insulate the top of the solidifying ingot and provide a degree of hot
Typical teapot ladle used to pour small- to medium-sized castings
topping, where the molten metal will fill the shrinkage occurring upon solidification lower in the ingot. Argon-oxygen decarburized electrodes also may be top poured, but the need to continually close the ladle tap, transfer the ladle, and reinitiate argon shrouding produces variable conditions from ingot to ingot. Should difficulty in closing the nozzle be experienced during top pour, then a safety problem arises. Thus, bottom pouring is the more generally used process. On the other hand, it is impractical to pour molten metal into small molds from a large bottom-pour ladle. The pressure head created by the metal remaining in the ladle delivers the molten metal too fast. Also, the time required to fill a small mold is short, thus requiring that a large bottom-pour ladle be opened and closed many times in order to empty it. This may cause the ladle to leak, although special nozzles have been developed to minimize leakage. Despite the fact that the size of bottom-pour
Schematic of automatic mechanized ladle
pouring
with
ladles could be scaled down for pouring smaller castings, this is unnecessary because of the almost equal ability of the teapot ladle to deliver clean metal. The teapot ladle incorporates a ceramic wall, or baffle, that separates the bowl of the ladle from the spout. The baffle extends almost four-fifths of the distance to the bottom of the ladle (Fig. 12). As the ladle is tipped, hot metal flows from the bottom of the ladle up the spout and over the lip. Because the metal is taken from near the bottom of the ladle, it is free of slag and pieces of eroded refractory. The teapot design is feasible in various sizes, generally covering the entire range of casting sizes that are below the minimum size for which the bottom-pour ladle is used. Lip-pour ladles resemble the teapot type but have no baffles to hold back the slag. Because the hot metal is not taken from the bottom of the ladle, this type of ladle pours a more contaminated metal and is seldom used to pour high-alloy castings. Nevertheless, it is widely used as a tapping ladle (at the melting furnace) and as a transfer ladle to feed smaller ladles of the teapot type. Methods of Automatic Pouring. There are two basic methods of automatic pouring: mechanized ladle pouring and direct pouring. Mechanized ladles generally have capacities matched to single mold requirements and include load cell weighing systems and controlled tilt rates (Fig. 13). In the case of direct pouring, metal is metered directly into the mold from the pouring vessel, with the appropriate controls to ensure the desired results. The use of a mechanized ladle allows for inladle alloying and inoculation before pouring. Dripping of molten metal onto the mold is eliminated when a mechanized ladle is used in combination with a stopper rod and nozzle
Transfer and Treatment of Molten Metal—An Introduction / 181
Fig. 14
Fig. 15
Automatic pouring into molds on continuously moving line. (a) Pouring into flaskless molds on an indexing line from bottom-pour unit with stopper rod. (b) Synchronized pouring on a continuously moving line with mechanized ladles capable of x-axis and y-axis travel and synchronization with the pouring line.
Rotary mechanical pourer with molten metal level detector and control
pouring unit. This is very important when pouring lost foam molds. Ladle tilting about an axis in line with the pouring lip can reduce moltenmetal stream travel during pouring. The capacities of mechanized pouring ladles are based on anticipated metal requirements. The dispensing of a metal stream from a ladle at a specific point requires that the nozzle or pouring lip not move, relative to the mold,
during pouring. In semiautomatic pouring units, bottom-pour ladles may be mounted above indexing molding lines to pour into molds indexed into position beneath the ladles (Fig. 14a). Synchronized pouring on a continuously moving mold line (Fig. 14b) may also require the use of a pouring vessel to measure predetermined weights of metal and either one or two (depending on line speed) mechanized
pouring ladles. Another approach is a mechanized rotary pourer with optical or laser monitoring for automatic control of melt level (Fig. 15). Direct bottom-pour ladles use either slide gates or stopper rods and nozzles to control and stop the flow of metal to the mold. Slide gates and nozzles are most commonly used to control flow rates on bottom-pour ladles for continuous and semicontinuous strand casting operations. Slide gates (Fig. 10b) offer more reliable shutoff than stopper rod valves; however, service life of the slide gate can be shorter than that of the stopper rod. Stopper rods and nozzles offer more control of pour rates than slide gate valves and are less expensive. For these reasons, bottom-mounted stopper rods and nozzles are more common in foundry pouring operations. Controls of Pouring. Satisfactory pouring results can be obtained in some automatic pouring applications without complex electronic control systems. Some requirements can be met by establishing a flow rate and simply timing the pour. A slightly more complex system includes dual pour rates and pour times. Simple infrared photosensors designed to detect the presence of molten metal can be used to end the pour cycle. If the exact weight of metal required for a mold is known and a mechanized ladle is weighed while filling, at least one variable is under control. If the movements of the operator controlling the pouring of a mechanized ladle are measured and then stored in a programmable controller, these movements and the results can be duplicated under automatic control. The programming of a controller by duplication of remote control is referred to as a teach-in system. The program is referred to as the pouring profile. Weighing of a bottom-pour ladle and pouring by weight can also be effective in some cases, although the relatively small capacity of the
182 / Molten Metal Processing and Handling ladle results in substantial variations in metal head pressure above the stopper and subsequent variations in pour rate. Increasing the capacity of the pouring vessel, facilitated by electric heating, reduces the variation in head pressure. The incorporation of a process control system to compensate for reduced head pressure through stopper positioning by feedback of weight, anticipated weight, or actual bath height is effective in some cases. If the rate of metal flow through a nozzle is constant for a constant head pressure, the repeated variations in the position of a stopper rod above the nozzle will yield repeatable and controlled pouring rates. A teach-in system is used to program the pouring profile. If slag buildup on the nozzle and variations in metal head pressure cause flow variations, the system pouring profile program must be corrected by a process control feedback loop. Feedback can be provided by a timing and metal-presence sensing system that checks the rate of mold fill or by a more sophisticated control system for mold-level measurement. Mold-level measurement and control systems fall into two categories: laser systems and vision systems. Both of these electronic process control systems monitor the level of metal in the pour cup and control the position of the stopper rod to maintain a nearly constant level throughout the pour. Optical sensors provide input to help regulate the electromechanical stopper drive in such a way that the liquid level in the filling head is kept at a constant level. This type of feedback system most closely duplicates the efforts of a skilled operator. Camera or laser level sensors are used depending on production factors. Manufacturers who mainly use a sprue cup diameter of 85 mm (3.3 in.) or larger will usually select the camera-based solution, whereas the laserbased solution will be used particularly where the focus is on minimizing the circulation material and a sprue cup diameter smaller than 80 mm (3.1 in.) is employed. The laser reflection angular detection system (Fig. 16a) determines the position of a lowpower laser beam reflected from the surface of the molten metal in the pouring basin. The laser beam is directed at the pouring basin from an angle. The distance to the molten-metal level will affect the incident and reflected angle and therefore the position of the beam on a sensing unit. The comparative value of the beam position on the sensor is used to control the position of the stopper rod; thus, the level of the metal in the pouring basin is controlled. A vision system uses a video camera to observe the pouring process (Fig. 16b). The video signal is displayed on a monitor for the operator and processed electronically in a manner that separates the brightness levels of the molten metal from the other portions of the picture. The molten-metal stream is electronically blanked from the composite picture of the metal in the pouring basin and the area of metal fill compared to the actual area of the pouring
Fig. 16
Optical sensors for pouring control. (a) Laser sensor. (b) Vision sensor
basin. The comparative value is used to control the position of the stopper rod, and the level of the metal in the pouring basin is controlled. The most sophisticated pouring systems may incorporate more than one set of controls. The use of integrated control schemes allows the system backup, improved safety, and increased flexibility that are required for pouring a variety of castings. For example, the benefits of moldlevel measurement and control are enhanced when used in combination with a teach-in duplicated pouring profile in some pouring applications. Pouring Furnaces. Radiant element, electric resistance-heated pouring furnaces are used in both nonferrous and ferrous pouring applications. For nonferrous, low-temperature applications, metal or silicon carbide heating elements are used. For ferrous high-temperature applications, graphite rod heating elements are used. Radiant-heated furnaces are typically bottompour units with stopper rods and nozzles mounted in the furnace bottom or in a detachable pour box. The furnace can be tilted to drain the pour box and to contain the molten metal beneath the radiant heat source. Crucible-type electric resistance-heated and coreless induction-heated furnaces are most commonly used in nonferrous pouring applications. These applications vary from the simple dip-and-pour mechanized ladle systems to the underpour stream discharge, axial-pour, and trunion-pour furnaces. Underpour stream
discharge in aluminum pouring for continuous and semicontinuous casting processes uses the oxide skin on the metal surface to protect the metal stream from further oxidation while transferring metal from the melting furnace to the automatic pouring ladle. Channel-type electric induction-heated pouring furnaces are sealed and designed to use gas pressure to discharge metal from the furnace vessel. Pouring can be accomplished over the lip of a pour spout or through a nozzle, with or without a stopper rod, in a pour box mounted on the furnace vessel. Vessel gas pressure can be varied to control the rate of pouring or to control the level of metal over the pouring nozzle (Fig. 9b). Control of the level of metal in the pour box establishes a constant static head pressure above the nozzle and ensures consistent pour rates. Venting of the pressure in the vessel causes the metal in the pour box to drain back into the furnace. Inert gas can be used as the pressure medium in applications where the metal is particularly susceptible to oxidation.
Dosing Furnaces Dosing furnaces are replacing ladle furnaces in many applications, especially in aluminum high-pressure die casting (see the article “Automation in High-Pressure Die Casting” in this Volume). Aluminum castings traditionally have been produced by either direct pouring from the melting furnace into the mold, or from a
Transfer and Treatment of Molten Metal—An Introduction / 183 transfer ladle in which further treatment has been performed just prior to pouring the mold, or from a holding furnace dip-well from which the casting is poured either by automatic robotic ladle or by manual means. Increasingly, automated and controlled direct pouring is becoming popular, especially in high-pressure die casting operations where the dosing furnace technology can be meshed with the automated die casting machine controls. A number of different units exist commercially that are comprised of holding furnaces with discharge capabilities directly into the mold cavity. These devices provide a very precise dosage of molten metal and are thus known as dosing furnaces. The dosing operation (Fig. 17) begins with the introduction of compressed air into the pressure-sealed furnace, thereby causing the ascension of the melt in the riser tube. With further increase in the pressure, the melt is discharged from the riser tube and flows via the dosing channel into the shot sleeve of the pressure die casting machine. Flow of the melt is stopped by rapid discharge of the compressed air, the molten metal returns to the level of that in the dosing furnace, and the die casting machine can now execute the casting operation. There has been constant further development in dosing with the dosing furnace to fulfill casting demands. The following describes several of the key technologies that represent the current state of the art. The following demands are placed on the overall process: High and reproducible dosing accuracy Short main time and minimum waiting time No loss of time with fluctuations in the cycle
cycle times without further adjustments, thereby excluding human influence. It is only necessary to enter the melt temperature and dosing weight. The dosing furnace automatically synchronizes itself with the die casting machine, thus always resulting in an optimized dosing cycle. Detection of Melt Level. It has been and still is necessary to be able to detect the level of the melt in front of the discharge edge at the rising tube. This was previously achieved by an electrode that generated an electric signal on contact with the surface of the melt (Fig. 18a). In order to avoid deposits of aluminum, this electrode had to be rapidly swung out after contact by means of a pneumatically operated device. Heat, oil, spraying agents, and sometimes open fire often led to failures in this component. A pneumatic scanning system (Fig. 18b) today (2008) replaces the electrode for detection of the melt level. Pneumatic scanning employs no moving components, requires no adjustment, and is especially advantageous with regard to high operational reliability and the absence of maintenance in a particularly harsh environment. The scanning accuracy is a precondition for exact dosing and is very well fulfilled by this system. The high degree of reliability in a pneumatic scanning system allows a level-scanning process known as the “top-stop-position” that has been used in the dosing furnace technology since approximately 2005. In this process, the height of the melt in the riser after dosing is not reduced to the level of the melt in the furnace but is held a little below the discharge
edge (Fig. 19). The furnace is thus always ready for dosing, so that there is no delay in discharge of the melt on receipt of the signal from the die casting machine. This compensates for fluctuations in the cycle of the die casting machine, and there is no loss of time through nonsynchronized process sequences. As a rule, this system achieves up to an 8% savings in the cycle time. Because the dosing furnace is a closed vessel that cannot be looked into, the filling level previously had to be determined by a weight load cell. This method was inaccurate since external influences on the furnace were also detected and could falsify the reading. With the development of new technology, the level in the furnace is precisely detected during the dosing process by means of the dosing contact and the system, so that the result is free from external influences. Level scanning is used to ensure that the furnace also remains ready for dosing while it is being refilled with molten metal (Fig. 20). Level scanning is started after opening of the filling cover, the level and the required amount of refill being semicontinuously detected in short cycles with indication of the current level. Furnace Recharging and Melt Holding. The quality of the melt in the transport ladle is only conditionally decisive for the casting quality. The decisive factor is the quality of the melt in the holding or dosing furnace and, finally, in the mold itself or the shot sleeve of the pressure die casting machine. If the quality of the melt in the shot sleeve is to be maintained or increased, closer consideration must
Fig. 18
Fig. 19
time Controlling and Monitoring the Dosing Operation. A digital control unit provides for real-time pressure control in the dosing furnace, and process information is monitored on a computer display with a large-surface keyboard for the operator. Precision hardware ensures adherence to dosage accuracy. The dosing program for the cold chamber pressure die casting machine optimizes the dosing to the shortest
Fig. 17
Dosing furnace
Level detection in dosing furnace. (a) Contact electrode. (b) Pneumatic scanner
Dosing sequence. (a) Melt hold, ready. (b) Melt flow (casting). (c) Melt again held, ready
184 / Molten Metal Processing and Handling
Fig. 20
Detection of melt level by level scanning
Fig. 21
Fig. 22
Recharging a dosing furnace
Dosing furnace with porous plugs
be given to the recharging operations and melt holding in the dosing furnace. Each filling or refilling operation is associated with a reduction in the quality of the melt. A correctly implemented filling of the dosing furnace can thus also substantially restrict the loss in quality. Figure 21 shows a correctly implemented filling operation. The filling funnel is kept full, comparable with the filling of a die in gravity die casting. This results in a lower formation of oxides, which remain on the surface, do not enter the dosing furnace, and can be skimmed off at the end of the filling operation. Holding of an aluminum melt in the dosing furnace allows for settling of inclusions to a major extent. Measurements confirm that the density index or other quality parameters usually improve during quiescent holding periods. The quality of the aluminum melt thus continuously improves with the holding time, and the melt quality is not negatively influenced by the dosing process itself with dry compressed
air (protection by the oxide skin). One option to provide further improvement in melt quality within the dosing furnace is through the installation of porous plugs (Fig. 22). Introduction of inert gas, preferably argon, at the lowest point in the furnace enables reliable maintenance of a hydrogen content in the furnace to acceptable low levels. ACKNOWLEDGMENT The editors extend thanks to the following for contributions to this article: Edward S. Szekeres, Casting Consultants
Incorporated, and Robert D. Pehlke, University of Michigan, on tundishes in continuous casting of steel Martin Reeves, Striko-Dynarad, on dosing furnaces Lee Gouwens, CMI Novacast, on electromagnetic mold-filling pumps Dave Neff and Paul Campbell, Metaullics Systems Division, Pyrotek, Inc.
REFERENCES 1. I.V. Samarasekera and J.K. Brimacombe, The Continuous Casting Mould, Continuous Casting, Vol 2, Heat Flow, Solidification and Crack Formation, ISS of AIME, 1984, p 33–44 2. Y. Sahai and Toshihiko, Melt Flow Characterization in Continuous Casting Tundish, ISIJ Int., Vol 36 (No. 6), 1996 3. M.M. Collur, D.B. Love, and B.V. Patil, “Use of Flow Modifiers to Improve Performance of a Tundish,” ISS Steelmaking Conference, April 1997 4. M.L. Lowry and Y. Sahai, Investigation of Steel Flow in a Continuous Casting Tundish with Multiple Baffles Using Mathematical Models and Tracer Studies, 1989 Steelmaking Conference Proceedings, p 71–79 5. A.M. Peel, Electromagnetic Pumps and Stirring Solutions for the Aluminum Industry, Alum. Times, Vol 7, 2005, p 32–33
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 185-193 DOI: 10.1361/asmhba0005353
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Degassing David V. Neff, Pyrotek, Inc.
GAS POROSITY is a major factor in the quality and reliability of castings, and there are various sources of gas porosity in castings. One source is the reaction between the casting mold and the molten metal. For example, the moisture in sand molds can react with the elements in the cast iron (carbon, silicon, aluminum, or iron itself) to form dissolved gases such as hydrogen or carbon monoxides (see the article “Gases in Metals” in this Volume). In addition to reactions with the mold wall, porosity also occurs from entrapped air during filling, centerline shrinkage that occurs during the final solidification, and blow holes from unvented cores (Ref 1). The other major cause of gas porosity in castings is the evolution of dissolved gases from melting, and dross or slag containing gas porosity. For example, liquid cast iron may have dissolved hydrogen and nitrogen. The solubility of these gases in the solid may be less than that in the liquid, and the gases may therefore be evolved during solidification. Whether or not this will occur depends on the amount of hydrogen and nitrogen present, the alloy being cast, chemical kinetics, and the surface tension of the alloy. Similarly, hydrogen can dissolve in liquid aluminum, but its solubility is much lower in solid aluminum and can cause porosity. For copper, the evolution of hydrogen, water vapor, or carbon monoxide could cause gas porosity. Gases can be a major problem in the casting of iron, aluminum, magnesium, and copper. Dissolved hydrogen is a particular concern because of its solubility in these metals (see the article “Gases in Metals”). The most common method for removing hydrogen from aluminum, copper, and magnesium is inert gas flushing, as described in this article. Inert gas flushing can also be used for removing hydrogen and nitrogen from cast iron (see the article “Gases in Metals”). Degassing of steels is described in the article “Steel Melt Processing” in this Volume.
Aluminum Foundry Degassing Degassing of aluminum alloy melts is a critical melt treatment technology necessary to provide readiness of the melt for casting into many
applications (Ref 1–8). Hydrogen is the only gas that is soluble in aluminum, as monatomic hydrogen. Figure 1 exhibits the hydrogen solubility of typical aluminum-silicon foundry alloys. Clearly, a substantial amount of hydrogen can be present in the liquid state, but as the melt solidifies, there is a substantial reduction in the amount of hydrogen that can exist in the solid state. How this change takes place will determine whether porosity will form. If not removed or reduced to an acceptable level determined by prior process and quality studies, the gas dissolved in molten aluminum, if not treated, can result in detrimental properties to castings. Figure 2 is a classic example of gas porosity developed during solidification in an aluminum alloy casting. Loss in mechanical properties, such as tensile strength, elongation, fatigue strength, and fracture toughness, are common results. In addition, interconnected porosity can result in “leakers” during pressure-tightness testing and blistering upon subsequent heat treatment as the pores grow and open up; this is especially a problem on the surface when surface finish/cosmetic properties are required. However, gas can arise not only from entrapped hydrogen dissolved during the melting process and evolving as gas during the liquid-solid transition, but also gas can result from aspirated air or gas evolution from molding materials (i.e., cores, binders, etc.). Thus, it is not totally clear in Fig. 2 that the porosity is due to evolution of hydrogen gas. Observable porosity may also very well be caused by lack of feeding, that is, shrinkage, during solidification. In fact, Fig. 2(a) depicts porosity that follows very well the interdendritic structure and hence is very likely shrinkage, not gas. In Fig. 2(b), the rounded porosity would suggest gas evolution during final stages of solidification and hence could very well be caused by hydrogen, However, further microscopic investigation may be necessary, since this porosity may also be caused by entrapped air or lubricant decomposition. In actual practice, it is therefore best to degas the melt to an acceptable level before casting. A typical target value, below which hydrogen-caused porosity is not normally a problem in the final casting, is that the degassing process should achieve a
level of 0.15 mL H2/100 g Al, the usual standard of measurement in a quantified test such as Alscan (Alcan Aluminum Ltd.). If one is instead using the reduced-pressure test as a measure of minimal porosity potential, it is not possible to accurately predict lack of hydrogen-caused porosity by this means alone. However, reaching a high percentage value in the reduced-pressure test related to the specific alloy theoretical density is a good measure, for instance, 2.62 in A356, whose theoretical density is 2.68 g/cm3.
Hydrogen Sources Given the propensity of aluminum to readily absorb hydrogen, there are many ways that hydrogen can and will always be present in a melt to a certain degree. Fuel-fired furnaces (natural gas, oil) have hydrogen available from fossil fuel decomposition. Water vapor is almost always present in the atmosphere, even at high temperatures directly in contact with the melting furnace. Indeed, high humidity conditions exacerbate not only hydrogen dissolution in the melt but also can retard the ability of any melt
Fig. 1
Hydrogen solubility in aluminum casting alloys. Source: Ref 2
186 / Molten Metal Processing and Handling
Fig. 3
Fig. 2
Typical porosity configurations in aluminum-silicon casting alloys. (a) Shrinkage pore found within a casting. (b) Gas pore in an Al-8% Si alloy. (c) Microporosity (gas plus shrinkage). (d) Microporosity (gas plus shrinkage). Source: Ref 2
treatment to remove hydrogen, because a degassed melt can then readily redissolve additional hydrogen. Ingot, additives, and fluxes are other common sources of hydrogen. In the case of fluxes, fluxing compounds are usually mixtures of various salts, including halides, and become very hygroscopic when exposed to air. In addition, the tools used to add and work in fluxes—rakes, spoons, and so on—will often retain flux and oxide debris after use, which will then adsorb moisture from the air as they cool down. Reuse of these tools, without intermediate cleaning, not only poses a potential safety hazard but also can reintroduce hydrogen. Scrap castings, machine turnings and borings, die cast trim press scrap, gates and risers, and sand and other nonmetallic molding material debris are additional sources of hydrogen.
Aluminum Degassing There are many means available to minimize hydrogen pickup or to degas a melt. Noting the increased solubility with temperature, as depicted in Fig. 1, merely minimizing the melting or holding times and temperatures will lessen the hydrogen absorption. Reducing the temperature will result in natural outgassing, given time, which is not always available. Therefore, additional degassing measures are often needed.
The basic principle of practical foundry degassing is that any gaseous component has the ability of collecting hydrogen gas, since all gases are insoluble in aluminum, with the exception of hydrogen, which exists as a single atom. Thus, a purge or process gas can function as in Fig. 3, depicting gas fluxing, which not only collects the hydrogen but also removes particulate matter, that is, inclusions, by flotation. The first application of this principle was mainly through the decomposition of a solid product that reacted in molten aluminum. Historically, the degassing material commonly employed was a hexachloroethane tablet (C2Cl6). When submerged and plunged into the melt, the hexachloroethane decomposed into its respective components. The chlorine component immediately forms a metastable aluminum chloride gaseous compound, which is insoluble in the melt and served as the purge gas. With any purge gas, the monatomic hydrogen diffuses into the purge gas bubble, combines with another hydrogen atom to form molecular hydrogen, and the bubble rises to the surface, where both gaseous species are released. The fluxing components in the hexachloroethane tablet, which are often halide salts, served not only to bind the tablet, but these salt components do aid in wetting of the gas bubble in the aluminum, reducing the interfacial energy
Schematic of gas fluxing in molten aluminum
and allowing easier transmittal of the hydrogen atoms into the collector aluminum chloride gas bubble. This use of hexachloroethane tablets has given way to other degas methodologies, in most instances because of environmental concerns and more efficient technologies. However, some foundries still employ this technique, especially on small crucible furnaces. The second most popular degassing technique in historical use—and still used today (2008)—is a simple flux wand or static lance. This can be a steel pipe or, more commonly, a graphite lance. A purge gas such as nitrogen or argon is injected under low pressure (less than 207 kPa, or 30 psi) into the melt. The insoluble nitrogen or argon gas bubble then rises within the melt, and the hydrogen atoms diffuse to this process gas bubble and form a hydrogen molecule. When the bubble reaches the surface of the melt, the hydrogen is released. If the relative humidity in the atmosphere is particularly high, this will be a dynamic process, because new hydrogen may become absorbed into the melt. Fundamentally, the lance degassing technique is relatively inefficient, as is seen later in this section, in relation to newer technologies. The reason for this is the relatively large gas bubble that emanates from the bottom of the lance and then agglomerates and rises quickly to the surface. The rate of rise increases with bubble size, and the bubble accelerates as it climbs toward the surface. Fig. 4(a) depicts a water model showing that a static lance only influences a small volume of fluid surrounding the lance, unless it is manually moved about. While this is possible in relatively small-diameter melting and holding vessels, it is impractical in larger ones. In Fig. 4(b), the use of either a disperser on a static lance or a rotor on a rotating shaft achieves a much finer bubble size, more dispersion, and better mixing of the gas bubbles and the liquid throughout a greater vessel volume. Visually, when performing this evaluation, one would see that the rate of rise of these smaller bubbles is measurably slower than with the larger gas bubbles in the static lance. This greater residence time increases the probability of reaction between the purge gas bubble and the hydrogen, thus increasing the efficiency. The efficiency of a degassing process is directly
Degassing / 187
Fig. 4
Water model with (a) static lance and (b) disperser or rotating shaft/rotor
Fig. 5
Fig. 6
to manipulate the gas injection into the system, and the shaft/rotor assembly itself. Equipment can be as simple as a well-mounted drive unit or more complex, with a fixed-mast drive-up ladle treatment station or a mobile unit that can be moved from furnace to furnace. Many manufacturers offer these various equipment designs, together with varying degrees of automated controls, and options to inject flux (see the article “Aluminum Fluxes and Fluxing Practice” in this Volume). For the operating foundry person, several important variables must be considered in developing the degassing process in the foundry to meet specific casting requirements. The parameters that must be integrated include:
Relative degassing efficiencies. Source: Ref 4
influenced by the bubble size, as shown in Fig. 5. A classic depiction of the fundamental difference in degassing efficiency between a static lance, a rotating injector, and a porous plug (i.e., a fine pore disperser device) is shown in Fig. 6 (Ref 4). It is clear that high hydrogenremoval efficiencies are possible with a disperser, but a rotating injector with an impeller achieves a much faster rate of reduction. Rotary Impeller Degassing. This technique has been borrowed from the chemical process industry and greatly improves mixing capability between the aluminum and the process gas. This technique was introduced to the aluminum
Degassing efficiency as a function of purge gas bubble size. Source: Ref 5
foundry industry in the mid-1980s and has become the preferred technology of choice for use in degassing wells in melting furnaces, transfer ladles, and in continuous-flow launder systems. In this technique, the purge gas is injected through a bored shaft, usually made of graphite. The incipient gas volume is then fractured into many small bubbles, that is, bubble shear, by the rotor attached to the bottom of the shaft, as the gas leaves the shaft. Many different rotor designs, sizes, and configurations are now employed throughout the industry. The basic components of a rotor degassing system include a mechanical drive unit to rotate the shaft/rotor assembly, a gas source, controls
Initial hydrogen level versus desired final
hydrogen level as determined by evaluation. This is a dynamic, ever-changing process where the initial hydrogen level is influenced by charge materials, melting practice, and relative humidity. Vessel (well or transfer ladle) size and volume. In a continuous-flow launder process, the metal volume residence time must be considered. Available time in a transfer ladle for melt treatment. The metal must then be delivered to the casting station(s) without excessive temperature loss, considering the relationship between rotor type and revolutions per
188 / Molten Metal Processing and Handling minute (rpm), gas volume being injected, avoidance of surface effects (vortexing, splashing), and time necessary and available to achieve the desired degassing results. The interplay among these variables must be determined on a case-by-case basis by the individual foundry and often on the specific casting(s) to achieve the optimum combination of process and equipment parameters. In general, the optimum result desired is to achieve the necessary degassing specification in as short a time as possible, at the lowest cost, and without excessive turbulence. As a general rule, most well-developed foundry degassing practices should achieve the desired degassing level in 5 min when treating a vessel with 455 kg (1000 lb) volume and when optimizing the foregoing parameters of rotor type, rpm, and purge gas flow rate. Selecting a Process Gas. Since all gases except hydrogen are insoluble in molten aluminum, there are theoretically many choices of which gas to use. Nitrogen is the most commonly employed and is the least expensive. Nitrogen creates a wet dross; that is, the dross that is formed is high in metallic aluminum content. Furthermore, since the density of nitrogen is comparable to air overall, there is a tendency to pick up moisture in nitrogen lines as fittings deteriorate and humid air is aspirated into the delivery lines. For this reason, all nitrogen sources should include an in-line dryer. Nitrogen purity should be at least 99.99% for any critical casting applications. Argon, while significantly more expensive than nitrogen, produces a less metal-rich dross. Argon is more inert, and, being heavier, provides a protective cover over the melt during degassing, precluding not only hydrogen absorption but further oxidation as well. It is also easier to keep the argon supply drier and cleaner as hoses and fittings deteriorate. While argon is the preferred process gas for in-line degassing treatment in the production of wrought alloys, it is not extensively used in foundry casting because of the higher cost. Active Halogens. Historically, a small percentage of chlorine (5 to 10%) was used with an inert gas when degassing with a static lance. The chlorine enhanced the reaction between the melt and the hydrogen atoms because it lowered the interfacial tension at the purge gas bubble surface (through the formation of aluminum chloride, a metastable gaseous phase) and facilitated the diffusion of the hydrogen atom into the purge gas bubble. Several commercial mixtures of halogen gases—Freon (CCl2F2) (Dupont) being the most popular—were employed in the past. With the growing concern and difficulties with the environmental impact of chlorine use, sulfur hexafluoride (SF6) gained prominence 20 years ago as a replacement for chlorine. However, subsequent experience showed that SF6 also posed environmental problems with ozone depletion, and the gas is
very corrosive, even in small amounts, to foundry metal structures. Fortunately, with the advent of rotor degassing and the resultant better kinetics of reaction and mixing, it has been found that chlorine does not enhance the degassing rate or efficiency when the purge/process gas bubbles are sufficiently small, as they are with rotor impeller degassing. Rather, when chlorine is used, there is better wetting of the nonmetallic inclusions during the degassing process, so that metal cleanliness is improved, and the burden on any subsequent particulate removal process is lessened, which in turn creates a better-quality cast product. Other Considerations for the Degassing Process. There are other issues that must be taken into account during the degassing process. Temperature. It is well understood that hydrogen absorption in the melt increases with increasing temperature. Conversely, hydrogen solubility and partial pressure of hydrogen gas in the melt decrease rapidly as the melt temperature is decreased. It follows that minimizing the temperature will result in lower hydrogen value. Given sufficient time, allowing the temperature to drop will result in natural outgassing. However, the injection of a purge gas at normal standard temperature and pressure creates an endothermic reaction, and heat will therefore be lost at an accelerated rate. A general ruleof-thumb is that degassing a metal volume of 455 kg (1000 lb) in an unheated refractory vessel for 5 to 10 min will result in a temperature drop of nearly 28 C (50 F). Whatever the actual magnitude of the temperature drop due to gas injection/treatment time, this reduction must be accommodated in relation to the necessary temperature that must be maintained as the treated melt is transferred into the mold or into the casting station(s). A degassing process should nevertheless be carried out at the lowest practical temperature possible. Grain Refining. Most aluminum melts in sand, permanent mold, and low-pressure foundry production require grain refinement to achieve the best mechanical properties in the finished casting. In most instances, grain refiner is added during the degassing process, because the increased mixing and agitation facilitates the dissolution and reaction of the grain refiner itself and provides more uniform dispersion throughout the melt or treated volume. Modification. Historically, the element sodium was used to provide modification of the eutectic silicon component in aluminum casting alloys—a necessary practice, together with grain refining, to produce desired mechanical properties. In current practice, sodium has given way to a strontium addition for many reasons, including safety in handling and in the environment, better consistency, and reduced fade, which is the loss of modification with time. Because both sodium and strontium are prone to oxidation (although strontium vapor pressure is substantially lower than sodium),
and the degassing mixing creates more convection during treatment, degassing should take place before modification, so as not to lose the benefits of the dissolved modifier due to oxidation. If chlorine is used, even in small quantities, in the degassing process, both sodium and strontium will react chemically to form solid chloride reaction products. Dross Formation and Skimming. The action of the purge gas injection process by whatever means—lance, rotor, or porous plug/disperser—will result in flotation of debris and some degree of particulate inclusion content within the melt. Naturally, the degassing process creates dross. This must be thoroughly skimmed off the surface to the greatest practical extent possible to avoid transfer into the casting mold or casting station. Determining Hydrogen Content. To control hydrogen content, one must be able to measure the hydrogen content in the melt and the subsequent result of any hydrogen-removal treatment. Laboratory evaluations are possible using vacuum fusion (Leco technique) and subvacuum fusion on previously solidified test specimens. Obviously, laboratory measurements cannot be used to make real-time decisions but rather are useful in process development. Real-time determinations are limited to the simple reduced-pressure test (RPT) and its variants, and in situ techniques, such as Alscan, continuous hydrogen analysis by pressure evaluation, and others. The RPTs are the simplest, but these tests are a measure of molten metal cleanliness and not absolute hydrogen content per se. The presence of inclusions in the melt during an RPT evaluation will create exaggerated porosity nucleated during solidification and a resultant lower density value. The Alscan test and others of this type do generate a real-time result through the principles of recirculating gas of a known standard hydrogen content and the thermal conductivity of dilute solutions.
Degassing of Magnesium Magnesium behaves similarly to aluminum in its ability to dissolve hydrogen easily in the molten state (Fig. 7). It too will suffer from porosity when cast if the hydrogen is not removed or naturally expelled during solidification. Aluminum and zinc, which are typical alloying elements in magnesium alloys, lower this solubility somewhat. Figure 8 indicates that much more gas is evolved from magnesium alloys containing zinc and aluminum than from pure magnesium upon solidification. Gas porosity occurs in solidified castings when the residual hydrogen exceeds 15 mL/100 g. Zirconium is also a major constituent of certain magnesium alloys and has a marked effect on hydrogen solubility. Hydrogen solubility decreases as the zirconium content of the alloy increases (Ref 9). Because of this low solubility even in the liquid state, hydrogen porosity is minimal in castings of magnesium-zirconium
Degassing / 189 alloys. However, excess hydrogen in the melt readily precipitates as zirconium hydride, thus depleting the alloy of zirconium in solid solution. It is therefore advisable to minimize hydrogen absorption into the melt and/or to employ degassing methods even with this alloy. Sources of gas in magnesium alloy melts are the same as those for aluminum—atmospheric conditions, fluxes, tools, and so on. Flux covers or antiburning agents do not necessarily preclude hydrogen absorption during melting. Like aluminum alloys, molten magnesium metal and alloys can be degassed with hexachloroethane, chlorine, or nitrogen. Hexachloroethane has the added advantage of supplying carbon through the decomposition of the hexachloroethane, with the carbon serving as a grain refiner (see the article “Magnesium and Magnesium Alloy Castings” in this Volume). Chlorine
or chlorine/nitrogen gas purging is optimally performed at temperatures of 725 to 750 C (1335 to 1380 F). In this range, liquid MgCl2 will also be formed, which will retard burning on the surface. Liquid MgCl2 wets the oxides and the magnesium melt surface, facilitating hydrogen escape. Pure chlorine provides for more liquid MgCl2 formation than either gas mixtures or hexachloroethane and therefore creates better degassing. Fluxes higher in MgCl2 that are used as covers during degassing also aid in hydrogen desorption at the flux/metal interface and in subsequent expulsion to the atmosphere. The degassing of magnesium alloys in crucible or pot-type furnaces should be carried out with clean, dry lances, preferably graphite, with thorough mixing for 5 to 10 min. Pressures and volumes should not exceed those required to achieve a mild, gentle roll or small bubble (25 mm, or 1 in., in diameter) on the melt surface. Rare-earth-containing magnesium alloys should only be degassed when required with nitrogen or argon, because chlorination will readily remove these expensive alloying elements.
Degassing of Copper Alloys In the melting and casting of many copper alloys, hydrogen gas absorption can occur because of the generous solubility of hydrogen in the liquid state of these alloys. The solid solubility of hydrogen is much lower; therefore, the gas must be rejected appropriately before or during the casting and solidification process to avoid the formation of gas porosity and related defects (excessive shrink, pinholing, blow holes, and blistering). These defects are almost always detrimental to mechanical and physical properties, performance, and appearance. Different copper
alloys and alloy systems have varying tendencies toward gas absorption and subsequent problems. Other gases, particularly oxygen, can cause similar problems (see the article “Copper and Copper Alloy Castings” in this Volume).
Gases in Copper The type and amount of gas absorbed by a copper alloy melt and retained in a casting depend on a number of conditions, such as melt temperature, raw materials, atmosphere, pouring conditions, and mold materials. Table 1 lists the gases that can be found in a number of copper alloy systems. Hydrogen is the most obvious gas to be considered in copper alloys. Figure 9 (Ref 11) shows the solubility of hydrogen in molten copper, tin, and copper-tin alloys. As in aluminum and magnesium alloys, solubility is reduced in the solid state; therefore, the hydrogen must be removed prior to casting or rejected in a controlled manner during solidification. Alloying elements have varying effects on hydrogen solubility (Fig. 10). There are many potential sources of hydrogen when melting copper alloys. For the most part, these are identical to the sources of hydrogen in aluminum melting. Oxygen also presents a potential problem in most copper alloys, but oxygen alone does not cause problems in copper alloys. In the absence of hydrogen, oxygen alone may not cause problems, because it has limited solubility in the melt. However, it forms a completely miscible liquid phase with the copper in the form of cuprous oxide (Fig. 11). During solidification, the combination of cuprous oxide and hydrogen can give rise to casting porosity resulting from the steam reaction (discussed later). Sulfur gases have significance in primary copper through the smelting of sulfide ores
Fig. 7
Solubility of hydrogen in molten magnesium at 101 kPa (760 torr). Source: Ref 9
Fig. 8
Effect of alloying elements on hydrogen solubility in liquid magnesium. AC, as-cast; HT, heat treated. Source: Ref 9
190 / Molten Metal Processing and Handling Table 1 Summary of gases found in copper alloys Alloy family
Pure copper Copper-tin-leadzinc alloys Aluminum bronzes Silicon brasses and bronzes Copper-nickels
Gases present
Water vapor, hydrogen Water vapor, hydrogen, CO Water vapor, hydrogen, CO Water vapor, hydrogen Water vapor, hydrogen, CO
Remarks
Approximate hydrogen/water vapor ratio of 1. Higher purity increases the amount of water vapor and lowers hydrogen. Lead does not affect the gases present. Higher tin lowers total gas content. Increased zinc increases the amount of hydrogen, with a loss in water vapor. The presence of 5 wt% Ni in alloy C95800 causes CO to occur rather than water vapor. Lower aluminum leads to higher total gas contents. Approximate hydrogen/water vapor ratio of 0.5. Increased zinc decreases hydrogen and increases water vapor. All three gases are present up to 4 wt% Ni, after which only CO and hydrogen are present. Hydrogen increases with increasing nickel up to 10 wt% Ni but is decreased at 30 wt% Ni.
Source: Ref 10
Fig. 11
Fig. 9
Solubility of hydrogen at 1 atm (760 mm Hg) in pure copper, pure tin, and copper-tin alloys. Source: Ref 11
and in the remelting of mill product scrap containing sulfur-bearing lubricants. Sulfur dioxide is the most probable gaseous product. Foundry alloys and foundry processing usually do not experience sulfur-related problems unless high-sulfur fossil fuels are used for melting. Carbon can be a problem, especially with the nickel-bearing alloys. The nickel alloys have extensive solubility for both carbon and hydrogen. Carbon can be deliberately added to the melt, along with an oxygen-bearing material such as nickel oxide. The two components will react to produce a carbon boil, that is, the formation of CO bubbles, which collect hydrogen as they rise through the melt (Ref 13). If not fully removed, however, the residual CO can create gas porosity during solidification. Nitrogen does not appear to be detrimental or to have much solubility in most copper alloys. However, there is some evidence to suggest that nitrogen porosity can be a problem in cast copper-nickel alloys (Ref 14).
Fig. 10
Effect of alloying elements on the solubility of hydrogen in copper. Source: Ref 12
Water vapor can exist as a discrete gaseous entity in copper alloys (Ref 10, 15, 16). In many copper alloys, the most important gas causing porosity is water vapor (Ref 11). Water itself is not dissolved, but dissolved atomic oxygen and hydrogen in the melt combine to form water until the oxygen becomes depleted; hydrogen is produced as a separate species. Water vapor is evolved from solidifying copper alloys, which always have some residual dissolved oxygen. Thus, gas porosity can occur even with low hydrogen content when oxygen is present. Gas porosity occurs when the gas pressure exceeds 1 atm (760 mm Hg). To prevent such porosity, the solid solubility limits must not be exceeded, and the sum of individual gas partial pressures must be less than 1 atm (760 mm Hg) through the solidification process (Ref 11). Testing for Gases. There are essentially three ways to determine the presence of gas in a copper alloy melt. The easiest and simplest is a chill test on a fracture specimen. In this
Copper-oxygen phase diagram. Source: Ref 12
method, a standard test bar is poured and allowed to solidify. The appearance of the fractured test bar is then related to metal quality standards (Ref 17). The second method is an RPT similar to the Straube-Pfeiffer test used for aluminum alloys. In recent years, considerable development and definition have been given to this test as adapted for copper alloys, because the different classes of copper alloys have different solidification characteristics that affect the response to the test (Ref 10, 18–20). The objective of the test is to use a reduced pressure only slightly less than the dissolved gas pressure so that the surface of the sample mushrooms but does not fracture (Fig. 12). The test apparatus is shown schematically in Fig. 13. The key to establishing controllable results, that is, unfractured but mushroomed test sample surfaces, is to provide wave-front freezing. This is accomplished by proper selection of vessel materials to achieve the correct solidification characteristics. Because the freezing ranges of copper alloys vary from short to wide, different materials must be used. The various material combinations that have proved successful are described in Ref 10.
Degassing Methods There are several kinds of treatments which may be employed for the improvement of melt quality in copper alloys, such as oxidationdeoxidation practice, zinc flaring, inert gas fluxing, and the use of solid degassing fluxes. Metallurgical principles for common copper alloy classes are discussed in detail in Ref. 21. Oxidation-Deoxidation Practice. The steam reaction previously mentioned is a result of both hydrogen and cuprous oxide being present in the melt. These constituents react to form steam, resulting in blow holes during solidification as the copper cools: 2H þ Cu2 O ! H2 OðsteamÞ þ 2Cu The proper deoxidation of a copper alloy melt will generally prevent this steam reaction,
Degassing / 191
Fig. 12
Effect of pressure on the appearance of copper alloy reduced-pressure test samples containing the same amount of gas. (a) Pressure of 7 kPa (55 torr) results in surface shrinkage. (b) At 6.5 kPa (50 torr), a single bubble forms. (c) Boiling and porosity occur at 6 kPa (45 torr).
Fig. 13
Fig. 14
Schematic of the reduced-pressure test apparatus used to assess amounts of dissolved gas in copper alloys
Hydrogen/oxygen equilibrium in molten copper. Source: Ref 10
Fig. 15
Influence of zinc content on boiling point or vapor pressure in copper alloys. Source: Ref 12
although excessive hydrogen alone can still cause gas porosity if it is not expelled from the casting before the skin is completely solidified. Fortunately, there is a mutual relationship between hydrogen and oxygen solubility in molten copper (Fig. 14). Steam is formed above the line denoting equilibrium concentration, but not below. Consequently, as the oxygen content is raised, the capacity for hydrogen absorption decreases. Therefore, it is useful to provide excess oxygen during melting to preclude hydrogen entry and then to remove the oxygen by a deoxidation process to prevent further steam reaction during solidification. This relationship gave rise to the Pell-Walpole oxidation-deoxidation practice for limiting hydrogen, especially when fossil-fuel-fired furnaces were used for melting. The melt was deliberately oxidized using oxygen-bearing granular fluxes or briquetted tablets to preclude hydrogen absorption during melting. The melt was subsequently deoxidized to eliminate any steam reaction with additional hydrogen absorbed during pouring and casting. Zinc Flaring. The term zinc flare applies to those alloys containing at least 20% Zn. At this level or above, the boiling or vaporization point of copper-zinc alloys is close to the usual pouring temperature, as shown in Fig. 15 (Ref 12). Zinc has a very high vapor pressure, which precludes hydrogen entry into the melt. Zinc may also act as a vapor purge, removing hydrogen already in solution. Furthermore, the oxide of zinc is less dense than that of copper, thus providing a more tenacious, cohesive, and protective oxide skin on the melt surface and increasing resistance to hydrogen diffusion. In zinc flaring, the melt temperature is deliberately raised to permit greater zinc vapor formation. Because some zinc loss may occur, zinc may have to be replenished to maintain the correct composition. Inert Gas Fluxing. With gas fluxing, an inert collector or sparger gas such as argon or nitrogen is injected into the melt with a graphite fluxing tube. The bubbling action collects the hydrogen gas that diffuses to the bubble surface, and the hydrogen is removed as the inert gas bubbles rise to the melt surface. The reaction efficiency depends on the gas volume, the depth to which the fluxing tube or lance is plunged, and the size of the collector gas bubble generated. Again, finer bubble sizes have higher surface-area-to-volume ratios and therefore provide better reaction efficiencies. Figure 16 depicts the amount of purge gas necessary to degas a 450 kg (1000 lb) copper melt. Figure 17 shows the response of an aluminum bronze alloy to nitrogen gas purging. The curve for oxygen purging is also shown. Solid Degassing Fluxes. Other materials can be used to provide inert gas purging. A tabletted granular flux such as calcium carbonate (CaCO3), which liberates CO2 as the collector gas upon heating, has been successful in degassing a wide variety of copper alloys. This type of degassing is simpler than nitrogen but must
192 / Molten Metal Processing and Handling
Fig. 16
Amount of purge gas required to degas a 450 kg (1000 lb) copper melt. Source: Ref 20
Fig. 17
Fig. 18
Comparison of the effectiveness of solid degassing flux versus nitrogen purging. Source: Ref 21
be kept dry and plunged as deep as possible into the melt with a clean, dry plunging rod, preferably made of graphite. This type of degasser may also have the advantage of forming the collector gas by chemical reaction in situ. This results in an inherently smaller initial bubble size than injected gas purging for better reaction efficiency. Such an advantage can be inferred from Fig. 18, although after 5 min of degassing, the end results are the same. Other solid degassers in the form of refractory metals or intermetallic compounds, such as calcium-manganese-silicon, nickel-titanium, titanium, and lithium, are effective in eliminating porosity due to nitrogen or hydrogen by their ability to form stable nitrides and hydrides. Again, maintaining dry ingredients, deep plunging, and stirring or mixing will enhance their effectiveness. For best results and optimum reaction efficiency, such degassers should be wrapped or encapsulated in copper materials to provide controlled melting when plunged. Alternatively, copper master alloys containing these elements are available.
Effect of nitrogen flow rate during purging on the residual gas pressure remaining in an aluminum bronze melt. The cure for oxygen purging is also shown. Source: Ref 19, 22
For both inert gas fluxing and solid degassing additions, the sparging gas reaction should not be so violent as to splash metal and create an opportunity for gas reabsorption. Furthermore, a melt can be overdegassed; an optimum amount of residual gas remaining in the melt helps to counter localized shrinkage in longfreezing-range alloys such as leaded tin bronzes. Vacuum degassing is not generally applied to copper alloys, although it can be very effective. However, the cost of the equipment necessary is relatively high, and there may be a substantial loss of more volatile elements having a high vapor pressure, such as zinc and lead. Furthermore, significant superheat (up to 150 C, or 300 F) may be required to accommodate the temperature drop during the degassing treatment, further aggravating the vapor losses. Auxiliary Degassing Methods. There are other ways to degas a copper alloy melt than using a specific treatment. The technique of lowering the melt superheat temperature, if possible, and holding the melt (in a dry, minimal gas environment) provides outgassing simply by lowering the equilibrium liquid-state solubility of the gas. During casting, the use of chills to provide directional solidification, particularly for the long-freezing-range alloys such as the red brasses and tin bronzes, results in less tendency for gas porosity. In the mold-metal reaction, previously degassed and deoxidized metal containing excess phosphorus or lithium deoxidants can react with green sand containing moisture. Hydrogen is liberated and absorbed as the metal comes into contact with the sand. This can be
minimized by correct deoxidant additions. In addition, finer facing sands or mold coatings (inert or reactive, such as sodium silicate or magnesia) can be used; these confine any reaction to the mold/metal interface area and retard hydrogen penetration into the solidifying skin of the cast metal.
REFERENCES 1. R. Monroe, Porosity in Castings, Trans. AFS, Vol 113, 2005, p 519–546 2. J.E. Gruzleski and B.M. Closset, The Treatment of Liquid Aluminum-Silicon Alloys, AFS, 1990 3. D. Apelian, Clean Metal Processing of Aluminum, Proceedings, Materials Solutions Conference ’98 on Aluminum Casting Technology, ASM International, 1998 4. G. Sigworth and T. Engh, Chemical and Kinetic Factors Related to Hydrogen Removal from Aluminum, Metall. Trans., Vol 138, 1982, p 447 5. G. Sigworth, A Scientific Basis for Degassing Aluminum, Trans. AFS, Vol 95, 1987, p 73 6. D.V. Neff, Understanding Aluminum Degassing, Mod. Cast., May 2002 7. D. Apelian and M.M. Makhlouf, Ed., High Integrity Aluminum Die Casting, NADCA, April 2004 8. M.M. Makhlouf, L. Wang, and D. Apelian, Measurement and Removal of Hydrogen in Aluminum Alloys, AFS Special Report, 1998 9. E.F. Emley, Principles of Magnesium Technology, Pergamon Press, 1966
Degassing / 193 10. P.K. Trojan, T.R. Ostrom, and R.A. Flinn, Melt Control Variables in Copper Base Alloys, Trans. AFS, 1982, p 729; “High Conductivity Copper Alloys,” Bulletin 43, Foseco, Inc., 1972 11. Casting Copper-Base Alloys, 2nd ed., American Foundry Society, 2007, p 128–131 12. R.J. Kissling and J.F. Wallace, Gases in Copper Base Alloys, Foundry, 1962, 1963 13. M.P. Renatv, C.M. Andres, G.P. Douglas, and R.A. Flinn, Solubility Relations in Liquid Copper-Nickel-Carbon-Oxygen Alloys, Trans. AFS, 1976, p 641
14. B.N. Ames and N.A. Kahn, Gas Absorption and Degasification of Cast Monel, Trans. AFS, 1947, p 558 15. T.R. Ostrom, P.K. Trojan, and R.A. Flinn, Gas Evolution from Copper Alloys—A Summary, Trans. AFS, 1981, p 731 16. T.R. Ostrom, P.K. Trojan, and R.A. Flinn, The Effects of Tin, Aluminum, Nickel, and Iron on Dissolved Gases in Molten Copper Alloys, Trans. AFS, 1980, p 437 17. Casting Copper Base Alloys, American Foundrymen’s Society, 1984 18. T.R. Ostrom, R.A. Flinn, and P.K. Trojan, Gas Content of Copper Alloy Melts—Test
19. 20. 21. 22.
Equipment and Field Test Results, Trans. AFS, 1977, p 357 T.R. Ostrom, P.K. Trojan, and R.A. Flinn, Dissolved Gases in Commercial Copper Base Alloys, Trans. AFS, 1975, p 485 R.J. Cooksey and R.W. Ruddle, New Techniques in Degassing Copper Alloys, Trans. AFS, Vol 68, 1960 D.V. Neff, Fluxing, Degassing, Deoxidation of Copper Base Alloys, Trans. AFS, 1989, p 439 P.K. Trojan, S. Suga, and R.A. Flinn, Influence of Gas Porosity on Mechanical Properties of Aluminum Bronze, Trans. AFS, 1973, p 552
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 194-205 DOI: 10.1361/asmhba0005351
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Molten-Metal Filtration D.V. Neff, Pyrotek Inc. Robert P. Pischel, FOSECO Metallurgical Inc.
FILTRATION has become a primary means of removing solid impurities from nonferrous molten metal. Aluminum alloys, especially with their tendency toward oxidation, often contain impurities (inclusions) that are deleterious to physical, mechanical, electrical, and aesthetic properties. In the handling of molten aluminum, it is fairly common to use filters as a part of the melting unit (in-furnace filtration) and in the gating and/or riser system (in-mold filtration). Zinc and copper alloy melts also are filtered but to a much lesser extent than aluminum. In almost all instances of metal casting, the presence of dross, inclusions, or particulate impurities detracts from casting performance and can be a major contributor to scrap castings and lost productivity and profitability for the foundry. For example, the presence of dross, inclusions, or particulate impurities detracts from the natural fluidity of any melt (Fig. 1) (Ref 1). Mechanical properties of the cast product can also be degraded, particularly in elongation/ductility and fatigue strength, which are most sensitive to the presence of inclusions. Inclusions also are harder than the surrounding matrix and give rise to tool chatter, tool wear, and ultimately, tool breakage during machining operations. In addition, inclusions may result in poor continuity, such as creating leakers in
pressuretight castings by nucleating porosity formation during solidification or by causing obvious discontinuities and blemishes in cosmetic surfacefinishing operations such as powder coating, plating, painting, polishing, and anodizing. Thus, inclusion-related defects should be minimized to the greatest extent possible to prevent scrap castings, lost productivity, and lower profitability. This article describes the methods of in-furnace and in-mold filtration, with emphasis on the filtration of molten aluminum. Moltenmetal filtration is an ongoing area of interest in aluminum foundry practice (Ref 2–6). For example, the following table comprises typical quality improvements in several aluminum foundries using in-furnace filtration: Foundry
Al alloy
Result from in-furnace filtering
380 319 380 319 380 518
Hardspots eliminated Leakers reduced 75% Tool life extended 600% Scrap reduced 95% Tool breakage reduced 95% Scrap eliminated
A B C D E F
Sources of Inclusions Undissolved particulate contaminants (inclusions) include slag (ferrous), dross (nonferrous), nonmetallic impurities (oxides, sulfides, nitrides), or insoluble intermetallic precipitates (see the article “Inclusion-Forming Reactions” in this Volume). Inclusions arise from many sources. Charge materials can contain nonmetallic inclusions from unfiltered primary or secondary ingot. Scrap materials may contain molding sand, debris from back-charged gates and runners, metalworking lubricants, paint and coatings, plastics, or unsorted and unwanted heavy-metal contaminant oxides (iron, manganese, chromium, and so on). Furnace tools, such as skimmers, puddlers, and rakes, can also add solidified oxide skins back into the melt if they are not cleaned before use. Factors that influence the formation of inclusions include: Excessive
Fig. 1
Effect of metal cleanliness on fluidity of an aluminum die casting alloy. Source: Ref 1
temperatures during melting accelerate the oxidation of aluminum and alloying elements. Where fluxes are used,
incomplete skimming or separation of fluxing salts often results in contamination and salt-wetted oxide inclusions. Grain-refiner additions can be a source of inclusions. Intermetallic compound precipitation may also contribute to inclusions. Iron, chromium, and manganese form intermetallics with aluminum, and the formation of intermetallic sludge is a common problem in die casting alloys. Refractory inclusions from the erosion and wear of melting and holding vessels may add alumina, spinel, silica, and silicon carbide inclusions to nonferrous alloy melts. Aluminum Inclusions. Inclusions in molten aluminum foundry processing arise from several sources. These inclusions exist as oxides present as skins on charge materials; from oxidation of the melt itself; from erosion of refractory vessels used for melting, transfer, and holding of molten aluminum; from additives such as master alloys; and grain refiners and modifiers as well. The various types of solidphase inclusions in molten aluminum alloys include:
Oxides (Al2O3, MgO) Spinels (Mg2AlO4) Borides (TiB2, VB2, ZrB2) Carbides (Al3C4, TiC) Intermetallics (MnAl3, FeAl3) Nitrides (AlN) Exogenous refractory inclusions (oxides and/ or carbides of iron, silicon, aluminum, etc.)
The size and shape of the inclusions vary from globular (oxides and borides) to stringerlike (oxides and spinels). The latter are particularly insidious in their effect on castability and the properties of the casting. Aluminum oxide is present in virtually all aspects of aluminum alloy casting production, because aluminum has a high oxidation potential. Prime metal in either liquid or solid form, alloy additions, and so on will always have an aluminum oxide skin. Melting in nearly all processes is done in air, thus providing another source for aspirated air to enter, especially in closed or even
Molten-Metal Filtration / 195 flue-operated larger furnaces that are not maintained under positive pressure. Foundry returns— an essential part of the foundry economic charge mix—also contain not only aluminum oxide skin but also can contribute sand debris, machining lubricant residues, die casting release agents, plunger lube, and other fluids that may contain some level of solids content. In addition, alloy and metallic impurity components in excess of the solubility limit form intermetallic compounds that do not dissolve during processing and hence can become inclusions in the solidified metal. In high-pressure die casting alloys and most aluminum-silicon shape casting alloys, an intermetallic compound known as sludge can form with sufficient constituent concentration (iron, manganese, chromium) and low temperature to form such an intermetallic. Additionally, segregated heavy metals (zinc, copper, iron, manganese) can also form a segregated agglomerate in the melting furnace, which has the potential to be transferred out and into a casting. During fluxing operations in a melting furnace or transfer ladle, residual flux inclusions may not become fully skimmed or removed from the surface nor have time to either float out or sediment. If these flux inclusions do become entrapped within the casting, they create sources for galvanic corrosion to take place. Typical solid-phase inclusions that may be found in aluminum casting alloys are listed in Table 1. Oxide-based inclusions are varied and will exhibit a spectrum of densities—some with specific gravity lower than the melt (nominally 2.4 g/cm3, or 0.087 1b/in.3) and some higher than the melt. Given sufficient time, separation of the solid particulate from the liquid is possible to a degree, based on the density difference. Thus, lighter particles could float to the surface; heavier particles will sink or sediment to the bottom of a melting or casting vessel or furnace. As Fig. 2 demonstrates, however, this phenomenon is time-dependent as well as size-dependent. In casting alloys, a discrete particulate soon becomes an agglomerate. Surface tension/interfacial energies usually dictate that complexes of inclusions have a bulk density nearer that of the melt, and hence, separation purely by gravity means becomes retarded. Liquid-phase inclusions can occur in molten aluminum alloys. The most prevalent is MgCl2 arising from the chlorination of magnesium-containing melts. These inclusions are not generally removed to a great extent by conventional filtration techniques and therefore must be avoided. Preventing excessive and inefficient chlorination of the melts during degassing and thorough skimming at least help to avoid such inclusions. The nature of these and other inclusions in aluminum alloys is discussed in more detail in Ref 7 and 9 to 12. Inclusions in Magnesium Alloys. The inclusions found in magnesium melts are similar in many cases to those found in aluminum—oxides, spinels, furnace refractory particles, and so on. In addition, alloying elements (zirconium and the rare earths) lead to oxides of these elements
as well as intermetallic compound precipitates. If the fluxing process is used to clean magnesium, fluxing salt inclusions may also be present. Inclusions in Copper Alloys. In addition to the usual inclusions arising from oxides, fluxing salts, and intermetallics, copper oxide inclusions and phosphorus pentoxide (from deoxidation) may be present in copper alloys if the melt is not allowed to settle or if it is inadequately skimmed before pouring and casting. Inclusions in Zinc Alloys. In the zinc foundry and die casting alloys containing aluminum, nonmetallic oxide inclusions are
usually less important than intermetallic inclusions. Iron-zinc and iron-aluminum intermetallic compounds may also be present, especially above 480 C (900 F) in melting in iron-base vessels. These particles may be small crystals less than 5 mm (200 min.) in size or may agglomerate into millimeter-sized aggregates. Other heavy metals present, such as lead, cadmium, nickel, manganese, and chromium, can also lead to intermetallic compound precipitation (for example, Mn2Al5, Ni2Al3, and CrAl4). These compounds (often called aluminides) can form when the concentration of impurity exceeds
Table 1 Common types of inclusions in aluminum foundry melts Inclusion type
Chemical formula
Density Form
g/cm3
Size range lb/in.3
mm
min.
8–1200 400–2 105 4–200 400–2 105 4–200 400–2 105 20–200
Oxides Alumina
Al2O3
Particles, skins
3.97
0.143
Magnesia
MgO
Particles, skins
3.58
0.129
Spinel
MgAl2O4
Particles, skins
3.60
0.130
Silica
SiO2
Particles
2.66
0.096
0.2–30 (particles) 10–5000 (skins) 0.1–5 (particles) 10–5000 (skins) 0.1–5 (particles) 10–5000 (skins) 0.5–5
Salts Chlorides Fluorides
Varies Varies
Particles Particles
1.98–2.16 1.98–2.16
0.072–0.078 0.072–0.078
0.1–5 0.1–5
4–200 4–200
Carbides Aluminum carbide Silicon carbide
Al4C3 SiC
Particles Particles
2.36 3.22
0.085 0.116
0.5–25 0.5–25
20–1000 20–1000
Nitrides Aluminum nitride
AlN
Particles, skins
3.26
0.117
10–50
400–2000
Borides Titanium boride
TiB2
4.50
0.163
1–30
40–1200
Aluminum boride
AlB2
Agglomerated particles Particles
3.19
0.115
Sludge Various
Al(FeMnCr)Si
Particles
>4.0
>0.145
Source: Ref 7, 8
Fig. 2
Sedimentation/flotation of inclusions in aluminum melts
0.1–3
4–120 ...
...
196 / Molten Metal Processing and Handling the liquid solubility level of the element, which is often quite low (<0.02% for manganese, silicon, chromium, and nickel in die casting alloys). The intermetallic particles can be lighter than the base alloy and therefore rise to the top dross. Fluxing salt inclusions and exogenous refractory particles can also be present, as with the other nonferrous alloy systems. Oxide inclusions in zinc are somewhat limited because of the high positive vapor pressure of zinc. However, alloys containing aluminum can be subject to greater oxidation of the aluminum constituent. Oxide inclusions can also result during melting or casting when excessive turbulence is encountered; thus, mechanical stirring or molten metal pumping should be controlled. Other nonmetallic inclusions arise from refractory vessel erosion and include silica, alumina, and silicon carbide. There is an appreciable density difference between these oxides and the base melt, unlike in aluminum; for this reason, gravity separation and thorough removal are substantially easier.
Removal of Inclusions There are three basic methods of mechanically removing or separating inclusions from molten metal: Sedimentation Flotation Positive filtration
Fig. 3
Two modes of positive filtration. (a) Depth filtration. (b) Cake filtration
The first two processes—sedimentation and flotation—are based on density differences between the liquid and solid phases. Actual separation capability is based not only on the density differences but also particle size. For example, rotary degassing is a method that removes inclusions by flotation. It is usually effective in removing inclusions larger than 40 mm (1600 min.) in size. Figure 2 is an application of Stokes’s law governing phase separation as applied to inclusions in molten aluminum. The greater the density difference from molten aluminum and the larger the particle size, the faster the rate of rise (flotation to the melt surface) or fall (sedimentation to the bottom of the furnace or holding vessel). Thus, in many metal casting operations, a holding or quiescent period within a furnace will result in much of the inclusion content being able to be removed by skimming the surface, or avoidance of the sedimented inclusions by careful pouring so as not to entrain the inclusions within the pouring stream or into the transfer ladle or casting mold. However, two factors combine to render such gravity-related processes less than efficient—at least for light nonferrous metals such as aluminum. First, there is a high surface tension/interfacial surface energy between most oxide inclusions in aluminum and the melt, which makes actual separation difficult. Secondly, many oxide inclusions, even small ones, can be detrimental to casting properties if concentrated into the casting during solidification. It can be seen that small-sized inclusions have very slow rates of separation. Additionally, many inclusions and/or inclusion complexes will have bulk
Fig. 4
densities that are very similar to the density of the melt itself and thus will also not separate easily or quickly. Consequently, positive filtration will need to be employed to remove these. Positive filtration involves the passing of molten metal through a porous device (a filter) in which the inclusions contained in the flowing metal are trapped or captured by one or more filtration mechanisms. With any filter, there are two mechanisms of positive filtration (Fig. 3): Cake filtration Depth-mode (or deep-bed) filtration
Cake filtration occurs when there is a buildup of particulate (filtrate) on the inlet surface of a filter. Particles that are larger sized than the inlet openings on a filter will naturally accumulate at the surface to form a cake. In most instances, however, inclusion particles are considerably smaller than the filter pore openings, but the filter body itself connecting the pores serves to slow the molten-metal flow through the filter, such that the particles begin to adhere to the surface. Fortunately, even though these inclusions begin to build up, the accumulation remains quite lacy and permeable to continued metal flow, up to an extent. Depth-mode filtration occurs within the body of a filter. In this mode, inclusion particles flowing with the liquid are subject to changes in velocity and direction created by the internal complexity or tortuosity of the flow path within the filter. The inclusions become captured or adhere to the internal surfaces of the filter by several possible mechanisms (Fig. 4) of diffusion (important for inclusions less than 1 mm, or
Transport and capture mechanisms of depth filtration. (a) Diffusion. (b) Direct interception. (c) Inertia or sedimentation. (d) Fluid dynamics effects on particulate transport
Molten-Metal Filtration / 197 40 min.), direct interception, sedimentation, and fluid dynamics. Filtration is a matter of probability, and so, the greater the flow path through the filter, the greater the opportunity for an inclusion to be subject to changes in direction and velocity and come into contact with the filter internal surface for adhesion and capture. Depth filtration is the filtering mode most prevalent in bed filtration. Depth filtration is capable of removing particles much smaller than 30 mm (1200 min.) as they are carried in a molten-metal flow stream through a very tortuous path in a deep-bed filter (Fig. 5), which is commonly used in mill product casthouses. The tortuosity provides increasing probability for a two-step particle-capture mechanism involving transport from the liquid metal flow and attachment to the filter surface itself. Bed filters composed of tabular Al2O3 are quite common in the casting of mill products, where up to 1800 Mg (2000 tons) of a single alloy may be filtered in-line between holding furnace and casting station before replacement is necessary. Bed filters provide the best source of fine particle filtration, but their size, expense, and single-alloy use preclude their economic use in multiple-alloy wrought casting (ingot) or foundry casting operations. Bed filters are capable of providing nearly the highest filtration efficiency attainable, but because the bed is not rigid, movement and flow channeling can occur, which can dump inclusions.
Filter Characteristics Molten-metal filters obviously require hightemperature integrity in terms of strength, refractoriness, thermal shock resistance, and corrosion resistance. The filters may be placed directly in either the melting or casting furnace (or sometimes both) in the cases of in-furnace applications, or the filters may be in the running system of the casting mold (in-mold filters). In either application, it is important to understand the basic characteristics of the filters. The mechanical removal of undissolved particulates from molten metal follows the same
Fig. 5
basic physics as the filtration of particulate from any liquid. The basic principles of filtration are: Larger inclusions are more easily removed
from melts than smaller ones.
Lower flow rates of metal through a filter
result in greater efficiency of inclusions removal. Filters with larger inlet surface areas have greater inclusion-holding capability and longer filter life, for a given incoming inclusion concentration. Filter bodies with greater internal tortuosity or filter depth provide a higher probability of inclusion capture. However, some basic misunderstandings persist in understanding the characteristics of filters. One misconception is that filters can catch everything unwanted in the melt. As a basic physical principle, mechanical filtration can only remove the insoluble particulates in a liquid. Mechanical filtration cannot remove contaminants that are liquid to semiliquid themselves or dissolved constituents such as hydrogen gas, various low-melting fluxes, or alloying eutectic components such as silicon. In addition, it is critical to understand that reactive molten metals (especially aluminum) react with the environment for the entire time of its molten state. Aluminum is particularly reactive with the atmosphere and reacts easily with oxygen, forming solid particulate aluminum oxide, which is filterable, and hydrogen, which is soluble and therefore not filterable. This includes the time spent traveling in the runner system and filling the mold postfilter placement. Hence, some particulate contamination may still occur if the molten aluminum is allowed to travel in a turbulent state postfiltration. This would be especially true in sand molds where lower compaction, incomplete bonding, and turbulence of metal flow will create the possibility of sand erosion. Secondly, there are significant differences in the mechanism of various types of media described as a filter. The two most common
Example of one type of bed filter used for in-line processing
media found in today’s (2008) aluminum foundries are screens and refractory foam filters (Fig. 6, 7). Screens, for all intent and purpose, are two-dimensional structures with evenly spaced round or square holes. Regardless of composition, that is, metal or fiberglass being the most common, they function primarily as sieves. Sieves can only remove particulate larger than the opening in the sieve or, in this case, the screen. Foam filters, on the other hand, have two completely different and distinct modes of filtration that allow them to act as true filters rather than just simple sieves. Because foam filters are three-dimensional and have a random pattern of openings, they exhibit two completely different behaviors. On the surface, they behave similarly to screens in that particulate larger than the openings of the pores of the foam will be trapped on the inlet side of the foam. This is often referred to in the literature as cake filtration, as previously noted. However, the metal must continue to pass through the pores and cells that make up the internal honeycomb structure of the filter. Within the structure of the filter, due to rapid changes in pressure and metal velocity, particulate much smaller than the actual pores and cells themselves may become trapped. This is commonly referred to as deep-bed or depthmode filtration. It is important to understand that typically, sieving or cake filtration can be nearly 100% under normal conditions for large particulates.
Fig. 6
In-mold ceramic foam filters
Fig. 7
In-mold screens
198 / Molten Metal Processing and Handling Unless excessive force is applied, all particulates greater than the opening size of the holes cannot pass through the holes. On the other hand, deep-bed filtration can only function as a percentage-reduction filtration mechanism dependent on many variables: metal fluidity, velocity, foam filter porosity, foam filter thickness, particle size of the contaminant to be trapped, and so on. That is, some percentage of extremely small particulate will be able to escape even the best foam filter. It therefore follows that the choice of filter medium is based at least in part on the requirement for casting quality. For those castings where very small particulates do not contribute to quality issues, screens may be adequate. However, as the quality requirement of the casting increases, a point will be reached where screens will be incapable of providing the level of melt cleanliness required to meet the casting quality demands. At that point, a ceramic foam honeycomb filter structure becomes the best option to achieve the required quality. The one characteristic sometimes missing in the discussion of aluminum filtration is the function of refractory material in the foam structure. Unlike iron, steel, and copper alloys, the majority of particulate impurities in aluminum alloys is aluminum oxide, the most prevalent impurity, which exhibits very little physical or chemical affinity for most common refractory materials. Unlike iron, for example, where sticky silicacious slags will physically stick to many suitable refractories, aluminum oxide (alumina) in aluminum is “dry” and does not exhibit a particular physical affinity to any refractory. For this reason, there has never been any reliable data presented that support a claim of one refractory being superior over another in nonmetallic reduction in molten aluminum alloys, all other attributes of the foam filter being equal. The refractory composition can be responsible for other filter behaviors, however, such as strength (both cold and hot), friability, thermal shock, wettability, thermal capacity (based on density of the material), thermal cycling (remelt behavior), and soon. For these aforementioned reasons, alumina, silicon carbide, mullite, and graphite have been used as media in foam filters for aluminum filtration. Although they offer different physical refractory performance for any given porosity and thickness, they all behave similarly in particulate-retention capability. Once the appropriate refractory is chosen, the most important physical characteristic of the foam filter is its porosity. At this time (2008), there is no true, accurate universal definition in the industry to describe the porosity of a foam filter. The nomenclature of pores per inch (ppi) is a carryover from the foaming process that is used to make all kinds of foam products, such as cushioning materials, packaging, air filtration substrates, carpet padding, and so on, as well as the reticulated polyurethane foam used as substrate for refractory foam filters.
In general, there are three basic classes of foam porosity most commonly used in the aluminum foundry industry: Coarse (10 ppi) Medium (20 ppi) Fine (30 ppi)
However, the porosity of the foam substrate is only a portion of the final story. No matter what foam is used, the refractory slurry applied to the foam can have different effects on the final foam filter flow characteristics, all else being equal. Therefore, comparing one manufacturer’s 10 ppi filter to another’s 10 ppi filter or even 15 ppi filter leads to confusion on the part of the foundry. Regardless of the ppi rating, what a foundry is really purchasing is the pressure drop that the filter provides the metal stream between the entrance face of the filter and the exit into the balance of the running system. For any given porosity rating and filter thickness, the pressure drop should be consistent, with minimum expected variation. As the pressure drop increases (as is the case with finer-grade filters), the filter has a greater ability to retain a higher and higher percentage of small and smaller particulates. Conversely, as the pressure drop increases, the expected flow rate in pounds per second will decrease. This often creates a basic dilemma. A common mistake made in filter printing is not allowing enough of an increase in the area of the filter versus the runner it is placed in in order to compensate for the filter pressure drop and correspondingly to sustain the necessary flow rate. Under normal circumstances, the rule of thumb is to allow for a 4-to-1 ratio of the usable filter area (excluding the print) versus the runner area in order to compensate for the filter pressure drop. When using coarser filters (10 ppi), this ratio can be somewhat reduced. When finer filtration is required (20 or 30 ppi) along with fast flow rates, the only alternative is to increase the usable surface area of the filter. Although most people focus on the particulate-removal capabilities of refractory foam filters as their primary function, it is, in fact, only a part of the contribution to overall casting quality improvement that they provide. In particular, unlike screens, properly printed foam filters have the capability to reduce the amount of turbulence and entrained air in the liquid metal stream, further minimizing the likelihood of the formation of regenerated oxides and dissolved hydrogen postfilter. In fact, an important additional function of an in-mold ceramic foam filter is to provide flow control. Unlike foam filters, screens actually split the metal stream into up to hundreds of small streams exiting the screen (Fig. 8, 9). This is particularly important in the first few seconds of mold filling prior to steady-state flow being achieved. The formation of the small streams dramatically increases the surface-area-tovolume ratio of the liquid metal, likewise increasing the surface area exposed to the air.
The more liquid metal surface area available to the atmosphere, the greater the potential for reoxidation to occur. Additionally, it can be shown that air bubbles can be easily aspirated into the mold through screens. Unlike screens, refractory foam filters inhibit the transmission of large air bubbles through the foam structure. An aluminum foundry should treat an air bubble as the encapsulated air it truly is. As damaging as an air bubble can be if it remains trapped in the solidified casting, it is equally damaging as a source of oxygen and hydrogen, which can
Fig. 8
Screen water model, split flow
Fig. 9
Ceramic foam filter, convergent flow
Molten-Metal Filtration / 199 react, forming aluminum oxides or dissolving in solution. One way to look at an air bubble in molten aluminum is as if it were a balloon on a string, the air bubble being the balloon and the string a trail of aluminum oxides following along wherever the balloon travels. Because refractory foam filters, via their unique pressure drop, can remove a significant portion of the energy from the metal stream as well as smooth the metal flow free of air bubbles, the metal exiting the refractory foam filter will be substantially less turbulent. Once this turbulence reduction is achieved, it will not remain so unaided by the foundry. The metal stream is the most nonturbulent when exiting the refractory foam filter immediately at the exit face. If anything is done to disrupt the metal stream as it travels away from the filter to the actual mold cavity, turbulence will increase. For this reason, anything that would increase the metal stream velocity postfilter should be avoided. This can include, but is not limited to, a choke postfilter, a further metal drop, a high-aspect-ratio runner (i.e., a runner that is substantially taller than it is wide), or any configuration that would tend to induce splash.
Types of Molten-Metal Filters Filter selection often depends on ease of use. The simplest filters are the strainer cores used in down sprues in conventional foundry sand casting. While serving to choke or control metal flow, they trap some particulates by surface attachment and a degree of cake filtration. For in-mold foundry applications, ceramic foam filters are common. Ceramic foams are lighter, easier to preheat, and provide easier flow capability than bonded-particle filters. In the primary and mill products casting industry, filter selection will depend on economics, space constraints, auxiliary metal treatment capabilities, desired filtration efficiency, and end-product application. Metal or Fiberglass Screens. Steel mesh screens and fiberglass cloth are also commonly used for aluminum. Molybdenum mesh is used for copper alloys. These materials act as planar filters. They provide moderate filtration, particularly of larger particles and inclusions exceeding 100 mm (0.004 in.) in size. The mesh screen, fiberglass cloth, and strainer cores are, of course, single-use filters. Bonded-particle filters consist of a refractory grain (either Al2O3 or SiC) bonded together to form a rigid structure. Bonded-particle filters provide a relatively tortuous path for metal flow through the filter (good depth filtration). These filters are often used vertically to separate melting or holding furnace hearth areas from the dip-out well. Silicon carbide vertical gate filters are quite durable and offer the potential of low thermal gradients between the hot face and cold face because of the relatively high thermal conductivity of silicon carbide compared to other materials, such as oxide ceramics. Temperature
gradients between inlet/outlet sides of the filter may start as low as 5 C but will build with use to gradients of 20 to 40 C (40 to 75 F) have been reported for these filters, versus 35 to 70 C (60 to 130 F) for oxide ceramic foams. Bonded- particle filters possess less porosity than for a ceramic foam (38 versus 75%, respectively) and greater tortuosity; therefore, they function more as depth filters. Cartridge Filters. One interesting application of the rigid filter concept is the cartridge filter. It is simply a shell-and-tube-type construction similar to a heat exchanger. This type of filter provides a great amount of surface area and usually has very fine pore size. Therefore, although a substantial priming head (up to 0.3 m, or 1 ft) of metal may be required before flowthrough begins, and although the flow rate is not very high, filtration efficiency is excellent. Cartridge filters are now receiving attention for such premium-quality applications as aircraft plate, computer disks, and foil. Efficiencies of 95% or better in removing particles less than 5 mm (200 min.) in size have been reported. Ceramic foam filters are the most common types of filters currently used in nonferrous casting. They are produced by slurry coating a reticulated polyurethane cellular foam, followed by drying and firing to burn out the precursor foam. This results in a ceramic replica of the original organic foam structure. Suitable ceramics for these filters are alumina, zirconia, and mullite. Some ceramic filters also employ chromic oxide. These filters are approximately 75% porous and are produced in a variety of pore sizes, as measured by pores per lineal inch (ppi). Ceramic foam filters are used in flat, platelike form. In the primary aluminum industry, plate sizes may range from 300 by 300 to 585 by 585 mm (12 by 12 to 23 by 23 in.), depending on flow rate (up to 680 kg/ min, or 1500 lb/min) and casting size (up to 34,000 kg, or 75,000 lb). Smaller filters are used in the die casting and foundry industries. Dual-structure or multilayered ceramic foam filters have been employed in the past but do not find present application. These use a finer pore size for greater inclusion-capture capability, coupled with a coarser pore size to permit greater throughput before the filter becomes saturated or clogged.
In-Mold Filtration Reduction of entrapped particulates in a casting is governed first by proper melting and handling practices and then by management of the casting process as the metal is poured, fed to the mold, and solidified. In-mold filtration is an important technology in aluminum foundry casting production. Filter application in the mold must follow sound principles to balance the pressure drop accompanying the resistance of a filter within the gating system with the necessary flow rate to minimize turbulence and therefore to achieve a sound, solidified cast structure. Proper selection of the specific filter
type and grade, sizing, and placement of the filter within the gating system will achieve significant particulate removal as well as metal flow control. Filter Selection and Application in Sand Gating. To achieve maximum filtration and flow modification benefits, the filter must be properly sized, placed, and coordinated with the other elements in the gating system. A filter that is too small can impede the flow of metal, increase fill times, and result in short pours if it blocks too soon. Also, remember that a filter is not an excuse to neglect good melting, treatment, and pouring practices. The selection of a filter size and porosity will depend on a variety of factors, including the pour weight of the casting, desired mold filling rate, alloy type, initial metal cleanliness, and pouring temperature. The most basic of guidelines for sizing based on predicted filter blockage for foam filter/flow modifiers is 0.2 to 0.4 g/cm2 (15 to 30 lb/in.2) of filter face before blockage. For flow rate, using the most common porosity foam filter/flow modifiers, the typical range for a 10 ppi filter is 50 to 100 g/cm2/s (0.7 to 1.4 lb/in.2/s) and for a 20 ppi filter, 30 to 80 g/cm2/s (0.4 to 1.1 lb/in.2/s). The filter (or filters in a multifilter system) should be large enough (total exit area and porosity) that the fill time is the same as it would be in an unfiltered casting with the same gating system dimensions. Also, the system should be unpressurized, and the filter should not become the choke in the system. If the filter becomes the choke, the metal will accelerate as it exits the filter and create undesired turbulence. If the ingate is the choke (a pressurized system), the flow modification effect of the filter will be lost as the metal sprays into the casting cavity. When blockage ultimately occurs from buildup of inclusions on and within the filter, it is not gradual but almost immediate. The selected filter must be large enough to maintain normal flow throughout the pour. As a starting point, the filter frontal area should be at least four times the cross-sectional area of the runner. Depending on foundry practice, it may be possible to reduce the filter size as casting experience and confidence are acquired. Filters may be placed horizontally, vertically, or at some intermediate angle within the runner system, depending on the molding method, pattern layout, or other molding restrictions. Filters in the vertical position are best placed just beyond the sprue well that will absorb the energy of the initial falling metal, as shown in Fig. 10. Filters in the horizontal position are best placed just off of the sprue well (Fig.11). Filters placed at the base of the sprue, where direct impingement of the metal stream will force the metal through the filter at high velocity until the sprue fills, may create turbulence in the initial metal entering the casting. Although the metal can be forced through the foam structure at these higher velocities, it can be defeating of the flow-modifying capability of the foam to do so.
200 / Molten Metal Processing and Handling
Fig. 11 Fig. 10
Horizontal installation of filter off of the sprue well
Vertical installation of filter in a mold feeding system
Within the basic guidelines for in-line filter application in sand molds, there is considerable latitude in positioning and placement (Fig. 12a–c). A single centrally located filter may be used, a filter may be placed in each runner, or a filter may be placed at each ingate for maximum effectiveness. Filters may be inserted manually or with core-setting equipment. Filter Selection and Application in Permanent Mold Casting. Most of the gating system design issues critical to sand casting of aluminum apply to permanent mold casting as well. There is even greater emphasis on the need for repeatable flow and tranquil filling characteristics and precisely controlled thermal properties to assure progressive solidification in the direction of the risers and feed heads. Proper gating design is vital to minimize turbulence as the mold fills and to avoid superheating the mold at any one point. The high density of the mold and total lack of permeability, compared to sand molds, put special emphasis on proper venting. Both horizontal and vertical gating and risering systems benefit from use of a tapered sprue to avoid air entrainment, an enlarged basin at the base of the sprue to absorb the kinetic energy of the falling metal, and an extended runner beneath the mold cavity to absorb entrained air and trap initial dross entering the runner system. The choke is normally at the base of the sprue, resulting in an unpressurized system. Quality issues with permanent mold castings may arise from entrained oxides and/or the reoxidation of treated metal, as is the case with sand castings. Filter screens are commonly used in an effort to intercept inclusions in permanent mold castings. However, screens have minimal filtration effect and damage the metal by splitting the stream and creating turbulence.
Development efforts are currently underway to apply foam filters to permanent mold castings to trap inclusions and minimize turbulence and reoxidation in vertical, horizontal, tilt, and low-pressure permanent mold applications. Filter selection is based on the same criteria (pour weight, desired fill rate, alloy type, initial metal cleanliness, pouring temperature) as sand casting. Direct Pouring. The primary purpose of runners and gating systems is to deliver metal to the casting cavity in a nonturbulent fashion. However, even though gating and running system components may be recycled, gating is not free. Significant costs of gating include: Cost of poured metal not turned into saleable
castings, including melt conversion cost, handling, and melt loss Total cost of remelting removed gating metal, including energy, melt loss, alloy loss, and other treatment costs Total cost of molding, including consumption of pattern and mold space, extra sand and binder, more metal, and runner maintenance due to runner erosion Pouring temperature cost, including superheat to compensate for loss of heat in runner system and oxidation aggravated by higher temperatures and turbulence Cleaning costs, including gating-removal labor, equipment and supplies, and handling and storage of removed gating Compromised thermal profile resulting from coldest metal ending up in the riser, and unnecessarily large risers needed for adequate thermal mass to accommodate shrinkage Oxidation and air entrainment resulting from turbulence in runners (rolling-back wave generation and air bubbles)
Fig. 12
Placement of filters (shaded). Filters may be placed in (a) each ingate, (b) one central location, or (c) each runner.
Molten-Metal Filtration / 201 By pouring directly through a sprue attached to the top of the casting cavity or mounted on a side entry close to the casting, many of the costs of a conventional gating system can be reduced or eliminated. Direct pouring will certainly reduce poured weight and increase yield. However, for automotive and other castings with stringent physical requirements, some means must be used to control filling turbulence and prevent mold erosion and/or core damage from direct impingement of the molten metal. A direct-pouring unit for sand or permanent mold casting (Fig. 13) has been developed and patented (FOSECO Metallurgical Inc.) that incorporates an insulating pouring sleeve (that may also act as a riser) with a foam filter to trap inclusions and modify the flow of the metal as it enters the casting cavity. Filter media other than foam have not been proven adaptable to the direct-pouring process. It is essential that turbulence reduction in the metal flow be possible from any device used to attempt direct pouring. In direct pouring, filtration takes a back seat to reduced turbulence in terms of primary operational characteristics.
Fig. 13
Patented direct-pouring filter. Courtesy of FOSECO Metallurgical Inc.
This direct-pouring system replaces all the traditional gating system elements and reduces feeding requirements. By eliminating the runner system and ingates, ingate hot spots requiring feeders are eliminated. Also, with metal entry through the top of the casting, the hottest metal ends up in the direct-pouring unit sleeve, reducing thermal mass requirements if the sleeve serves as a feeder as well. Example of In-Mold Filtration for Scrap Reduction in Sand Casting of Automotive Intake Manifold. Prior to the introduction of filtration/flow modification, a sand cast manifold experienced unacceptably high scrap rates resulting from oxides and other inclusions. Mold-filling turbulence was identified as a contributing factor. Minor modifications to the existing gating system enabled the addition of a single foam filter/flow modifier that led to a reduction of machining scrap and other oxiderelated scrap from 4.6 to 0.8%. With the casting valued at approximately $25 (United States) when the scrap was discovered, this scrap-reduction level represents a gross savings per casting of $0.95. Based on the typical production rate of 325,000 castings per year, this results in an annual gross savings of $308,750. Example of In-Mold Filtration for Scrap Reduction in Casting Oil Cooler Extension. Machining scrap and leakers were responsible for a combined scrap rate of 4.73%. Before the part could be leak tested, it required machining. At the point of machining, the part was valued at $16.75 (United States). After leak testing, the part was valued at $20. Placing a foam filter/flow modifier in the conventional gating system reduced the machining scrap rate from 1.81 to 0.4% and the leaker rate from 2.92 to 0.28%. The combined scrap rate reduction (machining and leakers) results in a gross savings to the foundry of $0.76 per casting or $45,600 based on the annual production rate of 60,000 pieces. Example of Direct-Pour Filtration in Casting Large Commercial Pump Body. Initially, this 30 kg (65 lb) casting included a conventional runner system and had a pour weight of 42 kg (93 lb) for a yield of 70%. By replacing the components of the runner system with a direct-pouring system, the pour weight was reduced by 9 kg (19 lb) and the yield increased to 88%. At $0.10/lb (United States), this accounted for a $1.90 savings per casting. The elimination of the conventional running and gating system also resulted in a reduction in cleaning room time from 6 to 3 min (removal of risers, flash, etc.), saving an additional $1.35 per casting. The filtration capability and flow-modifying effects of the foam filter in the direct-pouring unit reduced the (machining) scrap rate from 7 to 2%. With an approximate casting value of $75 at the time scrap was identified, this resulted in an additional savings of $3.75 per casting. With a total savings of $7.00 per casting and a production rate of 3000 castings per year, this resulted in a $21,000 gross
savings. At the selling price of $110, this represents a gross profit improvement of 6.4%.
In-Furnace Filtration of Molten Aluminum As described in the preceding section, filters are also placed in the feeding system for casting. However, it may not always be possible to employ direct filtering within the mold system. Consequently, in-furnace filtering is an alternative, such as with high-pressure die casting. Other processes, such as gravity permanent mold, low pressure, and sand casting, can also make use of filters directly into a melting or casting furnace. In-furnace filtration provides a very helpful means to the foundry in removing inclusions and thereby increasing the molten-metal quality to produce castings. In-furnace filters can be placed directly in either the melting or casting furnace, or sometimes both. In-furnace filtering products can be applied to most of the typical melting and holding/casting furnaces in use today (2008), with proper evaluation. The benefits of in-furnace filtration can be demonstrated using various analytical techniques on the molten metal itself prior to casting, or by monitoring and evaluation of machining and finishing operations and overall casting production, to determine reduction in the number of scrap castings that can be shown to be inclusion-related.
In-Furnace Filter Types There are several types of filter media that have been employed in a furnace. Historically, the earliest and simplest form of filter was to use a fiberglass sock at the tap-hole or outlet spout on the furnace to catch the dross or other debris. Additionally, use of a ceramic-media bed in a box just off the outlet allowed the metal to flow through and the ceramic media to collect not only the dross or other debris but also to filter some inclusions from the pouring metal itself. While these fiberglass filters at a tap-hole of a melting furnace are still employed and are indeed important, newer furnace filtering schemes have taken over in many instances. For in-furnace applications, either a ceramic foam filter or a bonded-particle filter is often used in modern practice. Many factors tend to favor the latter. The bonded-particle filter is an aggregate of silicon carbide granular particles together with a proprietary ceramic binder in a high-temperature firing process to give body to the filter. The forming process for the bonded-particle filter allows many different sizes, grades, and three-dimensional filter shapes to be produced for specific applications. This is especially useful in that there are many variants in size, refractory construction, and configuration in foundry melting and casting furnaces employed today (2008).
202 / Molten Metal Processing and Handling This type of filter is not suitable as a singleuse, in-mold filter but rather is useful and economical only in applications that allow for sustainable usage. A bonded-particulate filter has a close-packed structure, with just 40% nominal porosity. The open-pore volume at this level is insufficient to permit normal rules of solidification (Chvorinov’s rule) to be followed without an excessively-sized print in the gating to accommodate the flow rate required through the filter to achieve proper solidification in a casting. The bonded-particle filter provides outstanding longevity—often two to three months or more—in many aluminum casting operations because of its chemical stability afforded by the binder, which possesses excellent molten aluminum resistance. The inherent higher mass/density and strength of a bonded-particle filter versus a ceramic foam filter—which replicates a reticulated polyurethane foam structure—favors the application of the former in furnaces where long filter life is desirable. The bonded-particle filter is used either as a vertical gate filter in a melting furnace in front of a tap-hole or placed in front of a transfer pump to provide first-stage filtering. This firststage filtering, which lessens the burden on downstream filtering processes, minimizes entrainment of sludge and other debris associated with melting foundry returns and other materials with a high inclusion content. The molten metal is usually further delivered by ladle for subsequent treatment, such as degassing and/or flux injection, and then transferred to casting furnaces in the permanent mold or high-pressure die casting process or direct poured into sand or precision-sand molds. Figure 14 depicts a typical application of a vertical gate filter in a melting furnace in front of a tap-hole. In the dip-well of a casting furnace, further filtration is advantageous, and this can be accomplished either by another vertical gate filter (Fig. 15) or by an all-media box filter (Fig. 16), from which the casting can be poured by using either a manual ladle or an automatic ladle dipping from the interior. The box filter provides the additional advantage that the filtered metal is always available without regard to buildup and breakaway of inclusions/accretions on the sidewalls of the refractory. Both vertical gate and box filters can be produced in many sizes and geometric shapes to meet specific furnace and ladle dimensional requirements. The long-life potential of a bonded-particle filter allows for its long-term use as an in-furnace filter, complementing any further filtering that may be performed in-mold. In addition to long life, the silicon carbide composition provides greater thermal conductivity than an oxide-based filter, which is important in maintaining temperature in most casting processes. Additionally, the multifaceted particle structure, together with a fully interconnected pore structure, provides a tortuous metal-flow path through the filter. This tortuosity, coupled with
Fig. 14
Vertical gate filter in front of a tap-hole of a reverberatory melting furnace
Fig. 15
Vertical gate filter in a holding or casting furnace. (a) Filter separates heated chamber from dip-well. (b) Installation of a vertical gate filter in a new furnace dip-well
Fig. 16 dip-well
Box filters. (a) All-media bonded-particle box filter. (b) All-media box filter employed in a
Molten-Metal Filtration / 203 the chemical and physical affinity of the binder system, enhances the probability of inclusion capture and hence contributes to enhanced filtration efficiency.
In-Furnace Filter Installation Filter installation into reverberatory furnaces and use of in-furnace filters can be accomplished with careful effort. Measurements of the dip-well dimensions, or other areas in the furnace where the filter will be installed, should be taken carefully, with emphasis on the clean refractory, removing any buildup prior to taking the measurements. For a vertical gate filter, gaskets may be used on the filter edges to provide a seal and take up some degree of irregularity in the refractory sidewalls. It is necessary to fully scrape the sidewalls and bottom of the furnace vessel to provide a good fit. It is also very advantageous to create accepting slots in the furnace refractory for the filter, which can be done during a furnace relining. For either a vertical gate filter or a box filter, it is necessary to allow sufficient space for the ladling operation. Preheat of the filter is necessary to remove any moisture condensation and to prevent thermal shock, but with a thermally-conductive silicon carbide filter, this is straightforward and only requires that the filter be preheated to a temperature of 200 to 300 C (400 to 575 F). This may be accomplished by placing the filter on the furnace sill plate, covered with an insulating blanket, against a furnace hot wall or in a nearby oven. The filter may then be placed directly into the molten metal in the furnace as quickly as safety will allow, without the need to drain the furnace. When a melting or casting furnace is being originally manufactured, relined, or rebuilt, it is often desirable to create slots in the refractory sidewalls to accommodate a filter. In this case, the face-edges of the filter may have a compressible fiber gasket attached to further create a seal in the contact area between the filter medium and the sides or shoulder of the slotted construction. Crucible Furnace with Filter Baffle. Filter installation into crucible furnaces is somewhat similar in principle to installation into a reverberatory furnace, but there are specific differences. The filter can be viewed as a baffle separating the crucible into two sections, which may or may not be equal. Filter baffles have also been placed in the pouring ladle (Fig. 17). If a crucible furnace is used for both charging and degassing (and/or other melt treatment), space must be allowed for casting. In that instance, it is important to maintain both temperature and molten-metal level stability by matching intake with take-out volume, in order to avoid sludge formation and to maximize filtering surface area effectiveness. Further, the contours of the crucible place an extra imposition on the fabricated filter shape to become
sealed to the crucible wall to avoid metal bypass from unfiltered to filtered section, usually demanding extra gasketting. In addition, if due to space/volume constraints the filter must be placed offset to the main diameter, stops are necessary to avoid filter movement into the more open diameter. Figure 18 portrays a ceramic bonded-particle filter placement in a crucible. As an alternative to a baffle filter, there may be sufficient space available in casting crucible holding furnaces, especially larger ones, to accommodate a filtering vessel, such as an allmedia filter or one constructed of refractory materials with a smaller filter that can be replaced when spent. Low-pressure furnaces readily employ filters directly in or onto the stalk or riser tube. Both ceramic foam filters and bonded-particle filters are used in this application. In the lowpressure furnace, application of air forces metal up the stalk or riser tube and into the mold cavity with each cycle. Excess metal (the riser volume) drains back through the filter at the completion of each cycle. The filter medium must not create an undue restriction to metal flow, which would affect the pressurization cycle, nor the backflow of metal from the riser
during the drainage. The lower-porosity ceramic foam filters are often found to be more advantageous in this regard, with a disc shape being easily affixed into or onto the end of the riser tube. Therefore, bonded-particle filters with substantially lower porosity tend to be supplied in more three-dimensional shapes (Fig. 19) in this application to provide more filtering surface area and hence metal-flow capability. To accommodate a larger filter, there must be a sufficient gap or space between the end of the riser tube and the floor of the furnace, as shown in Fig. 19.
Factors That Determine Filter Life There are many factors that determine filter life, including overall incoming metal quality on the inlet side, the specific filter grade, metal flow rate and total throughput through the filter, and metal-level drawdown (in which the filter loses its prime). Other than simply gaining experience and knowledge about how long an in-furnace filter can be used, there are two observations that can be made. As the filter begins to clog, there will be a metal-level difference between the upstream and downstream
Fig. 17
Filter baffles in (a) a crucible and (b) a pouring ladle
Fig. 18
Crucible baffle filter
Fig. 19
Schematic of low-pressure furnace with bonded-particle filter
204 / Molten Metal Processing and Handling sides of the filter. In addition, there will often be a measurable and increasing temperature difference between the upstream and downstream metal. This is caused by the increasing concentration of oxide inclusions that are lower in thermal conductivity than both the molten metal and the silicon carbide filter. The following factors are extremely important in achieving optimum performance of any in-furnace filtering application: Intrinsically, the finer the filter grade and the
lower the filter inherent porosity, the better will be the filter particulate-capturing ability, the quicker the filter will clog, and therefore, the shorter the filter life. By contrast, the larger the filter, the more surface area is available to accommodate both metal flow and inclusion buildup, by both cake and depth-mode filtration. Any ceramic filter must be preheated to at least a minimal extent before being inserted or exposed to molten metal in order to resist thermal shock and premature failure due to mechanical/thermal damage. Excessive metal-level drawdown will lessen the filter useful life, since it will lose the filtering capability of the exposed surface. Heat will be lost as it wicks from an exposed filter surface, unless the filter is covered with an insulating material. In a casting furnace, the molten metal itself is a heat source coming through to a dipwell. When not casting, the furnace should be kept covered. Otherwise, heat will be lost, and the oxide buildup on the inlet side of the filter may not allow sufficient heat to be recovered through to the casting side when casting again commences. Naturally, the inclusion concentration in the molten metal being delivered to the filter in a given time, and the volume of molten metal passing through the filter, will affect filter life.
Fig. 20
Data comparison from K-mold technique in measuring melt quality of filtered and unfiltered molten aluminum ready for die casting
Dirtier metal will clog a filter faster and therefore require more frequent filter change-out.
Filtered Metal Quality It is often desirable to be able to demonstrate the differences in quality between filtered and unfiltered molten metal. There are several means that are possible to evaluate this, including a few real-time, shopfloor methods as well as more extensive laboratory methods. After initial molten-metal quality improvements have been demonstrated by the application of filtration, the true measure of improvement must be shown in product-quality enhancements on the cast product. Performance results of in-furnace filtering systems often show significantly reduced scrap, resulting in greater quality, productivity, and profitability. Specifically, one of the biggest benefits is improvement in machinability, since inclusion removal minimizes hardspots in castings that cause tool wear and tool breakage. Inclusions also nucleate porosity formation during solidification, which can cause leakers; filtering removes these and therefore improves
pressuretightness. Mechanical properties are also improved, particularly elongation and fatigue strength, which are quite sensitive to the presence of inclusions. Surface finish operations—polishing, anodizing, powder coating, and so on—are also enhanced, because inclusion-causing blemishes are minimized or eliminated. Beneficial effects of in-furnace filtration can be demonstrated by several different methods of measuring molten-metal quality. Figure 20 shows the difference between unfiltered and filtered molten metal using the real-time shopfloor K-mold technique. In this test, molten metal is carefully sampled several times in a notched-bar fracture test, and the fracture surfaces are examined macroscopically or with the aid of low-power magnification to observe any surface discontinuities, which will relate to inclusions, dross, or porosity. An index value can be determined; Fig. 20 shows that, of 20 sample surfaces investigated, filtration lowered the observed discontinuity count dramatically, to as little as 1 observation in 20 (one-twentieth, or 0.05). Another measurement method of melt quality is the Prefil Footprinter technique (N-Tec Ltd.). It has been used to compare quality relative to
Fig. 21
Prefil results on A356 benchmarked to other data (bars)
Fig. 22
Porous disc filtration analysis results on 380 die cast alloy. Flux injection in transfer ladle followed by settling and filtration in casting furnace
Molten-Metal Filtration / 205 data for several casting foundries using the same alloy. Prefil results are shown in Fig. 21 on benchmarked A356. Figure 21 shows the impact of in-furnace filtration, providing acceptable molten-metal quality suitable for a high-integrity casting process. In many aluminum foundry casting processes, metal is melted in a melting furnace, then transferred via transfer ladle to the casting furnace or to the mold. Metal treatment, such as fluxing, degassing, flux injection, grain refining, and modification, is often made in this transfer ladle. Figure 22 presents porous disc filtration analysis (PoDFA) (Alcan International Ltd.) results on a die casting alloy at several steps of the process. There is a proportionate improvement in metal quality, as evidenced by the metallographic inclusion average as obtained by image analysis, at each stage of the process. Even after settling time in the casting furnace, there is a further reduction in inclusion concentration after the metal has flowed through an in-furnace vertical gate filter. While the foregoing represents analytical techniques to assess the effectiveness of in-furnace filtration on molten-metal quality, it is necessary for the foundry to determine the beneficial effects and cost-effectiveness of filtration on the ultimate cast product. Destructive testing can be done on random or selected samples, especially to determine solidification structure and other microstructural features in the critical areas of a given casting, where engineering performance
criteria must be met. Separately-cast test bars do offer some measure of physical metal quality but do not reflect the cast part per se. Two mechanical properties that are most sensitive to the presence of, or absence of, inclusions are tensile elongation/ductility and fatigue strength. The presence of inclusions can lower the ductility and fatigue strength cycles to failure. However, good tensile results do not necessarily imply good casting quality, and a substantial number of test bars must be evaluated to create statistical significance. There are three foundry production methods to evaluate casting quality as affected by in-furnace filtration. Overall measurement of scrap quality and identification of those scrap castings that are due to observable inclusion-related defects, improved machining performance through elimination or minimization of hard spots, and leak testing (inclusions nucleate porosity formation) are three of the best means to ascertain the benefits of filtration. REFERENCES 1. D.E. Groteke, A Production Filtration Process for Aluminum Melts, Trans. Soc. Die Cast. Eng., 1983 2. D.V. Neff, Continuous, Sustained, ‘ReUsable’ Filtration Systems for Aluminum Foundries and Diecasters, Proceedings, Fourth International Conference on Molten Aluminum Processing, AFS, 1995, p 121
3. D.V. Neff, Using Filters and Evaluating Their Effectiveness in High Pressure Diecasting, Trans. NADCA, Sept 2002 4. D.V. Neff, The Role of Melt Treatment in the Production of High Integrity Aluminum Castings, Proceedings, International Molten Aluminum Conference, AFS, 2006 5. D.V. Neff, Ensuring Quality Prior to Pouring, Mod. Cast., June 2006, p 27 6. D.V. Neff, The Exclusion of Inclusions, Die Cast. Eng., Jan 2008, p 48 7. C.J. Simensen and G. Berg, A Survey of Inclusions in Aluminum, Aluminum, Vol 56, 1980, p 335 8. L.J. Gauckler and M.M. Wacher, Industrial Application of Open Pore Ceramic Foam for Molten Metal Infiltration, Light Metals, The Metallurgical Society, 1985, p 126 9. J. Langerweger, Nonmetallic Particles as the Cause of Structural Porosity, Heterogeneous Cell Structure and Surface Cracks in DC Cast Al Products, Light Metals, The Metallurgical Society, 1981, p 685 10. C.J. Simensen, The Effect of Melt Refining Upon Inclusions in Aluminum, Metall. Trans., 1980 11. C.J. Simensen and U. Hartvedt, “Analysis of Oxides in Aluminum by Means of Melt Filtration,” Center for Industrial Research, Trondheim, Norway, 1984 12. C.J. Simensen, Sedimentation Analysis of Inclusions in Aluminum and Magnesium, Metall. Trans. B, 1981, p 733
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 206-229 DOI: 10.1361/asmhba0005199
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Steel Melt Processing MELT PROCESSING of steels can be broadly classified as either primary steelmaking or secondary steelmaking. The two most important primary steelmaking processes are the electric arc furnace (EAF) process and the basic oxygen furnace (BOF) process or Linz-Donawitz process. The EAF process has been increasingly used, gradually overtaking the BOF as the primary steelmaking furnace in the United States. Some other processes, such as the open-hearth furnace, are still practiced in a few countries to produce special steels. Primary melt refinement also includes the uses of converters, which are furnaces in which oxygen is blown through a bath of molten metal, oxidizing the impurities and maintaining the temperature through the heat produced by the oxidation reaction. A typical converter in steel refining is the argon oxygen decarburization (AOD) vessel. In most modern steel operations, the steel is melted and only roughly refined in one furnace. Then, it goes to a secondary operation where it is further processed to yield the final desired chemistry. This secondary operation can be carried out while the steel is in the furnace, ladle, the degasser unit, or, in the case of stainless steel, the liquid metal can be poured into an AOD vessel. After secondary refining, the molten metal is transported to a tundish or a bottom pure ingot making facility for teeming. The melting furnaces used in steel foundries are essentially the same as those used for primary steelmaking, except that most foundry melting units are smaller. Although the equipment is the same, there are also important process differences. For example, one major difference between melting practice in a steel foundry and a steel mill is the higher tapping temperature used for foundry melting to attain better fluidity of the molten steel. Refractory materials in steel foundries are also different than those in primary steel production. Electric arc and induction melting are common methods of melting steels. After furnace melting, the molten steel may be refined in a converter (outside the melting unit) or in a ladle. Thus, steel melt refinements are classified as treatments by either converter metallurgy or ladle metallurgy. Converter metallurgy includes melt refinement in AOD vessels and vacuum oxygen decarburization in a converter vessel. Ladle metallurgy refers to melt treatments in a ladle refining station after melting in the
furnace/converter unit. Ladle metallurgy is used for deoxidation, decarburization, and adjustment of chemical composition, including gases. Methods of ladle metallurgy include: Vacuum induction degassing Vacuum oxygen decarburization Vacuum ladle degassing
These vacuum methods of melt refinement are briefly reviewed in the section “Ladle Metallurgy” in this article. Nonvacuum processes (carried out at atmospheric pressure) for production of special and clean steel and without any supplemental reheating include various argon bubbling processes (such as AOD), electroslag remelting (ESR), and ladle injection methods. Argon bubbling methods are discussed in the section “Converter Metallurgy” in this article. Melt refinement by ESR furnaces is covered in the article “Electroslag Remelting” in this Volume.
Furnaces and Refractories Electric arc and induction melting are the most common methods of melting steels. Coreless induction furnaces are also used for melting carbon, low-alloy steels, stainless steels, high-alloy steels, or nickel-base heat-resistant alloys. The type of furnace, refractories, and melt processing equipment depends on a number of factors that include the quality of the
Fig. 1
Refractory lining of ladles
charge material to be melted, chemistry control requirements, capital equipment costs, manpower, and energy and operational costs.
Refractories Several different types of refractories can be considered for the melting furnaces, ladles, and tundishes for holding and transporting molten metal. Usually, placement of refractory linings involves a practice called zoning, where refractories with different properties are used in different areas of the same vessel to resist the attack of the steelmaking process in that specific area. For example, several wear mechanisms occur in a ladle. At the slagline, wear proceeds primarily by corrosion, and magnesia-carbon or magnesia-chrome bricks are commonly used as the hot-face liner (Fig. 1). In the bottom and lower sidewalls of the ladle, wear spots can develop by erosion of bricks or by the impingement of the molten steel stream on the refractory. It has been reported that 70% alumina, dolomite, and alumina-spinel can be used as the hot-face liner in this area (Ref 1). Monolithic refractory use has progressed in accordance with the need for more mechanized lining installation and repair. Corrosion, oxidation, and thermal cycling contribute to lining wear. Properties are derived from the major refractory constituents, additives, microstructure, grain size, and fabrication process (bricks or
Steel Melt Processing / 207 castable). Changing these factors imparts specific characteristics that include melting temperature, creep resistance, oxidation resistance, corrosion resistance, mechanical strength, and erosion resistance. Lining technology is constantly being updated by better and more costeffective refractories. Refractories are classified as basic, acid, or neutral by an index of the “basicity” of the slag formed on the metal surface. The compatibility of this basicity index with a refractory can give an indication of its potential to corrode a refractory. Basicity can be represented by different indexes, such as the B2-ratio (the weight ratio of CaO to SiO2 in the slag) and the V-ratio (the weight ratio of CaO + MgO to SiO2 in the slag) (Ref 2): CaO ðwt%Þ SiO2 ðwt%Þ CaO þ MgO V¼ SiO2
B2 ¼
Steel slags typically contain five major oxides: CaO, MgO, SiO2, FeO, and Al2O3. Slags with B2- or V-ratios less than 1 are defined as acid slags, while basic slags are more than 1. The vast majority of steel is produced with a basic slag, while most iron production uses an acid slag. That is, most steelmaking refractories are composed of MgO and/or CaO, while ironmaking refractories consist of Al2O3 and SiO2. The use of these or other slag indexes may vary from plant to plant and from heat to heat, depending on the grade of steel being made. Foundries that have access to a good grade of scrap, and therefore do not need to reduce the phosphorus and sulfur contents of the steel to meet specifications, usually prefer to use a furnace lined with silica refractories (acid lined). Foundries that do not have a good source of low-phosphorus, low-sulfur scrap use refractories such as magnesite and dolomite for furnace lining (basic lining). The compositions of certain steels, such as austenitic manganese steels, require that they be made in a furnace with a basic lining. Basic Refractories. Basic linings are principally magnesium oxide, although some linings contain chromium oxide, calcium oxide, and other oxides in controlled amounts (Ref 3). They resist temperatures several hundreds of degrees higher than acid linings. Basic refractories are magnesite-based and are either natural or synthetic magnesite, brucite, and dolomite. Dense synthetic magnesia is produced from seawater and has a purity of 95 to 99% MgO. Natural magnesite is present as a hydroxide or a carbonate. The double carbonate of dolomite (CaCO3MgCO3) is found in abundance. Hightemperature calcining (firing in kilns to dissociate the carbonates) is required in all the cases to produce a dense material with minimum impurities. Magnesia-chrome refractory spinels are also commonly used without firing in combination with magnesia refractory. Neutral refractories are typically alumina refractory bricks, as high-alumina silicates or
pure sintered alumina. Neutral refractories are very resistant to chemical attack and are more tolerant to intermittent melting than acid or basic linings.
Direct Arc Melting The direct arc furnace consists essentially of a metal shell lined with refractories (Fig. 2). This lining forms a melting chamber, the hearth of which is bowl shaped. Three carbon or graphite electrodes carry the current into the furnace. Steel, either solid or molten, is the common conductor for the current flowing between the electrodes. The metal is melted by arcs from the electrodes to the metal charge—both by direct impingement of the arcs and by radiation from the roof and walls. The electrodes are controlled automatically so that an arc of proper height can be maintained. For basic electric furnace melting, the furnace lining is a basic refractory such as magnesite or dolomite. The charge is usually composed of purchased scrap steel and foundry returns. During the melting period, small quantities of lime are occasionally added to form a protective slag over the molten metal. Iron ore is added to the bath just as melting is complete. The slag is then highly oxidizing and in the correct condition to take up phosphorus from the metal. Shortly after all of the steel has melted, this first slag is taken off (if a two-slag process is to be used), and a new slag composed of lime, fluorspar, and sometimes a little sand is added. As soon as the second slag is melted, the current is reduced, and pulverized coke, carbon, or ferrosilicon, or a combination of these, is spread at intervals over the surface of the bath. This period of furnace operation is known as
Fig. 2
the refining period, and its purposes are to reduce the oxides of iron and manganese in the slag and to form a calcium carbide slag, which is essential to the removal of sulfur from the metal. The refining slag has approximately the following composition: Element
CaO SiO2 FeO CaF
Composition, %
45–55 15–20 0.50–1.5 5–15
Adjustments are made in the carbon content of the bath by the addition of a low-phosphorus pig iron. After the proper bath temperature is obtained, ferromanganese and ferrosilicon are added and the furnace is tapped. Aluminum is generally added in the ladle as a final deoxidizer.
Induction Melting The high-frequency induction furnace is essentially an air transformer in which the primary is a coil of water-cooled copper tubing and the secondary is the metal charge. The circular winding of copper tubing is placed inside the shell. Firebrick is placed on the bottom of the shell, and the space between that and the coil is rammed with grain refractory. The furnace chamber may be a refractory crucible or a rammed and sintered lining. General practice is to use ganister rammed around a steel shell that melts down with the first heat, leaving a sintered lining. Basic linings are often preferred; a rammed lining of magnesia grain or a clay-bonded magnesia crucible can be used. The process consists of charging the furnace with steel scrap and then passing a high-frequency current through the primary coil, thus
Schematic cross section of a typical electric arc furnace showing the application of different refractories
208 / Molten Metal Processing and Handling inducing a much heavier secondary current in the charge, which heats the furnace to the desired temperature. As soon as a pool of liquid metal is formed, a pronounced stirring action takes place in the molten metal, which helps to accelerate melting. In this process, melting is rapid, and there is only a slight loss of the easily oxidized elements. If a capacity melt is required, steel scrap is continually added during the melting-down period. Once melting is complete, the desired superheat temperature is obtained, and the metal is deoxidized and tapped. Because only 10 to 15 min elapse from the time the charge is melted down until the heat is tapped, there is not sufficient time for chemical analysis. Therefore, the charge is usually carefully selected from scrap and alloys of known composition in order to produce the desired composition in the finished steel. Composition can be closely controlled in this manner. In most induction furnaces, no attempt is made to melt under a slag cover, because the stirring action of the bath makes it difficult to maintain a slag blanket on the metal. However, slag is not required, because oxidation is slight. The induction furnace is especially valuable because of its flexibility in operation, particularly in the production of small lots of alloy steel castings. Induction melting is also well adapted to the melting of low-carbon steels, because no carbon is picked up from electrodes, as may occur in an EAF. Coreless Induction Furnaces. Both channel- and coreless-type induction electric furnaces are used in the steel industry. The coreless induction furnace melts by passing an alternating current through a water-cooled coil that surrounds the refractory linings of the furnace. The current generates a magnetic field that induces currents in the charge and molten metal. Coreless induction furnaces provide operational flexibility when processing different alloys, because the furnace can be easily emptied and recharged. Vacuum Induction Melting (VIM). In the VIM remelting process, the charge material has basically the same chemical analysis as the final melt. No major refining or metallurgical work is performed other than lowering the residuals and inclusions. When remelting alloy steel, certain elements that readily vaporize or oxidize must be added under controlled conditions. As a rule of thumb, alloying elements that oxidize should be added as late as possible, and difficult-to-dissolve elements should be added early and during high stirring periods. The furnace lining is dependent on the nature of the slag and the desired finish temperature. Oxygen, hydrogen, and, to some extent, nitrogen can be removed during vacuum melting. Carbon can also be lowered to produce some types of stainless steels. Pressures as low as 5 mm of mercury are used in these furnaces. The control of pressure and composition of gas over the melt makes it possible to deoxidize the melt with carbon or hydrogen, because they produce gaseous deoxidation products, preventing the formation of nonmetallic inclusions. Use of low
pressures also eliminates nitrogen pickup by steel. The volatility of certain alloying elements such as chromium, aluminum, and manganese can result in high losses that can be minimized by replacing the vacuum with an inert gas atmosphere over the melt during additions. Vacuum induction melting is often employed as a remelting operation for very pure steels. It is also used as a first stage in conjunction with vacuum arc remelting (VAR) furnaces and ESR furnaces for duplex (VIM-VAR) or triplex (VIMESR-VAR) melting. Some modified types of VIM equipment (such as the Therm-I-Vac process) use either a vacuum induction melter under low pressure or as a stream degasser. The system can be used for cold charging of steel in the induction furnace and melting, refining, and teeming of steel under low pressure. The steel could also be hot charged under low pressure in the vacuum furnace during which stream degassing occurs, reheated to compensate for the losses during degassing, adjusted for alloy additions to meet the chemical composition specifications, and teemed, all under low pressure. Superheating of steel is not required if hot charging is performed. The steel can be partly deoxidized before charging into the vacuum induction melter. This equipment is mostly used as a vacuum induction unit to melt a cold charge.
Converter Metallurgy Converters are used in secondary metallurgy to refine melts outside the primary metallurgical melting unit. Various converters are available that apply the bottom blowing of oxygen/inert gas mixtures. These bottom-blowing converters use different types, numbers, and arrangements of injection nozzles and the following gases: Argon as a cooling inert gas with a purity
ranging from 85 to 99.99%
Nitrogen as a cooling inert gas with a purity
of 99 to 99.9%
Argon-oxygen and nitrogen-oxygen mix-
tures in the case of the AOD converter
Dry air as a diluted reactive gas Oxygen as a reactive gas with a purity of 80
to 99.5%
Steam as a cooling reactive gas Carbon dioxide as a diluted reactive gas
Bubbling methods allow quick and uniform mixing of alloys, temperature homogenization, adjustment of chemical composition, and partial removal of nonmetallic inclusions. These functions are accomplished by lowering a refractory-protected lance close to the ladle bottom or by blowing through porous refractory plugs at the bottom or sidewall of the ladle or vessel. This section describes the converter designs for degassing: Argon oxygen decarburization Oxygen top and bottom blowing converter,
which is an extension of AOD technology
Vacuum oxygen decarburization converter
In addition to AOD in a converter unit, argon bubbling processes also include capped argon bubbling and composition adjustment by sealed argon bubbling. These bubbling methods are described with other ladle injection methods in the section “Ladle Metallurgy” of this article.
Argon Oxygen Decarburization Some foundries have installed AOD units to achieve some of the results that vacuum melting can produce. These units look very much like Bessemer converters with tuyeres in the lower sidewalls for the injection of argon or nitrogen and oxygen. They are processing units that must be charged with molten metal from an arc or induction furnace. Up to approximately 20% cold charge can be added to an AOD unit; however, the cold charge is usually less than 20% and consists of solid ferroalloys. The continuous injection of gases causes a violent stirring action and intimate mixing of slag and metal, which can lower sulfur values to below 0.005%. The gas contents approach or may be even lower than those of vacuum induction melted steel. The dilution of oxygen with inert gas, argon, or nitrogen causes the carbon-oxygen reaction to go to completion in favor of the oxidation reaction of iron and the oxidizable elements, notably chromium in stainless steel. Therefore, superior chromium recoveries from less expensive high-carbon ferrochromium are obtained compared to those of electric arc melting practices. Argon oxygen decarburization is a secondary refining process that was originally developed to reduce material and operating costs and to increase the productivity in production of chromium-bearing stainless steels. A premelt is prepared in an electric arc furnace by charging high-carbon ferrochrome, ferrosilicon, stainless steel scrap, burned lime, and fluorspar and melting the charge to the desired temperature. The heat is then tapped, deslagged, weighed, and transferred into an AOD vessel, which consists of a refractory-lined steel shell mounted on a tiltable trunnion ring (Fig. 3). As shown in Fig. 3, process gases (oxygen, argon, and nitrogen) are injected through submerged, sidemounted tuyeres. The primary aspect of the AOD process is the shift in the decarburization thermodynamics that is afforded by blowing with mixtures of oxygen and inert gas as opposed to pure oxygen. The heat is decarburized in the AOD vessel to 0.03% C in stages, during which the inert gas-to-oxygen ratio of the blown gas increases from 1-to-3 to 3-to-1. During the blowing, fluxes are added to the furnace and a slag is prepared. Following the decarburization blow, ferrosilicon is added, and the heat is argon stirred for a short period. The furnace is then turned down, a chemistry sample is taken, and the heat is deslagged. Additional alloying ele ments are added if adjustments are necessary,
Steel Melt Processing / 209
Fig. 3
Schematic of an argon oxygen decarburization vessel
and the heat is tapped into a ladle and poured into ingot molds or a continuous casting machine. With the AOD process, steels with low hydrogen (<2 ppm) and nitrogen (<0.005%) can be produced with complete recovery of chromium. In addition to its economic merits, AOD offers improved metal cleanliness, which is measured by low unwanted residual-element contents and gas contents; this ensures superior mechanical properties. The AOD process is duplexed, with molten metal supplied from a separate melting source to the AOD refining unit (vessel). The source of the molten metal is usually an EAF or a coreless induction furnace. Foundries and integrated steel mills use vessels ranging in nominal capacity from 1 to 160 Mg (1 to 175 tons). Although the process was initially targeted for stainless steel production, AOD is used in refining a wide range of alloys, including:
Stainless steels Tool steels Silicon (electrical) steels Carbon steels, low-alloy steels, and highstrength low-alloy steels High-temperature alloys and superalloys Argon oxygen decarburization units are used in the production of high-alloy castings, particularly of grades that are prone to defects due to high gas contents. Carbon and low- alloy steels for castings with heavy wall sections may be subject to hydrogen embrittlement and are also processed in these units with good results. Fundamentals. In the AOD process, oxygen, argon, and nitrogen are injected into a molten metal bath through submerged, side-mounted tuyeres. The primary aspect of the AOD process is the shift in the decarburization thermodynamics that is afforded by blowing with mixtures of oxygen and inert gas as opposed to pure oxygen. References 4 to 6 contain
Fig. 4
Carbon-chromium equilibrium curves
detailed discussions of the thermodynamics and the dilution principle and its influence on the refining of chromium-bearing steel. To understand the AOD process, it is necessary to examine the thermodynamics governing the reactions that occur in the refining of stainless steel, that is, the relationship among carbon, chromium, chromium oxide (Cr3O4), and carbon monoxide (CO). The overall reaction in the decarburization of chromium-containing steel can be written as: / Cr3 O4 þ C Ð 3/4 Cr þ COðgÞ
14
(Eq 1)
The equilibrium constant, K, is given by: K¼
ðaCr Þ3=4 ðPCO Þ ðaCr3 O4 Þ1=4 ðaC Þ
(Eq 2)
where a and P represent the activity and partial pressure, respectively. At a given temperature, there is a fixed, limited amount of chromium that can exist in the molten bath that is in equilibrium with carbon. By examining Eq 2, one can see that by reducing the partial pressure of CO, the quantity of chromium that can exist in the molten bath in equilibrium with carbon increases. The partial pressure of CO can be reduced by injecting mixtures of oxygen and inert gas during the decarburization of stainless steel. Figure 4 illustrates the relationship among carbon, chromium, and temperature for a partial pressure of CO equal to 1 and 0.10 atm (1000 and 100 mbar, or 760 and 76 torr). The data shown in Fig. 4 indicate that diluting the partial pressure of CO allows lower carbon levels to be obtained at higher chromium contents with lower temperatures. In refining stainless steel, it is generally necessary to decarburize the molten bath to less than 0.05% C. Chromium is quite susceptible to oxidation; therefore, prior to the introduction
of the AOD process, decarburization was accomplished by withholding most of the chromium until the bath had been decarburized by oxygen lancing. After the bath was fully decarburized, low-carbon ferrochromium and other low-carbon ferroalloys were added to the melt to meet chemical specifications. Dilution of the partial pressure of CO allows the removal of carbon to low levels without excessive chromium oxidation. This practice enables the use of high-carbon ferroalloys in the charge mix, avoiding the substantially more expensive low-carbon ferroalloys. Figure 5 compares the refining steps in the two processes. Equally important, however, are the processing advantages offered by submerged blowing. These include excellent slag/metal and gas/ metal contact, superior decarburization kinetics, and 100% utilization of the injected oxygen by reaction with the bath. In addition, submerged injection allows the accurate control of endpoint nitrogen and the removal of dissolved gases (nitrogen and hydrogen) and nonmetallic inclusions. Equipment. The processing vessel consists of a refractory-lined steel shell mounted on a tiltable trunnion ring (Fig. 3). With a removable, conical cover in place, the vessel outline is sometimes described as pear shaped. Several basic refractory types and various quality levels of the refractories have gained widespread acceptance (Ref 7, 8). Dolomite refractories are used in most AOD installations; magnesite-chromium refractories are predominant in small (<9 Mg, or 10 ton) installations and are used almost exclusively in Japan. As seen in Fig. 3, process gases are injected through submerged, side-mounted tuyeres. The number and relative positioning of tuyeres are determined in part by vessel size, range of heat sizes, process gas flow rates, and the types of alloys refined. Process gases are oxygen, nitrogen, argon, and, in some cases, carbon dioxide (CO2). The most recent AOD installations include the use of top-blown oxygen (one-third of all AOD installations have this capability). Oxygen top- and bottom-blowing converters are described in the discussion “Oxygen Top and Bottom Blowing” in this section. The specific volume of AOD vessels ranges from 0.4 to 0.8 m3/Mg. The gas control system supplies the process gases at nominal rates of 0.5 to 3 m3/min/Mg. The system accurately controls the flow rates and monitors the amount of gas injected into the bath to enable the operator to control the process and keep track of the total oxygen injected. Normally, a shop has three interchangeable vessels. At any given time, one of the vessels is in a tiltable trunnion ring refining steel, a second vessel is at a preheating station, and the third vessel is being relined. The vessel in the trunnion ring can be replaced with a preheated vessel in less than 1 h. The control of process gases, vessel activities, and ancillary equipment can range from manual to fully automated. Most installations
210 / Molten Metal Processing and Handling
Fig. 5
Composition changes in refining type 304-L stainless steel using electric arc furnace practice and argon oxygen decarburization (AOD)
Fig. 6
Schematic of stainless steel refining cycle for small (<27 Mg, or 30 ton) vessels showing the relationship between carbon and chromium contents as a function of time and temperature
are equipped with a computer to assist in process control by calculating the required amount of oxygen as well as alloying additions. Some installations have computer control systems capable of sending set points and flow rates to the gas control systems. AOD Processing of Stainless Steels. Charge materials (scrap and ferroalloys) are melted in the melting furnace. The charge is usually melted with the chromium, nickel, and manganese concentrations at midrange specifications. The carbon content at meltdown can vary from 0.50 to 3.0%, depending on the scrap content of the charge. Once the charge is melted down, the heat is tapped, and the slag is removed and weighed prior to charging the AOD vessel. In the refining of stainless steel grades, oxygen and inert gas are injected into the bath in a stepwise manner. The ratio of oxygen to inert gas injected decreases (3:1, 1:1, 1:3) as the carbon level decreases. Once the aim carbon level is obtained, a reduction mix (silicon, aluminum, and lime) is added. If extralow sulfur levels are desired, a second desulfurization can be added. Both of these steps are followed by an argon stir. After reduction, a complete chemistry sample is usually taken and trim additions made following analysis. Figure 6 illustrates the relationships among carbon, chromium, temperature, and the various processing steps for refining a typical type 304 stainless steel using AOD. Another critical aspect of AOD refining is the ability to predict when to change from nitrogen to argon to obtain the aim nitrogen specification. Table 1 lists the process parameters, carbon removal efficiencies, and carbon removal rates for the AOD refining of type 304 stainless steel. The point during refining when the oxygen-to-inert-gas ratio is lowered is based on carbon content and temperature. The ratios and carbon switch points are designed to provide optimal carbon removal efficiency without exceeding a bath temperature of 1730 C (3150 F). AOD Processing of Carbon and Low-Alloy Steels. The refining of carbon and low-alloy steels involves a two-step practice: a carbon removal step, followed by a reduction/heating step. The lower alloy content of these steels eliminates the need for injecting less than a 3:1 ratio of oxygen to inert gas. Once the aim carbon level is obtained, carbon steels are processed similarly to stainless steels. Figure 7 illustrates the carbon content and temperature relationship for the AOD refining of carbon and low-alloy steels. Because the alloy content of these grades of steel is substantially lower than that of stainless steel and because the final carbon levels are generally higher, there is no thermodynamic or practical reason for using an oxygen/inert gas ratio of less than 3:1. Oxidation measurements indicate that all of the oxygen reacts with the bath and that none leaves the vessel unreacted. By monitoring and recording the oxygen consumption during
Steel Melt Processing / 211 Table 1 Processing parameters for argon oxygen decarburization refining of type 304 stainless steel Total gas flow rate: 1.5 m3/min/ton Gas ratios
Partial pressures, atm
Carbon level, %
O2
N2
Ar
Ar
N2
CO
3.0–0.7 1–0.25 0.25–0.12 0.12–0.04 0.04–0.01
4 3 1 1 1
1 1 1 ... ...
... ... ... 3 8
... ... ... 0.83 0.96
0.14 0.20 0.53 ... ...
0.86 0.80 0.47 0.17 0.04
Temperature CRE(a), %
80 65 45 30 15
CRR (b), ppm/min
1200 830 430 133 26
C
1510 1700 1700 1730 1730
F
2750 3100 3100 3150 3150
(a) CRE, carbon removal efficiency. (b) CRR, carbon removal rate
Fig. 7
Schematic of carbon and low-alloy steel argon oxygen decarburization refining
Table 2 Composition control of argon oxygen decarburization-refined high-strength low-alloy steel Element
Aim
Mean
Carbon Manganese Silicon Phosphorus Sulfur Molybdenum Chromium Nickel Aluminum
0.27 0.80 0.50 ... ... 0.28 0.50 ... 0.05
0.264 0.83 0.49 0.019 0.002 0.28 0.56 0.13 0.043
Standard deviation
0.006 0.03 0.04 0.002 0.001 0.01 0.07 0.03 0.009
refining, very close control of end-point carbon is achieved. Because the oxygen and inert gases are introduced below the bath and at sonic velocities, there is excellent bath mixing and intimate slag/metal contact. As a result, the reaction kinetics of all chemical processes that take place within the vessel are greatly improved. Composition and Property Improvement with AOD. It has been well documented that
AOD-refined steels exhibit significantly improved ductility and toughness, along with impact energy increases of over 50% (Ref 9–11). These improved properties result from a decrease in the number and size of inclusions. The capability to produce low-gas-content steel with exceptional microcleanliness, along with alloy savings, is the primary factor for the growth of the AOD process for refining stainless, carbon, and low-alloy steels. Decarburization. In both stainless and lowalloy steels, the dilution of oxygen with inert gas results in increased carbon removal efficiencies without excessive metallic oxidation. In stainless grades, carbon levels of 0.01% are readily obtained. AOD Chemistry Control. The excellent compositional control of AOD-refined steel is indicated in Table 2 for a ten-heat series of highstrength low-alloy steel. The injection of a known quantity of oxygen with a predetermined bath weight enables the steelmaker to obtain very tight chemical specifications. Desulfurization with AOD. Sulfur levels of 0.01% or less are routinely achieved, and levels less than 0.005% can be achieved with single-
slag practice. When extralow sulfur levels are required, a separate slag treatment for 3min is sufficient. Slag Reduction. During oxygen injection for carbon removal, there is some metallic oxidation. Efficient slag reduction with stoichiometric amounts of silicon or aluminum permits overall recoveries of 97 to 100% for most metallic elements. Chromium recovery averages approximately 97.5%, and the recovery of nickel and molybdenum are approximately 100%. Nitrogen Control. Degassing in AOD is achieved by inert gas sparging. Each argon and CO bubble leaving the bath removes a small amount of dissolved nitrogen and hydrogen. Final nitrogen content can be accurately controlled by substituting nitrogen for argon during refining. Nitrogen levels as low as 25 to 30 ppm can be obtained in carbon and low-alloy steels, and 100 to 150 ppm N can be obtained in stainless steels. The ability to obtain aim nitrogen levels substantially reduces the need to use nitrided ferroalloys for alloy specification, and this also minimizes the use of argon. Hydrogen levels as low as 1.5 ppm can be obtained. Oxygen top and bottom blowing (OTB), also referred to as combined blowing, is an extension of AOD technology for refining steel. During OTB, oxygen is injected into the molten steel through a top lance during carbon removal, and inert gases, such as argon, nitrogen, or carbon dioxide, are injected with oxygen through submerged tuyeres or alternate forms of gas injectors such as canned bricks, porous bricks, or thin pipes set in the refractory brick. The top-blown oxygen can react with the bath, reducing refining times, or with CO. The combustion of CO above the surface of the bath increases the thermal efficiency of the refining process and decreases the quantity of silicon and aluminum required for reduction of metallic oxides. If the top-blown oxygen system is designed so that more than 65% of the oxygen reacts with the bath, the system is referred to as hard blown; if less than 65% of the oxygen reacts with the bath, it is referred to as soft blown. As a general guideline, a top-blown oxygen system installed in AOD vessels with a nominal capacity greater than 45 Mg (50 tons) will be hard blown; smaller vessels will be soft blown. In hard-blown top oxygen systems, the flow rate of oxygen through the top lance will be between 50 and 150% of the oxygen injected through tuyeres or injectors. In soft-blown systems, the top oxygen flow rate is between 50 and 100% of the oxygen flow through the tuyeres or injectors. Because foundry AOD vessels have a nominal capacity less than 45 Mg (50 tons), the top-blown oxygen systems are designed to be soft blown. This design maximizes the thermal efficiency of the smaller AOD vessels, reducing the quantity of fuel and reduction material required during refining. Top-blown oxygen flow rates in foundry AOD vessels range from 50% of the bottom oxygen flow rate to 120%.
212 / Molten Metal Processing and Handling Figures 8 and 9 compare OTB with various refining processes. More detailed information on OTB technology can be found in Ref 12 and 13.
Vacuum Oxygen Decarburization (VODC) Vacuum oxygen decarburization can be carried out in a converter vessel (VODC) or in the ladle (VOD). The parameters that favor the choice of ladles as reaction vessels include the following: Tapping, treatment, and teeming are done in
the same reaction vessel. Thus, there are no temperature losses due to any final transfer of the melt, and the high level of cleanliness achieved during the treatment can be preserved up to teeming. The use of electric power as an inexpensive energy source permits the highest flexibility in the melt shop. The ladle unit can act as a time buffer between the melting unit and the casting stand. Information on VOD ladle technology can be found in the section “Ladle Metallurgy” in this article. The primary reasons for selecting the converter as the VOD reaction vessel are as follows: Initial carbon contents are as high as possi-
ble, together with high oxygen blow rates.
Ease of deslagging and strong stirring action Shop restrictions such as a limited number
of ladles, prohibited use of slide gates, and restricted crane capacity in the casting bay make converter technology more attractive. This section reviews both the equipment and the processing parameters for VOD, which is used for the production of low-alloy steels, stainless steels, and superalloys.
Fig. 8
Equipment. Vacuum oxygen decarburization converters are similar to AOD and OTB converters in terms of design and the tilting device used. Bottom blowing is, however, restricted to the introduction of small amounts of inert gas through simple pipes, thus avoiding the special erosion-resistant refractory material used around tuyeres. Flue gas handling is easier and is incorporated into the vacuum system. In terms of vessel design, the conical converter top is closed by a vacuum hood with an oxygen lance feedthrough and vacuum addition lock (Fig. 10). Because the VODC system is closed and no air enters the vessel, permanent control of the decarburization rate and the carbon level in the bath can be maintained and monitored with a flue gas analyzing device (Ref 14, 15). Pollution control for carbon monoxide (CO) and dust is also incorporated into the system. VODC Processing Parameters. Figures 8 and 11 show the relationship among CO partial pressure, gas blow rates, and carbon and chromium oxidation rates for the VODC, VOD, AOD, and OTB processes. As shown, low chromium losses and high carbon removal rates are possible using VODCs. The high carbon removal rates are due to the higher oxygen yield of the VODC and the fact that oxygen dilution is not required. The amount of bottom-blown inert gas is only approximately 2 to 5% of the oxygen volume. With a specific oxygen blow rate of 0.7 m3/ min/Mg, the decarburization rate exceeds 600 ppm/min at high carbon levels (Fig. 8). In smaller VODC heats, higher blow rates of the order of 10 m3/min (350 ft3/min) are maintained, leading to a carbon removal of 1500 ppm/min in 5 Mg (5.5 ton) units. Figures 8 and 11 clearly demonstrate the low chromium losses and the high ratio of carbonto-chromium removal rate due to the very low
Carbon removal rates and chromium oxidation rates for the vacuum oxygen decarburization converter (VODC), oxygen top blowing (OTB), and argon oxygen decarburization (AOD) processes
CO partial pressure achieved in the VODC process. As in any vacuum process, the lowest carbon levels (from 50 to 500 ppm) are reached in stainless steel heats very quickly (6 to 10 times the carbon removal rate of AOD) without applying excessive heat and excessive amounts of inert gas. As shown in Fig. 8, temperatures of approximately 1650 C (3000 F) are common. Inert gas volumes are generally less than approximately 40 L/min/Mg. Figure 9 shows the relationship between the oxygen blow rates and the inert gas addition for various refining techniques, as well as the resulting chromium and carbon levels achieved at specific process steps. Figure 9 also shows the extent to which decarburization is accompanied by chromium oxidation. Carbon and nitrogen levels below 100 ppm are achieved with corresponding chromium yields of over 95% before reduction and 98% after reduction (Ref 14, 17). Sulfur levels are at 10 to 30 ppm for any grade of steel without any sacrifice in the degassing effect. Hydrogen levels are below 2 ppm. The operating conditions necessary for achieving low levels of sulfur and oxygen are the same as those in all strong stirring converter processes (AOD, OTB, VODC). A disadvantage of the VODC, as with all converter processes, is that gas pickup of the melt is encountered during tapping into the teeming ladle. Optimal cleanliness and castability therefore require a subsequent ladle treatment.
Ladle Metallurgy Ladle metallurgy follows furnace melting or refining in a converter. A primary reactor is emptied into a tapping ladle, and then the steel can be transferred into other ladles, furnaces, or vessels for treatment. Alternatively, the steel can be treated in the tapping ladle before being poured into molds. Ladle metallurgy is used for deoxidation, decarburization, and adjustment of chemical composition, including gases. Refining times in the furnace can often be shortened and production rates increased with an efficient secondary steelmaking practice. Ladle metallurgy also allows the control of teeming temperature, especially required for continuous casting. Steels with as low as 0.002 wt% S can be produced free of oxide or sulfide inclusions. Inclusion shape control is also performed in ladles to improve mechanical properties. There are various types of ladle treatment. Table 3 lists the primary methods and equipment used for ladle treatment and their relative capabilities. These methods can be categorized under four major groups (Fig. 12): stirring, injection, vacuum, and heating processes (Ref 18). Accurate assessment of temperature, chemical composition, and quantity of steel in the ladle must be known for efficient secondary processes. The use of high-quality refractories in the liquid steel handling facilities are desirable to achieve lower oxygen activity when
Steel Melt Processing / 213
Total gas flow Technique
m3/min/Mg
ft3/min/ton
Inert gas, %
Electric are furnace (EAF)
1.2
4.6
Lance bubbling equilibrium (LBE)
2.5
95
1
Oxygen top blowing (OTB)
2.0
77
10
Argon oxygen decardurization (AOD)
1.0
40
20--90
Vacuum oxygen decarburization converter (VODC)
0.7
27
2--5
Vacuum oxygen decarburization (VOD-ladle)
0.4
15
0.5--4
Fig. 9
0
Processing steps for achieving proper chromium and carbon levels in stainless steel using various refining techniques with varying gas blow rates and inert gas additions
steel is contained in ladles or tundishes. Ladle preheating and temperature control is also an integral part of ladle metallurgy. Vacuum treatments can be broadly classified into vacuum stream degassing, ladle degassing, and recirculation degassing (Table 3 and Fig. 12). Vacuum processes are designed to reduce the partial pressure of hydrogen, nitrogen, and carbon monoxide gas in the ambient atmosphere to enhance the kinetics of degassing, decarburization, and deoxidation. Alloy additions can also be made during the vacuum treatment. Alloy additions susceptible to oxidation, for example, iron-niobium, are added after deoxidation. The recovery of expensive alloys is improved by ensuring complete dissolution.
Injection treatment in ladles with inert gas, inert gas/oxygen mixtures, or pure oxygen is practiced. Injection of powdered agents of CaSi, CaC2, CaAl, rare earths, and magnesium with an inert gas carrier, or of submerged cored wire with a prior synthetic slag treatment, allows for simultaneous deoxidation, desulfurization, and sulfide shape control.
Metallurgical Objectives The fundamental objective of melt processing is the removal or addition of specific elements to achieve desired steel properties. Residual solid as well as gaseous elements in steel, such as
sulfur, phosphorus, oxygen, nitrogen, and hydrogen, are present due to thermodynamic and kinetic limitations in the primary and secondary steelmaking. Except for specific grades of steel, a lowering of these elements is preferred. Other elements, such as carbon, silicon, and manganese, can be controlled (lowered or raised) to the desired level. Other alloy additions, such as chromium, nickel, titanium, copper, and so on, may be present as residuals from the scrap used in steelmaking or as intentional additions to manufacture special alloyed steels. Degassing. Liquid steel absorbs gases, which can cause embrittlement, voids, inclusions, and so on in the steel when solidified. Proper melting practice and, to a lesser degree, proper heat treatment can limit the gas content of conventionally melted steel to acceptable levels, regardless of the type of furnace equipment employed. Failure to control oxygen, hydrogen, and nitrogen contents may result in porosity or a serious decrease in ductility or both. Gas content is largely adjusted during the oxygen boil. After the cold charge is melted and the bath is in the temperature range of 1510 to 1540 C (2750 to 2800 F), oxygen is introduced into the molten metal, usually by means of a piping arrangement. The oxygen combines with the dissolved carbon in the steel to form bubbles of carbon monoxide. As the bubbles form, dissolved hydrogen and nitrogen are caught up in the bubbles in much the same way that dissolved oxygen finds its way into bubbles of boiling water. Thus, the bubbles of gas contaminants are boiled out. Oxygen is the principal refining agent in steelmaking and plays a role in determining the final composition and properties of steel and influences the consumption of deoxidizers. The production of rimmed, semikilled, or killed grades of steel essentially refers to oxygen management in liquid steel before solidification. If the oxygen content of molten steel is sufficiently high during vacuum degassing, the oxygen will react with some of the carbon to produce carbon monoxide. The evolved carbon monoxide is removed by the created vacuum, known as vacuum-carbon deoxidation. In undeoxidized steel, the carbon and oxygen contents approach equilibrium at a given temperature and pressure according to: C þ O $ CO ðgasÞ
eq K1600 C ¼
p CO ðatmÞ wt%C wt% O
(Eq 3)
(Eq 4)
where Keq is the thermodynamic equilibrium constant, and pCO is the partial pressure of carbon monoxide. Figure 13 shows the contents of carbon and oxygen in equilibrium at 1600 C (2910 F) as a function of partial pressure of carbon monoxide. As the pressure is lowered by vacuum treatments, more and more carbon reacts with oxygen to establish the new carbon monoxide partial pressure, thus making the oxygen unavailable for
214 / Molten Metal Processing and Handling
Fig. 10
Fig. 11
Schematic layout of a vacuum oxygen decarburization converter
Comparison of CO partial pressure for the argon oxygen decarburization (AOD), vacuum oxygen decarburization converter (VODC), and vacuum oxygen decarburization (VOD) processes
inclusion formation with the added deoxidizers. Strong deoxidizers, such as aluminum, titanium, and silicon, react with oxygen with greater affinity than carbon at atmospheric pressure (Fig. 14), so carbon cannot react with oxygen when vacuum degassed. If not floated properly, oxides of deoxidants can result in inclusions when solidified. Carbon is a stronger deoxidizer than aluminum, titanium, or silicon below a pressure of 0.01 atm (10 mbar, or 7.6 torr). Hydrogen causes bleeding ingots, embrittlement, low ductility, thermal flaking, and blowholes. Hydrogen is usually picked up from the moisture in the charge and the environment and can be lowered by an effective vacuum treatment, according to Sievert’s law: ½%H ¼ KðpH2 Þ1=2
(Eq 5)
The amount of hydrogen dissolved in molten steel, [%H], is related to the partial pressure of
hydrogen gas above molten steel, pH2 . The equilibrium constant K is equal to 0.0027 at 1600 C (2910 F). Hydrogen removal during vacuum degassing is affected by factors such as surface area exposed to vacuum, degassing pressure, steel composition, extent of prior deoxidation, and hydrogen pickup from alloy additions, slags, and refractories. The normal hydrogen content in acid-melted steel is in the range of 2 to 4ppm. Hydrogen content resulting from basic practice is slightly higher. Austenitizing and tempering treatments serve to lower the hydrogen content further in sections of average thickness. However, these treatments may not suffice when very heavy sections are involved. The alternative is to degas the molten steel under vacuum, particularly when pouring heavy castings that will be subjected to severe dynamic loading. Nitrogen is particularly harmful for low-carbon steels intended for drawing applications
and should be lowered as much as possible. Although nitrogen follows Sievert’s law, its removal by argon flushing or vacuum degassing is limited due to the tendency of nitrogen to form stable nitrides. Primary control of nitrogen is attempted during steelmaking practices. The lowest nitrogen contents are obtained with either acid or basic open-hearth practice. Nitrogen contents in electric arc melted steels range from 3 to 10 ppm. Deoxidation and Alloying. Steels are required to meet certain temperature and chemical requirements for proper ladle metallurgy. Toward the end of tapping, some amount of furnace slag is carried over into the ladle. Deoxidation and alloying additions can be made during tapping. However, if ultralow carbon steel grades are to be produced, deoxidation is usually not performed at the time of tapping, since dissolved oxygen is later required to react with carbon. The primary role of deoxidation is to prevent pinholes due to carbon monoxide formation during solidification. Oxygen content should be less than 100ppm to prevent porosity. Silicon and manganese are mild deoxidizers; they are added to stop the carbon boil and to adjust the chemistry. Manganese (>0.6%) and silicon (0.3 to 0.8%) additions are limited by other alloy effects and are normally inadequate alone to prevent pinholes. Aluminum is the most common supplemental deoxidizer used to prevent pinholes. As little as 0.01% Al will prevent pinholes. Aluminum is normally added at the tap—approximately 1 kg/ Mg (2 lb/ton) (0.10% added) with a recovery of 30 to 50% for a final content of approximately 0.03 to 0.05%. This is normally supplemented at the pour ladle with additional deoxidation, which could be more aluminum or calcium, barium, silicon, manganese, rare earths, titanium, or zirconium. More than the minimum amount of deoxidizer required for preventing porosity is needed to maintain sulfide inclusion shape control, but excessive amounts can cause intergranular failures or dirty metal. The second addition at the pour ladle is normally 1 to 3 kg/Mg (2 to 7 lb/ ton). The commonly used elements for deoxidation are (in order of decreasing power) zirconium, aluminum, titanium, silicon, carbon, and manganese. The primary role of aluminum as a deoxidant is to react with the dissolved oxygen and float up as alumina. However, some of the finer alumina particles are retained in the steel as inclusions. Making low-inclusion steels requires further treatment to float the fine oxide particles. Slag carry-over is minimized to control the hydrogen and nitrogen pickup by the steel as well as the phosphorus reversal due to reduction of P2O5 by deoxidizing agents. Several types of devices are available as slag stoppers. Desulfurization. Ladle processes allow for desulfurization by judicious selection of an agent. Oxygen is a deterrent in sulfur removal and therefore should be pre-deoxidized to the lowest practicable oxygen content, preferably with aluminum. Simultaneous deoxidation and
Steel Melt Processing / 215
Fig. 12
Processes for secondary steelmaking. Stirring process: (a) bottom injection and (b) lance injection. Injection processes: (c) powder injection and (d) wire feeding. Vacuum processes: (e) stream degassing, (f) Ruhrstahl-Heraeus degassing, and (g) Dortmund-Horder-Huttenunion degassing. Heating processes: (h) Vacuum oxygen decarburization process, (i) vacuum arc degassing process, and (j) ladle furnace process
Table 3
Relative efficiency of secondary steelmaking processes
0, none; 1, good; 2, better; 3, best Metallurgical functions Steelmaking processes(a)
Stream degassing D-H degassing R-H degassing RH-OB degassing Ladle refining furnace, reheat only Ladle refining furnace, reheat and vacuum degassing Argon oxygen decarburization Argon bubbling, CAS Argon bubbling, CAB Lance powder injection Cored-wire injection REM canister
Composition control
Temperature control
Deoxidation (O2)
Degassing (H2)
Decarburization
Desulfurization
Microcleanliness
Inclusion morphology
1 2 2 2 3 3
0 2 2 2 3 3
1 2 2 2 3 3
3 2 2 2 0 2
1 2 2 3 0 2
0 0 0 0 1 3
1 2 2 2 3 3
0 1 1 1 0 1
1 2 2 1 2 1
2 1 1 0 0 0
2 1 2 2 2 1
1 0 0 0 0 0
3 0 0 0 0 0
2 0 2 3 1 1
1 1 2 2 2 1
1 0 2 3 3 3
(a) D-H, Dortmund-Horder-Huttenunion; R-H, Ruhrstahl-Heraeus; RH-OB, Ruhrstahl-Heraeus oxygen blowing; CAS, composition adjustment by sealed argon bubbling; CAB, capped argon bubbling; REM, rare earth metals
desulfurization is also commonly performed. For specialty steels, the ladle furnace is used to desulfurize the melt after vacuum degassing. Efficiency of desulfurization is lowered by the carry-over of oxidizing furnace slags, since the oxides are unstable during the sulfur removal treatment. In addition, FeO lowers the desulfurization efficiency. Similarly, oxides in the ladle lining can affect the efficacy of sulfur removal. Because desulfurization involves slagmetal reactions, intimate mixing of steel and
desulfurizing agents is a prerequisite. In addition to the previous requirements, the initial and final levels of sulfur are critical according to the law of mass action. Thus, steels with less than 0.006% S are difficult to make. Decarburization. It is difficult to produce steel by conventional steelmaking with <0.03% C. Carbon can be lowered to 0.01% from 0.04 to 0.06% levels by simply exposing it to a vacuum, as shown in Fig. 13. Particular attention must be paid to the amount of carbon
in alloying additions made after decarburization of steel. Sulfur and Phosphorus Control. It has been emphasized that the primary removal of sulfur is effected during ironmaking in the blast furnace, where reducing conditions prevail. However, in steelmaking, direct oxidation at the gas-metal interface in the jet impact zone causes 15 to 25% of dissolved sulfur to directly oxidize into the gaseous phase due to the turbulent and oxidizing conditions existing in the impact zone.
216 / Molten Metal Processing and Handling
Fig. 13
Carbon-oxygen equilibrium relationship in liquid steel at 1600 C (2910 F) for various partial pressures of CO gas above liquid steel. If carbon and oxygen levels are at 0.04 and 0.05 wt%%, respectively, at the start of vacuum treatment (region A), then enough oxygen is available to lower the carbon content to 0.01 wt% at pCO of 0.01 atm. However, for carbon levels greater than 0.05 wt% (region B), an external oxygen source is necessary before vacuum treatment to raise the oxygen level to region O, if 0.01 wt% C is desired in the steel at pCO of 0.01 atm. Note that regions A and B are above and below the CO gas stoichiometric line, respectively.
with carrier gas has been described using a reactor model (Ref 20). Ladle refining to produce ultraclean steels and efficient desulfurization are achieved when steel is treated under a basic, nonoxidizing slag. These synthetic slags have low oxygen potential, low melting point, moderate fluidity, and high solubility for alumina and sulfur. If the synthetic slag used in ladle metallurgy has a high sulfide capacity, sulfur removal can be accomplished to ultralow levels of 10 ppm S. If the activity of alumina in the slag is low, then the deoxidation product alumina can be absorbed in the top slag. Thus, calcium-aluminate-type synthetic slags (typically 50 CaO:7 SiO2:43 Al2O3) can deoxidize and reduce total oxygen in the steel as well as desulfurize, simultaneously. Several empirical relationships have been developed to relate the slag composition to the sulfide capacity (Ref 21). Melting and mixing of slag is as important as composition. Premelted slags are commonly used. Vigorous argon stirring of the metal-slag emulsion is necessary for proper deoxidation and desulfurization. Top slag heating is also performed in practice to prevent a thermal gradient in the slag layer. Alternatively, an exothermic mixture of burned lime, hematite, and aluminum powder is used, which helps lower the sulfur content of steel to 0.003 wt% or less, when strongly stirred with argon. The partitioning of phosphorus between the slag and metal is very sensitive to process conditions in steelmaking. An empirical relationship is used to define the phosphorus partitioning parameter, kps (Ref 22, 23): kps ¼ ð%P2 O5 Þ ð1 þ ð%SiO2 ÞÞ=½%P ¼ fðð%FeOÞ; BÞ
Fig. 14 Equilibrium relationship between the contents of total oxygen and various deoxidizing elements in a steel bath Regression equations based on operational data are employed to predict the end-point sulfur within acceptable limits (Ref 19): ð%SÞ=½S ¼ 1:42B 0:13ð%FeOÞ þ 0:89
(Eq 6)
where B is the slag basicity defined as the ratio of the weight percents of basic to acidic slag components. For example, (CaO + MgO)/ (SiO2 + Al2O3), (%FeO), and (%S) are FeO and sulfur concentrations in the slag, respectively, and [S] is the sulfur concentration in the metal. Desulfurization and inclusion morphology control are accomplished through injection metallurgy in secondary steelmaking. In essence, a powdered reagent, such as calcium-silicon, CaC2, magnesium, CaO + CaF2, or CaO + Al2O3, is introduced into the liquid steel with the help of a cored wire or a carrier gas. Desulfurization during powder injection
(Eq 7)
where B is the slag basicity; (%P2O5), (%SiO2), and (%FeO) are the respective oxide contents in the slag; [%P] is the level of phosphorus in the metal; and f is the function of FeO and B. The parameter kps is independent of basicity for B > 2.5. The distribution of phosphorus is also linked to the turn-down carbon level. In general, high basicity and low temperature favor dephosphorization. The addition of fluxes, for example, ilmenite, fluorspar, lime, and soon, to lower slag viscosity, also helps to lower the metal phosphorus content. Phosphorus can be essentially captured in the slag during early stages of steelmaking through judicious adjustment of both the blowing practice and hot metal silicon level (Ref 24). Thus, phosphorus control in secondary steelmaking refers to the control of phosphorus reversal occurring through the use of synthetic slags, which is dependent on the activity of P2O5 in the slag and the amount of aluminum and silicon in the metal. The Perrin process was developed for ladle dephosphorization using appropriate synthetic slags. It is important to do a slag-free tapping if synthetic slag-aided deoxidation with aluminum or silicon is carried out. The kinetics of
deoxidation by aluminum in an argon-stirred vessel or a Ruhrstahl-Heraeus degasser has been reported in literature (Ref 25). Inclusion Shape Control. Nonmetallic inclusions that form during solidification depend on the oxygen and sulfur contents of the casting. Deoxidation decreases the amount of nonmetallic inclusions by eliminating the oxides, and it affects the shape of the sulfide inclusions. High levels of oxygen (>0.0012% in the metal) form FeO, which decreases the solubility of manganese sulfides and causes the manganese sulfides to freeze early in solidification as globules. The use of silicon deoxidation alone normally causes the formation of globular (type I) sulfides. Decreasing the oxygen to between 0.0008 and 0.0012% in the metal, through the use of aluminum, titanium, or zirconium, increases the solubility of the manganese sulfides so that they solidify last as grain-boundary films. These grain-boundary sulfide films (type II inclusions) are detrimental to the ductility and toughness of the steel and cause increased susceptibility to hot tearing. If the oxygen content is very low (<0.0008% in the metal), then the deoxidizer forms complex sulfides that are crystalline and form early in solidification. These type III inclusions are less harmful than type II but are more harmful than type I. The actual sulfide shape depends on the type and amount of deoxidizer. There are complex interactions, so care must be taken to avoid porosity and type II sulfides. If a high level of rare earth metals (RE/S > 1.5) is added, then galaxies of sulfides (type IV inclusions) form; these are only for rare earth additions. The most common inclusion modifier is calcium. Methods of addition include injection (see “Ladle Injection Methods” in this article).
Argon Stirring Clean steel melt processing can be carried out at atmospheric pressure without any supplemental reheating by various methods that include ESR, ladle injection methods, and various argon bubbling processes such as AOD, capped argon bubbling (CAB), and composition adjustment by sealed argon bubbling (CAS). The CAB and CAS methods for argon bubbling were developed to make controlled additions to the ladle as well as to improve the steel refining capability. Stirring treatment with argon gas, also known as argon rinsing, can be done to homogenize the bath and to promote decarburization by lowering the carbon monoxide partial pressure. Bubbles agitate the bath and help in alloy dissolution, deoxidation, quick and uniform mixing of alloys, temperature homogenization, adjustment of chemical composition, and partial removal of nonmetallic inclusions. These functions are accomplished by either blowing argon through a refractory-protected lance lowered to within 300 mm (12 in.) of the ladle bottom or by blowing argon through porous refractory plugs inserted in the bottom or sidewall of the ladle.
Steel Melt Processing / 217 Methods include argon bubbling, stirring, trimming, and rinsing: Argon bubbling often supplements other
secondary steelmaking operations by promoting bulk movement of steel in the ladle for chemical and thermal homogeneity, enhancing the flotation of inclusions, and promoting intimate metal-slag mixing for refining operations, such as desulfurization and deoxidation. Argon stirring is the most vigorous of the bubbling treatments, injecting the highest flow rates of up to 0.3 m3/min (10 scfm) through liquid steel. Stirring, which usually follows tapping, is used to mix the slag and metal and for temperature and chemical homogenization. The high flow rate breaks through the top synthetic slag layer and creates an opening for alloy and deoxidant additions. Sometimes, radiant heat losses are promoted during stirring to prepare the heat for continuous casting. Carry-over of slag from the converter should be minimized. Argon rinsing follows the stirring to help float the inclusions into the slag, and the gas flow rate is below 0.15 m3/min (5 scfm). During argon rinsing, the gentle flow rate of argon prevents the formation of new heatradiating surfaces. If the steel composition requires adjustment through ferroalloy additions, it is carried out during argon trimming, which occurs between the stirring and rinsing steps. Just enough gas flow (0.15 to 0.3 m3/min, or 5 to 10 scfm) during the trimming period allows the ferroalloys to mix in the steel and not become lost in the slag layer. The specified rates in this description are somewhat dependent on the ladle size. Stirring is simple purging or rinsing of liquid steel by inert argon gas passed either through a porous plug or a submerged lance (Fig. 15) to achieve bath homogenization and oxide flotation. Deoxidation and desulfurization by enhancing slag-metal reactions can also be achieved by argon purging if a synthetic slag is used. Argon rinsing also helps in lowering the nitrogen and hydrogen content, as well as enhancing alloy dissolution and slag-metal reactions due to stirring effects. Exchange treatments with synthetic (low-FeO) slags (Ref 26) or conditioned (high-CaO) slags (Ref 27) can be carried out during or before tapping, followed by argon rinsing to promote deoxidation, desulfurization, and inclusion modification. Methods for argon rinsing are shown schematically in Fig. 15. Although a stopper-rod arrangement is shown in Fig. 15 for teeming the ladle, use of slide-gate systems is more prevalent now. The CAB process uses a conventional ladle with a cover and requires a synthetic slag over the steel surface after tapping (Fig. 16a) to act as a sponge for the absorption of nonmetallic
Fig. 15
Schematic illustrating methods for using the argon stirring process. (a) The lance method. (b) The porousplug method
Fig. 16
Schematic arrangement of equipment used in argon bubbling processes. (a) Capped argon bubbling process. (b) Composition adjustment by sealed argon (CAS) bubbling process
inclusions. The introduction of argon into the covered ladle through a porous bottom plug stirs the metal vigorously and creates a slagmetal mix, unlike other argon bubbling methods, which require an intact slag layer to protect the melt from oxidation. The cap on the ladle prevents any air from affecting the metal. The slag-metal emulsion is useful in enhancing microcleanliness, chemical homogenization, desulfurization, and deoxidation. The CAS process (Fig. 16b) uses a refractory-lined snorkel, which is lowered inside the melt during argon stirring so the steel inside the snorkel is slag free. This allows the addition of ferroalloys and deoxidizers without any slag interference and therefore is an effective secondary steelmaking process for achieving compositional control.
Ladle Injection Methods Alloy additions or deoxidation reagents are introduced either in a powder form through a
submerged lance with the help of a carrier gas or by encasing the materials in a wire by injection metallurgy. The primary objectives of introducing powdered reagents into liquid steel are for desulfurization, deoxidation, alloy dissolution, and inclusion shape control. Powders are mechanically mixed or premelted and granulated for injection with an inert gas carrier. Submerged cored wire with a prior synthetic slag treatment is the other injection method. If ultralow carbon steel grades are to be produced, deoxidation is usually not performed at the time of tapping, since dissolved oxygen is later required to react with carbon. A wide variety of powders (CaSi, CaC2, CaAl, rare earths, and magnesium) are used in industry for simultaneous deoxidation, desulfurization, and sulfide shape control. Calcium silicide treatment is most prevalent for simultaneous sulfur and inclusion shape control. Lime, alumina, and fluorspar mixtures are also injected to form the synthetic slag. Alloy additions that are usually made through injection include iron-silicon,
218 / Molten Metal Processing and Handling CaCN2, graphite, NiO, MoO2, lead, tellurium, and ferroalloys of iron-boron, iron-titanium, iron- zirconium, iron-tungsten, and so on. Several commercial systems are available for powder injection through a lance as well as for cored-wire feeding. Ultralow levels of sulfur (<0.005%) can be obtained with this technique. Powder injection and wire injection operations are shown schematically in Fig. 17. The basis of reactive metal injection desulfurization is the chemical combination of the reactive metal with sulfur dissolved in the liquid steel. Because the reactive metals are strong deoxidizers, it is necessary to have a low oxygen level in the liquid steel before the start of the injection process. Otherwise, the reactive metal would combine with the dissolved oxygen and not the dissolved sulfur. Aluminum is used to obtain the necessary low oxygen levels. It is also necessary to have a slag present to hold the sulfur removed by the reactive metal. If such a slag is not present, the metal sulfide will be oxidized, and the sulfur will return to the liquid metal. When a powder is fed through the lance in the steel melt, particles are accelerated by the high velocity of the carrier gas upon entering the gas stream. For example, in gas-borne powdered injection with the help of a submerged lance, 0.6 to 1.0 mm (0.02 to 0.04 in.) calcium-silicon particles are injected at a carrier gas pressure of 5 to 15atm. At the lance tip, the gas expands as a result of a pressure drop and a rise in temperature. In most steelmaking applications, jetting usually occurs during the injection, which is in the form of a gas-jet cone. On leaving the nozzle, the particles travel in a straight line downward due to their high momentum, until the jet breaks up into a swarm of rising gas bubbles. The majority of the powder particles are trapped in the rising gas bubbles. If the powder particles are wetted by liquid steel, they can break away and be recirculated in the melt. A slagmetal-gas emulsion thus forms near the top surface. Wire feeding allows the introduction of additions at the bottom of the ladle and therefore a better alloy recovery. Injection metallurgy is also used for addition of deoxidants, such as aluminum, in the form of wire feeding as well as lance injection. The initial capital investment for
Fig. 17
a powder injection installation is much more than that for a wire feeder, but cored wires are more expensive than ordinary metal or alloy powders. Figures 12(c) and (d) show a schematic of the injection processes. The calcium treatment, primarily carried out for sulfur control, also has a modifying effect on inclusion morphology. A complete modification of type II manganese sulfide inclusions takes place when the sulfur content is <0.07%. In practice, 0.5 to 0.7 kg per metric ton of steel is necessary for inclusion shape control (Ref 26). These reagents could be wire fed or powder injected. However, for wire feeding, recovery of calcium is three to six times greater than powder injection. Rare earth metals (REM) are very efficient in morphology control (stringer type to globular type), although they only have a minor role in desulfurization. Rare earth metal additions are usually accompanied by a magnesium addition to provide the agitation needed for uniformly distributing the REM throughout the heat.
Ladle Heating Processes Steadily increasing requirements for steel cleanliness have led producers to adopt several novel refining technologies and process routes, many of which involve increased use of the ladle as a refining vessel. Such procedures require longer holding times in the ladle, which necessitate either increased superheats in the furnace or external heating in the ladle in order to avoid early solidification. Higher superheat, in addition to requiring excessive energy expenditure, can contribute to the problem of dirty steel. The preferred solution is to supply heat to the ladle, maintaining the steel at minimal superheat during refining. The loss in temperature during secondary treatment can be compensated by an additional reheating process. The ladle furnace is normally used to heat up, to hold, and to finish all kinds of metal melts (see the section “Ladle Furnace and Vacuum Arc Degassing (LF/VD-VAD)” in this article). Depending on metal, ladle size, and expected heating rate, electrical power of up to several megawatts may be applied. Heating can be done by graphite electrodes, plasma torches, or induction coils. However, graphite electrodes
Schematics showing powder injection and wire injection methods used for desulfurization. (a) Powder injection. (b) Wire injection
cannot be used for numerous steel grades, especially those with extremely low carbon contents, because of the risk of carbon pickup. High-power plasma torches have been considered as an attractive nonconsumable heat source for ladle design. Prior to use in ladle metallurgy stations, plasma torches had been used primarily for scrap melting and heat loss compensation before and during continuous casting. Ladle heating with plasma torches is used for horizontal and vertical continuous casting. It also has application for ladle treatment in foundries. In principle, the plasma ladle furnace operates with arcs in exactly the same way as the conventional graphite-heated ladle furnace, but with two major differences: The electrodes produce no carburization. Arc stabilization with argon results in the
active creation of an inert furnace atmosphere. Depending on the application, the systems that can be operated with direct or alternating current power are equipped with either watercooled metal torches or special graphite electrodes. The plasma arcs are stabilized with argon. Where several torches are employed, the floor electrode can normally be eliminated. Torches can be positioned vertically or at an angle. An inclined torch arrangement also permits contact-free heating with nontransmitted arc. Because arcs can also be operated without melt contact, the plasma ladle furnace can be used for slag treatment, because even electrically nonconducting alloy elements can be melted without contact and superheated under total control. It is therefore well suited for desulfurization. A plasma ladle reheater with a graphite return electrode is shown in Fig. 18. It consists of a refractory-lined ladle cover with one or
Fig. 18
Schematic of a plasma ladle reheater
Steel Melt Processing / 219 more plasma arc torches mounted on top. Also mounted on top of the unit is a return electrode mechanism. The return electrode can be a graphite bar that is immersed in the melt from the top, or it can be another plasma arc torch. In an alternate design, the return electrode consists of a conductive layer of carbon-magnesite bricks (Fig. 19). When the cover is placed on the ladle, a copper plate makes contact with the electrically conducting bricks and serves to conduct the plasma current to the grounded cover. Induction furnaces equipped with a gasstabilized graphite electrode have also been investigated to control oxygen and inclusion levels and enhance desulfurization afforded by the transferred arc plasma. A schematic example of plasma torch heating with argon stirring is shown in Fig. 20 from a joint development effort by the Electric Power Research Institute Center for Materials Production and Maynard Steel Casting Company. Molten steel from the arc furnace is refined by the combined action of superheating (which increases the slag-metal reactivity) and electric currents at the slag-metal interface, which assist the electrochemical reactions. In addition, the entire bath of molten metal is stirred by argon passing through a porous plug to improve chemical and temperature homogeneity. The high heat capacity of the arc allows close control of the molten metal temperature and superheating of the slag cover. Ladle refiners with plasma heating can refine commercial grades of cast steel to very low levels of sulfur and oxygen. Oxygen can be reduced on the order of 30 ppm, with reduction of sulfur to 10 ppm. Additional advantages of ladle refining include reduced melting energy use, precise chemistry control, and improved shop logistics.
Vacuum Ladle Degassing Vacuum degassing processes expose liquid steel to a low-pressure environment to decrease gaseous content. The major gases to be eliminated are oxygen, hydrogen, and nitrogen. The effectiveness of these methods is directly linked to
the surface area of molten steel exposed. As described earlier, hydrogen removal is dependent on the partial pressure and diffusion, whereas oxygen removal primarily relies on the chemical reaction between carbon and oxygen and the partial pressure of carbon monoxide gas. These processes also accomplish composition and temperature control, decarburization, microcleanliness, and inclusion morphology control. Vacuum degassing is accomplished primarily by three methods: Stream degassing Recirculation degassing Vacuum degassing in the ladle
Because heat losses during degassing occur through radiation and absorption of heat by ladles or vacuum vessels, and because heat losses vary for different processes, the tapping temperature of the steel must be adjusted for compensation of temperature losses. Ways of conserving heat during vacuum processing are the retention of a slag layer on the molten steel surface, use of radiation shields, vessel preheating, supplemental heating arrangements, and optimization of processing time. The vacuum in most modern degassing systems is created by steam ejectors installed in series to provide multistage pumping capability. A greater number of pumping stages are used for low-carbon steels in the partly deoxidized condition than for deoxidized high-carbon steels when treated at the same rate. A reheating-type vacuum induction melting process uses up to six stages, whereas a recirculating-type Ruhrstahl-Heraeus (R-H) or Dortmund-Horder-Huttenunion (D-H) degasser uses five. A four-stage steam ejector is common for other vacuum degassing systems. Moreover, vacuum degassing processes allow the addition of deoxidizers and alloy additions to the molten steel while under reduced pressure. These additions are optimized for the capability of the system to manage the heat and mixing. Stream degassing methods include ladle-toladle, ladle-to-mold, and tap degassing. Ladle-
to-ladle degassing achieves hydrogen removal as well as oxygen removal by carbon deoxidation. A preheated teeming ladle with either a slide-gate mechanism or a nozzle and stopperrod arrangement is placed inside the tank, and the opening in the tank is covered with an aluminum disk. The tank is evacuated to 0.01 kPa (0.1 mbar, or 0.1 torr), and a superheated, partly deoxidized steel in the tapping ladle is placed on top of the evacuated tank. The stream of molten steel from the tapping ladle melts the aluminum disk cover and flows into the teeming ladle while other deoxidants are added through the addition hopper under vacuum. The pressure stabilizes at 0.11 kPa (1.1 mbar, or 0.8 torr) during degassing. Nitrogen is used to bring the system back to atmospheric pressure after degassing. In the ladle-to-mold degassing, an ingot mold is placed on a base, fitted with a hot top, and preheated using burners. The vacuum tank is placed around the base and makes a vacuumtight seal. A transfer or “pony” ladle, fitted with stopper rod and nozzle assembly, is placed in line with the mold inside the tank. A preheated refractory collar is fitted into an opening in the top of the tank to direct the pouring stream from the pony ladle nozzle into the mold. The tank is evacuated to 0.01 to 0.03 kPa (0.13 to 0.27 mbar, or 0.10 to 0.20 torr), and the tapping ladle containing the superheated steel is placed above the pony ladle for bottom pouring of steel. Once the pony ladle is two-thirds full, the stream of steel is allowed to enter the tank through the collar. The stream breaks up into tiny droplets after entering the vacuum chamber and attains a very large surface area, enhancing the degassing process. Several tapped steel heats can be vacuum cast into large ingots by this method. Nitrogen or air is allowed into the tank to break the vacuum.
Fig. 20 Fig. 19
Schematic of conductive ladle and cover with carbon-magnesite bricks as the counterelectrode
Direct-current plasma ladle refiner with argon flowing through the electrode to form and stabilize the plasma arc
220 / Molten Metal Processing and Handling Tap degassing consists of a special ladle with a flange at the top and an O-ring seal plate for the nozzle opening at the bottom, a ladle cover, and a stoppered tundish having an aluminumdisk seal just below the nozzle. The nozzle is sealed, and the special ladle is covered with the lid. An alloy addition hopper is inserted through the ladle cover. The preheated stoppered tundish is placed on top of the ladle cover, and the entire assembly is positioned such that the tapping stream from the steelmaking furnace can be captured in the tundish. The ladle is evacuated to 0.03 to 0.07 kPa (0.3 to 0.7 mbar, or 0.2 to 0.5 torr), and the stopper rod of the tundish is pulled when the tundish is full with the superheated steel tapped from the furnace. The required deoxidizers are added through the hopper during degassing. After degassing, the ladle is restored to atmospheric pressure and teemed. Recirculating degassing methods use atmospheric pressure to force the liquid steel from a ladle into an evacuated chamber, where it is exposed to low pressure and then flows back into the ladle. The metal can recirculate 40 to 50 times to achieve degassing, where the vacuum is used to recirculate the metal stream as well as to degas. A refractory-lined vacuum vessel with a “snorkel” tube, an electric-resistance graphite heating rod, an alloy addition hopper, a mechanism to lower or raise the vacuum vessel or ladle, and a steam-ejector system comprise the D-H degassing method (Fig. 12g). The vacuum vessel is internally preheated to the desired temperature, and the alloying additions are placed in the hopper. The superheated steel is tapped in an unoxidized or partially oxidized condition in a ladle and placed beneath the vacuum vessel. Either the ladle or the vacuum vessel are moved vertically. The snorkel is immersed in the molten steel, and the vacuum vessel is evacuated, causing the steel to be sucked into the vessel. Once the steel is degassed, the snorkel is raised, causing the steel to partially run back into the ladle. The gases are sucked out of the vessel at this stage. Several cycles of this operation degas the steel to the desired level. Alloy additions are subsequently made before the vacuum vessel is returned to atmospheric pressure. Some snorkels are designed to allow the passing stream of steel to be exposed to argon or oxygen. Two legs or snorkels are used in the R-H process, where one leg is fitted with a gas-inlet pipe (Fig. 12f ). The preheated vessel, similar in construction to the D-H degasser, is lowered in the molten steel ladle containing the superheated steel until 150 mm (6 in.) of both the legs are immersed. The vacuum chamber is evacuated, and an inert gas is introduced through the gas-inlet pipe. The lower density of the gas-liquid mix causes the steel to rise and flow into the vacuum vessel. The metal is deoxidized and returned to the ladle through the other leg under gravity. Several cycles of this process at a pressure of 0.04 to 0.08 kPa (0.4 to 0.8 mbar, or 0.3 to 0.6 torr) degas steel
to the desired level before conventional casting. A variation of the R-H process, known as the Ruhrstahl-Heraeus oxygen blowing (RH-OB) process, is used primarily to remove hydrogen and carbon from steel. The difference between the R-H and the RH-OB processes lies in the gases injected into the system during the process of degassing. The RH-OB process also introduces oxygen, in addition to an inert gas, into the system to enhance decarburization. Inert gas also helps in stirring the melt. Vacuum ladle degassing processes include induction stirring and vacuum oxygen decarburization (VOD). In the induction stirring process, induction coils are used to induce eddy currents in the molten steel to produce a stirring effect. Steel is contained in a ladle with shell sections made of nonmagnetic austenitic stainless steel and lined with high alumina or basic refractory. Superheated, fully deslagged, nondeoxidized steel is tapped at a relatively higher temperature than other vacuum degassing processes. The ladle is placed inside the induction coil positioned inside the vacuum chamber, a heat radiation and splash shield is placed on the vacuum tank, and the chamber is evacuated slowly to 0.01 kPa (0.1 mbar, or 0.1 torr). During degassing, the induction coils are energized to stir the steel. Alloy additions, primarily ferrosilicon, aluminum, and carbon, are added through a hopper toward the end of degassing. Stirring is continued for a few minutes to mix the alloys. After the vacuum tank is returned to atmospheric pressure, a synthetic slag is placed on top of the steel surface to conserve heat and prevent reoxidation. The VOD method was designed for the production of stainless steel. A stainless steel heat is prepared by charging and melting ferrochromium, stainless steel scrap, ferrosilicon, burned lime, and fluorspar and then heating it to the tapping temperature. The stainless steel is tapped in a preheated basic-lined ladle equipped with porous plugs for blowing argon. The ladle is placed inside a vacuum chamber and evacuated to low pressures while oxygen is lanced above the steel bath and argon is bubbled through the bottom plugs in the ladle (Fig. 12h). After the oxygen blowing period, further decarburization and deoxidation can be accomplished by additional vacuum time. Deoxidants could also be added during vacuum if desired. A desulfurizing slag can be added to the heat after the vacuum treatment, enabling the removal of up to 50% of the initial sulfur. Argon blowing through bottom plugs can be replaced with induction stirring.
The ladle, containing argon injection or induction coils, can be placed into an independent vacuum vessel (Ref 28), or, if argon injection/ induction coils are needed to stir the melt (Fig. 21), a vacuum ladle with a directly applied vacuum hood can be used. The vacuum vessel can be either stationary with a vacuum connection (Fig. 22a) or movable on a buggy (Fig. 22b). If additional heating of the melt is needed in the ladle, ladle furnaces or vacuum arc degassing are used (see the section “Ladle Furnace and Vacuum Arc Degassing (LF/VD-VAD)” in this article). Plasma arc torches are also used for heating during ladle refinement. The vessel is equipped with a ladle stool for quick and safe placement of the ladle into the vessel. The vessel also has an inert gas connecting device for the ladle. The ladle has the required freeboard (Fig. 23) and is covered by a ladle hood, which retains metal splashes and heat. There is enough space left below the ladle bottom to receive the molten metal in the event of a breakthrough of the ladle. The vacuum vessel cover is equipped with the necessary control devices for adding slag constituents and alloying and cooling material and for conducting temperature measurement, sampling, and observation to ensure uninterrupted molten metal processing. For optimal metallurgical results, all of these activities must be executed without breaking the vacuum. Devices are also available for the addition of bulky cooling material and for the injection of aluminum, calcium, calcium-silicon, graphite, and ferroalloys. Vacuum degassing is used to melt heats ranging in weight from 15 to 350 Mg (16.5 to 390 tons). Its success depends on treatment
Vacuum Ladle Degassing (VD) Static ladle vacuum degassing technology has developed from argon injection devices and the progress in ladle metallurgy with porous plugs, tuyeres, refractory material, heating technologies, slide gates, and so on. Two basic setups can be used during ladle degassing in vacuum.
Fig. 21
Flow pattern of inductive and inert gas injection stirring in ladles. Left: regular flow (no particle agglomeration). Right: irregular flow (increased particle agglomeration)
Steel Melt Processing / 221 reactions. Thus, chromium is protected against oxidation. This applies, to a limited extent, to manganese and silicon if the oxidation rates are kept sufficiently low. Modern VOD plants consist of the following equipment: Basic-lined ladle with a freeboard of 1 to
Fig. 22
Schematics of (a) stationary and (b) movable vacuum vessels
time, which in turn depends on the previous soaking of the ladle lining and the surface-tovolume ratio of the melt. The mean temperature decrease during vacuum treatment varies from 6 C/min (11 F/min) for 15 Mg (16.5 ton) heats to 0.7 C/min (1 F/min) for 300 Mg (330 ton) heats during the time the ladle lining is in thermal equilibrium. It has not been possible to achieve satisfactory degassing results with the presence of slag. However, this has changed because of: The progress made in the fabrication of
basic linings and in the corresponding lining and heating technology The availability of low-silica active slag mixtures The development of vacuum technology that prohibits foaming of the slag and permits the injection of large quantities of inert gas up to 10 L/min/Mg to optimize stirring Such improvements have led to simultaneous desulfurization rates of 95%, hydrogen removal of over 85%, nitrogen removal of 60%, and aluminum contents below 200 ppm. Figure 24 (Ref 29, 30) shows the removal of sulfur, hydrogen, and nitrogen as a function of gas circulation rate for a 180 Mg (200 ton) VD treatment. Vacuum treatment results for several low-alloy steel heats are given in Table 4. Vacuum Degassing Stirring Procedures. There are two vacuum degassing stirring procedures. In the first, hydrogen, nitrogen, and sulfur are eliminated by strong stirring, that is, with 5 to 10 L/min/Mg. In this case, the oxygen level is decreased only to 30 to 50 ppm total
oxygen, with the major portion of this being the suspended oxidic inclusions. In the second procedure, the total oxygen level is reduced to 10 to 25 ppm by a soft stirring phase in combination with a previously conducted aluminum deoxidation procedure. This requires a very low stirring rate, a strongly basic slag, and sufficient temperature reserve for an extended treatment. The distribution of stirring energy should be uniform. For strong stirring, inductive stirring combined with inert gas stirring is required. Inductive stirring alone is sufficient for gentle stirring if the flow pattern is irregular enough to create turbulences that will produce inclusion agglomeration. The differences in the flow patterns created by inductive and inert gas stirring are shown in Fig. 21. Vacuum oxygen decarburization differs from vacuum degassing, which is conducted under reducing conditions, in that decarburization is an oxidizing process that not only increases dissolved oxygen content in the melt by solid oxygen carriers but also introduces gaseous oxygen through a top-blowing lance that permits higher decarburization rates. The VOD process is primarily used for decarburizing stainless steel melts or other high-chromium-bearing alloys; however, a VOD plant can be used for all types of steel grades and alloys when decarburization, degassing, desulfurization, or chemical heating with oxygen is required (Ref 31–33). Chemical heating with oxygen is also called vacuum oxygen heating. The VOD process is based on the fact that decarburization is strongly favored under reduced pressure with respect to other oxidizing
1.5 m (3 to 5 ft), a porous plug in the bottom for inert gas injection, a slide gate, and refractory-lined ladle hood tightly sealed on the ladle Vacuum vessel with ladle stool, an inner lining, a safety basin for receiving a metal breakthrough, a vacuum connection, and a vessel cover with accessories and watercooled splash protection Top-blowing oxygen lance Probing device for in-process temperature measurement and sampling Movable guide pipe for feeding either aluminum wire or wire filled with such materials as calcium, calcium-silicon, titanium, graphite, and fluxes Observation port with television camera Vacuum pump set (steam ejectors, water ring pump, or mechanical pumps) with a suction capacity such that during oxygen blowing the vacuum is held below 10 kPa (100 mbar, or 75torr), and oxidizing as well as reducing vacuum treatments can be performed below 0.1 to 1 kPa (1 to 10 mbar, or 0.75 to 7.5 torr). The pump must also be adjusted to handle carbon monoxide gas safely and to remain operational under a heavy dust load of the carried gas of up to 0.4 kg/m3 (0.4 oz/ft3).
Oxygen blow rates vary from 5 to 25 m3/min (175 to 900 ft3/min), depending on heat size and initial carbon content. The specific blow rate ranges from 0.2 to 1.4 m3 O2/min/Mg, the highest specific blow rates are achieved with the smallest heat sizes. The inert gas blow rates vary strongly with the process and the desired results; that is, absolute rates range from 0.1 to 1.2 m3/min (3.5 to 45 ft3/min), and specific rates are 2 to 7% of oxygen flow. According to equilibrium data for the Fe-Cr-C-O system, the lower the chromium losses, the lower the partial pressure of the formed CO + CO2 (Fig. 4). However, the main oxygen-blowing period of the VOD process is performed at an intermediate vacuum pressure range of 3 to 8 kPa (30 to 80 mbar, or 20 to 60torr), which is governed by the suction capacity of the installed vacuum pump and the freeboard available in the ladle. The chromium loss at given carbon levels and temperatures depends on the prevailing kinetic conditions at the blowing spot. With an intense inert gas stirring from the ladle bottom, the chromium loss can be reduced to nearly zero by the strong stirring VOD process, as described in Ref 34. In most VOD plants, the chromium yield at the end of blowing is 90 to 95% and is increased to 98 to 99% during the subsequent reducing phase by the combination of low
222 / Molten Metal Processing and Handling vacuum pressure and the addition of reducing agents such as aluminum and silicon. The yield of chromium also depends on the initial and the final carbon contents of the melt. The limited freeboard of VOD ladles and the strongly exothermic effect of the process dictate that the starting carbon content and the oxygen blow rate be limited. A typical sequence of operations during a VOD treatment is shown in Fig. 25. Achievable results reported from several plants that use VOD capability are summarized in Table 5.
Ladle Furnace and Vacuum Arc Degassing (LF/VD-VAD)
Fig. 23
Freeboard requirements in secondary metallurgy
Fig. 24
Degree of refining as a function of the circulation rate in a 185Mg (205 ton) vacuum degassing plant. Source: Ref 29, 30
As noted in the section “Ladle Heating Processes,” secondary metallurgy techniques are energy-consuming, and any extended ladle treatment leads to heat losses. As a result, the goal has always been to perform such operations as vacuum degassing, desulfurization, and alloying with simultaneous heating so that the ladle treatment could be carried out with no evident heat loss. Simultaneous heating and degassing is difficult and previously was possible only by vacuum induction degassing or partially by vacuum arc degassing (VAD); however, the combination of an arc-heating ladle furnace (LF) with a vacuum degassing (VD) station meets operational requirements and produces excellent metallurgical results, as is summarized in this section. For vacuum vessels, there are two principal designs. The first is the stationary vessel, and the second is the track-mounted movable vessel. The stationary vessel is enclosed in a mobile vacuum cover or a swiveling ladle furnace cover. For tandem stationary vessels, each is enclosed in a moving vacuum cover, while both are served by a common swiveling ladle furnace, a common vacuum pump, and common alloying equipment. This design permits brief tap-to-tap times of approximately 30 to 45 min. A schematic of a combined LF/VD plant is shown in Fig. 26. The movable vacuum vessel moves with the vacuum ladle car to several treatment platforms that contain the vacuum cover, arc-heating ladle furnace, plasma torches (in the case of ladle plasma heating units), injection stands, and alloying and wire feeding station. There are single-stand layouts that permit both arc heating and vacuum degassing in addition to other secondary metallurgy treatments, such as alloying, injection, and stirring, at the same treatment station. The movable vacuum vessel is transported to this treatment station, where arc heating and vacuum degassing can be performed by ladle furnace vacuum degassing or vacuum arc degassing. For combined ladle furnace and vacuum degassing, arcing is carried out at the pressure created by a pollution-control device; this underpressure is 0.5 kPa (5 mbar, or 3torr). The protective atmosphere is created under a partially closed vessel lid, leaving only the openings for three electrodes, alloying, sampling, process control, and flue gas suction. The atmosphere composition is
Steel Melt Processing / 223 Table 4
Parameters and results of vacuum ladle degassing of low-alloy steel
EAF, electric arc furnace; see Fig. 23 for explanations of other abbreviations. Parameter
Initial hydrogen, ppm Final total hydrogen, ppm Initial nitrogen, ppm Final nitrogen, ppm Oxygen activity, ppm Final total oxygen, ppm Total aluminum, ppm Dissolved calcium, ppm Initial sulfur, ppm Final sulfur, ppm Temperature loss, C ( F) Alloying additions, kg/Mg (lb/ton) Slag additions, kg/Mg (lb/ton) Total amount of slag, kg/Mg (lb/ton) Argon volume, m3/Mg Duration of treatment, min
EAF/VD, 35–40 Mg (38–44 ton)
LF/VD, 40–42 Mg (44–46 ton)
LF/VD, 75–80 Mg (82–88 ton)
3–14 0.4–1.0 100–150 30–60 5–10 20–25 300–450 ... 100–200 40–50 50–60 (90–110) 2–6 (4–12) 20–25 (40–50) 20–25 (40–50) 0.03–0.05 12–16
5–10 0.6–1.5 120–150 50–100 ... 15–25 200–400 ... 300–400 100–200 30–40 (55–70) 0–0.5 (0–1.0) 0 (0) 20–25 (40–50) 0.02–0.03 15–16
5–6 1.0–1.5 80–120 45–90 ... 20–40 100–300 ... 130–200 20–30 30–40 (55–70) 0.5–1.0 (1.0–2.0) 1 (2) 18 (30) 0.15–0.18 20–30
BOP/VD, 180–190 Mg (200–210 ton)
3–7 0.9–1.6 50–70 25–35 1–2 10–20 100–400 20–30 150–200 5–10 35–50 (65–90) 1–3 (2–6) 7–9 (14–18) 15–22 (30–44) 0.12–0.15 25–30
step at 60 to 90 kPa (600 to 900 mbar, or 450 to 700 torr), followed by a vacuum procedure carried out at below 100 Pa (1 mbar, or 0.7 torr) with the electrodes lifted. Ladle Furnace Design. The basic design and equipment of any arc-heated ladle furnace are similar to those of electric arc melting furnaces, except that no tilting movement is required, and ladle furnaces are frequently in a gantry design (Fig. 27). Electrode regulation is accomplished by either electromechanics or hydraulics. The hydraulic fluid is a water emulsion, water glycole, or oil. Direct-current ladle furnaces are not manufactured, but if very small ladle furnace units are required, a monophase alternating-current design with two electrodes is available. Stirring, by inductive coils or inert gas, is necessary in any ladle furnace design. However, the ladle furnace has requirements that are very different from those of the electric arc melting furnace. Ladle furnace requirements include the following: Controls
governing electrode movement must be more sensitive because of the inert gas bubbling through the ladle bottom. The arc must be immersed in the foaming slag; this requires short arcs and therefore a low voltage, a low voltage- to-current ratio, and a low total inductivity.
Fig. 25
Typical sequence of operations during vacuum oxygen decarburization (VOD) ladle treatment. VCD, vacuum carbon deoxidation; VD, vacuum degassing
controlled by the carbon monoxide formed by the reaction of graphite electrodes with oxidic slag and the inert gas introduced through porous plugs in the ladle bottom or under the ladle furnace roof. For vacuum treatment, the electrodes are lifted above the vessel lid, and the lid opening is closed by a vacuum cover. When vacuum arc degassing is used, complete protection of the electrodes is provided by telescopic water jackets and inflatable top seals that permit simultaneous arcing, vacuum treatment, and inert gas stirring. However, arcing under reduced pressure is restricted to approximately 30 kPa (300 mbar, or 225 torr).
Below this pressure, the gas atmosphere would be ionized, resulting in severe damage to the equipment. It should be noted that 30 kPa (300 mbar, or 225 torr) arcing is feasible only if the arc voltage is kept below 100V and the electrode regulation is reliable. The arc length increases with the arc voltage and with the reduced pressure. Therefore, pressures are not applied below 80 kPa (800 mbar, or 600 torr), because the extreme arc length cannot be fully immersed in the slag. Submerged arcing is one precondition for the high thermal efficiency of this process. Therefore, the VAD process is also divided into two steps: one arcing
The refractory wear of immersed arcing is concentrated at the wall/slag zone. It is proportional to the amount of specific arc power brought into the slag and the intensity of slag cooling during treatment. For ladle furnaces, secondary voltages of 180 to 280 V, combined with a 2 to 3 mO furnace plus transformer reactance, lead to arc voltages below 100 V. The resulting relatively high current density is 10 A/mm2 (6500 A/in.2) in rigid bars and pipes, 7 A/mm2 (4500 A/in.2) in flexible water-cooled cables, and 0.40 A/mm2 (250 A/in.2) in graphite electrodes. The response time resulting from electrode control is defined as the time required to dissolve a short circuit by lifting the electrode approximately 10 mm (3/8 in.). This response time may vary between 100 and 400 ms, according to the electrode control system applied. This value is important because shorter response times result in less carbon pickup and higher thermal yields. The heating rate is also affected by heat transfer, particularly during the first 10 to 30 min of treatment. Heat losses from the melt to the different furnace components are illustrated in Fig. 28, which shows that the heating rate depends on the heat flow from the arc (area 1) to the different parts surrounding the melt (area 2). Even in the case of thermal equilibrium with the lining (area 3), permanent thermal and electrical losses (area 4) are present such that the best thermal yield under a steady state of heat transfer is only 50 to 60%. The thermal balance reached at a near steady state for arc heating (VAD/LF) and induction heating (vacuum induction melting and degassing, or VIM/VID) is shown in Fig. 29. Figure 30 shows
224 / Molten Metal Processing and Handling Table 5 Process parameters and metallurgical results in vacuum oxygen decarburization (VOD) ladle plants EAF, electric arc furnace; see Fig. 23 for explanations of other abbreviations. EAF-AOD-VOD Parameter
Chromium, % Nickel, % Molybdenum, % Initial carbon, % Final carbon, ppm Initial nitrogen, ppm Final nitrogen, ppm Chromium yield, % Titanium yield, % Manganese yield, % Oxygen blow rate, m3/min (ft3/min) Specific oxygen rate, m3/min/Mg Inert gas rate, m3/min (ft3/min) (ft3/min) Ar/O2 ratio, % Ladle freeboard, m (ft) Slag addition, kg/Mg (lb/ton) Alloy addition, kg/Mg (lb/ton)
EAF/VOD, 15 Mg (16.5 ton)
EAF/VOD, 27 Mg (30 ton)
EAF/VOD, 30 Mg (33 ton)
LF/VOD, 40 Mg (44 ton)
18 10 ... 1.1–1.9 200–600 ... ... ... ... ... 5–8 (175–280) 0.53
18 ... 2 0.6–1.1 80–180 300–430 40–80 97.6 40–60 60–80 5–14 (175–495) 0.52
18 11 2 0.8–1.0 300–400 ... 140–240 96.6 74.5 85.0 5–10 (175–355) 0.34
12 ... ... 0.5–0.6 50–100 150–200 50–85 96.3 ... ... 5–15 (175–530) 0.37
18 10 ... 0.6–0.8 10–20 400–800 50–70 99.2 ... 70.0 10–15 (355–530) 0.23
17 ... ... 0.7–0.9 10–30 300–600 50–80 99.0 ... 80.0 10–16 (355–565) 0.25
0.03
0.05–0.25
0.05–0.10
0.1–0.2
0.3–0.6
0.4–0.7
(1.0) 0.5 0.8–1.0 (2.5–3.0) 20–25 (40–50) 12–20 (24–40)
(1.75–9.0) 0.3–5.0 1.0–1.2 (3.0–4.0) 10–12 (20–24) 10–15 (20–30)
(1.75–3.5) 0.5–2.0 1.2–1.5 (4.0–5.0) 30–40 (60–80) 15–25 (30–50)
(3.5–7.0) 0.7–4.0 1.5 (5.0) 10 (20) 5–8 (10–16)
(11–21) 2.0–6.0 1.6 (5.2) 25 (50) 10–12 (20–24)
(14–25) 2.5–7.0 1.6 (5.2) 25 (50) 7–10 (14–20)
65 Mg (72 ton)
65 Mg (72 ton)
Equipment and Processing. Vacuum induction degassing units consist of a simple induction furnace, designed for vacuum operation, that is equipped with a bell-type vacuum cover (Fig. 31). The furnace crucible is either rammed or lined with brick and has a porous plug in the bottom. A frequency converter power supply provides the energy for the induction coil to control the temperature of the melt. With the vacuum cover removed, the furnace is charged with either liquid or solid. For vacuum treatment, the vacuum cover is placed on the furnace body by crane. A mechanical pumping system is used to achieve an operating pressure of approximately 0.01 to 0.1 kPa (0.1 to 1 mbar, or 0.075 to 0.75 torr). The process is primarily applied to tool and die steels, bearing steels, chromium steels, nickel-chromium steels, and nickel-base alloys to remove dissolved gases such as hydrogen and nitrogen and to remove oxide inclusions from the melt. Figure 32 shows the results of VID treatment on the hydrogen, nitrogen, and oxygen contents of a die steel. In Fig. 32(a), the use of VID lowered the hydrogen content from more than 2.2 ppm to 0.60 ppm (hydrogen removal of 50%). With proper metallurgical treatment, the nitrogen content can also be decreased. Figure 32(b) shows that the nitrogen content was lowered from more than 370 ppm to approximately 50 ppm (nitrogen removal of 50%). With VID, oxide inclusions can be removed by applying metallurgically active slags as an absorbing buffer. In addition, VID units are equipped with argon bubbling systems, which accelerate and improve the removal of nonmetallic inclusions. Changes in oxygen content as a result of VID, with and without argon bubbling, are shown in Fig. 32(c). The final total oxygen content when using VID with argon bubbling is 30% lower than VID without argon bubbling, which demonstrates the benefits of soft argon rinsing.
ACKNOWLEDGMENTS Contents adapted from:
Fig. 26
Schematic of a combined ladle furnace (LF) and vacuum degassing (VD) plant. VOD, vacuum oxygen decarburization
a typical treatment cycle and associated parameters for a 75Mg (80 ton) VAD-heated melt. Independent of thermal yield, the physical boundary for heat transfer by arc, plasma, electron beams, and oxygen gas streams is that the melt evaporates at the heat-transfer spots. Table 6 (Ref 28, 35, 36). compares the various types of heating processes described in this section, in addition to vacuum induction degassing, plasma heating, and vacuum oxygen heating, which are described elsewhere in this article.
Vacuum Induction Degassing (VID) Thermal balance is extremely critical for the secondary metallurgical treatment of small
heats. In most cases, excessive chemical heating through aluminum additions is the only method of compensating for temperature losses during the secondary metallurgical treatment of melts below 15 Mg (16.5 tons). This technique, however, has various disadvantages, such as increased aluminum oxide (Al2O3) inclusion content in the final product and the high equipment costs associated with aluminum heating and oxygen (O2) blowing. In such cases, VID, a new secondary steelmaking technology, also permits economical vacuum treatment for small heat sizes ranging from 1 to 15 Mg (1.1 to 16.5 tons). Table 7 compares the metallurgical applications of VID with other heating and melting operations.
W. Burgmann, Vacuum Ladle Degassing
(VD), Leybold AG, West Germany
G. Kienel, Degassing Processes (Converter
Metallurgy), Casting, Vol 15, ASM Handbook, ASM International, 1988, p 426–431 B. Mishra, Steelmaking Practices and Their Influence on Properties, Metals Handbook Desk Edition, 2nd ed., ASM International, 1998
REFERENCES 1. J.P. Bennett, K.-S. Kwong, G. Oprea, M. Rigaud, and S.M. Winder, Performance of Refractories in Severe Environments, Corrosion: Materials, Vol 13B, ASM Handbook, ASM International, 2005
Steel Melt Processing / 225
Fig. 27
Schematic of a vacuum arc degassing plant layout. 1, transformer; 2, bus bars; 3, flexible cables; 4, electrode arms; 5, temperature-measuring and sampling device; 6, alloying lock; 7, electrode housing; 8, electrode masts; 9, operating platform; 10, lifting device for platform; 11, vacuum cover; 12, ladle heat shield; 13, electrodes; 14, slag; 15, melt; 16, argon plug; 17, vacuum vessel; 18, vessel drive; 19, emergency pit
4.
5.
6.
7.
Fig. 28
Thermal flow pattern of a ladle furnace. 1, arc; 2, melt; 3, slag, electrode tip, and lining; 4, vessel, cover, busbars, and environment
2. V.D. Eisenhu¨ttenleute, Ed., Slag Atlas, 2nd ed., Verlag Stahleisen, GmbH, D-Du¨sseldorf, Germany, 1995 3. M. Blair and S. Thomas, Steel Casting Handbook, 6th ed., Steel Founders Society
8.
9.
of America and ASM International, 1995, p 14–2 F.D. Richardson and W.E. Dennis, Effect of Chromium on the Thermodynamic Activity of Carbon in Liquid Iron-Chromium-Carbon Metals, J. Iron Steel Inst., Nov 1953, p 257–263 R.J. Choulet, F.S. Death, and R.N. Dokken, Argon-Oxygen Refining of Stainless Steel, Can. Metall. Q., Vol 10 (No. 2), 1971, p 129–136 R.B. Aucott, D.W. Gray, and C.G. Holland, The Theory and Practice of the ArgonOxygen Decarburizing Process, J. W. Scot. Iron Steel Inst., Vol 79 (No. 5), 1971–1972, p 98–127 D.A. Whitworth, F.D. Jackson, and F.F. Patrick, Fused Basic Refractories in the Argon-Oxygen Decarburization Process, Ceram. Bull., Vol 53 (No. 11), 1974 D. Brosnan and R.J. Marr, “The Use of Direct Burned Dolomite Brick in the AOD Vessel,” Paper presented at the Electric Furnace Conference (Detroit, MI), Iron and Steel Society, Dec 1977 P.A. Tichauer and L.J. Venne, AOD, A New Process for Steel Foundries, Promises Better Properties, Less Repair, 33 Magazine, April 1977, p 35–38
10. D. Heckel, J.P. Wiencek, and N. Netoskie, “Metallurgical Aspects of the AOD,” Paper presented at the Electric Furnace Conference (Detroit, MI), Iron and Steel Society, Dec 1977 11. F.J. Andreini and S.K. Mehlman, “Quality Aspects of Foundry AOD,” Paper presented at the Electric Furnace Conference (Detroit, MI), Iron and Steel Society, Dec 1979 12. S.K. Mehlman, Ed., Mixed Gas Blowing, Proceedings of the Fourth Process Technology Conference, Iron and Steel Society of AIME, American Institute of Mechanical Engineers, 1984 13. L.G. Kuhn, Ed., Mixed Gas Blowing—A New Era of Pneumatic Steelmaking, Iron Steelmaker, Aug 1983 14. R. Heinke and S. Ko¨hle, Computer Control of the VOD Process, Proceedings of the Seventh International Conference on Vacuum Metallurgy (Tokyo), Section 34–2, Iron and Steel Institute of Japan, 1982, p 1356–1363 15. L. Tolnay, I. Sziklavari, and G. Karoly, Method of Regulating the Endpoint of the Vacuum Decarburization Process Developed for the Production of Acid Resistant Steel Grades with Extra Low Carbon Content in ASEA-SKF Ladle Metallurgy Plant, Proceedings of the Eighth International Vacuum Metallurgy Conference (Linz), M1.3, Eisenhu¨tte Osterreich (Vienna), 1985, p 606–615 16. P. Tennilae, M. Kaivola, and M. Walter, Systems Technology and Operating Results of VODC Converter, Proceedings of the First Electr. Steel Congress (Aachen), F4, Verein Deutscher Eisenhu¨tterleute (VDEh, Du¨sseldorf), 1983, p 1–12 17. W. Burgmann, O. Wiessner, and G. Reese, Vacuum Converter Technology for the Production of Superalloys, Stainless and Low Alloy Steels, Proceedings of the Vacuum Metallurgy Conference on Special Metals and Melting Process (Pittsburgh, PA), Iron and Steel Society, 1985, p 3–10 18. B. Mishra, Steelmaking Practices and Their Influence on Properties, Metals Handbook Desk Edition, 2nd ed., ASM International, 1998 19. B. Deo, Models for Predicting Sulfur and Phosphorus Distribution Ratios in Oxygen Steelmaking, Trans. IIM, Vol 41, 1988, p 475–479 20. L.E.K. Holappa, Ladle Injection Metallurgy, Int. Mater. Rev., Vol 27, 1982, p 53–76 21. E.T. Turkdogan, Physicochemical Aspects of Reactions in Ironmaking and Steelmaking Processes, Trans. ISIJ, Vol 24, 1984, p 591–611 22. F.D. Richardson, Physical Chemistry of Melts in Metallurgy, Vol 2, Academic Press, 1974 23. A. Bergman, The Application of Optical Basicity to Dephosphorization Equilibria, Steel Res., Vol 61, 1990, p 347–352
226 / Molten Metal Processing and Handling
Fig. 29
Thermal energy balances for arc-heated and induction-heated processes. VAD/LF, vacuum arc degassing/ ladle furnace; VID/VIM, vacuum induction degassing/vacuum induction melting
24. H. Ono et al., Removal of Phosphorus from LD Converter Slag by Floating Separation of Dicalcium Silicate During Solidification, Trans. ISIJ, Vol 21, 1981, p 135–144
25. K.H. Karl, H. Abratis, H. Maas, and M. Wahlster, Reaktionen Sauerstoffasffiner Begleitelemente des Eisens mit Eisenoxidhaltigen Basischen Schacken, Arch. Eisenhuttenwes., Vol 45, 1974, p 9–16
26. K.W. Lange, Thermodynamic and Kinetic Aspects of Secondary Steelmaking Processes, Int. Mater. Rev., Vol 33, 1988, p 53–89 27. T.H. Bieniosek, B.S. Kupchik, and K.C. Ahlborg, Ladle Slag Deoxidation with Aluminum-Limestone Mixtures, Steelmaking Conf. Proc., ISS-AIME, Vol 72, 1980, p 409–411 28. R. Schumann and W. Schwarz, Ladle Metallurgy Comparison of Vacuum Ladles and Vacuum Vessels, Metall. Plant Technol., Vol 4, 1985, p 10–18 29. K.H. Bauer and H. Wagner, Operational Experience with a Vacuum Ladle Degassing in a EOF Shop, Stahl Eisen, Vol 107, 1987, p 426–430 30. T. Hsiao, T. Lehner, and B. Kjellberg, Scand. J. Metall., Vol 9 (No. 3), 1980, p 10 31. H.G. Bauer, K. Behrens, and M. Walter, Review on VOD/VAD-Treatments, Proceedings of the Seventh International Conference on Vacuum Metallurgy (Tokyo), Section 33–1, The Iron and Steel Institute of Japan, 1982, p 1314–1331 32. W. Burgmann, J.-S. Yie, and M. Kaivola, Operational VOD Results in Ladles and Converters, Proceedings of the Seventh International Conference on Vacuum Metallurgy (Tokyo), Section 34-1, The Iron and Steel Institute of Japan, 1982, p 1348–1355 33. G. Pietschmann, G. Scharf, W. Burgmann, and M. Velikonja, The Freital Ladle Metallurgy Plant: Equipment and Results, Proceedings of the Eighth International Conference on Vacuum Metallurgy (Linz/ ¨ sterreich, 1985, Austria), Eisenhu¨tte O p 582–605 34. H. Kaito, T. Othani, and S. Iwaoka, Production of Super Ferritic Stainless Steels by S.S.-VOD Process, Tetsu-to-Hagane´ (J. Iron Steel Inst. Jpn.), 1977, p 139–156 35. R. Vasse, H. Gaye, J.P. Motte, and Y. Fautrelle, Thermal and Metallurgical Efficiencies in Ladle Induction Furnace Treatments, Proceedings of the Seventh International Vacuum Metallurgy Conference (Tokyo), Section 25–2, The Iron and Steel Institute of Japan, 1982, p 1125–1132 36. C. Lechevalier, J.P. Motte, and G. Forestier, Proce´de´ de Chauffage par Induction de L’acier Liquide en Poche, Rev. Me´t., Vol 10, 1980, p 791–797
Steel Melt Processing / 227
Fig. 30
Treatment cycle of a 75Mg (80 ton) vacuum arc degassing-heated melt. 1, tapping; 2, additions of C + Al + 0.5% CaO; 3, temperature check; 4, sampling and temperature check; 5, addition of 1.5% alloys + 0.5% CaO; 6, temperature check; 7, additions + 0.5% CaO; 8, degassing; 9, venting
Table 6
Data comparison for various ladle heating operations
VID, vacuum induction degassing; Vott, vacuum oxygen heating; see Fig. 23 for explanations of other abbreviations. Parameter
Method of heating Method of stirring Working pressure, mbar (torr)
VAD
LF
Arc Gas injection 500–800 (375–600) 10–20 (7.5–15)
Arc Gas injection 1020 (770)
ASEA/SKF(a)
VID
Arc Induction Induction plus gas Induction plus gas 1020 (770) 0.1–1 (7.5 102 to 7.5 101) O2 partial pressure, mbar (torr) 20–50 (15–38) 20–50 (15–38) 0.02–0.2 (1.5 102 to 1.5 101) Heating rate, K/min 3–5 3–5 2–4 5–15 Melt weight, Mg (ton) 30–150 15–300 15–250 0.5–15 (0.55–16.5) (33–165) (16.5–330) (16.5–275) Installed power, MVA 5–20 5–40 5–25 0.5–3 Lining thickness Bricked, mm (in.) 200–350 (8–14) 200–300 150–250 (6–10) 80–100 (3–4) (8–12) Rammed, mm (in.) ... ... ... ...
(a) Source: Ref 28. (b) Source: Ref 35, 36
VID, Calidus/IRSID(b)
Induction Induction 1–10 (7.5 101 to 7.5)
Plasma
Plasma Gas 1020 (770)
VOH
Oxygen + Al Gas 1–100 (7.5 101 to 75) 0.2–20 (1.5 101 to 15) 20–50 5–300 (5.5–330)
0.2–2 (1.5 101 to 1.5) 10–20 (7.5–15) 2–7 0.2–1 0.3–25 (0.33–27.5) 50–300 (55–330) 0.3–5 1–6 ... ... 200–300 (8–12) 200–300 (8–12) 100–150 (4–6)
...
...
228 / Molten Metal Processing and Handling Table 7 Comparison of applications and operational characteristics of various heating and melting processes Process(a) Benefit/application
VD
LF/VD
VOD
Melt improvement Hydrogen removal Nitrogen removal Deoxidation Vacuum carbon deoxidation Desulfurization Inclusion removal Extralow carbon Temperature adjustment
VODC
P P P P P P ... ...
Savings Basic materials Alloy material Operational capabilities Relief of primary melting system Increased production capacity Buffer function Collection of molten metal Split melts
VID
P P P P P P ... P
S P S P P S P ...
S P S P P S P ...
P P P P P P P P
S P
S P
P P
P P
S P
P P ... ... ...
P P P P P
P P ... ... ...
P P ... ... ...
P P P P P
(a) P, primary benefit or application; S, secondary benefit or application. VD, vacuum degassing; LF/VD, ladle furnace vacuum degassing; VOD, vacuum oxygen decarburization (ladle metallurgy); VODC, vacuum oxygen decarburization (converter metallurgy); VID, vacuum induction degassing
Fig. 31
Schematic of a vacuum induction degassing unit. (a) Top view. (b) Side view. 1, mold, die, or ladle; 2, charging device; 3, filter; 4, vacuum pumping system; 5, melt current supply; 6, vessel with vacuum bell; 7, gas-purging set; 8, cooling water manifold; 9, control cabinet; 10, hydraulic unit; 11, lid station; 12, transformer
Steel Melt Processing / 229
Fig. 32
Effect of vacuum induction degassing (VID) processing on the (a) hydrogen, (b) nitrogen, and (c) total oxygen contents of X 38 CrMoV 51 die steel (Fe-0.38C-1.0Si- 0.40Mn5.2Cr-1.3Mo-0.40V)
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 230-239 DOI: 10.1361/asmhba0005300
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Aluminum Fluxes and Fluxing Practice Rafael Gallo, Foseco Metallurgical David Neff, Pyrotek
IT IS WELL KNOWN that molten aluminum casting alloys have two inherent characteristics: the tendency to absorb hydrogen gas, and the ability to readily oxidize. Hydrogen is made available at the surface of molten aluminum alloys through the reaction of the molten bath with the water vapor present in the melting environment (Ref 1–4). The reaction between water vapor and molten aluminum yields not only hydrogen gas but also the formation (within milliseconds) of an amorphous aluminum oxide (Al2O3) film on the surface bath, which acts as a protective oxidation barrier for the molten metal underneath the film. The amorphous films have been referred to as young films (Ref 5). The two oxidation reactions that occur when molten aluminum contacts air proceed according to the following: 2Al þ 3H2 O ! Al2 O3 þ 3H2 4Al þ 3O2 ! 2Al2 O3
The hydrogen produced through the reduction of water in turn dissociates into its atomic form at the melt surface and then diffuses through the amorphous aluminum oxide film, from which it is quickly dissolved by the molten bath. The aluminum oxide films are an intrinsic part of the melting process; they protect the metal underneath the film from further oxidation. However, in actual foundry operations, the surface of the molten bath always has some movement due to one or more of the following melting practices:
Charging Skimming Cleaning Degassing Transferring Ladling
Any of these melting practices causes the thin aluminum oxide films to break and reoxidize, causing rapid aluminum oxide film thickening (oxide buildup). The constant metal movement and breaking of the aluminum oxide films cause the films to crumble, thicken, and encapsulate unoxidized molten aluminum, generating what is known as wet dross. A typical representation of this
phenomenon is well observed during the filling process of transfer ladles, such as the one shown in Fig. 1. As soon as the molten metal is tapped from the melting furnace into the transfer ladle, not only does the molten stream become disturbed (from the very beginning to the end), but also a significant amount of splashing between the molten stream and the rising metal surface promotes aluminum oxidation, causing the surface of the bath to thicken with wet dross. The aluminum content of wet drosses is typically reported to be on the order of 60 to 85%, while the remaining 40 to 15% is aluminum. The amount of trapped liquid metal in the dross varies according to the melting and/or handling molten practice. The aluminum oxide is a very stable compound that cannot be reduced to aluminum under ordinary melting conditions. However, the amount of suspended liquid metal could be reduced from the 60 to 85% range to 30% by proper fluxing and drossing techniques. Formation of dross is an intrinsic process when melting aluminum alloys. The dross is considered to be the main contributor in influencing the total metal loss during melting. Depending on the efficiency of the melting furnace, and melting practices, the amount of dross generated may be from 1 to 10% of the total metal melted. In addition to the formation and elimination of dross, another problem that aluminum foundries face in the molten metal bath is the nonmetallic and the metallic impurities that are suspended and floating in the bath. Impurities and aluminum oxides will remain suspended and floating in the liquid bath because they are porous and contain some gas adhering to and trapped in the pores. In this case, the oxide density is similar to that of the aluminum alloy. Nonmetallic and metallic impurities are introduced into the melt during the charging process, the molten metal treatment, and the handling operations. Even when melting primary ingot, nonmetallic impurities such as hydrogen and aluminum oxide are introduced. However, as would be expected, the majority of the nonmetallic and metallic impurities come from charging returns. The term returns means scrap castings, gates, risers, trimmings, and so on. According to the type and origin of the return
Fig. 1
Metal being tapped from a holding furnace into a transfer ladle, causing aluminum oxide films to crumble, thicken, and encapsulate unoxidized molten aluminum, generating wet dross
metal and the conditions of raw material storage, the metal can contain considerable amounts of both nonmetallic and metallic impurities. Hydrogen absorption, formation of dross, generation of bifilms, metallic and nonmetallic inclusions, and oxide buildup are inherent characteristics when melting and handling molten aluminum, regardless of the melting and/or holding furnace design or the energy used (gas or electricity). Therefore, it is of vital importance to understand and control such inherent characteristics, because they will greatly affect the quality of the molten metal, which in turn impacts the final casting quality with respect to porosity, shrinkage, oxides, and inclusions. While the existing types of inclusion material that would be present in melting and holding furnaces and/or transfer ladles will vary from foundry to foundry, their removal is essential for proper molten metal cleanliness. A number of commercially accepted melt treatment techniques are being used by aluminum foundries to remove and separate inclusions from the molten aluminum alloy prior to casting. These include various methods of fluxing, degassing, and filtration in the furnaces and the gating systems (Ref 3, 6–11). Any of these techniques will have an impact on the melt cleanliness of the molten aluminum alloy. However, the effectiveness of evaluating their removal would rely on the melt cleanliness measurement technique being used.
Aluminum Fluxes and Fluxing Practice / 231 The strengths and weaknesses of these methods with regard to equipment requirements, sampling, sensitivity, timing, and practical means of assessing inclusion levels on the foundry floor have been fully discussed and published in the literature (Ref 12–29). Technical articles, which specifically cover fluxing as a factor for ensuring metal cleanliness, have also been published (Ref 7, 30–35). In addition to aluminum, fluxing is also commonly done to some extent with virtually all nonferrous mill and foundry melting operations as the first step in obtaining a clean melt. Examples include: Fluxing of magnesium alloys using salt
fluxes or inert gas as a cover (see the article “Magnesium and Magnesium Alloy Castings” in this Volume). Fluxing of copper alloys to remove gas or prevent its absorption into the melt, to reduce metal loss, and to remove specific impurities and nonmetallic inclusions (see the article “Copper and Copper Alloy Castings” in this Volume) Fluxing of zinc alloys with chloride-containing fluxes that form fluid slag covers, which can be used to minimize melt loss if they are carefully skimmed from the melt before pouring (see the article “Zinc and Zinc Alloy Castings” in this Volume)
particulate (i.e., aluminum oxide inclusions) from the liquid metal. The principal function of gas fluxing is hydrogen removal, not dross treatment or recovery. Gas fluxing does not reduce dross but rather increases dross.
Solid Fluxes Solid fluxes can be broadly categorized as passive or active fluxes. Passive fluxes protect the surface of the molten aluminum from oxidation and prevent hydrogen pickup by the melt. Active fluxes react chemically with the aluminum oxides and clean the melt more effectively. Foundrymen recognize the importance of furnace type; furnace design; furnace maintenance; quality grade of smelter ingot; quality of the inhouse returns; condition of the skimmers, ladles, and the like; and the charging and melting practice on the starting quality of the molten metal to make a casting. Nonetheless, it seems that the foundrymen still have some concerns over the lack of general accepted information on conventional solid fluxes.
Aluminum Fluxing Fluxing is a term commonly used in foundries, especially within the melting department work force, to refer only to the addition of chemical compounds to clean molten aluminum alloy baths in either the furnaces (melting or holding) or the transfer ladles. Fluxing is the first step in obtaining clean molten metal by preventing excessive oxide formation, removing nonmetallic inclusions from the melt, and preventing and/or removing oxide buildup on furnace walls. In general, fluxes may be grouped into two classes: gaseous or solids. Gaseous fluxes may be a blend of an inert and a chemically active gas that is injected into the molten bath. Solid fluxes are blends of salts, which, at the present time (2008), are the most preferred type of fluxes used in foundries, especially since gaseous mixtures with chlorine have been almost completely eliminated because of environmental concerns.
Gas Fluxing Gas fluxing is the application of a gas treatment to the melt. The gas may be inert, such as argon or nitrogen, and the latter is more common because of economics. Gas mixtures may also be used, usually adding an active halogen component—chlorine or fluorine compounds— in smaller amounts, typically only 5 to 10%, to the base inert gas. The purpose of this addition is to provide a better separation of solid
Fig. 2
Equilibrium diagram of cryolite
Even though there are many different commercial fluxes on the market, the basic ingredients of such commercial fluxes are based on the very early work, conducted approximately 100 years ago, to improve the electrolytic production of aluminum by reducing the aluminum oxide using molten cryolite (Na3AlF6). Since then, both natural and synthetic cryolite and fluorides have been evaluated. Figure 2 shows the equilibrium diagram of cryolite. The removal of aluminum oxide by halides has its foundations in basic research and development on the systems NaF-AlF3-Al2O3 and NaF-AlF3. Studies on these systems have shown that the solubility of aluminum oxide increases with the sodium fluoride content. The reaction between cryolite and aluminum oxide in the presence of a surplus of NaF takes place according to the following reactions: 4NaF þ 2Al2 O3 ! 3NaAlO2 þ NaAlF4
or 2NaF þ Al2 O3 ! NaAlO2 þ NaAlOF2
232 / Molten Metal Processing and Handling In addition, halide and fluoride salts have been thoroughly researched by secondary (recycled) aluminum producers during the melting and processing of aluminum scrap and dross. The recovered aluminum results from the effectiveness with which nonmetallic and metallic impurities are removed or reduced to acceptable levels. In these operations, low-melting-point fluxes are typically used. They are basically sodium and potassium chloride with additions of fluorides and some other chlorides, such as calcium chloride and magnesium chloride. Salts are used in fluxes because of the following general characteristics:
Cost-effective Easy to use Combine easily with other ingredients Serve as fillers for active ingredients Have lower density than aluminum Give ability to cover the molten surface Allow a low-melting-point, high-fluidity product Have capability to absorb oxides and reaction products from the fluxing action Solid fluxes are basically blends of sodium chloride and potassium chloride salts, with or without the addition of fluorides. In addition, small quantities of oxidizing compounds, such as carbonates, sulfates, and nitrates, are added to promote exothermic chemical reactions (Ref 7, 29–31). Exothermic reactions are important because they prompt the coalescence of the trapped liquid aluminum particles in the dross. The mechanism of how salt fluxes work has been attributed to thermodynamic chemical reactions and surface tension effects between the aluminum oxide and the flux, the aluminum oxide and the molten metal, and the molten metal and the flux. This interaction between aluminum oxides, flux, and molten metal has been explained in earlier papers (Ref 11, 12, 31, 32). In addition, the effect of the flux on the liquid metal will depend on the chemistry of the flux used, morphology of the flux, total amount of flux added, molten metal temperature, flux contact time, rabbling (stirring) technique, and so on. Nevertheless, from the point of view of flux chemistry, it is important to understand that different combinations and proportions of ingredients will impart different flux properties, such as flux density, flux fluidity, flux wettability, and flux reactivity. These four different flux properties are responsible for the characterization of any particular flux. To better understand how these flux properties are accomplished, it is necessary to recognize that essentially each different ingredient provides different effects that directly influence the final property of a flux. Despite the fact that many ingredients are available, not all of them are necessary or found in each flux. Table 1 lists the characteristics of the materials used in fluxes. Generally, ingredients can be classified into four major groups based on their primary influence over the mixture.
Table 1
Characteristics of materials used in fluxes Melting point
Chemical
Molecular mass, g/mol
Solid density, g/cm3
LiCl NaCl KCl CaCl2 MgCl2 AlCl3 BaCl2 LiF NaF KF CaF2 MgF2 AlF3 Na3AlF6 LiNO3 NaNO3 KNO3 Li2SO4 Na2SO4 K2SO4 CaSO4 MgSO4 Li2CO3 Na2CO3 K2CO3 MgCO3 CaCO3
43.39 58.44 74.56 110.99 95.22 133.34 208.25 25.94 41.99 58.10 78.08 62.31 83.98 209.94 68.94 84.99 101.11 109.94 142.04 174.27 136.14 120.37 73.89 105.99 138.21 84.32 100.09
2.068 2.165 1.984 2.150 2.320 2.440 3.920 2.635 2.558 2.480 3.180 3.180 2.882 2.900 2.380 2.261 2.109 2.221 ... 2.660 2.610 2.660 2.110 2.532 2.420 2.960 2.710
C
605 801 770 782 714 190 963 845 993 858 1423 1261 ... 1010 264 307 339 859 897 1069 1450 ... 723 851 894 ... 1339
Boiling point F
1121 1474 1418 1440 1317 374 1765 1553 1819 1576 2593 2302 ... 1850 507 585 642 1578 1647 1956 2642 ... 1333 1564 1641 ... 2442
C
1325 1413 1500 1600 1412 177.8 1560 1676 1695 1505 2500 2239 1291(a) ... 600(b) 380(b) 400(b) High ... 1689 High 1124(b) 1310 High High 350(b) 850
F
2417 2575 2732 2912 2574 352 2840 3049 3083 2741 4532 4062 2356(a) ... 1112(b) 716(b) 752(b) ... 3072 2055(b) 2390
662(b) 1562
(a) Sublimes. (b) Decomposes. Source: Ref 32
Chlorides. Examples are aluminum chloride (AlCl3), barium chloride (BaCl2), calcium chloride (CaCl2), lithium chloride (LiCl), magnesium chloride (MgCl2), potassium chloride (KCl), and sodium chloride (NaCl). The melting point of these compounds, in their pure state, may range from 190 to 965 C (375 to 1765 F). However, they form low-temperature eutectics in combinations. In addition, it is very typical to have at least three compounds in a given flux, so that none of these compounds would be used as a single compound in a flux recipe. Chlorides are mainly used because of their fluidizing effects and because they are used as fillers. Fluxes based only on chloride salts should not react with molten aluminum, or at least the reaction should be negligible. In addition, these salts provide negligible effects on surface tension as compared to fluorides. Fluorides. Examples are simple fluorides such as aluminum fluoride (AlF3), barium fluoride (BaF2), calcium fluoride (CaF2), lithium fluoride (LiF), magnesium fluoride (MgF2), potassium fluoride (KF), sodium fluoride (NaF), and double-fluoride compounds such as sodium silicofluoride (Na2SiF6) and potassium silicofluoride (K2SiF6). The melting point of these compounds, in their pure state, may range from 845 to 1425 C (1553 to 2600 F). The high melting point of these compounds provides thickening effects on a flux. Fluoride salts have limited solubility for oxides. However, no fluoride salts can dissolve massive aluminum oxides. Fluoride salts act as surfactants, affecting surface tension forces
between flux, liquid metal, and aluminum oxides. As the flux wets the interface between the aluminum oxide particles and the liquid metal, the adhesion force between the liquid aluminum and the oxides decreases, promoting oxide separation and metal coalescence. Fluorides are still the most effective compounds being used in fluxes to improve aluminum recovery from wet dross. Oxidizing Compounds. Examples are: Nitrates such as potassium nitrate (KNO3)
and sodium nitrate (NaNO3) such as calcium carbonate (CaCO3), potassium carbonate (K2CO3), and sodium carbonate (Na2CO3) Sulfates such as potassium sulfate (K2SO4) and sodium sulfate (Na2SO4) Carbonates
The melting point of the nitrates ranges from 307 to 339 C (585 to 642 F); the melting point of the carbonates ranges from 851 to 1339 C (1564 to 2442 F); and the melting point of the sulfates ranges from 859 to 1450 C (1578 to 2642 F). Oxidizing compounds are used to promote exothermic chemical reactions. They react with the smallest molten aluminum particles that are present in the dross, yielding aluminum oxides as well as considerable heat. The purpose of the exothermic reaction is twofold: 1) It allows larger pockets of entrapped aluminum to coalesce and fall back into the molten bath, and 2) it facilitates reactions between aluminum and fluorides. The exothermic reaction continues until all of the fine aluminum particles are burned.
Aluminum Fluxes and Fluxing Practice / 233 Solvents of Aluminum Oxides. Examples are: Borax (Na2B4O7) Potassium borate (K2B2O4) Sodium cryolite (Na3AlF6)
Both the Hall and Heroult patents covered the electrolysis of aluminum oxide in a bath of molten halide salts. Since then, dissolution of aluminum oxides (Al2O3) by cryolite has been documented.
Flux Morphology The manufacturing of fluxes requires compliance with strict general quality requirements for the raw ingredients (commercial purity, grain size uniformity, and less than 1% moisture content) and for the manufacturing process (blending of the ingredients). Control over these variables assures consistent products and consistent performance. For most of the aluminum foundry history, powder fluxes have been used. Powder fluxes have been in commercial use for nearly 60 years. However, the basic formulations and compositions of the fluxes have not changed dramatically over this period. Present-day research and development in flux chemistry is still based on trial-and-error procedures. In the last ten years, research and development for possible new formulations and variations of traditional mixes has more closely evaluated the use of fluorides because of environmental emission concerns. It has been reported that if fluorides were removed from the formulation, it would render a powerless
Fig. 3
Powder flux being thrown inside the furnace
Fig. 4
Granular flux being thrown inside the furnace
and inefficient flux (Ref 7, 33). At the same time, it has been reported that emissions from fluxes containing fluorides could be reduced by at least half only if the morphology of the flux changed from powder to granular. It is still prudent to mention that emissions from powder fluxes are within standard environmental limits. Fluoride-containing salts are being subjected to stricter environmental regulations and constraints. Therefore, fluoride-free fluxes are available. However, it is important to realize that fluoridefree fluxes will never perform as efficiently as fluxes containing fluorides. On the other hand, fluoride-free fluxes are the best option to eliminate internal smoke and odor in the plant, as well as to treat metal in electrical resistance furnaces without damaging the electrical elements. A granular flux may or may not have the same formulation as a powder flux. However, because of the different grain morphology, granular fluxes offer other operational and product advantages over powder fluxes. Granular fluxes are less polluting to the atmosphere; emissions are reduced by at least 50%. Furnace tenders appreciate the fact that granular fluxes greatly reduce smell and smoke, facilitating working closer to the furnaces during drossing operations. Granular fluxes are easier to apply and to spread over the molten surface, since they are fines and dust free. Therefore, application rates can often be reduced. Additional product benefits are the better consistency in chemistry from grain to grain and the fact that there is no segregation during flux transportation or handling. Figures 3 and 4 depict the difference, which is very noticeable, when applying a powder or a granular flux. Figure 3 shows that, as the
Fig. 5
Powder flux grains
furnace tender throws the powder flux, the flux falls down in a big cloud as soon as it leaves the cup. On the other hand, Fig. 4 shows that the granular flux will travel farther before falling down into the melt. That difference in traveling distance facilitates spreading and directing the granular flux inside a furnace. Consistency in chemistry is based on the fact that, since powder fluxes are blends of salts, each grain will represent only one of the different chemical compounds from which the flux is made. On the other hand, and because of the manufacturing process, granular fluxes have 100% grain uniformity in chemistry, since each grain will represent the same chemistry as the flux recipe. Figures 5 and 6 are used to illustrate this fact.
Categories of Fluxes Years of casting experience have demonstrated that proper fluxing is a key element to cost reduction and quality improvement in the liquid melt. Fluxing is the best means of obtaining clean metal by preventing excessive oxide formation, removing nonmetallic inclusions from the aluminum melt, and preventing and removing oxide buildup from furnace walls. Historically, solid fluxes have been classified into four categories, depending on their use and function at the foundry operation. These categories are:
Cover fluxes Drossing fluxes Cleaning fluxes Furnace wall cleaner fluxes
Fig. 6
Granular flux grains
234 / Molten Metal Processing and Handling Cover, drossing, and melt-cleaning fluxes are usually manually spread, shoveled, or thrown onto the melt surface. However, drossing and melt-cleaning fluxes can also be added via rotary flux injection and/or lance flux injection. Furnace wall cleaner fluxes are typically blown (with a gunning device) onto the furnace walls at the melt line. The flux injection process introduces the powder flux in a finely dispersed stream to the bottom of the melt. To ensure that melt-cleaning and drossing fluxes will do their job effectively (regardless if they are either the powder or granular type), it is essential that, after manual addition, the flux be stirred into the melt to achieve as much contact as possible with the molten metal and dross layer. Based on the size of the furnace, the stirring can be accomplished manually by the furnace tender or mechanically by a forklift or mechanized vehicle. Typically, agitation and stirring of the flux is completed between 1 and 3 min in crucibles or transfer ladles and between 5 and 10 min in reverberatory furnaces with up to 41,000 kg (90,000 lb) of liquid capacity. If flux injection is used, there is no need for manual activation, since the mixing and activation would take place during the injection treatment cycle. It is also important to emphasize that the flux reaction efficiency depends on three interrelated factors: molten metal temperature, stirring, and activation time. Improper control of these factors will result in unreacted flux floating on top of the dross, as shown in Fig. 7. Cover fluxes are used, as the name implies, to cover the melt. These fluxes, which are comprised of salt compounds that form a liquid at normal aluminum melting temperatures, provide a molten barrier “blanket” on the surface of the metal to protect it against oxidation and hydrogen gas adsorption. Cover fluxes are mainly a blend of NaCl and KCl salts with small quantities of fluorides. Fluxes having these three compounds are suitable for almost all types of aluminum alloys, excluding hypereutectic aluminum alloys and aluminummagnesium alloys with over 7% Mg. Over the years, a great number of different commercial cover fluxes have been developed
with preferred additions and proportions of NaF, KF, Ca2F, and other compounds, such as CaCl2, for lowering the melting point of the flux, as well as for providing cleaning effects by agglomerating dirt and oxides. A typical base composition is 47.5% NaCl, 47.5% Kcl, and 5% Na3 AlF6, or cryolite. Other composition components are listed in Table 2. Cover fluxes, for smelting application may have slightly lower melting points than those used in foundry applications. Foundries use these types of fluxes during the melting of heavily oxidized foundry returns and machining chips, as well as when metal holding temperatures exceed the 770 to 790 C (1420 to 1450 F) range. In smelting operations, cover fluxes are mixed in the rotary furnace during the melting of fine scrap, turnings, sawings, fines, and so on. Typically, cover fluxes for smelting applications have lower melting points than those used for foundry applications (e.g., 424 versus 665 C, or 795 versus 1229 F). Drossing fluxes are designed to assist in separating the entrapped aluminum from the adherent oxide skin envelope, which surrounds the aluminum in the dross layer. A high percentage of useful metallics can be lost when wet drosses are removed from the furnace; thus, many operations attempt to reduce the metallic content by in-furnace treatment of the drosses with exothermic fluxes. These contain oxidants and fluorides to generate temperature and help promote a separation of the oxides from the metallics. The exothermic reaction does consume aluminum values while generating heat and can be the source of a loss of up to 20% of the contained metallics in the dross, if the flux is applied too liberally (see the article “Dross, Melt Loss, and Fluxing of Light Alloy Melts” in this Volume for more information on drossing fluxes). Drossing fluxes may be based on salt blends of chlorides, simple and/or double fluorides, and oxidizing compounds. Therefore, drossing fluxes are able to react exothermically, generating heat and improving flux wettability. The wetting action of the flux promotes coalescence, which brings the fine aluminum drops tighter to form larger drops that are much easier Table 2 Compound
Fig. 7
Unreacted flux
AlF3 CaCl2 MgCl2 MnCl2 KF NaF NaCl KCl CaF2 Na3AlF6 Na2SiF6 KNO3 C2Cl6 K2CO3 Na2CO3 K2TiF6 KBF4
to recover. The work of drossing fluxes is considered to be due to both the surface tension effects and the dissolution of aluminum oxides. There are wide commercial ranges of different flux compositions. Even though a drossing flux will always be considered as “hot or reactive flux” in furnace-tender terms, it is important to realize that the reactivity in a drossing flux is due to a combination of the oxidizers and the double-fluoride compounds. This is an important concept to understand, because even without double fluorides, the flux may look too reactive due to an excess of oxygen-bearing compounds that may be burning excessive amounts of good metallic aluminum without actually dissolving aluminum oxides in the melt. Some manufacturers may decrease the use of double fluorides, since they are more expensive compared to other compounds. At the same time, it is important to note that a lower-grade compound, which is cheap, may also decrease the effectiveness of the flux. The same concept will be true for any other type of flux; the quality of the chemical compounds will influence the price of the flux. Thus, a good drossing flux must be designed to reduce the rich metallic aluminum content of the dross. As the dross is treated with the drossing flux, it changes from a wet dross appearance (bright, shiny metallic color) to a dry dross appearance (dark, powdery). Proper flux treatment could reduce the amount of metallic aluminum content of the dross to 30%. Figure 8 shows the difference in appearance between untreated dross and dross treated with flux. Cleaning Fluxes. Melt-cleaning fluxes are designed to remove aluminum oxides from the melt. Melt-cleaning fluxes usually have similar chlorides and oxidizing compounds as the ones used in the drossing fluxes, but in different proportions. In addition, the composition of a meltcleaning flux will typically include only simple fluorides, compared to the simple and double fluorides that are present in drossing fluxes. Since a melt-cleaning flux is less reactive than a drossing flux, it will yield less dry dross than a drossing flux.
Typical fluxing compounds employed Fluidizer (F)/thickener (T)
Oxide surfactant
Chemically active
Exothermic
Gas released
Element added
F F F F F F F F T T T ... ... ... ... ... ...
... ... ... ... ... ... ... ... X X X ... ... ... ... ... ...
X ... ... X ... X ... ... ... ... ... X X X X X X
... ... ... ... ... ... ... ... ... ... ... X ... ... ... ... ...
... ... ... ... ... ... ... ... ... ... ... ... Cl2, AlCl3 CO2 CO2 ... ...
... ... ... ... ... Na ... ... ... ... ... ... ... ... ... Ti B
Aluminum Fluxes and Fluxing Practice / 235 Table 3
Typical manual addition rates for powder and granular fluxes
Application for:
Crucibles and ladles Reverberatory
Fig. 8
Untreated (left) and treated (right) dross
The work of metal-cleaning fluxes is considered to be only due to the surface tension effects, as previously stated. It is very important not to confuse a metal-cleaning flux with a furnace wall-cleaning flux. Furnace Wall-Cleaning Flux. Wall-cleaning fluxes are specifically designed for the softening and removal of excessive aluminum oxide buildup that occurs on the walls of melting furnaces, especially along the melt line. This type of flux helps keep crucible and furnace walls free of oxide buildup above and below the melt line. They are usually applied with a lance or flux gun or may simply be broadcast manually onto the walls. Slightly exothermic, these fluxes, under closed-door and high-fire conditions, increase the localized temperature of the oxide buildup and thereby soften or loosen the attached oxides, permitting easier mechanical removal and subsequent skimming. Wall-cleaning fluxes contain the highest amounts of double-fluoride compounds, such as Na2SiF6 and Na3AlF6. The exothermic reactions that occur because of the oxidizing compounds and the double fluorides enhance more penetration of the flux into the oxide buildup, facilitating the removal of the oxide buildup at the furnace wall. Fluoride-free fluxes cannot be used effectively as a wall-cleaning flux. If the oxide buildup has been converted over time to the insidious form of aluminum oxide called corundum, it cannot be removed chemically. However, in its early stages, sufficient loosening of the mildly adherent product can be effected with the combination of fluxing salts and hard work.
Foundry Practices Powder and granular fluxes can be added manually, either by weight percentage of the total metal charged or melted or by weight in pounds based on the molten surface area to cover. Manufacturer recommendations are the starting point. Typical recommendations are shown in Table 3. However, the exact quantity of flux will depend on a variety of operating conditions, such as furnace type, charging and melting practices, initial metal cleanliness,
Addition based on:
Powder flux
Granular flux
Percentage of total metal charged or melted Pounds of flux per square foot of molten surface area
0.25–0.50% 0.35–0.50 lb
0.20% 0.30 lb
molten metal temperature, and final metal cleanliness required. The rotary flux-injection process is typically used in crucible furnaces and transfer ladles, while lance flux injection is typically used for reverberatory furnaces. In both manual and flux injection processes, either powder flux or granular flux may be employed, provided that the composition is suitable for the required temperature/reactivity. In flux injection, the flux material must be flowable in the injection device. Typical recommendations for flux-injection rates are from 0.009 to 0.32 kg/min (0.02 to 0.70 lb/min). The amount of flux added is determined by the injection time, which will also depend on a similar variety of operating conditions mentioned previously. Molten metal cleaning with fluxes varies from foundry to foundry, with some performing a systematic and very careful fluxing and drossing procedure, and others skimming molten metal or cleaning furnaces without any flux at all. The difference in the furnace conditions and the scrap defects related to molten metal quality is very obvious. Although often overlooked, the furnace tender plays a key role in the success or failure of the use of any type of flux.
Flux Injection Foundry fluxing practice involves the use of either a manual flux addition, followed by mixing, stirring, or otherwise creating a reaction, or alternatively using a mechanical flux-injection process. Flux injection is becoming a common melt treatment for improving the cleanliness of an aluminum melt, and in a more efficient manner than manually adding flux, which can be very irregular and subjective (Ref 36). The general process involves the use of a specific fluxing salt that is applied through a delivery gas such as argon or nitrogen through an injection device. While chlorine has often been added to an inert gas in historical practice, today (2008), virtually all flux injection uses just an inert gas as the flux carrier. The purpose of flux injection is threefold in most foundry applications: hydrogen removal, a partial removal of inclusions by flotation, and, with the proper flux composition, treatment of the dross in situ for improved metal recovery. Static or stationary lances submerged into a furnace are used primarily in small melting furnace batch-type applications, such as small reverberatories, crucible furnaces, and holders. For larger furnace vessels, a rotary flux-injection device is used. The flux-injection method addresses the major drawback of conventional fluxing practices, namely the limited contact of flux salt surface with the unwanted impurities in the melt. Flux injection
overcomes this limitation by delivering predetermined amounts of powdered or granulated flux beneath the melt surface. The flux melts into small droplets that expose a large total specific surface area to the melt as they float to the surface, which accelerates flux-induced melt cleaning, that is, partial removal of inclusions. These inclusions then rise to the surface of the melt, where this residue or debris may be skimmed. When halogen-containing salts decompose in the metal, additional wetting occurs between the gas bubbles and the insoluble particulate within the melt. Inclusions are partially removed by attachment to the rising bubble to the melt surface, where they may be skimmed. Hydrogen is removed as well, because the dissolved monatomic hydrogen diffuses into the insoluble carrier gas bubble, then forms a hydrogen molecule. This double-gas bubble species, containing both the carrier gas plus the hydrogen, then continues to rise through the melt and reaches the top surface. Certain solid fluxing compound components themselves may be employed in the flux composition, which will decompose within the melt to generate an insoluble gas that further assists the carrier gas in this task. A typical static flux-injection system includes a dry powder or granulated material feeder that mixes the flux particulate into an inert gas stream, carrying it through a lance immersed in the melt, as shown schematically in Fig. 9. This can be done either into small melting furnaces or into melt transfer/treatment ladles. Flux injection provides many advantages, including delivering the flux into the body of the melt where it can truly react, minimizing the exposure of fluxing salt to the atmosphere and providing for better handling, and controlling the amount of flux to be used. However, there is one major drawback. The area of influence of the flux reaction in a static delivery system is confined to a very small area around the flux-delivery lance itself; very little twophase mixing between the melt and the carrier gas/flux mixture is feasible unless the operator moves the lance around. Consequently, static flux injection is principally suitable only for small furnaces. For larger furnaces, further evolution of the flux injection process has resulted in hybrid equipment that combines the best traits of flux injection and spinning nozzle rotor degassing, shown schematically in Fig. 10. Rotary flux injection is now commonly employed in many foundries and diecasters in furnaces as well as in batch-mode application, such as treatment in a ladle upon tapout from a melting furnace and before delivery to casting units. There are several manufacturers who offer this hybrid equipment; one example is depicted in Fig. 11.
236 / Molten Metal Processing and Handling
Fig. 9
Schematic of the flux-injection process
Fig. 11
Mobile unit for flux injection/rotor dispersion. Courtesy of Pyrotek Inc.
Fig. 12
Fig. 10
Metallic aluminum recovery during ladle flux injection with rotor dispersion. FIP, fluxinjection process
Schematic of flux injection plus rotor dispersion
The combination of flux injection coupled with rotor dispersion offers the following advantages: Fluxing salts are contained and protected from
the atmosphere, thereby minimizing personal exposure and health issues during use. Only measured or controlled amounts of fluxing salts are used, resulting in minimal spillage and waste. By introducing the flux in the submerged position within the molten metal, any emissions are minimized. Rotor dispersion increases the kinetics of mixing and two-phase contact, very frequently resulting in reduced treatment time.
By selecting the proper flux composition
being introduced, the dross resulting from the process is treated in situ, achieving increased recovery of the metallic aluminum contained in the dross. When performed regularly in a melting furnace or transfer ladle, the flux-injection process provides a strong cleaning action on the refractory walls when used wisely. The flux-injection process will provide efficient metal recovery in a batch-treat ladle process, with results analogous to furnace fluxing. Generally, a 455kg (1000lb) ladle may generate 7kg (15lb) of dross during a degas-alone melt treatment, and this dross will average 80 to
85% metallic aluminum content entrapped within a lacy matrix of aluminum oxide skins. Much of this metal is indeed recoverable. By treatment with the appropriate flux plus rotor dispersion, the resultant dross after the process may average just 3 to 4kg (6 to 8lb), and the metallic aluminum content may be reduced to 40% or less. Typical results are demonstrated in Fig. 12. Practical Application. Flux-injection results can be maximized or optimized with good process “housekeeping.” The operator must thoroughly skim the resultant flux-injection process dross from the bath in order to transfer good metal further down the line into casting furnaces. However, this treated dross in itself
Aluminum Fluxes and Fluxing Practice / 237 may still contain recoverable and economicallyviable metal units. Therefore, further opportunity exists to use an auxiliary dross-recovery process with other in-plant equipment. In many instances, it may not be worthwhile or a good use of foundry time to do this internally, but such residues can be sold off to an outside dross processor. It should also be noted that after a fluxing or flux-injection process, it is virtually impossible to remove all flux residues. In the first place, a certain residence time is required for all fluxing salt debris to be fully transported from within the melt to the surface. In a melt-transfer operation, where the treated melt must then be transported and delivered into holding furnaces, temperature loss must be minimized. Therefore, the ladle is usually emptied as quickly as possible to avoid too much temperature loss and the possibility of sludging, in which heavy metal segregation on intermetallic compound formation creates potential hard spots if they ultimately become entrained into a casting. Secondly, it is practically impossible to remove all fluxing debris from a ladle. Consequently, if a subsequent practice of good settling or sedimentation time in the holding/casting furnace is not used, or if final filtration is not employed, flux-based inclusions, such as those shown in Fig. 13, can become embedded into the casting, creating possible surface finish, corrosion, porosity, and mechanical property difficulties. Recent studies (Ref 35) have shown that rotary flux injection is the most effective process to reduce aluminum content in the dross and that manual fluxing is the worst process, leaving 84% metallic aluminum in the dross as compared to 6%. The second-best process to reduce metallic aluminum content in the dross is rotary degassing with a mechanized
Fig. 13
Flux inclusion entrapped in aluminum casting
flux-delivery system, yielding 17% Al in the dross. In addition, these studies showed that if flux were used during the degassing cycle, the aluminum content could be reduced by 77%. Combination degassing/flux-injection practice is gaining favor as a focus not only for improved metal quality but also for improved metal recovery from dross. Assessment and evaluations of fluxes and degassing equipment have shown that the reactivity and molten surface coverage by the powder or granular fluxes is greatly affected by the design of the rotor, operating condition of the equipment, and adherence to operating procedures.
Fluxing Applications Crucible Furnaces. Crucible furnace melting is commonly used at small foundries as well as at medium- and high-volume-production jobbing foundries in which melting versatility is a must. The most popular material for the crucible is silicon carbide. In these foundries, it is typical that the metal skimmings or dross be placed in a small metal bin located by the crucible furnace. Fluxing procedures in these shops are batch type and basically depend on the management philosophy about fluxing. It is very common that after the melting and alloying procedures, some degassing takes place. It is around the degassing stage when the fluxing and skimming operations occur. Typically, there are four variations in the use and application techniques of fluxes: Melting, degassing, and skimming without
any flux. With this procedure, it is common that the skimmings contain very rich metallic aluminum content, sometimes up to 85%.
The dross coming from the dross bin sometimes looks like an aluminum ingot. Melting, degassing, fluxing, and skimming. A very small number of foundries flux after the degassing operation. Proper fluxing can be accomplished after the degassing operation without introducing hydrogen gas. However, it is imperative that the added flux be dry, the skimmers be cleaned and dried, and the stirring of flux and the dross skimming be performed in a gentle manner. Nevertheless, the major drawbacks are that the metal temperature is usually at its lowest operating temperature, so the flux takes more hand stirring to become activated, and the metal treatment process takes longer. With this technique, strict discipline is a must. Melting and degassing while flux is applied on the metal surface, then skimming at the end of the degassing cycle. This is a better method than the first two described previously. If done properly, the operator could take advantage of the mixing capabilities and oxide removal of the rotary degassing equipment. However, with this procedure, it is very important to pay close attention to the degassing technology being used. Although this appears to be the preferred method used in the industry, it is not the most effective for metal cleanliness (Ref 35). Rotary flux injection. In this process, in which the flux is introduced (through the shaft) during the degassing process, the flux is also dispersed at the bottom of the melt by the rotary action. As the flux leaves the shaft, it liquefies, favoring the wetting of the inclusions, which in turn are carried to the melt surface by the degassing bubbles. Transfer Ladles. Transfer of molten metal by a ladle is a typical operation during the handling of molten aluminum in low-pressure permanent mold foundries and in diecasting foundries. At preset time intervals, the ladle is taken to the melting furnace, from which metal is tapped into the ladle. If degassing is part of the process, then the ladle is taken to the degassing station. If degassing is not part of the process, then the ladle is usually moved away from the furnace before performing the skimming operation. In any case, the molten metal cleanliness and drossing operations performed on the metal in the transfer ladle may include any of the four variations previously described for fluxing the molten metal in crucible furnaces. In addition, there are two supplementary variations with respect to degassing and fluxing at the same time. One variation is to add the flux at the bottom of the ladle before tapping metal into it. The other variation is to add the flux by directing it at the molten metal stream during the tapping of the furnace into the ladle. In both cases, not all the flux reacts when making contact with the stream of liquid aluminum. In the
238 / Molten Metal Processing and Handling majority of the cases, the flux tends to be pushed away from the molten stream. When the flux is added to the bottom of the ladle before tapping metal into the ladle, the only flux that reacts is that which becomes trapped below the initial surge of the molten metal cascading on top of it. The remaining flux is unreacted and tends to float and stay on the surface of the bath. Once the ladle is full, the flux must be stirred and rabbled to activate it before the degassing operation. Even in ladles in which an inert gas is introduced at the bottom of the ladle through a porous plug, the problem of unreacted flux exists. Some unreacted flux will float to the surface. The only way to react the flux is to create an excessive gas bubbling turbulence coming from the porous plug. Unfortunately, this would cause more molten metal oxidation, so that the benefits of fluxing are defeated. A technique not often used is to introduce the flux into the molten stream coming from the tapping of the reverberatory furnace. Again, this technique has the same negative effects with respect to activating flux. Only flux that becomes trapped in the molten stream is activated, while the flux falling down inside of the ladle remains unreacted. As mentioned previously, the unreacted flux will float to the surface of the molten bath. In addition, this process of adding flux is not a very safe practice, since splashing molten aluminum may burn the furnace tender more easily. Reverberatory furnaces are usually used as breakdown furnaces or holding furnaces. In some cases, melting and metal ladling are performed from the same furnace. In addition, some furnaces may have a dip-out and a degassing well at the front of the furnace. These wells may be divided by a ceramic baffle or a refractory wall. Basic use of fluxes in reverberatory furnaces is concentrated into two areas: molten metal and furnace sidewalls. Fluxing of the molten metal bath is typically accomplished after manual additions of powder and/or granular fluxes or by using the lance flux-injection process to deliver the flux. In this particular case, the powder flux is injected via a steel pipe submerged in the molten metal. Furnace sidewalls are susceptible to aluminum oxide buildup, especially along the sidewall areas that are exposed to combustion furnace atmospheres during the expected molten metal-level fluctuations inside the furnace. As the oxide buildup continues over time, it eventually will convert to corundum (the hardest phase of Al2O3. When this occurs, two things happen as the oxide builds up: Furnace efficiency and melting capacity are reduced, and it becomes too late for a wall-cleaning flux to be effective. Eventually, the furnace must be shut down for relining. However, having a routine maintenance schedule for cleaning furnace sidewalls can minimize the problem. Operational procedures for skimming and drossing of the molten metal and cleaning of the furnace sidewalls are normally divided into
exclusive sections according to the design of the furnace. It is very typical to have a specific procedure and frequency for the melting chamber, the charging well, the dip-out well, and the degassing well. Furnace sidewall cleaning should be done once a week, but it is normal to do it three times per week. Often, a good and conscientious furnace tender determines the frequency of using the wall-cleaning flux. Even though different foundries have different furnace-cleaning procedures, it is very common that good fluxing practices be based on well-written procedures and that furnace tenders adhere to proper flux addition, proper sequence by which the burners must be turned on and turned off and the furnace doors opened and closed, minimum flux activation time, and thorough scraping of the furnace sidewalls. An important consideration in the selection of fluxes for reverberatory furnaces may be derived from the fact that, as the furnace capacity increases, the furnace tender’s tolerance for discomfort decreases. A furnace tender who is subject to incessant heat, noise, and other working-related discomforts is less likely to be as effective and efficient as possible, especially if he has to deal with dust, smoke, and fumes from the flux. In such cases, a typical flux manufacturer should be able to recommend an environmentally-friendly flux for the application. Holding/Casting Furnaces. Generally speaking, foundries that employ furnaces used for holding or direct production of castings must take care to use flux and flux practices wisely. Fluxing salts will react and degrade refractory materials used in the furnace construction, and, as shown previously, flux debris can become entrained within castings that are poured directly from fluxed furnaces. If metal has been adequately processed previously, there is usually no need to flux a casting furnace, such as a holder from which metal is poured by automatic or manual ladle. The dross
Fig. 14
buildup or skim on such furnaces is considered to be new oxides, and periodic skimming can prevent these new oxides from becoming entrained into the casting. If a casting furnace is fluxed, flux must be used sparingly, and it is crucial to remove the fluxing debris as completely as possible. Other Considerations. Good management philosophy will make the furnace tender’s comfort an important consideration when selecting a flux to use. A furnace tender who is happy with the working conditions usually has the right attitude toward proper flux use. Perhaps, the final consideration that a foundry must take into account when dealing with fluxes and metal cleanliness concerns is the fact that, despite the weaknesses of commercial cleanliness techniques, each technique has a strength that could be used to assess the molten metal quality. Despite operational anomalies, analysis costs, and/or production constraints to use present-day metal cleanliness equipment, aluminum foundries could benefit from their use to set and monitor their melting and fluxing practices (Ref 34, 35). For example, the latest molten metal evaluations have shown, among other findings, that flux injection provides the cleanest metal; the dirtiest molten metal is obtained with no flux in the process. Any other process would produce molten metal qualities between curves 5 and 8, as shown by the pressure filtration results in Fig. 14. In addition, molten metal quality assessment by ultrasound techniques showed that the molten metal in the as-melted condition typically had average particle (inclusions) sizes between 200 and 500 mm (0.008 and 0.020 in., or 0.2 and 0.50 mm). After any rotary treatment, with and without flux, the particle sizes ranged from 30 to 160 mm (0.0012 to 0.006 in., or 0.03 to 0.16 mm). The great majority were in a range of 40 to 90 mm (0.0016 to 0.0035 in., or 0.04 to 0.09 mm) (Ref 35).
Typical pressure filtration curves representing metal cleanliness values as a function of different fluxing treatments. Dirtiest molten metal is given by curve 5 (rotary degassing without any flux usage). Cleanest metal is given by curve 8 (rotary flux injection).
Aluminum Fluxes and Fluxing Practice / 239 REFERENCES 1. R. Fuoco et al., Characterization of Some Types of Oxides Inclusions in Aluminum Alloy Castings, Trans. AFS, Vol 107, 1999 p 287–294 2. P. Crepeau, Molten Aluminum Contamination: Gas, Inclusions and Dross, Aluminum Alloy Castings, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 1–13 3. D. Apelian and S. Shivkumar, Molten Metal Filtration — Past, Present, and Future Trends, AFS Second International Conference on Molten Aluminum Processing, 1989, p 14–1 to 14–36 4. C. E. Eckert, The Origin and Identification of Inclusions in Foundry Alloys, AFS Third International Conference on Molten Aluminum Processing, 1992, p 17–50 5. J. Campbell, Castings, 2nd ed., Butterworth-Heinemann, Oxford, 1988, p 11–26 6. D. Apelian, How Clean Is the Metal You Cast? The Issue of Assessment: A Status Report, AFS Third International Conference on Molten Aluminum Processing, 1992, p 1–15 7. R. Gallo, Development, Evaluation, and Application of Granular and Powder Fluxes in Transfer Ladles, Crucible, and Reverberatory Furnaces, AFS Sixth International Conference on Molten Aluminum Processing, 2001, p 56–70 8. L.C.B. Martins and G.K. Sigworth, Inclusion Removal by Flotation and Stirring, AFS Second International Conference on Molten Aluminum Processing, 1989, p 16.1–16.28 9. D.V. Neff, Continuous, Re-Usable Filtration Systems for Aluminum Foundries, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 121–137 10. J.R. Schmahl, L.S. Aubrey, and L.C.B. Martins, Recent Advancements in Molten Aluminum Filtration Technology, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 71–102 11. L.C.B. Martins, Chemical Factors Affecting the Separation of Inclusion from Molten Aluminum, AFS Third International Conference on Molten Aluminum Processing, 1992, p 79–91 12. P. Crepeau et al., Characterization of Oxide Sludge, Dross and Inclusions in Aluminum Melting and Holding Furnaces, AFS Third
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23.
24.
International Conference on Molten Aluminum Processing, 1992, p 51–77 T.L. Mansfield and C.L. Bradshaw, Ultrasonic Inspection of Molten Aluminum, Trans. AFS, 1985, p 317–322 D. Doutre et al., Aluminium Cleanliness: Methods and Applications in Process Development and Quality Control, Light Met., 1985, p 1179–1195 R.I. Guthrie and J.E. Gruzleski, Quantitative Measurement of Melt Cleanliness in Aluminum Silicon Casting Alloys, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 15–28 A. Simard, F.H. Samuel, et al., Capability Study of the Improved PREFIL-Footprinter to Measure Liquid Aluminum Cleanliness and Comparison with Different Techniques, AFS Sixth International Conference on Molten Aluminum Processing, 2001, p 93–112 J.L. Roberge and M. Richard, Qualiflash Apparatus for Testing the Inclusion Quality of Aluminum Alloy Baths, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 29–42 E.L. Rooy and E.F. Fisher, Control of Aluminum Casting Quality by Vacuum Solidification Tests, Trans. AFS, Vol 76, 1968, p 237–240 D. Dispinar and J. Campbell, Critical Assessment of Reduced Pressure Test, Part 1: Porosity Phenomena, Int. J. Cast Met. Res., Vol 17 (No. 5), 2004, p 280–286 D.E. Groteke, Improve Your Vacuum Test, Mod. Cast. Mag., Aug 2007, p 28–31 N.D.G. Mountford, I.D. Sommerville, et al., Sound Pulses Used for On-Line Visualization of Liquid Metal Quality, Trans. AFS, 1997, p 939–946 N.D.G. Mountford, I.D. Sommerville, et al., Laboratory and Industrial Validation of an Ultrasonic Sensor for Cleanliness Measurements in Liquid Metals, Proceedings of International Conference on Cast Shop Technology, TMS-AIME, 2000, p 304–310 R. Gallo et al., Ultrasound for On-Line Inclusion Detection in Molten Aluminum Alloys: Technology Assessment, AFS International Conference on Structural Casting of Aluminum, 2003, p 29–42 R. Gallo, Measurement of Hydrogen Content in Molten Aluminum Alloys Using Crucibles and Reverberatory Furnaces for Sand and Permanent Mold Castings, AFS
25.
26.
27.
28.
29.
30. 31.
32. 33.
34. 35.
36.
Fourth International Conference on Molten Aluminum Processing, 1995, p 209–225 F. Painchaud and J.P. Martin, The New Alscan Analyzer: Easy to Use, Reliable, On-Line Measurement of Hydrogen in Liquid Aluminum Alloys, AFS Second International Conference on Molten Aluminum Processing, 1989, p 20.1–20.21 C. Schwandt et al., A Novel Electrochemical Hydrogen Analyzer for Use in Molten Aluminum and Its Alloys, AFS Sixth International Conference on Molten Aluminum Processing, 2001, p 149–157 D.E. Groteke et al., Foundry Methods for Assessing Melt Cleanliness, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 55–69 D. Dispinar and J. Campbell, “A Comparison of Methods Used to Assess Aluminum Melt Quality,” Shape Casting: TMS, The Second International Symposium, 2007 D. Apelian and S. Dasgupta, Interaction of Initial Melt Cleanliness, Casting Process and Product Quality: Cleanliness Requirements Fit for a Specific Use, AFS Fifth International Conference on Molten Aluminum Processing, 1998, p 234–258 K. Strauss, Applied Science in the Casting of Metals, Pergamon, 1970, p 253–256 M.H. Kogan, Design and Development of Fluxing Agents for the Aluminum Foundry Alloys, AFS Second International Conference on Molten Aluminum Processing, 1989, p 25–1 to 25–12 T.A. Utigard et al., The Properties and Uses of Fluxes in Molten Aluminum Processing, J. Met., Nov 1998, p 38–43 S.R. Sibley Granular Fluxes for Aluminum Alloys, Environmental and Technological Advances, AFS Fourth International Conference on Molten Aluminum Processing, 1995, p 417–430 D. Neff, Ensuring Quality Prior to Pouring, Mod. Cast., June 2006, p 27–29 R. Gallo, Cleaner Aluminum Melts in Foundries: A Critical Review and Update Transactions, 112th AFS Casting Congress, May 17–20, 2008 (Atlanta, GA) D. Neff, The Benefits of Flux Injection Combined with Rotor Dispersion for Producing Clean Aluminum Casting Alloys, Proceedings, Third International Conference on Molten Aluminum Processing, Nov 1992 (Orlando, FL), AFS
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 240-254 DOI: 10.1361/asmhba0005301
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Modification of Aluminum-Silicon Alloys* G.K. Sigworth, Alcoa Primary Metals
ALUMINUM-SILICON ALLOYS are widely used in components where good strength and light weight are required or where corrosion resistance and good castability are needed. However, the commercial application of these materials depends on the control of the structure of the silicon phase. This control is usually called modification. Because of its commercial importance, modification has been the subject of several hundred studies since its discovery in 1921 (Ref 1). This article does not attempt to review the full extent of the available literature. Instead, the focus is on aspects important for the commercial production of castings. An examination of the literature shows that modification is a complex and controversial subject. Many conflicting theories have been proposed. Hence, this article by necessity reflects the personal opinions and interpretations of the author. The reader desiring more information and other viewpoints on this fascinating subject is referred to a number of excellent surveys available in the literature (Ref 2–12). This article considers the modification process in hypoeutectic and eutectic alloys, which differ only in the relative volume fraction of primary aluminum and aluminum-silicon eutectic. The refinement of hypereutectic alloys is not considered. In hypereutectic alloys, phosphorus is added to produce finer primary silicon particles, but the eutectic is not modified (Ref 3, 13–17). An unmodified alloy contains large flakes of brittle silicon, which cause the casting to have poor ductility. Unmodified alloys often have elongations no more than a few percent, and the fracture surface is primarily brittle. With a successful modification treatment, the silicon assumes a fine, fibrous structure. These fibers appear to be small, individual particles on a polished surface, but etching some of the aluminum from the surface shows that silicon is connected in a seaweed- or coral-like structure. The elongation of the modified alloy is improved significantly, and an examination of
the fracture surface reveals a dimpled texture normally associated with ductile failure. Figures 1 and 2 show the structure of modified and nonmodified A356 alloy, as observed on polished and deep-etched surfaces (Ref 18). Fracture surfaces can be seen in Ref 19.
The effect of modification on as-cast tensile properties is given in Table 1 for several alloys. Modification has a significant effect on ductility. This is the basis for a test that “old timers” sometimes use in the foundry. It is possible to pour a small bar casting. One end of the bar is
50 mm (a)
Fig. 1
(b) Structure of nonmodified A356 alloy solidified at 1.5 to 2 C/s (2.7 to 3.6 F/s). (a) Polished surface. Original Magnification: 400. (b) Deep-etched surface. Original Magnification: 1000. Source: Ref 18
50 mm (a)
Fig. 2
(b) A356 alloy modified with 156 ppm Sr and solidified at 1.5 to 2 C/s (2.7 to 3.6 F/s). (a) Polished surface. Original Magnification: 400. (b) Deep-etched surface. Original magnification: 1000. Source: Ref 18
*Used with permission from American Foundry Society (www.afsinc.org). Copyright 2008 from the following: International Journal of Metalcasting (ISSN 1939-5981), Vol 2 (No. 2), 2008, p19–41, and AFS Transactions (ISBN: 978-0-87433-338-1), Vol 116, 2008, p115–139 of the article by G.K. Sigworth, “The Modification of Al-Si Casting Alloys: Important Practical and Theoretical Aspects.”
Modification of Aluminum-Silicon Alloys / 241
Yield strength, 0.2% Alloy
Time, hr
Reported mechanical properties in the as-cast condition Ultimate tensile strength
Cast product
Modification
MPa
Ksi
MPa
Ksi
Elongation, %
Reference
13% Si
Sand-cast test bar Permanent mold test bar
... ... ... ... ... ... 112 108 125 125 126
... ... ... ... ... ... 16.2 15.6 18.1 18.1 18.3
124 193 193 221 180 210 137 159 168 191 193
18.0 28.0 28.0 32.1 26.1 30.5 19.9 23.1 24.4 27.7 28.0
2 13 3.6 8 5.5 12.0 1.8 8.4 6.0 12.0 12.8
20
13% Si
None Na-modified None Na-modified None 0.007% Sr None Sr-modified None Sr-modified 0.06% Sr
359
Permanent mold test bar
A413
Sand-cast test bar
A413
Permanent mold test bar
A413
Sample from auto wheel
20 21 22
Heat Treated Castings An examination of Fig. 1 and 2 and Table 1 makes it abundantly clear that the mechanical properties of castings depend strongly on the size and shape of silicon crystals. One way to improve ductility is to employ a chemical modification, as described previously, by adding a small amount of sodium or strontium. Another way is to employ a thermal modification by subjecting castings to heat treatment. Few foundrymen understand that a thermal modification is possible or that standard heat treatment schedules are designed for nonmodified alloys. Consider the schedules recommended for the solution heat treatment of A356 alloy by the American Foundry Society (Ref 26): Permanent mold castings: 12 h at 538 C
(1000 F) for T6 temper; 6 to 12 h for T61 temper
4
2
1
Perm
anen
t mo
ld
400
Sand
350
Sr-modified Nonmodified
22 22
8
450
300 0
0.1
0.2
cast
0.3
t,−1/3
Fig. 3 clamped in a vise, and the free end is hit with a hammer. The energy required to break the bar and the appearance of the fracture surface indicate when the alloy is modified. Sodium was the modifier first used by foundrymen. It was employed for many years, especially in aluminum-silicon alloys low in magnesium. These alloys provide good elongation with modest strength in the as-cast condition and were used in applications requiring reasonable ductility, such as automotive wheels. However, it is difficult to add sodium, which has low solubility in molten aluminum and a high vapor pressure, so its recovery is erratic. Sodium is also rapidly lost to oxidation after the addition (Ref 23). These problems made it difficult to control sodium modification. In the 1970s, strontium began to replace sodium as the preferred modifier. Strontium is easily added via master alloys with nearly 100% recovery, and its loss to oxidation is slow (Ref 23). With strontium, the control of microstructure is easier and the mechanical properties in castings are more consistent (Ref 24, 25). Today (2008), the use of sodium as a modifier has almost completely disappeared. For this reason, the primary focus of this article is on strontium modification.
20
500
Q, MPa
Table 1
Sand castings: 12 h at 538 C (1000 F) for
Mechanical properties (Q) versus solution time (t) in A356-T6 alloy castings. Source: Ref 27
T6 temper A detailed study of the solution heat treatment process (Ref 27) has shown that only 60 min is needed to place the magnesium and silicon in solution in permanent mold castings. The additional 11 h recommended for the T6 solution heat treatment are used to modify (spheroidize) the shape of silicon crystals. The practical consequence is important: When aluminum-silicon alloys are chemically modified with sodium or strontium, the solution heat treatment time may be reduced. This observation was first noted by Levy in 1978 (Ref 28), who described steps undertaken at Reynolds Metals to increase ductility in lowpressure die-cast aluminum wheels. Prior to this time, most aluminum wheels in the United States were sold to the automotive aftermarket. The original equipment manufactures wanted an elongation of 10%, whereas 4% was the typical value in cast wheels. Levy found that 10% elongation could be achieved by:
for the quality index follow a straight line when plotted against the inverse of the cube root of solution time (minutes). The corresponding solution time (in hours) is also indicated at the top of Fig. 3. It is useful to consider what happens to the silicon phase during solution heat treatment and how the change in structure influences the quality of aluminum castings. A study of nonmodified alloys under a hot stage microscope (Ref 31) showed that silicon plates (in the ascast material) first break up into smaller pieces. This breakup happens rapidly, in an hour or two. There then follows a period where the small silicon pieces gradually coarsen and become more spherical. The same process occurs in modified alloys (Ref 27, 29). Boeing engineers conducted extensive studies of highquality aerospace castings. They examined the effect of solution heat treatment time along with other process variables. They finally concluded (Ref 32) that the quality of A357 alloy castings was well represented by the equation:
Maintaining a low iron content in the metal Modifying the metal (first with sodium, then
Q ¼ 101 N 0:5 129
later with strontium) Using a shorter aging time These procedures are now standard practice, and aluminum castings are used extensively for wheels in new automobiles. Levy also noted that the solution time with modified metal could be reduced from 8 to 4 h. The reduced solution time possible with strontium modification is now also a common practice. These considerations led to detailed studies (Ref 27, 29) on the solution heat treatment of nonmodified and strontium-modified A356 alloy castings. Some of these results are plotted in Fig. 3 for permanent mold (ASTM B108) and sand mold test bar castings. The mechanical properties of the castings in Fig. 3 have been characterized by the quality index developed empirically by Drouzy et al. (Ref 30): Q ¼ UTS þ 150 log Ef
(Eq 1)
where the quality (Q) and the ultimate tensile strength (UTS) are in MPa, and Ef is the elongation to fracture in a tensile test. The data points
þ
N 0:5 AR
1100 98 P 0:5 570 AR
(Eq 2)
where Q is the quality index (in MPa), N is the cell count (aluminum cells per 0.0001 in.2), AR is the average aspect ratio of the silicon particles, and P is the area percent of porosity found on the polished surface of test specimens. The magnitude of the coefficient of the third term (1100 MPa) shows the importance of silicon particle shape (AR) on casting quality and gives credence to theoretical studies that suggest that cracking of silicon particles controls tensile elongation and quality in aluminum-silicon alloy castings (Ref 33). Studies made by the author establish that the cost for solution heat treatment is approximately 0.5 cents/lb-h. With strontium modification, the solution time required can be cut in half. For permanent mold castings, this represents a saving of at least 4 h, or 2 cents/lb. This amount is significant in today’s (2008) competitive market and is nearly ten times the cost of the strontium addition. Moreover, halving the
242 / Molten Metal Processing and Handling solution time doubles the throughput of heat treatment furnaces, removing a production bottleneck in many foundries. These are important benefits that often justify the modification of heat treated castings. The aforementioned results show that, when long solution times are employed, non- and strontium-modified permanent mold castings have almost the same tensile properties. In fact, when examining the microstructure of castings subjected to a standard solution heat treatment, it is difficult to distinguish any difference in the silicon phase. (Consider, for example, Fig. 1 of Ref 8 or Fig. 10 of Ref 29.) Hence, one may ask: “Why modify with strontium?” One answer is that modification usually changes the distribution of porosity and shrinkage in a casting, so that strontium additions may cure shrinkage problems. This topic has been the subject of a great deal of controversy over the years and is considered in detail as follows.
Porosity and Shrinkage Formation In most aluminum foundries, one-half to threequarters of all scrap is caused by the undesirable formation of porosity or shrinkage or by problems directly related to these defects. Lack of pressure tightness, poor quality on machined surfaces, low elongation and loss of tensile strength, and poor fatigue strength are problems related to porosity and/or shrinkage formation. Unfortunately, there is no single recipe that foundrymen can use to produce good castings. A number of factors must be
considered and controlled, and the balance point or “sweet spot” that produces the best results in one casting may give poor results in another. The situation is complex, poorly understood, and subject to a good deal of controversy. This section first considers porosity and how modification changes porosity formation in a casting. Then, shrinkage is considered as well as how modification changes feeding and shrinkage formation. The amount of porosity in a casting depends on several factors, including:
Solidification rate Gas content Metal cleanliness Modification Grain refinement Pressure in the casting
Four of these factors were studied by Fang and Granger (Ref 34). They poured directionally solidified castings of A356 alloy, which had a water-cooled copper chill at the bottom and were solidified directionally upward from this chill. Because of the directional solidification, these castings had more than adequate feeding, and shrinkage formation was not an issue. The amount of porosity and the average size of the pores were measured by quantitative metallography at various distances from the chill. The solidification rate at these locations was determined by thermocouples placed in the mold during solidification. The gas content of the molten metal was measured directly by the Telegas instrument prior to solidification.
The porosity of the castings produced as a function of solidification rate and gas content is plotted in Fig. 4. Several important observations can be made from these results: 1. Solidification rate has an overriding influence on the amount of porosity in a casting. Castings that freeze quickly can tolerate high contents of hydrogen gas. Slowly cooled castings, however, easily form significant amounts of porosity. This is why low gas levels are needed for best quality in large sand castings, but gas is sometimes added to permanent mold castings. 2. Grain refinement tends to reduce the amount of porosity. (Compare the results for castings 1 and 3 in Fig. 4b). Grain refinement also reduces the size of pores. This occurs because hydrogen gas pressure does not usually become sufficiently high for pores to form until the last 5 to 10% of solidification. Hence, pores must fit into the spaces left between primary aluminum grains. Smaller pores produce a beneficial effect on fatigue life. 3. Modification tends to increase the amount of porosity formed, at least at the relatively high levels of gas (0.31 cm3/100 g) present in this study. (Compare the results for castings 2 and 5 in Fig. 4b). The question of porosity distribution in sodiumand strontium-modified castings was studied by Pechiney and Montupet engineers (Ref 25).
0.8 0.85 ⬚C/s 1 0.7
A356 Alloy
0.5
Castings 1, 2, 4 Nonmodified Grain refined
0.4
0.3
2.6 ⬚C/s
Pore volume fraction, %
Pore volume fraction, %
0.6
10−1
(5)
(2)
10−2
(5) (2) (1) (3) (4)
0.2
0.1
7.5 ⬚C/s
10−3 10−1
0 (a)
Fig. 4
0.2
Hydrogen content, cm3/100 g Pore volume fraction in A356 alloy castings. Source: Ref 34
0.31*** 0.31** 0.25* 0.25** 0.11**
(4)
No grain refiner Grain refined Grain refined and modified
28 ⬚C/s 0.1
(1) (3)
Casting No. H2content (cm3/100g)
0.3 (b)
1
10 Cooling rate, ⬚C/s
102
Modification of Aluminum-Silicon Alloys / 243 As part of their study, they poured an end-chilled sand mold bar casting. The design of the casting is shown in Fig. 5. The porosity was determined in this bar by measuring the density (specific gravity) in sections cut along its length. These measurements were made in sodium-, strontium- and nonmodified A356 alloys. In the sodium- and strontium-modified alloys, there was an increased amount of porosity near the riser. The results for the strontium-modified alloy are shown in Fig. 5. (The results for sodium-modified alloys were similar but are not duplicated here.) The increase in porosity is most pronounced at higher addition levels of strontium. In practice, this means that one should use the minimum amount of strontium needed to produce an acceptable level of modification. This is generally approximately 0.008 to 0.012% or 80 to 120 ppm Sr. There are a number of actions that foundrymen can take to minimize the increase in porosity associated with strontium-modification. Aside from keeping strontium levels at a minimum, one should consider: Lower gas in the metal reduces the amount of
porosity. Longer degassing times and low metal temperatures will help. The influence of gas content is illustrated well in a paper by Emadi and Gruzleski (Ref 35). They produced directionally solidified castings (similar to Fang and Granger in Fig. 4). Some of their results have been reproduced in Fig. 6. At a low gas content (0.13 cm3/100 g), the increase in porosity with modification is negligible, except at the slowest cooling rates (compare curves 1 and 2). At a higher gas content (0.26 cm3/100 g), there is a significant increase in porosity when strontium is added (curves 3 and 4). The importance of foundry degassing practices should be obvious from these results. Ingate 250
Increasing solidification rate always reduces
porosity, but this is especially true for strontium-modified alloys. If excessive porosity is found in part of a casting, this can often be removed by using chills in sand castings or conductive die coatings (or increased water cooling) in permanent mold castings. Grain refinement is beneficial and tends to counteract the porosity formation caused by strontium additions. For this reason, it is a good idea to use grain refinement in combination with strontium modification. Grain refinement is best accomplished by an addition of 10 to 20 ppm B in the form of 5% Ti-1%B or 3%Ti-1%B rod. Mold design is also an important consideration in new castings. It is important to keep sources of heat away from areas where porosity is not desired. For example, in a wheel casting, one would not want to gate metal into an area that will be machined to a high polish and coated. With the increased availability of powerful casting simulation programs, it is possible to determine with great precision where porosity may form before machining patterns or molds. In fact, when a casting is designed properly, an xray examination will show that the increased porosity formed by strontium additions is located in the base of the riser(s). This represents the ideal case. Filtration will remove inclusions, making it more difficult for gas pores to nucleate during solidification.
Modification Mechanisms At this point, it is useful to consider the mechanisms responsible for silicon modification and that cause the changes in porosity noted previously. When these mechanisms are understood, it is easier to see how strontium modification produces other effects:
37.5
Chills
1.2 Unmodified 1 3 Sr-modified 2 4
Riser
Specific gravity
2.68
0
Testpiece
Length along testpiece, in. 4 6 8 2
10
2.67
No Sr
2.66
0.01% Sr
2.65 0.07% Sr 0.10% Sr
2.64 2.63
0
Chill end
Fig. 5
0.8
3
Gas content 0.13 0.26 0.13 0.26
castings
Increased segregation of iron compounds Increased segregation of copper as blocky
CuAl2 in copper-containing alloys Cooling curves recorded during the solidification of aluminum-silicon casting alloys are obtained by pouring a liquid metal sample into a small crucible or sand cup, into which a thermocouple is placed. The interpretation of the temperature profile recorded during freezing is called thermal analysis. Thermal analysis can be a powerful tool, because the cooling curve of the alloy acts like a fingerprint to identify the phases forming during solidification. Thermal analysis has a long history as a laboratory tool for the study of aluminum-silicon alloys. Two outstanding early examples are the 1950 work by Thall and Chalmers (Ref 20) and the 1966 paper by Crosley and Mondolfo (Ref 3). The latter experiments, in particular, helped to lay the foundation for the present understanding of aluminum-silicon alloys. In the early 1980s, two versions of commercial thermal analysis equipment became available (Ref 36, 37), one of which was subsequently used widely as a foundry tool to evaluate grain refinement and modification treatments in the cast shop (Ref 24, 37). Unfortunately, no improvements have been made since that time, and the tool has fallen into disuse. It is rare to see thermal analysis equipment operating today (2008) in foundries. Also, researchers have proposed a number of advanced analytical techniques to determine silicon modification (Ref 18, 36, 38, 39), but none of these ideas has been transferred onto the foundry floor. There appears to be a useful opportunity for someone willing to convert this knowledge into a modern, inexpensive, and easy-to-use computer-based thermal analysis system. Researchers in Australia poured castings with a base Al-Si-Mg alloy and then made additions of strontium and phosphorus to the same metal (Ref 40). Thermal analysis curves for the three alloys are given in Fig. 7. Important clues to operating mechanisms are available from these curves. Before extracting the information, however, it is necessary to introduce a
0.6 2 0.4 1 0.2
0.12% Sr
2.62 2.61
Pore volume fraction, %
25
1.0
Improved feeding from risers A changed distribution of shrinkage in
200 50 100 150 Length along testpiece, mm
250
0 0.3
Distribution of porosity in an end-chilled sand casting. Source: Ref 25
0.5
1
2
3
4 5
Average cooling rate, ⬚C/s
Riser end
Fig. 6
Porosity versus freezing rate in an A356 alloy. Source: Ref 35
Fig. 7
Effect of strontium and phosphorus on cooling curves in an A356 alloy. Source: Ref 40
244 / Molten Metal Processing and Handling few definitions and some basic concepts. For this purpose, the cooling curve for a modified A356 alloy is employed, as shown in Fig. 8 (Ref 41). Figure. 8(a) shows the full curve. TC is the temperature at the center of a small crucible during freezing. dT/dt is the first derivative of the cooling curve, TC, and shows the instantaneous cooling rate in degrees per second. (The scale for the dT/dt curve appears on the right side of Fig. 8a.) Figure. 8(b) shows an enlarged version of the cooling curves for the solidification of the aluminum-silicon eutectic. (This is the area inside the box outlined by dashed lines in Fig. 8a.) Before the first eutectic forms, the cooling rate (dT/dt) is nearly constant at 0.6 C/s (31 F/s) (Fig. 8). At a temperature of approximately 570 C (1060 F), the first derivative starts to increase. This happens because the aluminum-silicon eutectic has begun to freeze, releasing latent heat into the sample. This point is called the eutectic nucleation temperature, TN. As more eutectic grains form and grow, the amount of the latent heat released becomes larger, until the cooling stops and a minimum temperature occurs, TMin. The sample ⬚C 630
1.5⬚C/s 1.3
620
1.1
Tc
0.9
610
0.7 0.5
600
0.3 590
0.1 −0.1
dT/dt
580
−0.5 −0.7 −0.9
560
−1.1 550 0
50
(a)
100
150
−1.3 200
Time, s
584
1
582
0.8
580
The eutectic nucleation temperature (TN) is
reduced.
The eutectic growth temperature (TG) is
reduced.
The nucleation temperature is lowered with
respect to the growth temperature (that is, TN TG). Phosphorus additions can be seen to have the opposite effects of those produced by strontium.
Eutectic Nucleation How does an addition of a few parts per million of strontium (or sodium) change the nucleation of the silicon eutectic? This is an interesting and controversial question that has been the subject of dozens of research papers and academic theses. In the author’s opinion, the most likely explanation is that strontium (and other modifiers) react with to remove or somehow neutralize AlP, which is an effective nucleant for the silicon phase. The supporting evidence for this hypothesis is outlined briefly as follows: Phosphorus is added to hypereutectic alloys
−0.3
570
then begins to heat up. (This is called recalescence.) The sample continues to heat until the steady-state growth temperature for the eutectic, TG. From the cooling curves in Fig. 7, strontium additions are demonstrated to produce the following effects:
Al-10%Si alloy. The sample was quenched from the semisolid state. The fine eutectic surrounding the larger silicon crystals formed during this quench. (The dashed lines in Fig. 9 indicate the outline of primary dendrites forming prior to eutectic solidification.) Some researchers have reported finding Al-P-O particles, but this observation appears to be an artifact. AlP is a reactive compound that combines readily with oxygen in the air and perhaps with water used in polishing media. The oxygen is therefore the result of contamination. This was shown by elegant (and complicated) ion milling techniques used to remove the contaminated oxygen-containing layers. The same study (Ref 43) also showed there is a definite epitaxial crystallographic relationship between AlP and the silicon crystal. Modifiers react with phosphorus. The elements that have been reported to produce modification (Ref 7, 41, 44) are indicated by the shaded areas in Fig. 10. These are all reactive metals that form compounds with phosphorus. As a consequence, they all have the capability to neutralize AlP particles in aluminum melts. This interpretation appears to be supported by thermal analysis data (Ref 41). Phosphorus additions poison modification. A number of researchers have shown that high phosphorus levels make it difficult to modify aluminum-silicon alloys. Some of the maps produced by these studies are shown later, when the effects of elemental additions on the modification process are
to reduce (or refine) the size of primary silicon. In 1966, Crosley and Mondolfo (Ref 3) found phosphorus-containing particles inside primary silicon crystals. Based on crystallographic data, they proposed that AlP has an epitaxial relationship with silicon and would be an effective nucleus. AlP nucleates silicon in hypoeutectic alloys. This has been observed by a number of researchers. An example is shown in Fig. 9. This particle was observed in a nonmodified
I
IIA
IIIB
H
Li
Be
Na
Mg
K
Ca
Sc
Rb
Sr
Y
Cs
Ba
La
0.6
Tc
0.4
578 dT/dt 576
0.2
574
0
572
−0.2
570
TG
TN
−0.4 −0.6
568 Tmin
566 564 80 (b)
Fig. 8 Ref 41
−0.8 −1
100
120
140
160
Time, s Thermal indicators in eutectic solidification. See text for an explanation of parameters. Source:
Fig. 9
Phosphorus-rich particle inside silicon crystal in an Al-10%Si alloy. Source: Ref 42
Fig. 10
Elements that have been observed to produce modification in aluminum-silicon alloys
Modification of Aluminum-Silicon Alloys / 245 discussed. One example of this poisoning effect is shown in Fig. 11. Strontium poisons AlP in hypereutectic alloys. The interaction between AlP and strontium is reciprocal and also occurs in hypereutectic alloys. In these alloys, phosphorus is normally added to nucleate and refine primary silicon. When strontium is added in sufficient quantities, and some time is allowed for the reaction between AlP and strontium to occur, the size of primary silicon increases significantly and the number of primary silicon particles is reduced (Ref 46, 47). Modifier additions change eutectic grain size. When a modifier is added, nucleants for the silicon eutectic are removed. The most significant consequence is a larger eutectic grain size and the change in porosity distribution discussed previously. Due to the importance of this subject, it is worthy of discussion in a separate section.
Fig. 12(b). In this case, no silicon eutectic grains nucleated in the liquid. Instead, the eutectic first formed at the outside of the casting (in contact with the mold) and grew toward the center as a planar solidification front. Unfortunately, no thermal analysis was conducted in this study. From the structures observed, however, it would be reasonable to assume that for nonmodified alloys, the eutectic nucleation temperature is greater than the growth temperature (that is, TN > TG). This may be inferred from the appearance of eutectic grains in the liquid phase ahead of the solid/liquid interface. Conversely, for sodium-modified
alloys, the nucleation temperature is lower than the growth temperature (TN < TG). These differences have an important effect on the formation of porosity and shrinkage. Inferences can also be made about the eutectic grain structure by anodizing a polished surface and looking at the sample in a microscope with polarized light. This procedure reveals the grain structure (size and relative orientation) of the aluminum phase in the eutectic. This procedure was used in the 1975 study by Jenkinson and Hogan (Ref 49). More recently, a number of studies have been undertaken by researchers in Australia.
Eutectic Grain Structure Most popular casting alloys are hypoeutectic. For example, A356 alloy contains 7% Si, compared to the eutectic composition of 12.7% Si. Consequently, the as-cast structure contains 50 to 60% primary aluminum grains surrounded by a layer of aluminum-silicon eutectic. It is difficult to discern the eutectic grain structure in this material by simple metallographic observation. One way to observe eutectic grains is to partially solidify the alloy and subject a sample to a rapid quench. This procedure is usually called an interrupted quench. (That is, solidification is interrupted by a rapid quench.) This procedure was used in the 1963 research paper by Kim and Heine (Ref 44). It is possible that they were the first to observe the change in the eutectic grain structure of aluminum-silicon alloys that occurs with modification. Unfortunately, their observation was not connected with the important issues at hand (porosity and shrinkage distribution in castings), and it was forgotten. In 1981, Flood and Hunt used interrupted quench experiments to study sodium modification of eutectic alloys (Ref 48). Their results clearly showed that sodium prevented the nucleation of eutectic grains. Figure. 12 is taken from their study. Figure. 12(a) shows the structure of the nonmodified alloy. The lighter (white) portion of the ingot was allowed to solidify normally. The darker-gray areas were quenched after approximately one-third of solidification was completed. It can be seen that numerous eutectic grains nucleated in this nonmodified alloy and grew into small, spherical colonies of silicon crystals. The arrows in this figure show some of these eutectic grains. It is noteworthy that these grains are isolated, or completely surrounded by liquid metal. The sodium-modified alloy is shown in
Fig. 11
Effect of strontium and phosphorus additions and holding time. Mechanical properties are for an as-cast Al-12%Si alloy. Source: Ref 45
Nonmodified alloy (a)
Fig. 12
Na-modified alloy (b)
Structure of aluminum-silicon eutectic alloy observed by an interrupted quench. Arrows show individual eutectic grains. Source: Ref 48
246 / Molten Metal Processing and Handling They developed special techniques to see eutectic grains in aluminum-silicon alloy castings. For the researcher interested in experimental detail, the new procedures are reviewed in Ref 50 and 51. For the purpose of this article, however, it is sufficient to extract the most important observations that have practical significance for foundrymen. The eutectic structures observed in aluminum-silicon alloy castings (Ref 42, 50–55) are summarized in Fig. 13. This figure shows schematically the structure of three Al-7%Si alloy castings at a point during solidification when the primary aluminum grains have formed and the eutectic solidification is just beginning. Figure 13(a) shows the nonmodified structure. The black areas indicate regions where aluminumsilicon eutectic has solidified. There is a shell of eutectic in contact with the mold wall, and numerous small, solid grains of eutectic are in contact with primary aluminum. In this case, eutectic grains nucleate and grow on existing aluminum dendrites. Figure 13(b) shows the strontium-modified alloy. In this case, the number of eutectic grains is smaller, and the eutectic grain size is much larger. In experiments conducted by McDonald et al. (Ref 53–55), the eutectic grain size increased by an order of magnitude when strontium was added. The eutectic grains were a few hundred microns in diameter in nonmodified alloys, and a few thousand microns in strontium-modified castings. Note also that the solid shell of eutectic is thinner when strontium is present. In fact, in large sand or investment castings there may be little solidified eutectic at portions of the mold surface. These changes have important ramifications for shrinkage formation. The structure found in sodium-modified alloys is shown in Fig. 13(c). In this case, no nucleation occurs in the liquid, and eutectic growth is from the mold walls toward the center. (This case was also shown in Fig. 12b for an alloy of eutectic composition.) In the discussion of Fig. 4, it was noted that grain refinement (of the primary aluminum
(a)
Fig. 13
(b)
phase) had a beneficial effect on the amount of porosity. The size of the pores was also reduced. From Fig. 13, it can now be seen that strontium additions have the opposite effect: The size of eutectic grains is increased. This change is almost certainly responsible for the increase in porosity found in most modified castings. To one familiar with the technical literature in this area, this simple conclusion may be surprising. There has been a good deal of controversy over the last 25 years, and nearly a dozen theories have been proposed to account for the change in porosity observed in modified alloys. Surprisingly, eutectic grain size was never identified as a possibility, presumably because no one could measure it. For those interested in the scientific history and background, the other theories and the literature are reviewed elsewhere (Ref 52, 54, 56–58). It may be useful to present a simple calculation. Fuoco and co-workers (Ref 59) duplicated the experiments of the Pechiney researchers summarized in Fig. 5 and also conducted interrupted quench experiments in small crucibles. Their melts were maintained at a gas content of 0.13 cm3/100 g (standard temperature and pressure, or STP). When strontium was added as a modifier, the amount of porosity increased, and the pore size increased to an average diameter of 300 mm. Now, the question is this: From what volume of liquid metal (eutectic grain size) did this pore receive its hydrogen gas? It is assumed that all of the gas originally present in the melt finds its way into a pore during solidification. If this is true, how much porosity forms? One hundred grams of metal, at a liquid density of 2.35 g/cm3, has a volume of 43 cm3. The gas volume at standard temperature and pressure is 0.13 cm3. However, at the eutectic temperature (and 101 kPa, or 1 atm, pressure), this gas will have a volume of 0.13 cm3 (843/ 273) = 0.4 cm3. Thus, the amount of porosity that corresponds to a gas content of 0.13 cm3/100 g (at STP) is approximately:
(c)
Structure of aluminum-silicon eutectic grains in three semisolid Al-7%Si alloys. (a) Unmodified. (b) Strontium modified. (c) Sodium modified. Source: Ref 42
% porosity ¼ Volume of gas=Volume of metal ¼ 0:4 cm3 =43 cm3 ¼ 0:0093 or 0:93%
(Eq 3)
For the material to have that amount of porosity, each pore must be surrounded by a volume of solid (pore-free) metal having an average size 107 times larger than the pore ([1 0.0093]/0.0093 = 107). From this is derived: ffiffiffiffiffiffiffiffi Dgrain p ffi 3 107 or 4:75; thus Dpore Dgrain ffi 1425 mm
(Eq 4)
In nonmodified alloys, the eutectic grains are approximately 250 mm in size. Using the same calculation as in Eq 4, it is found that the pore size is much smaller—approximately 50 mm. From this simple calculation, one can see how a large eutectic grain size in modified alloys may contribute to a larger pore size in castings. The previous calculation assumed that all hydrogen was able to diffuse to the pore, and it corresponds to the case of slow solidification. During rapid freezing, much of the hydrogen is frozen or trapped in the solid metal. In this case, the amount of porosity is less and the pores are smaller. Unfortunately, this simple picture is not the complete story. The experiments of Fuoco et al. (Ref 59) and Anson et al. (Ref 56, 57) show that, at medium to high gas levels, pores nucleate or form earlier during the solidification process when strontium is present. The change in eutectic grain size cannot account for this observation. One possibility is a different nucleus for the pores. It is well known that oxide films nucleate porosity in aluminum castings (Ref 60, 61). Strontium additions have been observed to change the composition of the oxide on the surface of a eutectic aluminum-silicon alloy and to increase the oxidation rate of the melt (Ref 62). Strontium-containing oxides have also been observed to nucleate pores in strontium-modified castings (Ref 63, 64). So, it is possible that a strontium-containing oxide film is a better nucleant for porosity. The discussion now shifts to other important changes that occur when an aluminum-silicon alloy is modified. Feeding and Shrinkage Formation. Considering once more the structures indicated in Fig. 13, it is evident that relatively large channels of liquid eutectic will exist in the semisolid, strontium-modified castings. These open channels change the feeding patterns. Because there is a thin shell of eutectic at the surface, these channels sometimes come to the surface. The author has seen large, strontium-modified sand castings that exhibited a “wormhole” on the surface. These holes may go deep into the casting where an unfed thick section sucked liquid away from the surface. This left an empty hole where the eutectic feeding channel was located. More common is the problem of a rough surface, usually called an “orange-peel” defect, sometimes seen in investment castings. Here again, poor feeding in a thick section pulls liquid away from the surface, leaving
Modification of Aluminum-Silicon Alloys / 247 a rough surface of exposed, dry aluminum dendrites. This defect has been described in the literature (Ref 65). Strontium has also been shown to promote surface slumping in unfed sections of permanent mold castings (Ref 52, 66). Sodiummodified castings did not exhibit surface slumping, presumably because the solid outer shell of solidified eutectic is thicker. The literature regarding the effect of modifiers on feeding is contradictory. Iwahori and co-workers (Ref 67) conducted an elegant series of experiments that showed feeding stopped in nonmodified alloys when the aluminum-silicon eutectic began to freeze. In Fig. 13 (a), it is tempting to say that the many small eutectic grains in this structure act to close off liquid feeding paths between aluminum dendrites. Iwahori and his colleagues later (Ref 68) repeated the same experiments in modified alloys. Sodium-modified alloys fed longer (later into solidification), which meant that smaller risers could be employed without creating shrinkage in their casting. On the other hand, strontium-modified alloys were observed to behave in the same manner as nonmodified melts. Unfortunately, these experiments were conducted with 1.7 kg (3.7 lb) melts, so the results cannot be scaled up directly to melt practices found in operating foundries. A third study (Ref 69) reported that strontium made feeding worse in a simple bar casting, whereas three other papers (Ref 70–72) reported that strontium modification improved feeding. The author also has experienced many instances of strontium effects on shrinkage. In the early 1980s, the author began to work for a company that made and sold aluminumstrontium master alloys. During one foundry visit, they were having problems with shrinkage located on the machined deck face of lost-foam cylinder heads in a 319-type alloy. When they modified the alloy with strontium, the shrinkage defect disappeared. Since that time, the author has seen the same thing occur in many other castings, and others (Ref 42, 66, 73, 74) have
(a)
Fig. 14
(b)
reported the same effect of strontium on shrinkage. An example of the change possible is given in Fig. 14. This shows the internal structure at a hot spot in a permanent mold casting. In the nonmodified alloy (Fig. 14a), a large, concentrated shrink is found. When strontium is added (Fig. 14b and Fig. 14c), the shrink disappears and is replaced by dispersed microporosity. The author thinks that strontium helps to cure shrinkage by a combination of better feeding and increased porosity. Both tend to displace shrinkage. By looking carefully at the tops of risers during solidification, it will become apparent that feeding usually improves when strontium is used as a modifier. (The shape of the cavity at the riser top changes, showing that feeding occurs longer when the metal is modified.) Grain refinement also assists feeding, which is another reason to use it in combination with strontium. In the sodium-modified casting (Fig. 14d), a concentrated shrink is found. However, it is smaller in size than in the nonmodified alloy. Perhaps the shrink was reduced by improved feeding. Copper and Iron Segregation. The larger eutectic grain size in modified alloys tends to increase copper and iron segregation in the same way it increases segregation of gas. Several studies have noted that strontium promotes the formation of large, blocky CuAl2 particles (Ref 55, 75–77). These particles take longer to dissolve during solution heat treatment (Ref 75). There also is an observable change in the copper eutectic cooling curves (Ref 75, 77), so thermal analysis could be used to help control CuAl2 morphology in modified Al-Si-Cu casting alloys. The effect of strontium on copper segregation in an Al-10%Si-2%Cu alloy is shown in Fig. 15. Grain refinement of the primary aluminum phase helps to disperse the blocky CuAl2 particles (Ref 75). A similar segregation effect may also occur with iron intermetallic compounds, depending
(c)
on the iron content of the alloy. Figure 16 shows the distribution of b plates in an Al10%Si alloy that contains 0.11% Fe and 140 ppm Sr. The iron-containing plates are found at the boundaries of the eutectic grains.
Recommended Foundry Practices Consider how strontium is best used in the foundry. Each process step is considered individually. Adding Strontium. Making additions of strontium is relatively easy. There are several master alloy compositions and different products available for this purpose. Their recovery and speed of dissolution have been studied in detail (Ref 24, 39, 78–80). Waffle products tend to be slow to dissolve. Higher metal temperatures generally promote more rapid dissolution and higher recoveries. The exception is the 90Sr-10Al product, which is exothermic and goes into solution best when the melt temperature is low. Strontium recoveries are generally high with all products—90% or more. In past years, an Al-14Sr-10Si master alloy was marketed by some companies, but this was a terrible product. It was hard to dissolve, and it always gassed the metal. This is perhaps because the product contained high levels of calcium and barium (Ref 78). This master alloy has not been available in North America for many years, but the author was recently told that an overseas company is once again making this product. Foundrymen should be aware of potential problems with this strontium-silicon master alloy. Several workers have noted that strontium is slow-acting, showing an incubation time of 1 or 2 h (Ref 10, 24, 39, 45, 72). Modification was observed to improve with time after the addition, even though part of the added strontium may be lost to oxidation. An example of this effect was shown in Fig. 11. More recent research has established that most of the delay
(d)
Cross section of hot-spot area (spoke-rim junction) of permanent mold Al-10%Si-Mg alloy casting. (a) Unmodified alloy. (b) 50 ppm Sr. (c) 100 ppm Sr. (d) Na-modified. Source: Ref 42
248 / Molten Metal Processing and Handling
Distribution of b-phase in an Al-10%Si alloy. Arrows show b plates located at eutectic grain boundaries. Source: Ref 55
Fig. 16
100 µm
(a)
(b)
Fig. 15
Binary image showing distribution of CuAl2 phase in an Al-10%Si alloy casting. (a) 2 wt% Cu, 0 ppm Sr. (b) 2 wt% Cu, 100 ppm Sr. Source: Ref 55
Fig. 17
Quality index of A356-T6 alloy castings at three cooling rates. Curve 1, 1.5 C/s; curve 2, 0.5 C/s; curve 3, 0.08 C/s. Source: Ref 80
is caused by the time needed for intermetallic particles in waffle products to dissolve. Strontium-containing rod products usually have small particles of SrAl4 and are fast dissolving and fast acting (Ref 24, 39). With strontium master alloy rod additions, no more than approximately 10 to 15 min is needed to reach full modification.
The melt loss of strontium is slow (Ref 23), so it will be convenient for many foundrymen to have strontium added to their alloys by their ingot supplier. How Much to Add. It is generally accepted that the best silicon modification in A356 alloy occurs at a level of approximately 0.012% Sr, or 120 ppm. An example of how mechanical
properties have been observed to vary with strontium content is shown in Fig. 17 (Ref 80). This plot shows the quality index of A356 alloy at three different cooling rates. In practice, however, good modification can be obtained at significantly lower concentrations, depending on the purity of the alloy. In primary metal, the phosphorus level is usually below 10 ppm, and other bad impurities are absent, so full modification can be obtained with strontium contents as low as 60 to 70 ppm. Good modification can also be obtained at higher levels of strontium, although this tends to result in more porosity. The reason for the small drop in quality as strontium becomes more than approximately 0.03% is that higher levels result in coarser silicon particles (Ref 18, 81). Gas Pickup and Degassing. It is generally agreed that adding strontium does not result in any appreciable gas pickup (Ref 82, 83). Sodium additions, however, do add gas to the melt (Ref 84). The evidence regarding the effect of modifiers on gas pickup from the atmosphere is less clear, however. Experiments with small melts held under atmospheres having a controlled humidity (Ref 62) show that strontium additions do increase the gas by changing the composition of the oxide on the melt surface. Other laboratory experiments (Ref 83) suggest that strontium has no effect. The author’s experience in the foundry is in agreement with the observations offered by Hurley and Atkinson (Ref 24). As long as large strontium additions and high melt temperatures (>760 C, or 1400 F) are avoided, there is no significant increase in gas pickup from the atmosphere. Degassing of modified melts should be done using nitrogen or argon, preferably with a rotary impeller. No chlorine or other reactive gas (nor salts or fluxes) should be employed, because these will tend to burn out dissolved strontium. Contrary to an observed loss of sodium (Ref 84), degassing with inert gases does not result in a significant loss of strontium. Quality-Control Testing. Under normal conditions, a spectrographic analysis for strontium is sufficient to control modification in a
Modification of Aluminum-Silicon Alloys / 249 foundry using primary metal. As shown subsequently, however, some elements can poison the effect of modifiers. These may become a concern when secondary metals or scrap are incorporated into the furnace charge. In this case, it would be wise to check modification levels by the routine use of thermal analysis. This measurement technique has been found to be an acceptable quality-assurance tool (Ref 24, 37). Mold Reactions. The addition of strontium has been observed to increase hydrogen gas pickup in green sand mold castings (Ref 84, 85). This can become a serious problem in large, complicated castings. To a certain extent, this gas pickup can be avoided by careful gate and runner design and by use of mold coatings.
Growth of Silicon Eutectic A careful examination of the literature makes it evident that the crystallization of silicon in aluminum-silicon alloys is an exceedingly complex affair. In particular, the presence (or absence) of impurity elements has a significant effect on the structure formed. In the earlier discussion of cooling curves, it was noted that modifiers change the nucleation of silicon. This effect was considered in some detail, to explain effects ancillary to silicon modification (changes in porosity formation, feeding, etc.). There is also, however, a change in the growth of the silicon phase during chemical modification with sodium or strontium. This is evidenced by a 6 to 10 C (11 to 18 F) depression in the eutectic growth temperature, TG. The formation of silicon crystals in liquid aluminum-silicon melts is influenced by the high latent heat (and entropy) associated with freezing of silicon (Ref 86). This complicates the kinetics of the silicon crystals growing from the melt. In fact, depending on the impurities present and the solidification conditions, silicon can form an amazing range of structures. The papers by Fredriksson et al. (Ref 87) and Day and Hellawell (Ref 88) show the many chameleon-like forms that silicon can assume. A number of theoretical and experimental studies have been offered to explain the mysterious growth of silicon. A detailed review of this work is beyond the scope of this article, whose focus is on practical aspects of modification. Instead, the reader is referred to a number of key papers and reviews that provide an introduction to this field of inquiry (Ref 9, 11, 12, 55, 86–92). There are also directional solidification studies that show how the aluminum-silicon eutectic grows from liquid melts (Ref 49, 93, 94). Nearly 20 years ago, a committee of foundrymen (within the American Foundry Society, or AFS) considered the range of silicon structures commonly observed in commercial castings. They established a modification rating system (Ref 95). It is available as a wall chart that
may be purchased from AFS. A number from one to six is assigned, depending on the silicon structure observed. For the sake of discussion here, an abbreviated form of the rating system is given in Fig. 18. Each micrograph in Fig. 18 is for a polished, unetched sample of A356 alloy. The silicon phase has a dark-gray color. Iron-containing intermetallic phases are light gray, and aluminum is white. The following is a brief description of each rating. Nonmodified Structure (1). The fully nonmodified aluminum-silicon eutectic consists of large, acicular silicon flakes. This is the structure observed in normal commercial alloys when no modifiers are present. However, very rapid solidification can produce some refinement and a structure similar to rating number 2 or 3. Lamellar Structure (2). Small amounts of strontium produce some refinement of the large, acicular plates, so that fine lamellae of silicon are present in the eutectic. This structure is preferable to rating number 1, since sufficiently long solution heat treatments break up the lamellae and produce rounded silicon particles. Transitional Lamellar Structure (3). At slightly higher strontium levels, the lamellar plates of silicon become even finer and start to break down into smaller segments. The arrows in Fig. 18(3) indicate two plates that are starting to transform into an aligned row of individual fibers. The reason for this change in structure is perhaps best seen in studies of directionally solidified aluminum-silicon alloys (Ref 49, 93, 94). In nonmodified alloys, the silicon plates grow faster than aluminum, so that new cells of aluminum nucleate on the protruding silicon plates. As strontium is added, the growth of silicon is retarded, until the aluminum phase starts to overgrow the silicon. (The discussion by Major and Rutter in Ref 94 on this point is particularly lucid.) This overgrowth is starting to occur in this structure, as indicated by the arrows. The result is that some of the lamellar silicon plates are starting to break down into smaller, fibrous segments. Mostly Modified Structure (4). At this point, sufficient strontium has been added to completely transform the silicon from acicular/ lamellar forms into a modified, fibrous structure. However, some larger silicon crystals are still found in the structure, as indicated by the arrows in Fig. 18(4). The larger silicon crystals in this micrograph appear to have formed at a location where two eutectic grains have collided. Between two eutectic cells, the local growth rate will slow considerably, and when the strontium level is marginal, there will be a local loss of modification. (See the modification maps in Ref 49, 53, and 55 for a more detailed explanation.) Fully Modified Structure (5). This is the fibrous silicon that is normally considered to be the fully modified structure. Super modified Structure (6). On rare occasions, an extremely fine, fibrous eutectic was observed. The individual fibers are so small that they cannot be seen in an optical
micrograph. (Figure 18(b) is shown at 800 magnification.) They are, however, distinguishable with a scanning electron microscope at higher magnifications. Overall Ratings. It is fairly common to find a mix of structures in a casting. The usual way to assign a rating to a mixed structure is to multiply the area percentage of each structure by the rating number and average the values. For example, in a casting that is approximately 20% rating 3 and 80% rating 4, the overall rating for the casting would be: 0:2 3 þ 0:8 4 ¼ 3:8 Similarly, a casting that is 50% structure 4 and 50% structure 5 would have an overall modification rating number equal to 4.5.
Effect of Other Elements A number of elements have a significant effect on the silicon structure, even when present in small quantities. These effects are not usually of concern in primary alloys. For the foundry using secondary alloys, however, it is important that impurity elements be considered. Each element is briefly considered as follows. Phosphorus is present in most primary alloys at a level of 5 to 10 ppm. Even this very small amount has a significant effect on the silicon structure in foundry alloys. Studies have been made with very high-purity alloys (Ref 4, 96). A zone-refined alloy containing 0.6 ppm P exhibited a refined eutectic structure similar to a strontium-modified alloy (Ref 4), whereas an alloy containing 6 ppm P had the normal nonmodified structure shown in Fig. 18(1). Directional solidification studies by Major and Rutter (Ref 94) showed that phosphorus changes the relative growth kinetics of aluminum and silicon during eutectic solidification and the morphology of the solid/liquid interface. A number of other researchers have considered the effect of phosphorus in aluminum alloys (Ref 4, 22, 40, 43, 96–100). It is well known that phosphorus poisons the modifying effect of strontium (Ref 4, 22, 97, 100) and sodium (Ref 4, 99, 100). The effect of phosphorus on strontium modification is shown in Fig. 19. This figure is a map of the strontiumphosphorus interaction in A356 alloy sand castings. The “F” represents the fibrous or fully modified structure. The “L” represents the region where lamellar silicon is found, and the “A” stands for acicular, nonmodified silicon. The results in this figure indicate that, for each 10 ppm increase in phosphorus, one must increase the strontium by 30 ppm to offset the poisoning effect of phosphorus. Antimony additions will produce a lamellar structure similar to the one indicated in Fig. 18 (2). This was the basis for the Calypso series of antimony-modified foundry alloys developed by Pechiney approximately 30 years ago (Ref 4, 25). Antimony does not increase porosity, as strontium and sodium sometimes do (Fig. 5).
250 / Molten Metal Processing and Handling
Fig. 18
Fig. 19
Modification rating system proposed by the American Foundry Society. See text for explanations of rating system. Source: Ref 95
Strontium-phosphorus interaction in sand-cast A356 alloy. F, fibrous or fully modified structure; L, lamellar silicon region; A, acicular, nonmodified silicon region. Source: Ref 100
For many years, antimony was added to aluminum-silicon casting alloys by French and Japanese wheel makers to produce a partial refinement (lamellar structure). For this reason, antimony sometimes turns up in secondary alloys made from scrap wheels. Antimony resides underneath phosphorus in group VA of the periodic table, so its chemical properties are similar in many respects. A number of researchers have documented that antimony poisons the modification effect of both strontium and sodium (Ref 100–105). The deleterious effects of antimony can often be counteracted with increased levels of strontium, as shown in Fig. 20. Unfortunately, once antimony is added, it remains in the scrap cycle. It cannot be removed by oxidation with oxygen or chlorine. Improved properties and shorter solution times are possible with strontium. The
widespread use of rotary impellor degassers now makes it easy to obtain a low hydrogen content in melts. Antimony is a heavy metal, and many are concerned about possible health risks associated with its use in the foundry. For all of these reasons, antimony is now rarely used. Bismuth is an element that is added to freemachining alloys. This scrap sometimes finds its way into secondary alloys such as 319. Because it is also a group VA element, it poisons the effect of strontium modification when present in significant quantities (Ref 97, 106, 107). Researchers and foundrymen agree that the strontium-bismuth ratio must be greater than 0.45 in order to produce a fully modified silicon eutectic in 319 alloy. Magnesium. When magnesium is added to aluminum-silicon alloys, there is a coarsening of the silicon eutectic (Ref 4, 108). This is perhaps because magnesium changes the phase relationships to produce a eutectic mushy zone. Boron. Some researchers have experimented with boron additions to grain-refine aluminumsilicon alloys. Boron forms the compound AlB2, which is an effective nucleant in foundry alloys, but boron also forms the compound SrB6. Thus, boron additions interfere with strontium modification (Ref 109). For this reason, grain refinement is best accomplished by additions of TiB2, via Al-Ti-B master alloys. Calcium. There have been many studies on the use of calcium as a modifier (Ref 41, 44, 62, 106, 110, 111), either alone or in combination with sodium or strontium. When calcium is added to strontium-modified melts, there is a coarsening of the silicon eutectic, similar to that which occurs with magnesium or excess quantities of strontium. The action of calcium is similar to strontium as a modifier, except calcium does not work as well at slow cooling rates. Calcium changes the surface oxide when dissolved in aluminum melts, making the metal extremely susceptible to gas pickup. Calcium sometimes finds its way into the foundry via silicon additions used to make up the alloy. It is best to keep calcium at less than approximately 30 ppm to avoid excessive gas pickup. When necessary, calcium is easily removed by additions of reactive salts or by degassing with a few percent of chlorine in the inert gas. Salts and chlorine also remove strontium, however, so it is best to first remove calcium, skim the melt to remove any reactive salts, and then make the strontium master alloy addition.
Summary In writing this article, the author was amazed by the volume of information available in the technical literature. However, there is a relative paucity of information in some areas, with possible future advances in two areas of modification technology. For these reasons, a few suggestions are made for areas where future research and development appear to be useful.
Modification of Aluminum-Silicon Alloys / 251 These nucleants contained titanium and vanadium. They then added titanium, vanadium and boron (in unspecified proportions) and were able to reduce the eutectic grain size. (Note Fig. 6 in Ref 51.) This is an exciting result. If it could be translated into commercial practice, it would be the most significant advance in modification technology since strontium was introduced more than 30 years ago. It is to be hoped that this work will continue.
500 F
450
Acicular (A) 400 Lamellar (L)
Sr analysis, pmm
350
Partial modification (Pm) Fibrous (F)
PM
300
ACKNOWLEDGMENT The author gratefully acknowledges Alcoa Primary Metals, whose support has made this study possible.
350 A
200
REFERENCES
150 L + PM 100
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0
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Sb analysis, pmm
Fig. 20
Strontium-antimony interactions in A356 alloy. Source: Ref 102
There have been studies on the effect of phosphorus or strontium additions on the formation of iron-containing compounds in foundry alloys (Ref 112–115). Unfortunately, the available information is conflicting and contradictory. Since iron is always a problem in aluminum castings, it would be worth exploring any method that offers promise of controlling or limiting the loss of elongation and tensile strength normally associated with higher iron contents. A detailed fundamental study would be extremely useful. Perhaps the best way to develop better information would be with directional-solidification experiments. Careful thermal analysis studies may also be helpful, and studying the effect of freezing rate on particle formation would also be informative. In Fig. 18(6), a supermodified structure was shown. If it were possible to produce a structure like this consistently in commercial casting processes, this would be a significant advance in casting technology. The mechanical properties of the castings should be superior to anything that is produced today, and it may allow for further reduction of solution heat treatment times. Two recent laboratory studies suggest ways in which this may be accomplished. The first is the excellent study of aluminum-silicon alloys produced by semisolid metal processing (Ref 18). In melts that were subjected to stirring during solidification, an extremely fine supermodified eutectic was produced. In the second
study (Ref 116), ultrasound produced the same structure. This question arises: What level of vibration or ultrasound (or metal movement) is required to produce a supermodified structure? It would be worth trying to answer this question. It also raises profound questions about the mechanisms operating to produce the modified structure. There was only one paper (Ref 117) that studied the effect of modification practice on fatigue properties. This study suggested that best fatigue properties are obtained when a grain refiner is used in combination with strontium. In view of the importance of fatigue life, further study appears to be justified. Likewise, there was only one paper on machinability (Ref 118). This paper suggested that sodium and strontium modification improves surface finish and reduces tool wear. An increase in shear angle was recommended in modified alloys for best results. Machinability is another subject that merits further study. Finally, there is the foundryman’s dream to somehow modify aluminum-silicon alloys without the undesirable increase in porosity. This may be possible. Researchers at the Cooperative Research Center for Cast Metals Manufacturing (CAST) in Australia have been studying nucleants for the silicon eutectic in aluminum-silicon alloys (Ref 51). In nonmodified alloys, they found AlP, but in strontiummodified alloys, other nucleants were observed.
1. A. Pacz, U.S. Patent 1,387,900, Aug 16, 1921 2. R.W. Smith, Modification of Aluminum-Silicon Alloys, Solidification of Metals, Iron and Steel Inst., London, 1968, p 224–237 3. P.B. Crosley and L.F. Mondolfo, The Modification of Aluminum-Silicon Alloys, Mod. Cast., March 1966, p 89–100 4. S. Bercovici, “Control of Solidification Structures and Properties of Al-Si Alloys,” Paper presented at the 45th International Foundry Congress (Budapest), Oct 1978 5. R.J. Kissing and J.F. Wallace, Refinement of Aluminum-Silicon Alloys, Foundry, Vol 91 (No. 4), April 1963, p 74–79 6. S.-Z. Lu and A. Hellawell, Modification of Al-Si Alloys: Microstructure, Thermal Analysis, and Mechanisms, J. Met., Feb 1995, p 38–40 7. G.K. Sigworth, Theoretical and Practical Aspects of the Modification of Al-Si Alloys, AFS Trans., Vol 91, 1983, p 7–16 8. M. Adachi, Modification of Hypoeutectic and Eutectic Al-Si System Casting Alloys, J. Jpn. Inst. Light Met., Vol 34 (No. 6), 1984, p 361–373 9. J.E. Gruzleski and B. Closset, The Treatment of Aluminum-Silicon Alloys, American Foundry Society, 1990, p 25–105 10. L. Backerud, G. Chai, and J. Tamminen, Solidification Characteristics of Aluminum Alloys—Volume 2: Foundry Alloys, AFS/ Skanaluminium, 1990, p 25–38 11. J.E. Gruzleski, The Art and Science of Modification: 25 Years of Progress, AFS Trans., Vol 100, 1992, p 673–683 12. M.M. Makhlouf and H.V. Guthy, The Al-Si Eutectic Reaction: Mechanisms and Crystallography, J. Light Met., Vol 1, 2001, p 199–218 13. M. Adachi, Modification of Hypereutectic Al-Si System Casting Alloys, J. Jpn. Inst. Light Met., Vol 34 (No. 7), 1984, p 430–436 14. A.P. Bates, Hypereutectic Aluminum-Silicon Alloys: A Review of Published Information, Metallurgia, Vol 61 (No. 364), Feb 1960, p 70–79
252 / Molten Metal Processing and Handling 15. J.L. Jorstad, The Hypereutectic Al-Si Alloy Used to Cast the Vega Engine Block, Mod. Cast., Vol 60, Oct 1971, p 59–64 16. Y.P. Telang, Process Variables in Al-21 Si Alloys Refinement, Mod. Cast., Vol 43 (No. 5), 1963, p 232–240 17. J.L. Jorstad and D. Apelian, Hypereutectic Al-Si Alloys: Practical Processing Techniques, Die Casting Eng., Vol 48 (No. 3), May 2004, p 50–58 18. S. Nafisi and R. Ghomashchi, Effects of Modification during Conventional and Semi-Solid Metal Processing of A356 AlSi Alloy, Mater. Sci. Eng. A, Vol 415, 2006, p 273–285 19. M. Helmanova, The Influence of Solidification Rate on the Spatial Distribution of Eutectic Silicon in Technical Silumins, Prakti. Metallogr., Vol 12, 1975, p 59–68 20. B.M. Thall and B. Chalmers, Modification in Aluminum-Silicon Alloys, J. Inst. Met., Vol 77, 1950, p 79–97 21. Z. Poniewierski, Effect of the Type of Modification on the Microstructure of Al-Si Eutectic, Rev. Int. Hautes Temp. Re´fract., Vol 14, 1977, p 253–260 22. F. Fommei, Modification Treatments of AlSi Alloys, Aluminio, Vol 46, 1977, p 121–135 23. J. Gobrecht, The Influence of Alloying Elements on the Duration of Modification of Na and Sr in Al-Si Cast Alloys, Giesserei, Vol 65, 1978, p 158–164 24. T.J. Hurley and R.G. Atkinson, Effects of Modification Practice on Aluminum A356 Alloys, AFS Trans., Vol 93, 1985, p 291– 296 25. M. Garat, G. Laslay, S. Jacob, P. Meyer, P.H. Gaerin, and R. Adam, State of the Art Use of Sb-, Na- and Sr-Modified AlSi Casting Alloys, AFS Trans., Vol 100, 1992, p 821–832 26. J.J. Jorstad and W.M. Rasmussen, Aluminum Casting Technology, 2nd ed., American Foundry Society, 1993, p 76–77 27. S. Shivkumar, S. Ricci, B. Steenhoff, D. Apelian, and G. Sigworth, An Experimental Study to Optimize the Heat Treatment of A356 Alloy, AFS Trans., Vol 97, 1989, p 791–810 28. S.A. Levy, Techniques for Improving Ductility of Low Pressure Die Cast Aluminum Wheels, AFS Trans., Vol 86, 1978, p 135–138 29. F. Paray and J.E. Grusleski, Factors to Consider in Modification, AFS Trans., Vol 102, 1994, p 833–842 30. M. Drouzy, S. Jacob, and M. Richard, Interpretation of Tensile Results by Means of a Quality Index, AFS Int. Cast Met. J., Vol 5, 1980, p 43–50 31. P.Y. Zhu, Q.Y. Liu, and T.X. Hon, Spheroidization of Eutectic Silicon in Al-Si Alloys, AFS Trans., Vol 91, 1983, p 609–614 32. D.L. McLellan and M.M. Tuttle, Aluminum Castings—A Technical Approach, AFS Trans., Vol 91, 1983, p 243–252
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Modification of Aluminum-Silicon Alloys / 253
63.
64.
65.
66. 67.
68.
69.
70.
71.
72.
73.
74.
75.
76.
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77. M. Djurdjevic, T. Stockwell, and J. Sokolowski, Effect of Strontium on Microstructure of Al-Si and Al-Cu Eutectics in 319 Aluminum Alloy, Int. J. Cast Met. Res., Vol 12, 1999, p 67–73 78. B. Closset and S. Kitaoka, Evaluation of Sr Modifiers for Al-Si Casting Alloys, AFS Trans., Vol 85, 1987, p 233–240 79. M.O. Pekguleryuz and J.E. Gruzleski, Conditions for Strontium Master Alloy Additions to A356 Melts, AFS Trans., Vol 96, 1988, p 55–64 80. B. Closset and J.E. Gruzleski, Mechanical Properties of A356 Alloy Modified with Pure Strontium, AFS Trans., Vol 90, 1982, p 453–464 81. R. DasGupta, C.G. Brown, and S. Marek, Analysis of Overmodified 356 Aluminum Alloy, AFS Trans., Vol 96, 1988, p 297–310 82. J. Gruzleski, N. Handiak, H. Campbell, and B. Closset, Hydrogen Measurement by Telegas in Sr-Treated A356 Melts, AFS Trans., Vol 94, 1986, p 147–154 83. F.C. Dinayuga, N. Handiak, and J.E. Gruzleski, The Degassing and Regassing Behavior of Sr-Modified A356 Melts, AFS Trans., Vol 96, 1988, p 83–88 84. X.-G. Chen, F.J. Klinkenberg, R. Ellerbrok, and S. Engler, Efficiency of Impeller Degassing and Regassing Phenomena in Aluminum Melts, AFS Trans., Vol 102, 1994, p 191–197 85. G.K. Sigworth, C. Wang, H. Huang, and J.T. Berry, Porosity Formation in Modified and Unmodified Al-Si Alloy Castings, AFS Trans., Vol 102, 1994, p 245–261 86. B. Gallois and G.K. Sigworth, An Analysis of Silicon Eutectic Modification, Thermal Analysis of Molten Aluminum, American Foundry Society, 1985, p 101–119 87. H. Fredriksson, M. Hillert, and N. Lange, Modification of Aluminum Silicon Alloys by Sodium, J. Inst. Met., Vol 101, 1973, p 285–299 88. M.G. Day and A. Hellawell, The Microstructure and Crystallography of Aluminium-Silicon Eutectic Alloys, Proc. R. Soc. (London) A, Vol 305, 1968, p 473–491 89. S.-Z. Lu and A. Hellawell, Mechanisms of Silicon Modification in Aluminum-Silicon Alloys: Impurity Induced Twinning, Metall. Trans. A, Vol 18, 1987, p 1721–1731 90. A. Hellawell, How Does a Modifier Work?, Proc. Second International Conference on Molten Aluminum Processing, American Foundry Society, 1989, p 9–1 to 9–17 91. R. Elliott, Eutectic Solidification Processing, Butterworth, London, 1983 92. R. Elliott, Growth Kinetics in Aluminum Silicon Eutectic Alloys, Proc. Second International Conference on Molten Aluminum Processing, American Foundry Society, 1989, p 10–1 to 10–24 93. L. Clapham and R.W. Smith, Partial Modification in Unidirectionally Solidified Al-Si Eutectic Alloys, Acta Metall., Vol 37 (No. 1), 1989, p 303–311
94. J.F. Major and J.W. Rutter, Effect of Strontium and Phosphorus on Solid/Liquid Interface of Al-Si Eutectic, Mater. Sci. Technol., Vol 5 (No. 7), 1989, p 645–656 95. G.K. Sigworth, Determining Grain Size and Eutectic Modification in Aluminum Alloy Castings, Mod. Cast., Vol 77, July 1987, p 23–25 96. S.D. McDonald, K. Nogita, and A.K. Dahle, Eutectic Nucleation in Al-Si Alloys, Acta Mater., Vol 52, 2004, p 4273–4280 97. C.R. Loper, Jr., H.-G. Seong, and J.-I. Cho, Interaction of Phosphorus and Bismuth (with Sr) in A356.2 Alloy, AFS Trans., Vol 109, 2001, p 1–13 98. C.R. Loper, Jr. and J.-I. Cho, Influence of Trace Amounts of Phosphorus in Aluminum Castings, AFS Trans., Vol 108, 2000, p 667–672 99. B. Gu¨nther and H. Ju¨rgens, Influence of Minor Impurities on Mechanical Properties, Structure and Solidification Form of Eutectic on Near-Eutectic Al-Si Casting Alloys, Giesserei, Vol 67, 1980, p 8–13 100. S. Bercovici, M. Garat, and R. Scalliet, “A Review of Recent French Casting Alloy Developments,” Report 13691, Pechiney Ugine Kuhlmann, April 1978 101. N. Handiak, J.E. Gruzleski, and D. Argo, Sodium, Strontium and Antimony Interactions During Modification of AS7G03 (A356) Alloys, AFS Trans., Vol 95, 1987, p 31–38 102. E.N. Pan, Y.C. Cherng, C.A. Lin, and H.S. Chiou, Roles of Sr and Sb on Silicon Modification of A356 Aluminum Alloys, AFS Trans., Vol 102, 1994, p 609–629 103. B.L. Tuttle, D. Twarog, and E. Daniels, The Effect of Trace Amounts of Antimony on the Structure and Properties of Aluminum Alloy A356.2, AFS Trans., Vol 99, 1991, p 7–16 104. W. Wang and J.E. Gruzleski, Interactive Effects During Sodium or Strontium Treatment of Antimony-Containing A356 Alloy, AFS Trans., Vol 98, 1990, p 227–234 105. Incompatibility of Sodium and Antimony in the Modification of Al-Si Alloys, Fondeur, Vol 272, 1976, p 12–13 106. S. El-Hadad, A.M. Samuel, F.H. Samuel, H.W. Doty, and S. Valtierra, Effect of Bi and Ca Additions on the Characteristics of Eutectic Silicon Particles in Sr-Modified 319 Alloys, Int. J. Cast Met. Res., Vol 15, 2003, p 551–564 107. C.J. Machovec, G.E. Byczynski, J.W. Zinel, and L.A. Godlewski, Effect of BiSr Interactions on Si Morphology in a 319 Type Aluminum Alloy, AFS Trans., Vol 108, 2000, p 439–444 108. G. Heiberg et al., Effect of Mg, Fe and Cu on Eutectic Solidification of Hypoeutectic Al-Si Alloys, AFS Trans., Vol 110, 2002, p 347–358 109. B. Golhahar, A.M. Samuel, F.H. Samuel, H.W. Doty, and S. Valtierra, Effect of Grain Refiner-Modifier Interactions on
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ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 255-262 DOI: 10.1361/asmhba0005302
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Grain Refinement of Aluminum Casting Alloys* G.K. Sigworth and T.A. Kuhn, Alcoa Primary Metals
GRAIN REFINEMENT of aluminum was first accomplished in the early 1930s by making titanium additions to the melt (Ref 1–4). The improvement in castability was considerable, and this quickly became the established practice. The resulting change in the structure of the casting is remarkable. This may be seen by cutting a piece from the casting and polishing the cut surface with a series of progressively finer sandpapers, until a mirror-smooth surface is obtained. After the polished surface is etched with acid, the underlying grain structure is revealed. An example of the transformation obtained by grain refinement is shown in Fig. 1. Grain refinement has a significant effect on the mechanical properties of castings, primarily because the distribution of second-phase materials is changed. Consider the casting shown in Fig. 1. The nonrefined casting (top in Fig. 1) has a small area near the chill where equiaxed grains were formed, but the rest of the metal has elongated, feathery grains. All brittle intermetallic phases and any porosity formed will be located between these large grains, so the elongation normal to the grains will be poor. The material shown at the bottom of Fig. 1 has a much better structure. The grain size is small (or fine, which is the basis for the term refinement). Hence, mechanical properties are isotropic, and the material is stronger. As is shown subsequently, grain refinement in casting alloys tends to reduce the amount of porosity and the size of the pores. This improves mechanical properties, especially fatigue strength. Feeding is also improved. It is for these reasons that most cast aluminum is grain refined.
procedure is more convenient with large samples, for example, when cutting through the middle of a casting. The sample is placed on a milling machine, and metal is cut away until a flat surface is obtained. Then, if a final light cut with a sharp tool is made, the resulting surface may be etched directly without polishing. For aluminum alloys low in copper, the test sample is placed in hot, nearly boiling water for a few minutes. When the sample is hot, it is removed from the water, and the surface to be etched is immersed briefly in a solution of Poulton’s etch. (This etchant contains the following ingredients: 60% HCl, 30% HNO3, 5% HF, and 5% H2O.) This acid is held at room temperature. Alternatively, the surface may be etched by wiping it with a piece of cotton wool or absorbent cloth that has been soaked in the etchant. For alloys that contain copper, it is more difficult to determine the grain size. One procedure that works reasonably well is to etch the sample at room temperature in a solution of 10% HF. This results in a black smudge on the sample, produced by the copper in the alloy. This smudge may be cleaned off by rubbing the
Average intercept distance Calculated average grain diameter ASTM grain size number Grains per unit area
Table 1 gives values in different units of measurement, for the sake of comparison. The easiest procedure is usually to measure the average intercept distance. All grain sizes in this article are reported in this way. The units employed are microns (mm, or 10 6 m).
Mechanisms of Grain Refinement Grain refinement dates from the early 1930s, when foundrymen began to use titanium additions to improve the structure of their castings. To understand this action of titanium, one must
Chill
Measurement of Grain Size To determine grain size, one obtains a sample and produces a smooth surface on one side. Thismay be done in the laboratory using standard metallographic techniques. However, another
surface of the sample under running water or by rinsing it in a dilute solution of nitric acid. There are various procedures established to measure grain size on the etched samples. These are outlined in ASTM E 112. The grain structure can be reported in one of several units:
Chill
Fig. 1
Grain structure in 3004 alloy castings. Top: no refinement; bottom: 10 ppm boron added as 5Ti-1B alloy
*Republished with the permission of American Foundry Society (www.afsinc.org). Copyright 2007 from the following: International Journal of Metalcasting (ISSN 1939–5981), Vol 1 (No. 1), 2007, p 31–40, and AFS Transactions (ISBN 978-0-87433-338-1), Vol 115, 2007, p 177–188
256 / Molten Metal Processing and Handling Table 1
Comparison of various grain size measurements (ASTM E 112)
1
Average intercept distance in.
mm
mm
Average diameter, mm
ASTM grain No.
Grains/cm2
Grains/in.2
0.008 0.015 0.031 0.047 0.079
0.2 0.4 0.8 1.2 2.0
200 400 800 1200 2000
252 504 1008 1512 2520
14.5 12.5 10.5 9.5 8
1890 496 112 56 20
12,200 3200 724 362 128
2
3 1000
1830
900
1650 L
4
L + TiAl 3
Temperature, ⬚C
700
1290
665
0.15
1.15
660.1 α + TiAl 3
α
600
Temperature, ⬚F
1420
800
5
1110
930
500
6 400
900 0
0.5
1.0
1.5
2.0
Titanium, wt%
Fig. 2
The aluminum-rich side of the aluminum-titanium phase diagram. Source: Ref 5
7 of the base alloy. Hence, the first nucleation (or formation) of solid aluminum will be at the surface of the aluminide particle, as shown by “3”. The aluminum crystal then grows around the surface of the aluminide (“4” and “5”). In the process, it consumes the dissolved titanium in the vicinity of the particle, and growth stops. As the metal cools further, dendritic growth begins (“6”) and continues (“7”) as solidification proceeds. This course of events may be confirmed in two ways. The first is to pour a sample of metal into a small cup and record the temperature during solidification with a thermocouple. The resulting cooling curves of refined and nonrefined metal are shown in Fig. 4. When titanium is added, the nucleation temperature, Tn, is often above the growth temperature, Tg. When no titanium is present, the metal cools several degrees below the growth temperature before nucleation occurs. Another confirmation of the mechanism proposed in Fig. 3 may be found by a careful examination of the microstructure of aluminum castings. When anodized in the proper way and viewed under polarized light, the titanium-rich areas corresponding to “6” and “7” of Fig. 3 are revealed. An example is presented in Fig. 5 (Ref 8). The darker areas located in the
Fig. 3
TiAl3 nucleates aluminum grain. Source: Adapted from Ref 6
1 Grain refiner added
2 3,4 Temperature
consider the aluminum-titanium phase diagram (Ref 5), shown in Fig. 2. The important thing to note is that titanium raises the melting point of aluminum. Pure aluminum melts at 660.1 C (1220.2 F); adding titanium increases the melting point to 665 C (1229 F). This effect has important ramifications for the nucleation and growth of solid aluminum grains. The best explanation was given in 1983 by Backerud (Ref 6). Figure 3 is taken from his paper and shows schematically what happens. Titanium additions are made via aluminum master alloys, which contain between 5 and 10% Ti. These materials contain numerous crystals of the titanium aluminide compound TiAl3. (Pictures of aluminide crystals are given in Ref 7.) When a master alloy is added, typically a few minutes before casting, millions of these microscopic particles are released into the melt. The presence of one is indicated by “1” in Fig. 3. When the TiAl3 crystal comes in contact with liquid aluminum, it starts to dissolve. This means the liquid metal at the surface of the particle becomes enriched in titanium, as shown by “2” in Fig. 3. From the relationships shown in the aluminum-titanium phase diagram, the titanium-rich metal in contact with the aluminide can begin to solidify at a temperature that is above the melting point
5 Tn
6
7
Tg Tn No grain refiner
Time
Fig. 4
Cooling curves during solidification. Tn, nucleation temperature; Tg, growth temperature
Grain Refinement of Aluminum Casting Alloys / 257
100 mm
Fig. 5
Titanium-rich areas in dendrites of Al-0.06%Ti alloy. Source: Ref 8
10,000
100
5000 2000
200
1000 400
500
1 2
Grain size, mm
Number of grains per cm3
20,000
3
200 100
800 0
0.10
0.20
0.30
Titanium, wt% 1 2 3
Fig. 6
Delamore and Smith Grebe and Grimm Crossley and Mondolfo
Refinement of 99.7% Al by titanium. Source: Ref 9
center of the grains were shown by microprobe measurements to be high in titanium. The aforementioned is a clear explanation of how titanium acts to grain refine aluminum, but there is an additional factor that must be considered: the dissolution of TiAl3 particles. A normal addition level of titanium is approximately 100 ppm (0.01% Ti). When this small amount of titanium is added to relatively pure aluminum at normal casting temperatures, it dissolves. When the added particles have dissolved, the grain-refining effect disappears. This loss of refining ability with time is called fade. The fade time depends on a number of factors. The most important are the titanium content of the metal, the metal temperature, and the type of master alloy used. The latter determines the TiAl3 particle size. Aluminum-titanium rod products tend to have smaller particles and dissolve faster—typically in 10
to 15 min. Waffle products are cast from a higher temperature and have larger aluminide particles. These may take 30 to 40 min to dissolve. The practical consequence in the foundry is relatively simple: If the titanium in the metal is significantly less than the solubility (0.15% Ti at 665 C or 1229 F), then it is not possible to grain refine consistently. Results from three early studies of grain refinement in high-purity aluminum are shown in Fig. 6 (Ref 9). The best refinement was obtained when the titanium added was more than 0.15%, the solubility limit of TiAl3 in aluminum. The picture painted thus far is fairly attractive and relatively easy to understand. This understanding of grain refinement has been around since 1930. During that time, it has entered deeply into the consciousness of foundrymen, much like Michelangelo’s Mona Lisa has entered the mainstream of consciousness forming our popular culture. For this reason, it needs to be well understood. This picture is essentially the source of, and basis for, most of the grain-refinement practices used today (2008) in aluminum casting alloys. The maximum limits for titanium are quite high in most alloys, and, in some cases, a minimum is specified, for example, 201 alloy. It is vital to grain refine this alloy to prevent hot cracking during casting. The specification for 201 alloy published by the Aluminum Association permits up to 0.35% Ti, and a minimum of 0.15% Ti is called for. Many of the other 2xx and 7xx casting alloys also have a minimum specified (0.10 or 0.15% Ti). With a few exceptions, aluminum casting alloys permit a maximum of at least 0.20% Ti. The reason for these large quantities of titanium may as well be engraved in stone. As shown subsequently, it should be a tombstone. In other words, the epitaph for this 75-year-old grain-refining practice should read:
High quantities of titanium are necessary to prevent dissolution of aluminum-titanium refiners. From an historical point of view, the aforementioned accurately describes the first 40 years of aluminum grain refinement. Then, in the early 1970s, the first modern Al-Ti-B grain refiners became available. The role of boron was studied in the laboratory in the early 1950s (Ref 10, 11). This work showed that boron in combination with titanium gave much better grain refinement, but the use of boroncontaining salts in the casthouse was an extremely messy and haphazard business. Finally, in the late 1960s, Kawecki discovered how to produce boron-containing master alloys by reacting a mixture of KBF4 and K2TiF6 with molten aluminum. The result was the first commercial Al-Ti-B master alloy—a fast-acting rod product that could be fed continuously into the metal transfer launder or holding furnace by wire feeders. The new boron-containing grain refiners were much more powerful and economical. They therefore replaced the use of the earlier aluminum-titanium products in most aluminum alloys and still form the basis for accepted commercial practices to this day. How do the modern refiners differ from the old ones? There are many differences, but the most important is this: Most of the titanium in modern Tibor refiners is present as TiB2 particles, and titanium diboride is virtually insoluble in molten aluminum. The consequence: High quantities of titanium are NOT necessary to prevent dissolution of Al-Ti-B grain refiners. However, that is only part of the story. In addition to low solubility, borides are extremely effective nuclei in foundry alloys. In fact, boron is better than titanium as a grain refiner for aluminum casting alloys.
Boron as a Grain Refiner In 1981, Lu, Wang, and Kung published a study of the grain refining of A356 alloy in the Journal of the Chinese Foundrymen’s Association (Ref 12). The results shown in Fig. 7 are taken from their study. Three different master alloys were added to A356 alloy: Al-5%Ti, Al-5%Ti-1%B, and Al-4%B. The results clearly showed that boron is more powerful than titanium as a grain refiner. This result ran completely counter to the prevailing viewpoint described previously. Everyone believed that titanium was needed for grain refinement, but the Chinese study showed titanium was only a weak grain refiner. It also suggested that boron is much better than titanium. Intrigued by these results, Sigworth and Guzowski (Ref 13, 14) began their own study. Laboratory tests confirmed that aluminumboron master alloys containing AlB2 particles gave excellent grain refinement. They then went to Stahl Specialty Company in Missouri, where foundry trials confirmed that aluminum-
258 / Molten Metal Processing and Handling With this understanding, it is now possible to define an optimized grain-refining practice for each of the common casting alloys.
2000
Best Practices for Grain Refinement Grain size, mm
1500
Before defining the best grain-refinement practices, it is necessary to clearly distinguish between the two forms of titanium.
Al-5%Ti 1000
Forms of Titanium 5%Ti-1%B 500 Al-4%B
0
0
0.02
0.04
0.06
0.08
0.10
Addition level, %Ti or %B
Fig. 7
Grain refining of A356 alloy by three master alloys. Source: Ref 12
1000 A356 alloy No TiB 2
Grain size, mm
800 600 400
10 ppm B 200 0
0
0.02
0.04
0.06
0.08
0.10
0.12
Soluble titanium, wt%
Fig. 8
Grain size as a function of titanium content. Source: Ref 16
boron master alloys were good refiners. However, the added boron reacted with dissolved titanium to produce a sludge in holding furnaces. Also, boron apparently reacted with strontium, since boron additions resulted in a loss of modification in some heats modified with strontium. A more recent study has confirmed the interaction between strontium and boron (Ref 15). This presented an interesting problem. If one compares the nucleation potency and other characteristics of the different crystals in Al-7%Si alloys, one finds: TiAl3 crystals are poor nuclei. TiAl3 crystals
also have a relatively high solubility in aluminum. For both reasons, a large amount of titanium must be added to produce consistently small grain sizes. TiB2 particles are excellent nuclei. Titanium diboride has almost no solubility in liquid aluminum. Thus, TiB2 particles produce good
refinement at small addition levels. The refinement is also long lasting when the particles are not allowed to sediment from the melt. AlB2 is the best nucleus. It produces the smallest grain size. However, AlB2 dissolves readily in aluminum, where it reacts with titanium and strontium in the melt. Longterm use would produce an undesirable sludge in furnaces. Thus, in spite of its potency as a grain-refining nucleant, AlB2 cannot be used in the foundry. Would it be possible to develop a grain refiner with the potency of AlB2 and without the problems associated with boron additions? After reflecting on this question for some time, Sigworth and Guzowski produced grain refiners containing a mixed boride (Ref 13, 14). The compounds AlB2 and TiB2 have the same crystal structure, and x-ray studies showed the two crystals form a continuous solid solution. Thus, a mixed boride has a composition intermediate between AlB2 and TiB2 and can be represented by the formula (Al,Ti)B2. Mixed borides appear to have a nucleation potency close to AlB2, with a low solubility similar to TiB2. Thus, the problems associated with straight boron additions were avoided. The first mixed boride refiner was an Al2.5Ti-2.5B alloy. More recently, it has been produced as an Al-1.7%Ti-1.4% master alloy, marketed under the name Tibloy. This is an excellent refiner but somewhat slow acting, requiring approximately 5 min to produce the best grain size. In 356 alloy, its performance does not usually justify the additional cost, but in 319 alloy, Tibloy refines significantly better than Al-5Ti-1B. (See Fig. 17 and 18 in Ref 13.) One could summarize the previous discussion by saying: Boron is the most powerful grain refiner in aluminum casting alloys.
Soluble Titanium. This is the TiAl3 compound, which is present as 10 to 20 mm-sized particles in aluminum-titanium master alloys. It is also present in Al-Ti-B alloys that have an excess of titanium required to produce TiB2. (The stochiometric ratio is 2.22 wt% Ti for each 1 wt% B.) The TiAl3 compound dissolves quickly in alloys that contain less than approximately 0.15% Ti. As noted previously, this particle is a poor refiner in aluminum-silicon casting alloys. Insoluble Titanium. This is the titanium present as TiB2, which has such a low solubility in aluminum that, for all practical purposes, it can be considered to be insoluble. (TiC is another form of insoluble titanium. The carbide is also a good grain refiner, but because it offers no advantage in casting alloys, Al-Ti-C refiners are not considered further in this article.) In the past, the single word titanium has been used to describe both its soluble and insoluble forms. This terminology appears to be the principal reason for confusion by foundrymen when trying to understand grain refinement. Thus, to establish a linguistic foundation for an improved understanding, the following terminology is proposed: The word titanium shall refer to the amount of titanium dissolved in an alloy; the word boron shall refer to the quantity of insoluble boride particles present. With this terminology in mind, the question can be posed: What combination of titanium and boron gives the best grain refinement? The answer is this: It depends on the alloy family under consideration. For this reason, each is considered separately.
Practices Aluminum-Silicon Casting Alloys. The AlSi and Al-Si-Mg casting alloys, such as 356 and 357, offer excellent castability and provide a range of attractive mechanical properties. Hence, they find widespread commercial application. These alloys have also been studied extensively. The grain refining of the (copperfree) aluminum-silicon casting alloys has been reviewed in a separate publication (Ref 16). Figure 8 is taken from this study. These results are for grain refinement of an Al-7%Si-0.35% Mg (A356) alloy. This alloy was modified with approximately 0.02% Sr and grain refined with a 10 ppm boron addition of Al-5Ti-1B master alloy. Castings were also poured without the
Grain Refinement of Aluminum Casting Alloys / 259 500 A356 alloy
Grain size, μm
400 5% Boral
2.5Ti-2.5B
300
200 5Ti-1B 100
0
0
0.05
0.10
0.15
0.20
Soluble titanium, wt%
Fig. 9
A356 alloy grain size from Aluminum Association test. Source: Ref 17
10,000
Grain size, μm
%Ti 0.0 0.05
1000
0.10 0.15 0.20
300 0
Fig. 10
5 10 15 Boron added as 3Ti - 1B, ppm
20
Grain refinement of 319 alloy. Source: Ref 20
boron addition. All samples were cut from a plate casting at a distance of 35 mm (1.4 in.) from a chill, where the cooling rate was similar to that obtained in permanent mold cast wheels. Grain sizes were measured in accordance with ASTM E 112. The average and standard deviations of the grain sizes are plotted. It is readily seen that when TiB2 particles are added via commercial grain refiners, the grain size of the casting is essentially constant and independent of the dissolved titanium-content in the alloy. This result was also confirmed in other experimental studies. A plot of results obtained with the Aluminum Association grain size test is given in Fig. 9. Three different boron-containing grain refiners were employed. It appears a small amount of titanium (approximately 0.02%) is desired when using Tibloy or aluminum-boron master alloys, but this is not necessary for TiB2-containing refiners, such as Al-5Ti-1B.
These results clearly point to improved grainrefining practices for the aluminum-silicon casting alloys, such as 356 and 357. Best grain refinement is obtained when 10 to 20 ppm of boron are added in the form of Al-5Ti-1B or Al-3Ti-1B rod. With this boron addition, it is possible to reduce the large amounts of soluble titanium now added to these alloys. Lower titanium does not appear to result in any loss of mechanical properties or castability, as determined by grain size (Ref 16). Lower titanium contents also offer the advantage that the tendency for sludge formation is eliminated. Sludge can become a significant problem when the titanium content is held over 0.10% for long periods of time (Ref 18, 19). Aluminum-Silicon-Copper Casting Alloys. The grain refinement of 319 alloy was studied in detail by Pasciak and Sigworth (Ref 20). A number of castings and solidification conditions were studied. The grain size observed in test castings, whose freezing rate was similar to that found in permanent mold castings, has been plotted in Fig. 10. Both the (dissolved) titanium and (insoluble) boron contents were varied in these trials. For some reason, the presence of copper in the alloy completely changes the grain-refining response to titanium. In this case, it is beneficial to have a minimum of approximately 0.10% Ti dissolved in the alloy. Without this quantity of titanium, it was not possible to produce a small grain size. The best grain size is obtained when 10 to 20 ppm of boron are added. The boron addition can be made in the form of Al-5Ti1B or Al-3Ti-1B rod. A Tibloy alloy can also be used and may be cost-effective, but, in this case, at least 5 min should be allowed for the refiner to become active. Although they have not been studied, it is reasonable to presume that other copper-containing aluminum-silicon casting alloys (such as 355 and 308) will behave similarly to 319. For this reason, it would be advisable to have approximately 0.10% Ti present in these Al-
Si-Cu casting alloys before making an addition of 10 to 20 ppm of boron. Aluminum-Copper Casting Alloys. The first casting alloy used in the United States was an Al8%Cu composition, known as No. 12 alloy. In 1909, Alfred Wilm discovered that an Al-4.5% Cu-0.5%Mn alloy would strengthen by aging after a quench from an elevated temperature. This alloy was called Duralumin and formed the basis for the aluminum-copper family of alloys used today (2008). These alloys have excellent mechanical properties, high-temperature strength, and fatigue life. They are difficult to cast, however. Their long freezing range makes them extremely susceptible to hot cracking. It has long been well known that grain refinement reduces the tendency for hot cracking in net-shaped castings. The study that shows this most clearly is the 1970 paper by Davies (Ref 21), who found the total length of hot tears in a casting to be proportional to the cast grain size. In other words, a good grain-refining practice is essential for the production of crack-free 2xx alloy castings. The specifications for 2xx alloys date from the time when titanium was used as a grain refiner. For this reason, a minimum content of 0.15% Ti is specified for 201, 202, 203, 204, and 206 alloys. More recently, however, it has been shown that lowering the titanium content reduces the grain size of the casting and improves hot crack resistance (Ref 22). The best grain size and hot crack resistance are obtained when the titanium is less than 0.05% and 10 to 20 ppm of boron are added as Al-5Ti-1B or Al-3Ti-1B rod. Aluminum-Zinc-Magnesium Casting Alloys. A number of aluminum casting alloys are based on the ternary Al-Zn-Mg system. These alloys naturally age to high strengths at room temperature. A high-temperature solution and aging treatment is not required. Consequently, there is a potential cost savings that can be realized by eliminating the heat treatment, especially with complex shapes that would tend to distort when quenched from a high temperature. Castings from these alloys may also be welded or brazed immediately after casting, and the assemblies will age to high strength at room temperature. This combination of properties makes the alloys promising candidates for a number of applications. In spite of their excellent properties, the 7xx casting alloys are seldom used, mainly because of their propensity for hot cracking. As a consequence, the Aluminum Association chemical composition limits established for titanium in these alloys have relatively high upper limits; in three alloys (712, 771, and 772), a minimum titanium content (0.10 or 0.15%) is specified. A recent study has shown that low titanium contents in these alloys reduce the grain size and improve hot crack resistance (Ref 23). The best grain size and hot crack resistance are obtained when the titanium is 0.02 to 0.05% and 10 to 20 ppm of boron are added as Al-5Ti-1B or Al-3Ti-1B rod.
260 / Molten Metal Processing and Handling Table 2
Composition of aluminum-magnesium alloys tested
Table 3 Grain sizes observed in 535 alloys
Composition, wt% Alloy
Cr
Cu
Fe
Mg
Mn
Ni
Si
Ti
Zn
535 535L
0.0 0.0
0.004 0.004
0.15 0.15
6.98 7.04
0.18 0.18
0.002 0.002
0.05 0.05
0.196 0.036
0.0 0.0
1000 mm
Fig. 11
Grain structure in an A356 alloy casting
Aluminum-Magnesium Casting Alloys. The 5xx casting alloys provide extremely good corrosion resistance and moderate strength without heat treatment. They are commonly used in marine and architectural applications. The alloys are not easy to cast. The high magnesium content produces a thick, sticky oxide film on the melt, and castings are susceptible to hot cracking and shrinkage. The Aluminum Association chemical composition limits established for these alloys have relatively high upper limits for titanium; in three alloys (516, 520, and 535), a minimum content of 0.10% Ti is specified. Sigworth (Ref 24) conducted a study of the grain refinement of Al-7%Mg (535) alloy. Two alloys were prepared, having the same composition except for the soluble titanium. The chemistry of the two alloys is given in Table 2. To each alloy, an addition of 30 ppm boron was made in the form of Al-3%Ti-1%B master alloy. Grain size samples were taken by using the test specified by the Aluminum Association. The resulting grain size, as measured by the average intercept distance, is shown in Table 3. Since the statistical (1s) error associated with the determination of grain size is approximately 10%, for all practical purposes, the two alloys have the same grain size. Thus, it appears the soluble titanium has little effect in 5xx alloys. This result is similar to that found previously for Al-Si and Al-Si-Mg alloys, such as 356 and 357. In other words, it should be
possible to reduce the titanium content of aluminum-magnesium alloys without affecting castability when boron is used as a refiner.
Benefits of Grain Refinement Grain refinement tends to reduce both the amount of porosity and the size of the pores in casting alloys. The reason for this may be seen by careful examination of the structure in a casting (Fig. 11). In 356 alloy, the first 50 or 60% of the solidification process is taken up by the formation and growth of aluminum grains. Then, a film of aluminum-silicon eutectic forms around the aluminum dendrites. Gas porosity typically forms during the last 5 to 10% of the solidification process. Thus, gas must fit into the spaces available between the aluminum dendrites and solidified eutectic. (The pores are seen as black areas between dendritic grains in Fig. 11.) When the grain size is reduced, the size of the spaces available for pores is also reduced. The consequence is a smaller pore size. Figure 12 shows a plot of measured pore size in directionally solidified A356 alloy castings (Ref 25). By comparing curves 1 and 3 of this figure, one can see that grain refinement reduces the size of pores formed. (Rapid solidification also reduces pore size.) The result is an improved fatigue life for the casting. The other major benefit of grain refinement is improved feeding and a reduced tendency for
Alloy
Grain size, mm
535 535L
82 86
shrinkage formation in the casting. The mechanism for this effect was first established by Guocai Chai (Ref 26), who placed a slowly rotating paddle in a solidifying cylindrical casting. By measuring the torque on the paddle, he was able to measure the point during solidification when the semisolid “mush” started to jell, or became stiff. He called this transition point dendrite coherency. Coherency indicates the time when solid grains start to connect with each other to reduce feeding. Chai’s measurements showed that grain-refiner additions delay the onset of dendrite coherence. Figure 13 presents results obtained for an Al-4%Cu alloy solidified at two different cooling rates. Without grain refinement, dendrite coherency occurs at 25% solid; with grain refinement, coherency occurs at 50% solid. This means that a grainrefined riser behaves like a liquid nearly twice as long as a nonrefined riser. The measurements of Chai were later duplicated in other alloys by Arnberg and Backerud (Ref 27).
Concluding Remarks This review of grain refinement in aluminum casting alloys has shown that the large titanium additions commonly used, and often specified, are primarily an historical artifact. Titanium was used by itself for nearly 40 years, and with this practice, large additions are needed for good results. Now that powerful Al-Ti-B refiners are available, however, there is no longer any reason (with one exception) for large titanium additions. Lower titanium contents eliminate the tendency for sludge formation, and, in some alloys, a smaller grain size and improved resistance to hot cracking are found. The first key to an improved understanding of grain refinement is to realize that titanium can be present in two different forms: as soluble and insoluble particles. The soluble form is called titanium in this article. The insoluble form is called boron, since insoluble titanium is present as titanium diboride. The second key is to realize that boron (boride particles) is an extremely powerful nucleant in all casting alloys. Boron is much more effective than titanium. Thus, it makes more sense to say that grain refinement is accomplished with boron. It then becomes a question of which combination of titanium and boron gives the best results. The answer depends on the alloy system. A brief summary is offered as follows: Low titanium is best: The Al-Cu and Al-Zn-
Mg alloys give best results when the soluble
Grain Refinement of Aluminum Casting Alloys / 261 As noted previously, titanium diboride particles are virtually insoluble in molten aluminum. They are more dense than the metal, however, and will settle from the melt. Sedimentation can be rapid if there is no stirring (Ref 28). The result will be a loss of grain refinement, usually called fade. For this reason, it is useful to provide a small amount of gentle stirring in holding furnaces. When stirring is present to prevent sedimentation, borides will provide grain refinement for long times in foundry alloys. In this article, the advantages and disadvantages of the various commercial Al-Ti-B master alloys have not been considered. The interested reader is referred instead to the excellent review by Spittle (Ref 29). He concluded that, when different refiners are compared on the basis of the amount of boron added, the popular Al-5Ti-1B alloy usually gives the most consistent refinement.
100 (5) (2) (1)
Equivalent average pore diameter (dp), μm
(3)
(4) 10 A356 alloy Casting #
H2 content (cc/100g)
(5) (2) (1) (3) (4)
0.31*** 0.31** 0.25* 0.25** 0.11**
1 0.1
* Non-grain-refined ** Grain-refined *** Grain-refined and modified
1
10
100
Cooling rate, ⬚C/s
Fig. 12
ACKNOWLEDGMENT The authors wish to acknowledge Alcoa Primary Metals, whose support made this study possible.
Pore size in A356 alloy castings. Source: Ref 25
REFERENCES Coherency fraction solid (fs*), %
60 55 50 45 40 35 30 25
0.5⬚C/s O 1⬚C/s
20 15 0.00
0.02
0.04
0.06
0.08
0.10
0.12
0.14
0.16
0.18
0.20
0.22
Grain refiner added
Fig. 13
Dendrite coherency in an Al-4%Cu alloy. Source: Ref 26
titanium content is kept low, preferably less than approximately 0.05% Ti. The grain size is reduced and the resistance to hot cracking is improved in most castings. This runs counter to the accepted view, since many of these alloys call for a relatively high minimum titanium level. Example alloys are 201, 203, 204, 206, 240, 242, 710, 712, 713, 771, and 772. Titanium has little effect: In many alloys, the titanium content has little effect on grain size when a powerful boron-containing refiner is employed. The Al-Si, Al-Si-Mg, and Al-Mg alloys fall into this category. Example alloys are 356, 357, 358, 359, 512, 520, and 535.
High titanium is best: Tests in 319 alloy have
shown that it is necessary to have a minimum of approximately 0.1% Ti in order to produce a small grain size. By analogy, other Al-Si-Cu alloys presumably would also benefit from high titanium additions. Example alloys are 308, 319, 354, and 355. In all of the aforementioned cases, the best grain-refinement practice is to make an addition of 10 to 20 ppm of boron, preferably in the form of Al-5Ti-1B or Al-3Ti-1B rod. A Tibloy alloy can also be used and may be cost-effective in hard-to-refine alloys (such as 319), but at least 5 min should be allowed for the refiner to become active.
1. J.A. Nock, Jr., Method of Making and Casting Aluminum Alloys, U.S. Patent 1,912,382, 1934 2. A. Pacz, Aluminum Alloy Casting and Process of Making the Same, U.S. Patent 1,860,947, 1932 3. J. Suhr, Casting Alloy AP-33, Rev. Aluminium, Vol 9, 1932, p 1669–1680 4. W. Rosenhain, J. Grogan, and T. Schofield, Gas Removal and Grain Refinement of Aluminum Alloys, Foundry Trade J., Vol 43, 1930, p 177–180 5. G.K. Sigworth, The Grain Refining of Aluminum and Phase Relationships in the AlTi-B System, Met. Trans. A, Vol 15, 1984, p 277–282 6. L. Backerud, How Does a Good Grain Refiner Work?, Light Met. Age, Oct 1983, p 6–12 7. M.M. Guzowski, G.K. Sigworth, and D.A. Sentner, The Role of Boron in the Grain Refinement of Aluminum, Met. Trans. A, Vol 18, 1987, p 603–619 8. M.A. Easton and D.H. St. John, The Partitioning of Titanium during Solidification of Aluminum Alloys, Mater. Sci. Technol., Vol 16 (No. 9), 2000, p 993–1000 9. G.S. Cole, J. Cisse´, H.W. Kerr, and G.F. Bolling, Grain Refinement in Aluminum and Aluminum Alloys, AFS Trans, Vol 80, 1972, p 211–218 10. A. Cibula and R.W. Ruddle, The Effect of Grain Size on the Tensile Properties of High Strength Cast Aluminum Alloys, J. Inst. Met., Vol 76, 1949–1950, p 361–376
262 / Molten Metal Processing and Handling 11. A. Cibula, The Grain Refinement of Aluminum Alloy Castings by Additions of Titanium and Boron, J. Inst. Met., Vol 80, 1951–1952, p 1–16 12. H.T. Lu, L.C. Wang, and S.K. Kung, Grain Refining in A356 Alloys, J. Chin. Foundrym. Assoc., Vol 29, June 1981, p 10–18 13. G.K. Sigworth and M.M. Guzowski, Grain Refining of Hypo-Eutectic Al-Si Alloys, AFS Trans., Vol 93, 1985, p 907–912 14. G.K. Sigworth and M.M. Guzowski, Grain Refiner for Aluminum Containing Silicon, U.S. Patents 5,055,256 and 5,180,447; U.K. Patent GB2174103A 15. K. Nogita, S.D. McDonald, and A.K. Dahle, Effects of Boron-Strontium Interactions on Eutectic Modification in Al-10 mass% Si Alloys, Mater. Trans., Vol 44, 2003, p 692–695 16. G.K. Sigworth, M.A. Easton, J. Barresi, and T.A. Kuhn, Grain Refining of Al-Si Casting Alloys, Light Met., 2007, p 691–697 17. D.K. Young, B.T. Dunville, W.C. Setzer, and F.P. Koch, A Survey of Grain Refiners
18.
19.
20.
21. 22.
23.
for Hypoeutectic Al-Si Alloys, Light Met., 1991, p 1115–1121 J.H. Sokolowski, C.A. Kierkus, B. Brosnan, and J.W. Evans, Formation of Insoluble Ti(Al,Si)3 Crystals in 356 Alloy Castings and Their Sedimentation in Foundry Equipment: Causes, Effects and Solutions, AFS Trans., Vol 108, 2000, p 491–495 X.-G. Chen and M. Fortier, Formation of Primary TiAlSi Intermetallic Compounds in Al-Si Foundry Alloys, Mater. Forum, Vol 28, 2004, p 659–665 K. Pasciak and G.K. Sigworth, Role of Alloy Composition in Grain Refining of 319 Alloy, AFS Trans., Vol 109, 2001, p 567–577 V. de L. Davies, The Influence of Grain Size on Hot Tearing, Br. Foundryman, Vol 63, April 1970, p 93–101 G.K. Sigworth and F. DeHart, Recent Developments in the High Strength Aluminum-Copper Casting Alloy A206, AFS Trans., 2003, p 341–354 G.K. Sigworth, M. Kaufman, O. Rios, and J. Howell, Development Program on
24.
25.
26.
27.
28. 29.
Natural Aging Alloys, AFS Trans., Vol 112, 2004, p 387–407 G.K. Sigworth, Grain Refining of Aluminum Casting Alloys, Sixth International AFS Conference on Melt Treatment of Aluminum, Nov 11–13, 2001 (Orlando, FL), p 210–221 Q.T. Fang and D.A. Granger, Porosity Formation in Modified and Unmodified A356 Alloy Castings, AFS Trans., Vol 97, 1989, p 989–1000 G. Chai, Dendrite Coherency during Equiaxed Solidification in Aluminum Alloys, Chem. Commun., (No. 1), Stockholm University, 1994 L. Arnberg and L. Backerud, Solidification Characteristics of Aluminum Alloys, Vol 3: Dendrite Coherency, American Foundrymen’s Society, 1996 P.L. Schaffer, A.K. Dahle, and J.W. Zindel, Grain Refiner Fade in Aluminum Alloys, Light Met., 2004, p 821–826 J.A. Spittle, Grain Refinement in Shape Casting of Aluminum Alloys, Int. J. Cast Met. Res., Vol 19, 2006, p 210–222
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 263-265 DOI: 10.1361/asmhba0005247
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Refinement of the Primary Silicon Phase in Hypereutectic Aluminum-Silicon Alloys John L. Jorstad, JLJ Technologies, Inc.
PRIMARY SILICON in hypereutectic Aluminum-silicon alloys is very hard, imparting improved wear resistance but also decreasing tool life during machining. To minimize machining issues while fully utilizing outstanding wear properties, primary silicon crystals should be controlled to uniform small size and have uniform spatial distribution. This can be accomplished by adding phosphorus to the melt, thus creating an abundance of AlP3 nuclei suitable for initiating primary silicon crystal growth, and by solidifying cast product as rapidly as feasible. The conventional high-pressure die casting process normally solidifies product rapidly enough that artificial nuclei are not needed, but even there, added AlP3 nuclei may offset the less-than-ideal melt temperature and melt flow conditions that sometimes occur. Phosphorus added for refinement and the elements added to modify eutectic silicon react together such that primary silicon refinement and eutectic silicon modification in hypereutectic alloys cannot be accomplished simultaneously, except to the extent that both are affected by solidification rate. Only the additive in excess of that reacting together is effective in accomplishing refinement (phosphorus) or modification (strontium, for instance).
Importance of Primary Silicon Refinement The most notable feature of a hypereutectic aluminum-silicon alloy microstructure is the primary silicon phase. Silicon, in either the primary or eutectic form (Fig. 1), is present in aluminum alloys as essentially pure silicon crystals. (Only a miniscule amount of aluminum is dissolved in the silicon, and other alloying elements dissolve to an even lesser degree.) Silicon crystals are harder than any other phase typically found in aluminum casting alloys. On the Mohs scale, silicon is 6.5 versus corundum at 7. (The corundum form of Al2O3 is a
common medium for abrasively cutting and grinding hardened metals.) On the Knoop microhardness scale, silicon particles have been measured at 1100 and above, whereas sludge (the complex Al15(MnFe)3Si2 phase often referred to in die castings as hard spots) typically measures less than 600, and common forms of Al2O3 and MgAl2O4 (spinel) skim and particulates measure even less. On the Knoop scale, even the fully hardened Al-Cu, Al-Mg, or AlCu-Mg alloy alpha matrices typically measure well under 200. The hardness of silicon generally provides two very dramatic effects in aluminum-silicon casting alloys: superior wear resistance and reduced tool life during machining (Ref 1). Both are affected by the volume fraction, size, morphology, and distribution of silicon crystals; thus, hypereutectic alloys (containing a much larger total volume fraction of silicon crystals as well as larger average silicon size) have the greatest impact. To minimize machining issues while maintaining favorable wear performance, the primary silicon must be reasonably small and uniformly distributed throughout a cast matrix, hence the importance of primary silicon refinement.
causing eutectic aluminum grains to nucleate first and forcing silicon to then grow between those aluminum grains, thus acquiring a fibrous, broomlike morphology. This controlled growth phenomenon is commonly termed modification.
Accomplishing Primary Silicon Refinement Primary silicon size is controlled both by nucleation and solidification cooling rate. Absent any artificial nuclei, primary silicon nucleates after relatively deep undercooling and, once initiated, can then grow rapidly to large size and very irregular shape (Fig. 2). However, providing an abundance of suitable nuclei on which primary silicon crystals can initiate and grow dramatically changes that overall microstructure (Fig. 3). One of the most effective means to assure abundant nucleation sites is to add phosphorus to the melt. Phosphorus is added to hypereutectic aluminum-silicon alloys in an amount sufficient to retain 0.0005 to 0.0025% (less that 50% of that added is usually retained). Phosphorus combines with aluminum to form tiny, insoluble AlP3
Refinement versus Modification Primary Silicon Phase Refinement. Controlling the size, shape, and distribution of the primary silicon phase in hypereutectic aluminumsilicon alloy castings (alloys with >12% Si) is based on maximizing the number of nucleation sites, augmented with rapid solidification cooling rates. This nucleation phenomenon is commonly termed refinement. Eutectic Silicon Modification. Controlling the size and morphology of the eutectic silicon in hypoeutectic or eutectic aluminum-silicon alloy castings (alloys with 12% Si or less) is, on the other hand, accomplished by chemically inhibiting the usual nucleation of eutectic silicon,
Fig. 1
Primary (the larger, more equiaxed dark phase) and eutectic (the smaller, more elongated dark phase) silicon crystals in a T-6 hypereutectic aluminumsilicon alloy cast microstructure. Both forms/phases are essentially pure silicon. Original magnification: 800
264 / Molten Metal Processing and Handling particles (Fig. 4) (Ref 2). Because of very similar crystal habits (both are diamond cubic) and lattice constants, AlP3 particles are effective nuclei on which primary silicon crystals initiate. Sources of Phosphorus. Over the years, a variety of means emerged for introducing phosphorus to melts. In the timeframe of 1960 to 1980, several salts containing phosphorus were
200 µ
Fig. 2
Irregular growth of primary silicon crystals after deep undercooling and rapid solidification
Fig. 3
More uniform size, shape, and distribution of nucleated primary silicon crystals. Original magnification: 200
popular. Another method in that timeframe was vaporized PCl5 bubbled into the melt entrained in N2 as a carrier gas. While they accomplished the desired refinement of primary silicon, those products fell out of favor because the fumes associated with breakdown of salts or vaporizing of PCl5 were unpleasant if not unhealthy. Another additive researched and reported in the 1960s (Ref 2) and continuing in use even today (2008) is the alloy of phosphorus and copper. Commonly called phos-copper, it is a product readily available in the form of extruded brazing rod or as shot used to deoxidize copper alloy melts. Brazing rod is usually eutectic phosphorus-copper containing nominally 8% P. Deoxidizing shot is more often hypereutectic phosphorus-copper and contains 15 to 20% P. Both are effective additives, but the eutectic phosphorus-copper is easier to use. Phosphorus-copper must dissolve in the melt, although the melting point of the eutectic grade is close to that of many hypereutectic aluminum-silicon alloy melts. Hypereutectic phosphorus-copper often has a tough oxide surface that renders it slow and perhaps difficult to dissolve, whereas the eutectic phosphoruscopper dissolves quickly and completely. The eutectic phosphorus-copper is therefore the preferred product, and manufacturers now produce eutectic phosphorus-copper in shot form, too. It is less expensive than extruded rod and handy to precisely weigh for charging in the relatively small quantities usually required. An AlCuP master alloy rod for primary silicon refinement was introduced by VAW in Germany in the late 1980s. This rod introduces preformed AlP3, and the product has proven quick and reliable for primary silicon refinement. Effect of Solidification Rate. Solidification rate alone can be sufficient to refine primary silicon, although only at the extreme rates often
(a)
Fig. 4
AlP3 nucleus in primary silicon. Original magnification: 1500
(b)
experienced during conventional high-pressure die casting (HPDC) does AlP3 seem unnecessary. Conventional high-pressure die casting probably requires no artificial nucleation. As noted in Fig. 5, HPDC can achieve very small primary silicon size even without benefit of specifically-added AlP3. This is likely the result of the exceptional solidification rates normal to HPDC but may be aided by turbulent flow affecting nucleation, too. Regardless of the reason, controls are still required to assure reasonable uniformity of primary silicon size and spatial distribution. Figure 6, for instance, compares a well-made HPDC compressor body with one in which the die was filled either from melt that was too cold (already below the alloy liquidus entering the die) or was flowing too slowly. Primary silicon should ideally nucleate in the die cavity, influenced by pressurized intimate contact between melt and die steel. However, if the melt is held too close to liquidus temperature, or if the shot is delayed too long in the shot sleeve, or if flow ahead of the die cavity is too slow, primary crystals can initiate and grow too early in the casting cycle and under less-than-ideal conditions existing in the part cavity. Because the ideal scenario can be difficult to assure, many die casters require phosphorus in their hypereutectic alloy melts and claim that somewhat better primary silicon microstructures result. The ideal refined primary silicon size thus depends on many factors, including process conditions and specific product requirements. The type of casting process itself is also a factor, because solidification rate plays a role. In general, the nominal primary silicon size that is desirable in many sand castings is 35 to 50 mm; in permanent mold, 25 to 40 mm; and in diecasting, where refinement may or may not be employed, 15 to 25 mm is sufficient. Ingot pretreated with phosphorus can be purchased and is quite suitable; however, melts
(c)
Effect of solidification rate. (a) Unrefined high-pressure die casting chain saw crankcase, >150 C/s (300 F/s). (b) Gravity permanent mold piston, 10 C/s (50 F/s). (c) Sand cast engine prototype, <1 C/s (34 F/s). (b) and (c) AlP3 refined. Original magnification: 75
Fig. 5
Refinement of the Primary Silicon Phase in Hypereutectic Aluminum-Silicon Alloys / 265
(a)
Fig. 6
(b) (a) Well-made high-pressure die cast microstructure versus (b) one where primary silicon nucleated both outside (large crystals) and within (tiny crystals) the die cavity. Original magnification: 100
(a)
(b)
Fig. 7
(a) AlP3 inclusion on a machined surface. (b) Agglomeration of tiny AlP3 particles. The phosphide reacts with atmospheric moisture to become the phosphate, which then protrudes from the machined surface. Original magnification: 1000
often contain a significant fraction of returns, and refinement fades over time. Thus, a more robust and reliable method is to treat each melt with phosphorus according to its specific need.
Loss of Refinement Over Time. The AlP3 particles have an affinity for each other and seem to lose effectiveness largely by agglomerating into clusters that are no longer effective as nuclei. Clusters migrate to melt surfaces,
oxidize, and are removed in dross. Agglomeration and clustering are hastened by agitation of melts, meaning that melt size (relative volume affected at any given time by stirring or ladling frequency, for instance) can influence the rate of loss of AlP3 effectiveness. As a rule, large melts (tons) lose the effectiveness of phosphorus refinement over a matter of many hours and even days, whereas small melts (a few kilograms) can lose it within 4 to 5 h. The AlP3 clusters are easily floated into the dross and removed during fluxing and degassing, especially when “active” (de-wetting) fluxes such as those containing chlorine are employed. Agglomerated particles can eventually collect into relatively large masses, and, if not removed during fluxing/degassing or filtering, AlP3 can become entrapped in castings to appear as inclusions (Fig. 7). While such inclusions are rare, their mere possibility is strong argument for not going overboard with refinement frequency or phosphorus amount and for properly maintaining melt surface cleanliness. Primary Silicon Refinement and Eutectic Modification Are Not Compatible. Artificial nucleation of primary silicon with phosphorus and modification of the eutectic with sodium, strontium, or other modifiers cannot be achieved simultaneously. Phosphorus and the elements used for modification react to form compounds that accomplish neither primary refinement nor eutectic modification. Only if one or the other of the intended additives is in excess of that needed to form the compounds will refinement or modification take place, but not both.
REFERENCES 1. J.L. Jorstad, The Hypereutectic AluminumSilicon Alloys 390 and A390, Trans. TMS/ AIME, Vol 242, 1968, p 1217–1222 2. F.L. Arnold and J.S. Prestley, Hypereutectic Aluminum-Silicon Alloy Phosphorus Refinement, Trans. Am. Foundrymen’s Soc., Vol 69, 1961, p 129–137
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 269-275 DOI: 10.1361/asmhba0005206
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Thermodynamics and Phase Diagrams Ursula R. Kattner, National Institute of Standards and Technology
THERMODYNAMICS is a branch of physics and chemistry that covers a wide field, from the atomic to macroscopic scale. In materials science, thermodynamics is a powerful tool for understanding and solving problems. The part of thermodynamics that concerns the physical change of state of a chemical system following the laws of thermodynamics is called chemical thermodynamics. The thermodynamic properties of individual phases can be used for evaluating their relative stability and heat evolution during phase transformations or reactions. Traditionally, the most common application of thermodynamics is in the form of phase diagrams. Phase diagrams are visual representations of the state of a material, generally equilibrium, as a function of temperature, pressure, and, in multicomponent systems, the concentrations of the constituent components. They are therefore frequently hailed as basic blueprints or roadmaps for alloy design, development, processing, and basic understanding. The importance of phase diagrams is also reflected by the publication of many handbooks (Ref 1). Thermodynamic quantities of pure substances are tabulated in a number of compilations (Ref 2), while only a few compilations of thermochemical data for solution phases exist (Ref 3, 4). Most of these compilations provide the data in tabular form, although some provide them in the form of analytical functions or in electronic form. The Scientific Group Thermodata Europe compilation of binary systems (Ref 4) provides selected thermochemical values in tabulated and graphical form. This compilation also includes phase diagrams calculated from the same thermodynamic functions as the tabulated thermochemical values. It is the first compilation that makes use of the correlations between thermodynamics and phase diagrams established by Gibbs.
Thermodynamics The thermodynamic quantities that are most frequently used in materials science are the enthalpy, in the form of the heat content of a phase, the heat of formation of a phase, or the latent heat of a phase transformation; the heat
capacity, which is the change of heat content with temperature; and the chemical potential or chemical activity, which describes the effect of composition change in a solution phase on its energy. All of these thermodynamic quantities are part of the energy content of a system and are governed by the three laws of thermodynamics. Detailed physical derivations of the relationships of these quantities can be found in many textbooks (Ref 5, 6). A system is considered to be in equilibrium when its energy is at a minimum. Under constant pressure conditions for a closed system, this energy is the Gibbs energy, G: G ¼ H TS
(Eq 1)
where H is the heat content or enthalpy, T is the temperature, and S is the entropy. The entropy is a rather abstract entity and can be envisioned as a measure of the disorder of atoms in the system. The temperature dependence of the Gibbs energy of a pure substance is usually expressed as a power series of T: G ¼ a þ bT þ cT lnðT Þ þ dT 2 þ eT 3 þ fT 1 þ gT 7 þ hT 9
(Eq 2)
where a . . . h are coefficients (Ref 7). Pressure and magnetic states result in additional contributions to the Gibbs energy. However, pressure dependence for condensed systems at nearatmospheric pressures is usually ignored. For multicomponent systems, it has proven useful to distinguish three contributions from the concentration dependence to the Gibbs energy of a phase, Gj: Gj ¼ G0 þ Gideal þ Gxs
(Eq 3)
The first term, G0, corresponds to the Gibbs energy of a mechanical mixture of the constituents of the phase: G ¼ 0
X
xi Gi ðT Þ
(Eq 4)
i
where xi is the mole fraction of each of the components, and Gi(T) is the Gibbs energy of the pure component. The second term in Eq 3, Gideal, is the contribution from the
configurational entropy from atoms substituting for each other on equivalent sites in a random mixture. In substitutional solutions, such as liquid or disordered solid solutions, it is given as: Gideal ¼ RT
X
xi ln xi
(Eq 5)
i
The third term in Eq 3, Gxs, is the so-called excess term. The excess Gibbs energy or Gibbs energy of mixing of a binary regular solution is given as: Gxs ¼ xA xB
(Eq 6)
In a regular solution, the parameter O is a measure of the enthalpy of mixing of the phase of interest and represents the deviation of the Gibbs energy from ideal behavior. It is negative if A-B neighbors are favored in the solution and positive if A-A and B-B are favored in the solution. The solution is called subregular when O is a function of the composition and quasi-regular when O is also a function of temperature. In general, these solutions are referred to as regular-type solutions. The formalisms needed for the description of Gxs of ordered phases with extended homogeneity ranges or liquid phases with short-range order are generally more complex, since an internal order parameter must be taken into account. The order parameter is zero for the completely disordered state and is at maximum for the perfectly ordered state. Under equilibrium conditions, the Gibbs energy must also have a minimum with respect to this internal parameter. For a given temperature and concentration, the equilibrium (minimum Gibbs energy) may be single phase or a mixture of phases. The Gibbs energy of a mixture of phases is the weighted sum of Gibbs energies of the individual phases: G¼
X
f j Gj
(Eq 7)
j
where fj is the phase fractions of the individual phases. The requirement of minimum Gibbs energy for a system in equilibrium requires that the chemical potential, mji , of component i in phase j must be equal in all phases:
270 / Principles of Solidification maA maB .. . maZ
¼ ¼ ¼
mbA mbB .. . mbZ
¼ ¼ ¼
mgA mgB .. . mgZ
¼ ¼ ¼
... ... .. _ .
(Eq 8)
...
The chemical potentials of the two elements in a binary regular solution phase are (Ref 6): mA ðT; xA Þ ¼ GA ðT Þ þ ð1 xA Þ2 þ RT ln xA mB ðT; xB Þ ¼ GB ðT Þ þ ð1 xB Þ2 þ RT ln xB ðEq 9Þ
The Gibbs energy for a given composition can also be calculated from the chemical potentials: G¼
X i
xi mi
(Eq 10)
The chemical activity is related to the chemical potential by: mi ¼ G0i þ RT ln ai
Pi Pi0
G
T1
T2
β
α
T
l
l
l+β
T2 β
T3
α
T4
B
T3
α
TB
α XB
G
T1
l+α
l
α l+α A
β
l
TA
α+β A
XB
B
l+α
l
l+β
β
XB
A G
B
T4
β
α
(Eq 12)
The vapor pressure of a component in solution, Pi, is reduced in comparison to that of the pure component, Pi0 . The vapor pressure of an individual component is the same over all phases in equilibrium.
Fig. 2
G
(Eq 11)
where G0i is the reference state of the pure component. In an ideal solution, the activity of a component is equal to its molar fraction, ai = xi. This relation is known as Raoult’s law. Real solutions, such as regular-type solutions, deviate from Raoult’s law. The deviation of the activity from the molar fraction of the component is described by the activity coefficient, gi, ai = gi xi; this relation is known as Henry’s law for dilute solutions. The activity, ai, of component i is defined as: ai ¼
(representing the line of equal chemical potential for all components) and the equilibrium compositions of the two phases are given by the tangent points. The line in the phase diagram connecting these two points is called a tie line. The relative position of the Gibbs energy curves varies with temperature and causes the concentrations of the tangent points to change as the temperature changes. The phase boundaries plotted in a phase diagram are the locii of these tangent points as a function of temperature. A special temperature can exist where three phases may occur in equilibrium, for example, T3 in Fig. 1. Although a phase diagram is usually determined from direct measurement of concentrations and temperature, it is
The requirement that the chemical potential of a component must be the same for all phases in equilibrium directly links the Gibbs energy to the phase diagram, as illustrated in Fig. 1. The Gibbs energy functions of three phases, liquid (l), a, and b, are plotted as functions of the concentration, xB (xA + xB = 1 for a binary system) for various fixed temperatures. The minimum of the Gibbs energy is given by the lowest of the individual Gibbs energy curves and tangents to the Gibbs energy curves. The equilibrium is a single phase if the minimum Gibbs energy is given by the curve of this phase. The equilibrium is a twophase mixture if the minimum Gibbs energy is given by the common tangent of two curves
l
β
α+β
β
α α A
Fig. 1
l+α
l+β XB
β
α B
A
XB
B
Relation between the Gibbs energy curves of the liquid (l), a, and b phases
Calculated phase diagrams for a liquid (L) and solid (S) phase using the regular solution model with various regular solution parameters OL and OS. (a) OL = OS = 0 (ideal solution). (b) OL = 10 kJ/mol; OS = 20 kJ/mol. (c) OL = OS = 20 kJ/mol. The transformation entropies of the pure components are 10 J/(mol K).
Thermodynamics and Phase Diagrams / 271 important to understand the underlying thermodynamic basis for these diagrams, since it creates rules that experimentally determined phase diagrams must obey. The Gibbs energy functions can be used to illustrate how the magnitude and particular shape of the excess functions affect the resulting phase diagram. Figure 2 shows that just by varying the parameter of regular solution of a liquid and a solid phase, OL and OS, lensshaped isomorphous (Fig. 2a), eutectic (Fig. 2b), or pertitectic (Fig. 2c) phase diagrams can be obtained. A positive value of Oj results in the formation of a miscibility gap (a region of phase separation) at lower temperature. However, since the contribution of Gideal is negative, the miscibility gap closes at the critical point at sufficiently high temperature. For a regular binary solution, the miscibility gap is symmetric and closes at xjA ¼ xjB ¼ 0:5 with the critical temperature, Tcj (Ref 5): Tcj ¼
j 2R
(Eq 13)
The relationship between the Gibbs energy functions and phase diagrams is the basis for the calculation of phase diagrams (CALPHAD) method (Ref 8). The calculation of phase diagrams has a number of advantages: (1) the asobtained phase diagrams and thermochemical data are self-consistent because they are calculated from the same set of Gibbs energy functions, (2) prediction of higher component systems is possible, (3) efficient storage of the information in the form of a thermodynamic database instead of a multitude of diagrams is possible, and (4) input data for other simulation software are provided. The Gibbs energy functions employed by the CALPHAD method are relatively simple mathematical functions that are based on physical models for atomic interactions in each phase. The adjustable parameters of these functions are obtained from fitting experimental thermochemical data (enthalpies, chemical potentials, etc.), experimental phase diagram data, or predictions from first principles calculations that are based on density functional theory (Ref 9). Phase diagrams for ternary and higher-order systems can be generated by the CALPHAD method. First, the thermodynamic descriptions of the constituent binary systems are derived. Thermodynamic extrapolation methods are then used to extend the thermodynamic functions of the binaries into ternary and higher-order systems. The results of such extrapolations can then be compared to experimental data, and, if necessary, interaction functions are added to the thermodynamic description of the higher-order system. In principle, this strategy is followed until all 2, 3, . . . n constituent systems of an ncomponent system have been assessed. However, experience has shown that, in most cases, no, or very minor, corrections are necessary for
reasonable prediction of quaternary or highercomponent systems (Ref 8). Since true quaternary phases are rare in metallic systems, assessment of most of the ternary constituent systems is often sufficient to describe an n-component system. This method has been successfully applied to many multicomponent systems and a number of commercial alloy systems.
Phase Diagrams Multicomponent phase diagrams are usually two-dimensional maps of phase stability regions for ranges of temperature and composition. Phase boundaries identify the location of phase changes. The liquidus boundary separates phase regions containing only liquid phase(s) and phase regions containing liquid and solid phase(s). This boundary represents the limit where the fraction of solid is zero. The solidus boundary separates phase regions containing only solid phase(s) and phase regions containing solid and liquid phase(s). This boundary represents the limit where the fraction of liquid is zero. The solvus is the boundary between phase regions containing only one solid phase and phase regions containing two or more solid phases. The solvus line represents the limit where the fraction of the other solid phase(s) is zero, that is, the limit where the solvent phase is saturated with the solute constituents, and higher concentrations result in the precipitation of other solid phases. The degrees of freedom, F, of a phase region are governed by the Gibbs phase rule: þF ¼Cþ2
(Eq 14)
where F is the number of phases in equilibrium, and C is the number of components in the system. The degrees of freedom are the externally controllable conditions of temperature, pressure, and, in a multicomponent system, composition. Most materials processes take place at constant pressure, which often is the atmospheric pressure, removing one of the degrees of freedom: þF ¼Cþ1
(Eq 15)
Equilibria with no degrees of freedom, that is, temperature and compositions of all fixed phases, occur where C + 1 phases simultaneously exist. These are important features of a phase diagram and are called invariant equilibria or invariant reactions. For example, in a binary system, an invariant three-phase equilibria exists where a liquid is in equilibrium with two solid phases. The invariant three-phase equilibria in a binary system can be divided into two groups: decomposition or eutectic-type reactions and formation or peritectic-type reactions. In a binary eutectic (Fig. 2b), the liquid decomposes on cooling into two solid phases, and in a peritectic (Fig. 2c),
the liquid reacts with a solid phase on cooling to form a second solid phase. Phase regions with C phases in equilibrium have, according to the phase rule, one degree of freedom (monovariant equilibrium). Only one of the variables, temperature or phase composition, can be chosen freely. An invariant binary equilibrium becomes monovariant in a ternary system because of the additional degree of freedom and generally retains its eutectic-type or peritectic-type character, although it is possible for its character to change at a certain temperature. An invariant equilibrium occurs at the intersection of C + 1 monovariant equilibria. The phase fractions, fj, in a multiphase equilibrium can be determined from: x0i ¼
X j
f j xji
(Eq 16)
where x0i is the concentration of component i in the alloy, and xji is the concentration of component i in phase j. In a binary system where xA + xB = 1, the phase fractions are: fa ¼
xbB x0B xbB
xaB
and f b ¼
x0B xaB xbB xaB
(Eq 17)
and the ratio of the two phase fractions becomes: f a xbB x0B ¼ f b x0B xaB
(Eq 18)
The same ratio exists by imagining a lever balanced at x0B , with weights replaced by the phase fractions and the lever arm lengths replaced by the concentration differences; the amount of a phase becomes larger the closer the alloy composition is to the phase boundary. This technique for determining phase fractions has come to be known as the lever rule.
Metastable Phase Diagrams Metastable equilibria occur when a phase is missing, for example, if the phase fails to nucleate. The failure of an equilibrium phase to form may result in the formation of a metastable phase. In the Gibbs energy diagram, the Gibbs energy curve of a metastable phase always is above the minimum Gibbs energy envelope of the equilibrium phases. The most well-known example for such a case is the iron-carbon system, shown in Fig. 3. Metastable cementite, Fe3C, results if graphite, or carbon, fails to form. The phase diagram shown in Fig. 3 is actually the superimposition of two phase diagrams, the stable equilibrium iron-carbon diagram (solid lines) and the metastable Fe-Fe3C diagram (dashed lines). In general, the absence of an equilibrium phase results in extended solubility in the remaining equilibrium phases. The liquidus of
272 / Principles of Solidification the metastable phase is always below the liquidus of the equilibrium phase that is absent.
Solidification Phase diagrams provide useful information for understanding the solidification of alloys. In addition to the liquidus and final freezing temperatures, important quantities for the mathematical treatment of solidification processes are obtained from the phase diagram. The ratio of the concentrations of the liquidus, xLi , and the solidus, xSi , at the same temperature is described by the partition coefficient or distribution coefficient, ki: xS ki ¼ Li xi
called the Scheil path), no diffusion in the solid and complete diffusion in the liquid are assumed. This case, where thermodynamic equilibrium exists only as local equilibrium at the liquid/solid interface, produces microsegregation with the lowest possible final freezing temperature and therefore represents the most extreme case of solidification with microsegregation. An alloy that solidified following the Scheil path is not in equilibrium after solidification. An expression similar to Eq 20 can be obtained for the Scheil path by equating the solute rejected during the formation of a small amount of solid, resulting in an increase of solid in the liquid (Ref 10):
xLB xSB df S ¼ 1 f S dxLB
(Eq 21)
(Eq 19)
Phase diagram information can also be used in two simple models to describe the limiting cases of solidification behavior. For solidification obeying the lever rule at each temperature during cooling, complete diffusion is assumed in the solid as well as in the liquid. Thus, all phases are assumed to be in thermodynamic equilibrium at each temperature during solidification. For the solidification of a single solid phase in a binary alloy, rewriting Eq 17 using f S = 1 f L and assuming a constant partition coefficient gives: xLB 1 ¼ x0B 1 ð1 kB Þf S
(Eq 20)
In comparison, for solidification following the Scheil-Gulliver path (for convenience, here
0
5
1600 1538⬚C (dFe)
10
Integration from xSB ¼ kB x0B at f S = 0 gives the so-called Scheil-Gulliver equation (Ref 10): xLB 1 ¼ x0B ð1 f S Þð1kB Þ
(Eq 22)
While the Scheil-Gulliver equation can also be derived for ternary and higher-component alloys, it requires knowledge of the tie lines. An alloy that is not in equilibrium after solidification can be the result of a phase failing to nucleate, incomplete diffusion, or solute trapping. The description of such solidification processes with equilibrium phase diagram information alone is not possible. However, for most alloys where the solidifying phase is a substitutional solid solution, predictions of the Scheil-Gulliver model are closer to reality than those of the lever model, while for
Carbon, at.% 15
20
25
0.09 0.53 L
1493⬚C 1400 0.16 1394⬚C
L + C(graphite) 1252⬚C 2.1
(gFe)
4.2
1153⬚C 2.14 1147⬚C
4.3
1000
Fe3C
Temperature, ⬚C
1200
912 ⬚C 800
0.65 0.021 0.022 0.76
740⬚C 727⬚C
600 (aFe)
400
Fig. 3
0 Fe
1
2
3 4 Carbon, wt%
5
6
Superimposed stable and metastable iron-rich part of the iron-carbon system. The metastable phase boundaries are shown with dashed lines.
interstitial solid solutions where the interstitial element diffuses rapidly, predictions of the lever model are closer to reality. Modeling of real solidification behavior requires knowledge of the temperature dependence of the partition coefficients and liquidus slopes and a kinetic analysis of microsegregation and back diffusion at each temperature. The solidification of a binary alloy is schematically shown in Fig. 4 (the sketches are oversimplified to emphasize the concentration gradients). The solidification of the alloy labeled with “x” begins at TL. In the case of lever solidification, the composition of the growing a phase changes homogeneously during cooling according to the tie lines (horizontal lines) until the temperature TP, the peritectic formation of the b phase, is reached. At this temperature, b phase is formed from the a phase and the remaining liquid phase, and solidification is complete. On further cooling, the compositions and phase fractions of the a and b phases continue to change according to the lever rule. In the case of Scheil solidification, only the composition of newly growing a phase changes according to the tie lines during cooling; the composition of a phase formed at higher temperatures remains unchanged because of the absence of diffusion in the solid phase, and the remaining liquid becomes enriched in solute. On reaching the peritectic temperature, TP, solidification of the b phase begins. The amount of a becomes frozen at this point. No formation of b phase from a phase reacting with the liquid phase occurs because of the absence of diffusion. Solidification of the b phase continues until all liquid is consumed or when the eutectic temperature, TE, is reached, where the remaining liquid solidifies in a eutectic reaction. An A-rich alloy that freezes as single a phase in lever solidification may show the formation of b phase in Scheil solidification. A B-rich alloy that freezes as two phases, b + g, in lever solidification may actually have all three solid phases, a + b + g, in Scheil solidification if its composition is on the A-rich side of the peritectic point. In some alloys, annealing at elevated temperatures but below the solidus temperature can be used to remove nonequilibrium phases and concentration gradients from microsegregations. Lever solidification in a ternary alloy is, in general, similar to that of a binary alloy, but, due to the added degree of freedom, a threedimensional representation is necessary. The phase boundaries that are crossed during cooling and the sequence of phase formation can be read from a temperature-concentration section (e.g., an isopleth, a vertical section through Fig. 5a or b). However, such a section provides only very limited information on the compositions of the individual phases or phase fractions, since tie lines are rarely in the same plane as the section. A temperature-concentration section is ill suited to obtain information for Scheil solidification; more information can be obtained from the liquidus projection, which
Thermodynamics and Phase Diagrams / 273 is a topological map of the liquidus surface, and isothermal sections. To illustrate Scheil solidification for different topologies of ternary systems, three-dimensional phase diagrams and liquidus projections of two hypothetical systems are shown in Fig. 5. Figure 5(a) shows a simple ternary system. The topmost surfaces represent the liquidus surface where various phases first solidify, depending on the alloy composition. Intersections of these primary solidification surfaces occur along lines resembling valleys. These lines are also referred to as lines of twofold saturation because the liquid is simultaneously in equilibrium with two solid phases. According to the phase rule, such a three-phase equilibrium is monovariant and therefore is also called a monovariant eutectic. At the intersection of three monovariant eutectics, a ternary invariant eutectic is formed.
During Scheil solidification of the alloy labeled with “y” in Fig. 5(a), primary b phase is formed first, and the liquid becomes enriched in solutes as b is formed until composition of the liquid coincides with one of the valleys. Solidification proceeds with descending temperature following the valley and the formation of two solid phases, b + g, until final freezing occurs at the temperature of the ternary eutectic equilibrium, shown as a dot on the liquidus projection in Fig. 5(c). The path of the Scheil solidification is marked with the thick line and arrows in Fig. 5. The hypothetical system shown in Fig. 5(b) is more complex. Not all intersections of the primary solidification surfaces resemble valleys; only a slope change occurs at the intersection of the L + g and L + d surfaces. The character of such a monovariant line is peritectic, and, as in the case of the binary
L TL T1
Temperature
TP
T2 T3 T4
α
TE
β γ
A
Lever solidification
Solid fraction TL
TL Total solid fraction
T1
T1
T2
T2
T3
T4
TE Fraction β
T3
T4
TE
B
Composition
X
L
α
Scheil solidification α
L
α
L
L
α
α L
L
α
α
α
β
β
β
β
α
Solid fraction TL
T1
Total solid T1 fraction
T2
T2
T3
T3
L
β
T4
α L
β
α
Fraction β
TE Eutectic (β+γ)
Fig. 4
TL
T4
TE Fraction γ
Hypothetical binary system. Phase fractions and oversimplified microstructures are shown for lever and Scheil solidification.
peritectic equilibrium, solidification “switches” from the formation of one solid phase to the formation of another solid phase. The solidification path of an alloy encountering such a monovariant peritectic is shown for alloy y in Fig. 5(b). The sequence of phase formation, shown on the liquidus projection in Fig. 5(d), during the Scheil solidification of this alloy is L!L+g!L+d!L+d+b!d+b+ a. However, the solidification of alloy z is different; its phase formation sequence is L ! L + g ! L + g + b ! L + d + b ! d + b + a. The solidification also begins with the formation of g until the liquid composition intersects with the valley of the L + b + g monovariant eutectic. Solidification proceeds with decreasing temperature following the valley until the invariant transition reaction occurs. In the transition reaction, the valley switches from the L + b + g monovariant eutectic to the L + b + d monovariant eutectic. The character of the invariant transitions reaction is closer to that of a peritectic reaction, and it is easy to understand why this invariant equilibrium is also called quasi-peritectic. Both alloys follow the L + b + d monovariant eutectic until final freezing occurs at the ternary eutectic. Although the methodology for obtaining solidification paths for lever and Scheil solidification is straightforward, the information needed for accurate predictions is, in many cases, incomplete. Due to their multidimensionality, the interpretation of diagrams for more complex systems can be quite cumbersome. For systems with more than three components, the graphical representation of the phase diagram in a useful form becomes a very challenging task. However, the difficulty of graphically representing systems with many components can be overcome by using thermodynamics for understanding reactions and phase transformations in multicomponent systems.
Calculation of Phase Diagrams and Solidification In recent years, the application of phase diagram information obtained from thermodynamic calculations to practical processes has increased significantly. Thermodynamic descriptions have become available for a large number of alloy systems and allow the calculation of the phase diagrams of multicomponent alloys. In addition to the phase equilibria data, the thermodynamic calculation also provides parameters that are needed for the simulation of casting processes, such as enthalpy data or driving forces for phase formation, as well as thermodynamic information that is needed for simulations considering kinetic factors. The calculation of the lever and Scheil path is a straightforward task. The change in temperature during cooling is approximated by a series of small decrements in temperature. For the lever path, the temperature is stepped, with the overall composition of the alloy being constant. The calculation of the Scheil path is based on
274 / Principles of Solidification z
y T
y
T TL γ
TL
β
δ
γ C α
α
β
B
C
A A
B
(a)
(b)
z
y γ
β
γ
δ
α
α (c)
Fig. 5
β
(d)
Solidication paths in hypothetical ternary systems. (a) Phase diagram of a system with a ternary eutectic. (b) Three-dimensional phase diagram of a system with a ternary transition reaction and a ternary eutectic. Solidification of the alloys begins at TL. (c) Liquidus projection of (a). (d) Liquidus projection of (b)
Fig. 6
the assumption that local equilibrium exists at the interface for each temperature step. The method used for calculating the Scheil path is illustrated in Fig. 4. At each temperature step, the fraction of solid phases formed is removed from the calculation, and the composition of the remaining liquid phase is used as the composition of the freezing alloy for the next temperature step. For example, the equilibrium is calculated at temperature T1 for alloy composition x, the amount of a that has formed is removed from the calculation, and the composition of the remaining liquid becomes the alloy composition for the calculation at the next temperature, T2. Simultaneously, the fractions of the solid phases that have been removed are summed to give the total fraction of solid formed during solidification: j j fnew ¼ fold þ f j
L j with f j ¼ fold fstep
(Eq 23)
j where fnew is the total accumulated fraction of j phase j, fold is the total fraction of phase j j is the fraction of from previous steps, and fstep phase j formed from the liquid in the current step. Stepping is usually stopped after a eutectic equilibrium has been encountered or a predetermined fraction of solid phase has formed. However, calculation of the enthalpy may be continued beyond solidification if results are needed for other calculations, such as process simulations. The calculation of the enthalpy versus temperature during Scheil solidification requires that
Phase fraction versus temperature curves for solidification of alloy Al-4.44Cu-1.56Mg-0.55Mn-0.23Fe-0.21Si-0.05Zn (mass fraction 100) using (a) equilibrium calculation and (b) nonequilibrium calculation
Thermodynamics and Phase Diagrams / 275 the variations in the compositions within the accumulated solid phases be considered. Instead of calculating the enthalpy for all compositions and increments of solid-phase fraction accumulated, the enthalpy can be calculated for the average composition of each individual solid at each temperature step (Ref 11). The average composition of each phase is obtained from: xji;new ¼
j xji;old fold þ xji;step f j j fold þ f j
(Eq 24)
The results obtained from the lever and Scheil calculations can be conveniently presented in graphical form in a phase fraction versus temperature diagram. Such diagrams are shown in Fig. 6 for lever and Scheil calculations for an alloy that is close in composition to the commercial aluminum alloy 243. The composition (in mass fraction 100) used for this calculation is 0.21Si-0.23Fe-4.44Cu0.55Mn-1.56Mg-0.05Zn, with the remainder being aluminum. For these calculations, the Al-DATA* database (Ref 12) was used. The difference between lever and Scheil solidification becomes most noticeable toward the end of solidification. For both solidification paths, the solidification begins with the precipitation of (Al) and Al6Mn. The lever solidification, shown in Fig. 6(a), continues with the formation of the a phase, Al20Cu2Mn3, Al7Cu2Fe, decomposition of Al6Mn and the a phase, and the formation of Mg2Si. Finally, after solidification is complete, precipitation of Al2CuMg begins. The final microstructure con-
*Trade names are used in this paper for completeness only, and their use does not constitute an endorsement by the National Institute of Standards and Technology.
sists mainly of (Al) and small amounts of Al20Cu2Mn3, Al7Cu2Fe, Mg2Si, and Al2CuMg. The Scheil solidification, shown in Fig. 6(b), continues with the formation of Al7Cu2Fe, Mg2Si, Al20Cu2Mn3 (this phase is not shown in Fig. 6b since the amount is extremely small), Al2CuMg, and Al2Cu. The microstructures obtained from these solidification paths are quite different and probably result in different mechanical properties. The actual microstructure found in castings can be compared to diagrams such as those shown in Fig. 6, and typical casting parameters, such as cooling rate or mold design, can be adjusted to obtain the desired microstructure by following a solidification path somewhere between these extremes. Examples of using the calculation of phase equilibria in solidification and casting simulations can be found in various chapters of this Handbook. REFERENCES 1. U.R. Kattner, The Thermodynamic Modeling of Multicomponent Phase Equilibria, JOM, Vol 49, Dec 1997, p 14–19 2. N. Jacobson, Use of Tabulated Thermochemical Data for Pure Compounds, J. Chem. Ed., Vol 78, June 2001, p 814–819 3. R. Hultgren, P.D. Desai, D.T. Hawkins, M. Gleiser, and K.K. Kelley, Selected Values of the Thermodynamic Properties of Binary Alloys, American Society for Metals, 1973 4. Thermodynamic Properties of Inorganic Materials Compiled by SGTE, Landolt-
5. 6.
7. 8.
9.
10. 11.
12.
Bo¨rnstein, New Series, Group IV, Vol 19, Subvol B, Springer, Berlin, Germany, 2002, onward C.H.P. Lupis, Chemical Thermodynamics of Materials, North Holland, New York, NY, 1983 M. Hillert, Phase Equilibria, Phase Diagrams and Phase Transformations: Their Thermodynamic Basis, Cambridge University Press, Cambridge, U.K., 1998 A.T. Dinsdale, SGTE Data for Pure Elements, CALPHAD, Vol 15, Dec 1991, p 317–425 N. Saunders and A.P. Miodownik, CALPHAD (Calculations of Phase Diagrams): A Comprehensive Guide, Pergamon, Oxford, U.K., 1998 J. Hafner, C. Wolverton, and G. Ceder, Toward Computational Materials Design: The Impact of Density Functional Theory on Materials Research, MRS Bull., Vol 31, Sept 2006, p 659–665 W. Kurz and D.J. Fisher, Fundamentals of Solidification, 3rd ed., Trans Tech Publications, Aedermannsdorf, Switzerland, 1989 W.J. Boettinger, U.R. Kattner, and D.K. Banerjee, Analysis of Solidification Path and Microsegregation in Multicomponent Alloys, Modeling of Casting, Welding and Advanced Solidification Processes VIII, B. G. Thomas and C. Beckerman, Ed., TMS, 1998, p 159–170 N. Saunders, Phase Diagram Calculations for Commercial Al-Alloys, Mater. Sci. Forum, Vol 217–222, 1996, p 667–672
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 276-287 DOI: 10.1361/asmhba0005207
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Nucleation Kinetics and Grain Refinement J.H. Perepezko, University of Wisconsin—Madison
NUCLEATION PROCESSES play a key role in the solidification of castings by controlling to a large extent the initial structure type, size scale, and spatial distribution of the product phases. During many solidification processes, the size scale of critical nucleation events is too small and the rate of their occurrence too rapid for accurate observation by direct methods. Nonetheless, nucleation effects in the solidification microstructure exert a strong influence on the grain size and morphology as well as the compositional homogeneity. The final microstructure is also modified by the crystal growth, fluid flow, and structural coarsening processes that are important in the later stages of ingot freezing. In large bulk castings, the solidification temperature corresponding to the onset of freezing is often close to but slightly less than the melting point or equilibrium liquidus temperature. The offset of the solidification temperature with respect to the equilibrium temperature is called the undercooling or supercooling, DT, and it plays a vital role in the overall description of the initial stage of the solidification that is controlled by nucleation. The level of melt undercooling at the onset of solidification is important to consider in developing an understanding of the variety of structural modifications and grain-refining practices in common casting alloys and is the basis of more recent solidification processing technologies using rapid solidification methods. In this article, selected highlights of thermodynamic relationships during solidification and nucleation kinetics behavior are discussed in connection with the basis of nucleation treatments, such as grain refinement and inoculation, to provide a summary of nucleation phenomena during casting.
and to account for the free energy changes involved in various crystallization processes. In most castings, a full global equilibrium associated with a uniform phase composition does not occur throughout the ingot. However, it is often valid to take the liquid and solid compositions at the solidification front to be represented by the liquidus and solidus compositions determined by the tie line end points at a given temperature on the phase diagram. This concept is called local interfacial equilibrium, and it has been successfully applied to the description of casting processes in which the rate of diffusion, particularly in the solid, is too sluggish to achieve uniform phase compositions (Ref 1). The consideration of nucleation and the level of melt undercooling introduce another type of departure from full equilibrium that is known as metastable equilibrium. Only when there is a departure from full liquid-solid equilibrium will solidification be possible. For solidification to occur, this departure brings the liquid into an undercooled state in which it is metastable because of the absence of one or more stable solid phases. The change of the liquid from a stable to a metastable state occurs in a continuous manner without an abrupt change in the physical properties, such as molar volume or heat capacity. Therefore, metastable states can exhibit a true reversible equilibrium. The thermodynamic description of solidification can be quantified by recalling that for a pure metal at the equilibrium temperature, Tf, the free energy change is:
Thermodynamics of Solidification
so that:
Macroscopic Solids. Throughout the analysis of solidification, thermodynamics is used to judge the alloy phase constitution, to describe the solidification path and composition changes in terms of partition coefficients and the slopes of the liquidus and solidus phase boundaries,
Gf ¼
Hf ðTf T Þ Hr T ¼ Tf Tf
(Eq 3)
when the small correction due to heat capacity terms is neglected. Based on Eq 3, there is a direct relationship between DGf and the undercooling DT. The relationship in Eq 3 is represented in Fig. 1, in which the relative stability of the two phases (stable a and metastable b) that form from the liquid without composition change is shown as a function of temperature. The behavior is also approximately followed for stable and metastable eutectics such as Al/ Al3Fe and Al/Al6Fe and Fe/C and Fe/Fe3C (Ref 2). An important feature illustrated in Fig. 1 is that the melting of a metastable phase always occurs at a lower temperature than the stable-phase melting. For alloys, the free energy of a phase is a function of both temperature and composition. When components are mixed together to form an alloy, the molar free energy change, DGm, can be estimated from thermodynamic solution models (Ref 3). The simplest approach is to assume that the heat of mixing, DHm, is zero and that the entropy change upon mixing, DSm, is given by the ideal entropy of mixing. This is the description of an ideal solution, and it is often a useful approximation for dilute alloys. For an ideal solution of composition CB:
GS GL ¼ ðHS HL Þ Tf ðSS SL Þ ¼ 0
or Gf ¼ Hf Tf Sf ¼ 0
Sf ¼ ðHf Þ=Tf
(Eq 1)
(Eq 2)
where GS and GL, HS and HL, and SS and SL are the molar free energy, enthalpy, and entropy for the solid, S, and liquid, L, respectively. At T, a temperature other than Tf, DGf 6¼ 0, and Eq 1 and 2 can be combined to yield:
Fig. 1
Molar free energy versus temperature curves for a single component with a stable a-phase and a metastable b-phase
Revised and updated from J.H. Perepezko, Nucleation Kinetics, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 101–108
Nucleation Kinetics and Grain Refinement / 277 Gm ¼ Hm T Sm ¼ RT ½CB ln CB þ ð1 CB Þ lnð1 CB Þ
(Eq 4)
With more concentrated alloys, a regular solution model is often applied in which the heat of mixing is given by the symmetrical term OCB(1 CB), where O reflects the net interatomic interaction. For regular solutions, DSm is taken as the ideal solution term in Eq 4. The description of equilibrium between solid and liquid phases over a range of alloy composition and temperature is represented in terms of the component chemical potentials mA and mB for each phase. The chemical potential is related to the molar free energy at a given composition by the standard thermodynamic definition (Ref 3): mA ¼ Gm CB
@Gm @CB T;P
@Gm mB ¼ Gm þ ð1 CB Þ @CB T;P
(Eq 5a)
(Eq 5b)
As illustrated in Fig. 2(a), the definitions in Eq 5(a) and Eq 5(b) indicate that a tangent to a molar free energy curve at composition Co intersects the pure component axes at mA and mB. With this description, alloy phase equilibrium is represented by an equality of chemical potentials for each of the components in the equilibrated phases so that mLA ¼ mSA and mLB ¼ mSB , as shown in Fig. 2(b). The geometric interpretation of alloy phase equilibrium is expressed in terms of the common tangent construction, which is also shown in Fig. 2(b). The molar free energy of the liquid or the solid phase is also related to the chemical potential by the expression: Gm ¼ ð1 CB ÞmA þ CB mB
(Eq 6)
and provides a useful form in representing the free energy changes involved in the formation of solid phases from an alloy liquid solution. Free energy versus composition diagrams provide a useful method of analyzing the free
Fig. 2
energy relationships for alloy solidification reactions. Some of the relationships are shown in Fig. 3 for a temperature below the eutectic. For temperatures below the eutectic, the (a + L) and (b + L) equilibria are metastable with respect to the stable (a + b) phase mixture. In the phase diagram, the metastable two-phase conditions can be examined in terms of an extension of the respective solidus and liquidus phase boundaries below the eutectic, as noted in Fig. 3. For large excursions below the eutectic, the phase boundary extrapolations may not follow a simple linear extension. The relationships illustrated in Fig. 3 also show the common characteristic for a metastable phase to exhibit a higher solubility for solute than the equilibrium stable-phase solution. Similar behavior is found for the cases of peritectic and monotectic solidification reactions (Ref 2). When one or both of the stable solid phases in a eutectic or peritectic reaction are suppressed, metastable solid phases may develop and yield metastable eutectics and peritectics (Ref 2). Figure 4 shows the modified phase diagram for a metastable peritectic. The stable Al/ Al3Ti and metastable Al/AlxTi peritectic reactions may be relevant to the operation of grain refining in aluminum alloys (Ref 4, 5). To provide a description of compositional segregation plus the influence of constitutional supercooling and solute redistribution during two-phase eutectic growth, phase diagram information such as the partition coefficient, k, or the liquidus slope, mL, is often used in alloy solidification analysis (Ref 6, 7). For dilute solutions, the van’t Hoff relation for liquidsolid equilibrium holds and relates k and mL as: k¼1
mL HfA RðTfA Þ2
Fig. 3
(Eq 7)
where HfA and TfA are the latent heat and melting point, respectively, of pure solvent, A. The partition coefficient can be evaluated separately for a dilute solution by: HfB 1 1 k ¼ exp R TfB TfA
where HfB and TfB are the latent heat and melting point, respectively, of component B. Microscopic Solids. The preceding discussion of solidification thermodynamics applies to cases in which the solid phases are of macroscopic size. There are several important situations during nucleation, dendritic solidification, or eutectic growth at high velocity in which the solid is of microscopic size or
Eutectic phase diagram (a) illustrating the metastable extensions of the liquidus and solidus phase boundaries to a temperature T1 below the eutectic. (b) The molar free energy versus composition diagram at T1 shows the metastable (a + L) and (b + L) equilibria by dashed common tangents and the stable (a + b) equilibrium by the solid common tangent.
(Eq 8)
Alloy free energy versus composition diagrams showing the chemical potentials of the liquid and solid components. (a) Variation of the molar free energy of mixing with composition for a liquid alloy solution. The chemical potentials, mLA and mLB , for the solution at composition Co are given by the intercepts on the pure component axes. (b) Molar free energy versus composition diagram for two-phase liquid-solid equilibrium in an alloy of composition Co. Solid of composition CS is in equilibrium with liquid of composition CL, and the chemical potentials of A and B are equal in the solid and liquid.
Fig. 4
Modified phase diagram for a metastable peritectic. When a stable A3B phase is suppressed, the stable peritectic reaction involving A and A3B may be replaced by a metastable peritectic reaction between A and AxB. The peritectic temperature Tp and a-phase solubility for the metastable peritectic are higher than for the stable peritectic case.
278 / Principles of Solidification exhibits a sharply curved boundary with the liquid. The free energy of a small particle increases inversely with its size or radius of curvature. The increased free energy for the microscopic solid compared to the macroscopic solid results in a decrease in melting temperature, as expected for metastable equilibrium behavior. For a pure component, the free energy increment due to curvature K is given by: Gk ¼ sKVm
(Eq 9)
where Vm is the molar volume, and s is the interfacial energy. For a sphere of radius, r, K = 2/r. In terms of Eq 9, the melting of a microscopic particle of pure component with radius r, Tf(r), is depressed from the melting point of a macroscopic particle by: Tf ðrÞ ¼ Tf
2sVm ðrSf Þ
(Eq 10)
For alloys, the influence of curvature can be treated by including a term based on Eq 9 to the chemical potential for the solid (Eq 6). This will shift the liquidus and solidus phase boundaries, but in a manner that results in a negligible change in k and mL (Ref 6).
Nucleation Phenomena Nucleation during solidification is a thermally activated process involving a fluctuational growth in the sizes of clusters of solids (Ref 8, 9). Changes in cluster size are considered to occur by a single atom addition or by removal exchange between the cluster and the surrounding undercooled liquid. At small cluster sizes, the energetics of cluster formation reveal that the interfacial energy is dominant, as can be observed by noting that the ratio of surface area to volume of a sphere is 3/r. For the smallest sizes, clusters are called embryos; these are more likely to dissolve than grow to macroscopic crystals. In fact, the excess interfacial energy due to the curvature of small clusters is the main contribution to the activation barrier for solid nucleation. This accounts for the kinetic resistance of liquids to crystallization and is manifested in the frequent observation of undercooling effects during solidification (Ref 10, 11).
Homogeneous Nucleation The principal features of nucleation phenomena and the kinetics of the rate process during solidification can be illustrated in the simplest terms by using the capillarity model to evaluate the kinetic factors (Ref 8). With this approach, it is useful to examine first the case of homogeneous nucleation in which solid formation occurs without the involvement of any extraneous impurity atoms or other surface sites in contact with the melt. As a further simplification, only the case of isotropic interfacial
energy is treated, but it should be recognized that anisotropic behavior can yield faceted cluster shapes. The energetics of cluster formation for a spherical geometry can be expressed in terms of a surface and a volume contribution as: GðrÞ ¼ 4pr2 s þ 4/3pr3 GV
(Eq 11)
where DG(r) is the free energy change to form a cluster of size r, and DGV = DHf DT/TfVm. The relationship in Eq 11 is characterized in Fig. 5, in which the activation barrier for nucleation, DGcr, is reached at a critical size rcr (that is, dDG(r)/dr = 0), as given by: 2s 2sTf Vm ¼ Hf T GV
rcr ¼
(Eq 12)
and
Fig. 5
Gcr ¼
3
16ps ¼ 3G2V
16ps Tf2 Vm2 3Hf2 T 2 3
(Eq 13)
At increasing values of undercooling, rcr is reduced (rcr a DT1), and Gcr is reduced more rapidly (DGcr a DT2). A cluster is often considered to reach the stage of a nucleus capable of continued growth with a decreasing free energy when the size rcr is achieved, but in fact, stable nucleus growth ensues when the cluster size exceeds rcr by an amount corresponding to (DGcr kT) in Fig. 5 (Ref 8). The relationship between cluster size and the number of atoms in a cluster, ncr, is expressed by (ncr Va Þ ¼ 4pr3cr =3, where Va is the atomic volume. To relate cluster energetics and fluctuational growth to the rate of nucleation, it is necessary to describe the cluster population distribution. As a liquid is undercooled, solid clusters begin to evolve over a transient period until the cluster size distribution attains a steady state. Because the mixture of clusters in an undercooled melt is a dilute solution, the metastable equilibrium concentration of clusters of a given size, C(n), can be expressed by (Ref 8): GðnÞ CðnÞ ¼ C‘ exp kT
(Eq 14)
where C‘ is the number of atoms per cubic meter in the liquid, and DG(n) is given by Eq 11 when r is converted to n, as noted previously. If solid nucleation is regarded as the growth of clusters past the critical size, then the resulting cluster flux or the nucleation rate I (in nuclei m3 s1) can be represented kinetically by the product:
I ¼ nSL Scr ZCðncr Þ
(Eq 15)
where nSL is the jump frequency associated with atom jumps from the liquid to join the cluster and can be estimated from the liquid diffusivity, DL, and jump distance, a, as in (DL/a2); Scr is the number of atoms surrounding a cluster that is approximately ð4pr2cr=a2 Þ; Z is the Zeldovich factor that accounts for the decay of
Free energy change for cluster formation as a function of cluster size. The surface contribution to DG(r) is 4pr2s, and the volume contribution is 4pr3 DGv/3; rcr occurs at 2s/DGv where DG(r) = DGcr.
supercritical clusters and is usually of the order of 0.1 to 1.0 (Ref 8); and C(ncr) is the concentration of critical clusters. The full expression for the steady-state nucleation rate is then: I¼
2 DL 4prcr Gcr ZC exp ‘ a2 a2 kT
(Eq 16)
For typical metals, C‘ 1028 m3 , DL 109 m2/s, and ao 0.3 109 m so that: I ’ 1040 exp
16ps3 Tf2 Vm2 3k Hf2 T T 2
(Eq 17)
and shows a rather steep temperature dependence, as illustrated in Fig. 6. At high temperature, the temperature dependence of I is dominated by the driving free energy term, which is contained within the exponential dependence on the activation barrier, and I can vary by approximately a factor of 5 per degree Celsius. Equation 16 or 17 indicates that the maximum nucleation rate occurs at (1/3) Tf. Over a wide temperature range, the liquid diffusivity cannot be taken as constant but can instead be represented as an Arrhenius function, DL = Do exp (Q/kT), where Do is a constant, and Q is the activation energy for liquid diffusion. At large undercooling or low temperatures, and especially for glass-forming alloys, the diffusivity is often obtained from the viscosity in the form DL = (kT/3paZo) exp [A/(T To)], where Zo, A, and To are constants. Often, To can be estimated by Tg, the glass transition temperature. In evaluating nucleation rates, accurate values for the activation barrier are important because of the exponential dependence of I on DGcr. The evaluation of DGV is based on thermodynamic data or reasonably accurate models (Ref 12). However, separate, independent measurements of liquid-solid interfacial energy, sLS, at the nucleation temperature are not available. Instead, estimates for sLS
Nucleation Kinetics and Grain Refinement / 279
Fig. 6
Steady-state nucleation rate as a function of undercooling below the melting point. At low undercooling, the nucleation rate is primarily controlled by driving free energy; at high undercooling, as the temperature approaches the glass transition temperature, Tg, nucleation is limited by diffusional mobility.
The interfacial energy, s, relationships among a planar nucleant substrate (n), a spherical sector solid (S), and the liquid (L). The interfacial regions are designated by the subscripts LS (liquid-solid), nL (nucleant-liquid), and nS (nucleant-solid).
Fig. 7
are obtained from correlations (Ref 13), calculations based on structural models (Ref 14), or, more recently, from ab initio atomistic computation analysis (Ref 15–17).
Heterogeneous Nucleation Homogeneous nucleation is the most difficult kinetic path to crystal formation because of the relatively large activation barrier for nucleus development (DGcr). To overcome this barrier, classical theory predicts that large undercooling values are required, but in practice, an undercooling of only a few degrees or less is the common observation with most castings. This behavior is accounted for by the operation of heterogeneous nucleation in which foreign bodies such as impurity inclusions, oxide films, or crucible walls act to promote crystallization by lowering DGcr. Because only a single nucleation event is required for the freezing of a
liquid volume, the likelihood of finding a heterogeneous nucleation site in contact with a bulk liquid is great. Indeed, it has been estimated that even in a sample of high-purity liquid metal there is a nucleant particle concentration of the order of approximately 1012 m3 (Ref 10). Only by using special sample preparation methods to isolate the melt from internal and external nucleation sites by subdivision into a fine droplet dispersion has it been possible to achieve undercoolings in the range of 0.3 to 0.4 Tf (Ref 11). The action of heterogeneous nucleation in promoting crystallization can be visualized in terms of the nucleus volume that is substituted by the existing nucleant, as illustrated schematically in Fig. 7. For a nucleus that wets a heterogeneous nucleation site with a contact angle y, the degree of wetting can be assessed in terms of cos y = (snL snS)/sLS, where the interfacial energies are defined in Fig. 7. As y approaches 0 , complete wetting develops; as y approaches 180 , there is no wetting between the nucleus and the nucleant (which is inert), and the conditions approach homogeneous nucleation. The energetics of heterogeneous nucleation can be described by a modification of Eq 11 to account for the different interfaces and the modified cluster volume involved in nucleus formation. In terms of the cluster formation shown in Fig. 7, the free energy change during heterogeneous nucleation is expressed by:
in applying the continuum model described previously for heterogeneous nucleation. As the contact angle decreases below approximately 20 to 30 , h approaches atomic dimensions, so that the continuum description of wetting is not valid. This issue is of importance in the analysis of heterogeneous nucleation during grain refinement where the observed undercoolings are small and the implied contact angles are below 20 . Following an analysis similar to that developed for homogeneous nucleation based on Eq 15, an expression can be obtained for the rate of heterogeneous nucleation. In this case, the value of the surface area, Scr, for a spherical cap geometry is estimated by ½2pr2cr ð1 cos yÞ=a2 . In addition, the concentration of critical clusters is represented in terms of the number of surface atoms of the nucleation site per unit volume of liquid, Ca, which is of the order of 1020 m3. The heterogeneous nucleation rate expression Ihet (m3 s1) is given by:
GðrÞhet ¼ VSC GV þ ASL sSL þ AnS snS AnL snL
where the prefactor represents an average value that can vary by approximately 1 order of magnitude for specific cases. The comparison between heterogeneous and homogeneous nucleation kinetics is illustrated in Fig. 9 by a schematic time-temperature transformation diagram. Although only a single transformation curve (that is, C-curve) is shown in Fig. 9 for heterogeneous nucleation, in reality there will be as many curves as the number of heterogeneous nucleation sites. Each curve for heterogeneous nucleation will be distinguished by a catalytic potency, that is, f(y), and a site density (Ref 11). To attain homogeneous nucleation conditions, it is clear that all heterogeneous nucleation sites must be removed or bypassed kinetically, such as during rapid quenching.
(Eq 18)
where VSC is the spherical cap volume, and ASL, AnS, and AnL are the solid-liquid, nucleant-solid, and nucleant-liquid interfacial areas, respectively. When the volume and relevant interfacial areas are expressed in terms of the geometry of Fig. 7, the evaluation of DGcr for heterogeneous nucleation yields: 16ps3LS 2 3 cos y þ cos3 y 4 3G2V (Eq 19) ¼ Gcr ðhomÞ½fðyÞ
Gcr ðhetÞ ¼
so that the barrier for homogeneous nucleation is modified by f(y), the shape factor, during heterogeneous nucleation. Over the variation in y ranging from complete wetting (y = 0 ) to nonwetting (y = 180 ), f(y) or [DGcr(het)/D Gcr(hom)] varies from 0 to 1.0, as illustrated in Fig. 8. It is important to note that the value of rcr for the curvature of the critical nucleus does not change in the classical analysis of heterogeneous nucleation, but the reduction in the activation barrier shown in Fig. 8 has a significant influence on the nucleation rate. This feature is also described in Fig. 8, in which the spherical cap size as measured by the ratio (h/r) is plotted as a function of y. The dimension h is the height above the substrate. A closer examination of the trend of (h/r) in Fig. 8 reveals that there is a problem
2 DL 2prcr ð1 cos yÞ 2 2 a a Gcr ðhomÞ fðyÞ Ca exp kT
Ihet ¼
(Eq 20)
or for typical metals: 16ps3LS Tf2 Vm2 fðyÞ Ihet ’ 1030 exp 3k Hf2 T T 2
(Eq 21)
Rapid Solidification Over the past 25 years, developments in rapid solidification have expanded the understanding of a number of central solidification microstructural characteristics, such as supersaturated solid solutions, cells, and dendrites, associated with growth kinetics and have expanded the product phase selection to include nonequilibrium structures such as metastable phases and metallic glass (Ref 12). Initially, the analysis of rapid solidification was viewed in terms of a high cooling rate; however, the understanding of the synthesis of metastable
280 / Principles of Solidification
Fig. 8
The variation in shape factor f(y) or [DGcr(het)/DGcr(hom)] and spherical cap size, h/r, as a function of the contact angle, y
Fig. 9
Comparison between heterogeneous nucleation (A) and homogeneous nucleation (B) in terms of the relative transformation kinetics below the melting point. The reduced temperature Tr = T/Tf and time t ∝ l1. Tg, glass transition temperature
phases and metallic glass requires a consideration of nucleation behavior under conditions of high undercooling. Indeed, the development of bulk metallic glasses, where melts of centimeter size can vitrify during slow cooling (Ref 18), demonstrates the importance of melt undercooling in rapid solidification behavior. Under usual conditions, the steady-state nucleation rate given by Eq 17 or 21 is constant. During rapid melt quenching and especially at high undercooling near Tg, atomic transport may not be adequate to maintain cluster equilibrium. Under conditions of sluggish
atomic motion, the initial cluster distribution can be different from the steady-state form, and the time required to reach steady state represents an effective delay time, t (Ref 8, 9). Essentially, the delay time is the period for cluster growth over the size range covered by (DGcr kT) in Fig. 5 and can be estimated from (bZ2)1, with b = nSL Scr. During this transient period, the time-dependent nucleation rate I(t) increases continuously and can be described approximately by I(t) = I exp[t/t] (Ref 19). When transient nucleation is significant, crystallization is delayed to facilitate glass formation. For a number ofmaterials, the delay time and the cooling rate, T , are inversely pro portional, so that T 1 K and may be used to estimate the relative importance of transient effects on nucleation (Ref 20). Current experience indicates that non-steady-state effects are most important for glass formation and during devitrification upon heating.
Inoculation Practice There is a wide commercial application of cast alloy treatments that modify the initial solidification characteristics to provide a means for effective control of grain size and morphology. A number of methods are available to control grain size, including electromagnetic or mechanical stirring, ultrasonic cavitation, and
directed fluid flow. However, inoculation appears to give the most consistent results and is the most economical approach. Inoculation involves the introduction of nucleating agents to a melt, either externally in the form of a fine dispersion or through internal means by phase or chemical reactions that result in the formation of a solid reaction product. Currently, grain-refining agents are usually fabricated in rod form, for example, as an aluminum-base master alloy containing insoluble particles of borides or carbides. The rod is introduced into the molten metal as it flows into a casting unit to facilitate rapid dissolution of the aluminum matrix and dispersion of the particles. Grain refinement effectiveness is often evaluated by the standard TP-1 test procedure (Ref 21). An effective inoculation treatment typically yields an equiaxed grain morphology with minimum grain sizes in the 100 to 150 mm range. In commercial practice, inoculants are used as additions to numerous molten alloys to yield fine grain sizes in castings. In cast iron, ferrosilicon is added to nucleate graphite; zirconium or carbon can be added to magnesium alloys; iron, cobalt, or zirconium is added to copper-base alloys; phosphorus is introduced to aluminumsilicon alloys to refine the silicon phase size; arsenic or tellurium can be added to lead alloys; and titanium is added to zinc-base systems. Other additives, such as sodium in aluminumsilicon alloys, are used to modify the growth morphology. In addition, there is a large body of experience on the grain refinement of aluminum-base alloys by aluminum-titanium, aluminum-titanium-boron, and aluminumtitanium-carbon melt additions. There is widespread experience with inoculation, and several models have been developed for the process, primarily (but not exclusively) in connection with aluminum alloy melts. It is not possible to discuss all of this work in this article; instead, the key features of inoculation are treated in relation to specific features of nucleation during solidification. To illustrate the principal features of inoculation and the physical mechanisms that have been developed to account for the effectiveness of inoculating agents, some of the experience with aluminum alloys is used as a basis for discussion. The basic features of inoculation in other alloy systems follow the pattern that appears to apply to aluminum alloys. The basic requirements for an efficient nucleant can be assessed from consideration of nucleation theory (Ref 22–24). To promote the formation of crystals on a nucleant, the interface between the nucleant and the liquid should be of higher energy than that between the nucleant and the crystal solid. A means of maximizing this condition is to provide a nucleant crystal relationship that is associated with a good crystallographic fit between the respective crystal lattices. In fact, the potency of a given catalyst is believed to be increased for decreasing values of relative lattice disregistry, with the degree of undercooling decreasing with a
Nucleation Kinetics and Grain Refinement / 281 lowering of lattice disregistry. In addition to the basic requirement of lattice compatibility, the melt should tend to wet the surface of the nucleant, and for a uniform refining, the solid nucleant should remain widely dispersed in the liquid. As an ancillary condition, the onset of nucleation should not be followed by a rapid and an isotropic crystal growth, because this will permit the full effect of the potential nucleants to be realized. In alloys, the growth restriction can be achieved as a consequence of solute segregation during freezing, which allows for nucleants throughout the melt to become effective and favors an equiaxed fine grain structure (Ref 24, 25). Aside from these general conditions for an efficient nucleant, it is important to note that for the typical undercooling values of less than 5 C (9 F) observed for grain refinement, the associated contact angle is well below 30 , and thus, the conventional heterogeneous nucleation model represented by Eq 20 cannot be applied for predicting nucleant performance. However, several advances in experimental observations and modeling efforts have yielded a new perspective on grain refinement, and, for the first time, a degree of predictive capability has emerged to provide a valuable basis for the optimization of grain refinement operations. Within the general requirements for effective nucleants, there appear to be two general classes of compounds that are effective in aluminum-base alloys. Similar compound particle types apply to other common casting alloys. The first group includes Al3Ti, Al3Zr, and Al7Cr and can be considered to be associated with a peritectic reaction. The second group comprises compounds such as TiC, TiB2, and AlB2, which are added either intentionally or which result from a chemical reaction in the melt with residual impurities. It is useful to note that the undercooling required to activate solidification is usually less than approximately 5 C (9 F). Despite the rather extensive discussion concerning the mechanism of action and the activity of these proposed nucleants, the strongest evidence for the operation of a compound nucleant is derived from the observed orientation relationships. Table 1 lists disregistry calculations along two directions, a and b, that are normal to each other and are located in the plane in which nucleation is considered to initiate. For the most part, the crystallographic relationships that have been observed involve planes of relatively close packing between the nucleant and aluminum, with disregistries of usually less than approximately 10%. The issue of specific nucleant identity is not insignificant, considering that in commercial melts the rather broad spectrum of background catalysts may often result in solidification at undercoolings of approximately 5 C (9 F). Indeed, only approximately 1 to 2% of all potential nucleant particles in a master alloy addition result in the formation of grains. Furthermore, there is some indication that these proposed nucleants may
not operate singly but rather in association with each other in the melt. Among the nucleant-forming elements listed in Table 1, titanium and, in particular, titanium associated with boron in the master alloy form with a typical composition (Al-5Ti-1B or Al-3Ti-1B) is the most effective grain refinement addition in use for aluminum, but other compositions, such as Al-3Ti-0.15C, are receiving increasing use. In the commercial application of grain-refining agents involving titanium and boron, it is well known that an optimum contact time (that is, the holding time necessary after the addition of refining agents to obtain maximum grain refinement) can vary considerably among different master alloy compositions. As shown in Fig. 10, this feature is generally observed to occur where the optimum contact time interval (period A-B) for minimum grain size is followed by contact time intervals that yield large grain sizes during the interval B-C, which is called a fading process. It is possible to recover from inoculant fading if it is due to nucleant particle settling by stirring the melt, as indicated by the interval C-D.
Grain Refinement Models The operation and underlying mechanism of grain refinement has been under active study for more than 50 years. During this period, numerous models have been proposed to account for the grain initiation and development process. For the most part, the early models served to rationalize existing observations at the time of their introduction but offered little predictive capability. This situation has been changed by the advances in experimental approaches and computational modeling methods so that there is currently a growing predictive capability that can be used for the design of effective grain-refining agents and the optimization of the grain refinement process. The available coverage in this article does not allow for a comprehensive review of all models, but two of the classical models are discussed briefly to provide a perspective for two of the contemporary models. It is generally viewed that Al3Ti crystals are mainly responsible for grain refinement at
alloy compositions above 0.15% Ti, that is, within the peritectic range. For compositions outside the peritectic range, where the Al3Ti phase is not stable under equilibrium conditions, there are basically two interpretations of the grain refinement effect: the carbide-boride view and the peritectic reaction theory. The carbide-boride model suggests that a compound such as AlB2, TiB2, or TiC is responsible for the nucleation of aluminum crystals, because the Al3Ti particles that are introduced into the master alloy are expected to dissolve rapidly (Ref 28, 29). The borides are hexagonal structures, and the carbide is a cubic structure with a relatively close correspondence (Table 1) with the aluminum lattice. Although it is considered that a boride or a carbide is the site of nucleation, the Al3Ti particles do play a role in the overall grain refinement process. The dissolution of Al3Ti is required, for example, to provide a constitutional deterrent (that is, excess titanium) to rapid crystal growth. Within this viewpoint, the establishment of an optimum contact time can be related to the initial dissolution of Al3Ti and to the possible conditioning of the nucleant surface. The subsequent fading reaction is related to the agglomeration and settling of nucleant particles (Ref 29).
Fig. 10
The influence of holding time after the introduction of a grain-refiner master alloy on the degree of grain refinement. The period A-B represents the contact time, the period B-C refers to a fading process, and the period C-D refers to a recovery process.
Table 1 Observed characteristics of potential nucleant compounds Lattice disregistry, % Compound
Crystallographic relationships
Al3Ti
(221)Alk(001)Al3Ti; [010]Alk[113]Al3Ti (001)Alk(001)Al3Ti; [100]Alk[100]Al3Ti (011)Alk(001)Al3Ti; [011]Alk[110]Al3Ti ð111ÞAlk(110)Al3Ti (111)Alk(001)AlB2 None observed (111)Alk(001)TiB2 (001)Alk(001)TiC
Al3(Ti,B)
AlB2 TiB2 TiC
a direction
b direction
0
0.25
5.2
Undercooling (DT)(a)
Ref
C
F
Ref
25
...
...
25
5.2
5,25
3–5
5–9
25
0.6
5.2
25
0
0
25
5.1 5.9 ... 5.9 6.5
14.3 5.0 ... 5.8 6.5
26 26 26 25 27
... 0.5–1.0 0–0.3
... 0.9–1.8 0–0.5
... 25 26
0
0
28
(a) For peritectic compounds, DT is referenced to Tp; for other compounds, DT is referenced to Tf.
282 / Principles of Solidification Although there is evidence for the stability of boride or carbide particles in aluminum melts and reports of observations of these particles in ingots, other observations have indicated a number of problems in attributing nucleation solely to the action of the boride or carbide particles. For example, a consistent orientation relationship is observed between Al3Ti and aluminum. For TiB2 and aluminum, the reported results differ on detection of a consistent orientation relationship, as noted in Table 1. In addition, boride particles have often been observed to be located at grain boundaries rather than within grain centers, which suggests limited effectiveness (Ref 26). Peritectic Reaction Theory. There is an alternate viewpoint of the nucleation behavior during grain refinement that was originally developed on the basis of a peritectic reaction, but this viewpoint has subsequently evolved into a model with more general scope (Ref 30–35). Although many types of particles are believed to be active in the nucleation of aluminum crystals at undercoolings less than approximately 5 C (9 F), the presence of Al3Ti is believed to offer a more active catalyst requiring little or no undercooling. Thus, Al3Ti crystals, when present, will predominate in the grain refinement nucleation. For peritectic alloys such as the one illustrated by the solid lines in Fig. 4, it is clear that Al3Ti can form upon cooling below the liquidus and can promote the formation of aluminum crystals at the peritectic temperature, Tp. Even for compositions in the peritectic range, a properitectic formation of aluminum has been observed at temperatures above Tp that appears to be associated with a metastable reaction (Ref 36–39). At alloy compositions below the peritectic range, a grain refinement effect is expected to be reduced by the loss of Al3Ti particles, which are added externally in master alloy form. Indeed, independent Al3Ti particles in liquid aluminum that is not saturated with titanium dissolve within several minutes. On the other hand, boride particles, that is, TiB2 or (Al,Ti)B2, are observed to be reasonably stable in melts and offer a fine particulate dispersion. To account for the beneficial effect of incorporating boron into the master alloy compositions and also the observation of the interaction between TiB2 and Al3Ti, it appears likely that TiB2 particles may provide an effective substrate for the nucleation of Al3Ti for compositions spanning the peritectic reaction. During cooling of the melt or holding (that is, the contact time period), the TiB2 particles will react slowly with liquid aluminum to form an (Al,Ti)B2 solution and to liberate excess titanium locally to provide for the formation of a sheath of Al3Ti as a coating on the (Al,Ti)B2 particles (Ref 5). This view is supported by observations on the behavior of boride particles following direct injection into an aluminum alloy melt (Ref 40) and the proposed hypernucleation model that deals with preferential adsorption (Ref 41). However, other work
(Ref 39) indicates that the interactions between titanium, boron, and aluminum are complex and there are thermodynamic issues that do not favor Al3Ti development from the melt (Ref 42). Overall, the applicability of the peritectic reaction model to the hypoperitectic composition range is unresolved and appears to be complex (perhaps involving preferential adsorption) compared to the case for the peritectic composition range. Free-Growth Model. In recognition of the difficulties in applying the conventional heterogeneous nucleation model to the conditions of low undercooling and low contact angle, Greer (Ref 43–49) has adapted a model proposed by Turnbull (Ref 50) to treat the evolution of aluminum crystals during grain refinement. The long-standing challenge of observing the early stage of grain refinement that is obscured in bulk castings was overcome by the use of glass-forming aluminum alloys containing a dispersion of grain-refining particles. With this approach, the initial stage of grain refinement could be captured for detailed analysis, as shown in Fig. 11. The basis for the model is derived from transmission electron microscopy observations that aluminum crystals develop initially on the (0001) planes of TiB2 hexagonal prism particles (Ref 45, 51). Moreover, there is evidence that between the aluminum crystal and TiB2, there is some Al3Ti with an orientation relationship between the individual phases, such that f0001gTiB2 kf112gAl3 Ti kf111gAl . The duplex nature of the particles suggests that TiB2 may not be the actual grain initiation site.
Fig. 12
The initial event leading to the aluminum crystal is not specified completely (i.e., it could be conventional nucleation or adsorption on the particles) (Ref 52), but the key tenet of the model is that for a platelet (modeled as a disc of diameter d) following lateral growth to cover the TiB2, further growth must be outward, as illustrated in Fig. 12. The solid aluminum takes
Fig. 11
Bright-field transmission electron micrograph showing a-aluminum crystallites formed on a TiB2 particle embedded in an amorphous Al88Y8Ni5Co2 (at.%) matrix and annealed. The labeled crystallite sits on a (0001) ledge on the ð1010Þ prismatic face of the hexagonal platelet particle that is coated with a thin layer of Al3Ti. Source: Ref 45
The bold curve indicates the undercooling necessary for grain initiation. The free-growth undercooling is calculated; the nucleation undercooling is schematic only. Inset (i) shows the classical spherical cap model for heterogeneous nucleation. Inset (ii) shows a cap of a-aluminum growing on an inoculant particle through the critical hemispherical condition.
Nucleation Kinetics and Grain Refinement / 283
Tfg ¼ 4sLS =Sf d
(Eq 22)
As the undercooling continues to increase, free growth can commence on particles of successively smaller size. This model represents a deterministic description of aluminum grain development that is fundamentally different from a stochastic nucleation description. In the implementation of this model, the cooling path is considered as a sequence of isothermal steps. For each step, the number of grain initiation events is determined as well as the latent heat release from growth of existing grains. The balance between the latent heat release and the external heat extraction yields the cooling rate and the temperature for the next calculation step. Once recalescence begins due to the release of the latent heat, the number of grains per unit volume is established, since there are no further grain initiation events. This approach is similar to the original model of Maxwell and Hellawell (Ref 25), except that they assumed conventional heterogeneous nucleation and a uniform nucleant particle size. At the low undercooling values that are typical for grain refinement, the growth of an aluminum crystal of radius r can be approximated by a spherical growth model to yield a growth rate v: v ¼ dr=dt ¼ ðDL Ts Þ=ðrQÞ
(Eq 23)
where DTs is the solutal undercooling, and Q is a growth restriction parameter defined by: Q ¼ mL ðk 1ÞCo
(Eq 24)
where Co is the alloy solute content. Sometimes, the growth restriction is related to the constitutional undercooling parameter, P = Q/k. A listing of the relevant parameter values for the main solutes of interest in aluminum grain refinement is given in Table 2. From the values given in Table 2, it is evident that titanium is the most effective solute in terms of restricting growth of aluminum grains. In the free-growth model, the measured nucleant particle size distribution is incorporated in the numerical analysis. For TiB2 particles in an Al-Ti-B grain refiner, a log-normal distribution was shown to be appropriate (Ref 45). In a standard grain-refining test based on the TP-1 procedure, the predictions of the free-growth model for grain size as a function of grain-refiner level, alloy solute content, and cooling rate are in good agreement with the observations. For example, in Fig. 13, the
Table 2
Relevant parameters for the main solutes used in aluminum grain refinement
Solute element
Ti Ta V Hf Mo Zr Nb Cr Si Ni Mg Fe Cu Mn
k
7–8 2.5 4.0 2.4 2.5 2.5 1.5 2.0 0.12 0.007 0.49 0.02 0.17 0.62
mL, K/wt%
mL(k 1) = Q/Co
mL(k1)/k = P/Co
Maximun concentration, wt%
Reaction type
26–33 70 10.0 8.0 5.0 4.5 13.3 3.5 6.8 3.3 5.7 3.0 3.4 1.2
156–236 105 30 11.2 7.5 6.8 6.6 3.5 5.9 3.3 2.9 2.9 2.8 0.46
22–29.5 42 7.5 4.7 3.0 2.7 4.4 1.8 49.8 4714 5.9 145 16.5 0.74
0.15 0.10 0.1 0.5 0.1 0.11 0.15 0.4 12.6 6 37 1.8 33.2 2
Peritectic Peritectic Peritectic Peritectic Peritectic Peritectic Peritectic Peritectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic
1013 Number of grains, m–3
the form of a growing spherical cap. At a given undercooling, growth ceases when the radius of curvature reaches rcr for the undercooling, as given by Eq 12. As the undercooling increases, rcr decreases and the crystal can grow accordingly. When rcr has decreased to equal d/2, the solid forms a hemisphere, and the solid can now grow freely to develop into a grain. The critical undercooling, DTfg, for the onset of free growth is given by:
Predicted 1012
Maxwell & Hellawell
1011
(A)
(B)
1010 1010
Measured 1014 1012 1013 1011 Number of particles, m–3
1015
Fig. 13
The number of grains per unit volume as a function of the number of refiner particles per unit volume, showing a general trend to lower efficiency at higher addition level. Data from grain diameters measured in TP-1 tests (closed circles) are compared with predictions of the free-growth model (open circles). Source: Ref 44. The predictions are qualitatively different from those of Maxwell and Hellawell (Ref 25) and are a much better fit to the data.
Grain diameter (mean linear intercept) for commercially pure aluminum inoculated with A1-5Ti-lB at various levels. The grain diameters measured in TP-1 tests (diamonds) are compared with free-growth model predictions based on an isothermal melt (circles) or on directional solidification (squares). Source: Ref 45
number of aluminum grains as a function of the number of nucleant particles is shown and demonstrates an agreement between the freegrowth model analysis and the measurements. Similarly, in Fig. 14, the observed grain size dependence on refiner addition level shows a similar trend with the free-growth model prediction, except at small levels. The free-growth model has also been applied to the grain refinement of magnesium-base alloys by SiC particles and again has revealed the capability to predict the resultant grain size (Ref 53). The agreement between the free-growth model and grain refinement measurements is a strong support for the model, since the input parameters are determined independently and are not adjustable. The apparent success of the free-growth model that is essentially an athermal treatment of grain initiation rather that the usual stochastic nucleation process provides a valuable means to treat grain development at low undercooling, where conventional nucleation does not apply. In fact, the behavior treated by the free-growth model is consistent with some of the modeling approaches used in simulating
the development of solidification microstructure (Ref 54–56). At the same time, the freegrowth model requires further development to treat large casting volumes, where significant liquid flow can contribute to grain initiation from dendrite fragments, and particle agglomeration and settling are important. Similarly, solute effects such as those involved in poisoning are treated only in terms of their effect on Q. However, in multicomponent alloys, strong chemical interactions can influence particle surfaces as well as growth restriction. Constitutional Undercooling Model. Along with the development of the free-growth model analysis of grain refinement, an alternate viewpoint has been proposed by Easton and St. John that incorporates the chemical effects of solute content on restricting grain growth and the structural characteristics of grain refinement as reflected by nucleant potency (Ref 57– 65). In this model, the analysis of constitutional undercooling, DTc, plays a major role in the evolution of the as-cast grain size. It is assumed that following the initial thermal undercooling that is generated by the mold chill, constitutional undercooling is the dominant form of
Fig. 14
284 / Principles of Solidification undercooling. As an additional assumption, it is considered that nucleation initiates when DTc = DTn, the nucleation undercooling, but again, the grain initiation mechanism on grain-refining particles is not specified. At this point, the value of the fraction solid, fs, corresponds to the final grain size, since further growth is arrested. Due to the shallow thermal gradients in the melt, a progression of nucleation events develops from the mold wall toward the casting center. Separate nucleation events in this progression initiate as the previously nucleated grains develop sufficient DTc to reach DTn, which is a measure of nucleant potency. To satisfy this condition, a small amount of growth is necessary and is evaluated by expressing the constitutional undercooling as: Tc ¼ ml c0 1
1 ½1 fs p
(Eq 25)
where p = 1 k. The initial rate of development of constitutional undercooling is then related to the growth restriction factor as: dTc ¼ ml c0 ðk 1Þ ¼ Q dfs
(Eq 26)
By rearranging Eq 25, an expression for the relative grain size (RGS) is obtained in terms of solute content and particle potency as: RGS ¼ fs;n ¼ 1
ml c0 ml c0 Tn
1=p (Eq 27)
where fs,n is the amount of growth required to develop sufficient constitutional undercooling to reach DTn and initiate grain nucleation. The application of Eq 27 to reveal the relationship between grain size and the combined effects of nucleant potency and solute content is illustrated in Fig. 15. It is evident that nucleant potency has a strong effect on grain size for all solute levels, as expected, while solute content is a significant factor at low levels and has little effect at high levels. As a further development in this analysis, the constitutional undercooling model for grain refinement was related to the semiempirical expression: D ¼ a þ b=Q
(Eq 28)
where D is the grain size, and a and b are constants. Since Q is related to the rate of development of constitutional undercooling by Eq 26, the grain size was taken as proportional to (DTn/Q), so that b = (b0 DTnm/Q), where m may be between 1 and 2 to reflect either plane front (m = 1) or dendritic (m = 2) growth. Moreover, it is apparent that D is dependent on the separation between active particles. By assuming that the grain density is proportional to the nucleant particle density, the parameter a was evaluated as [rf ]1/3, where r is the nucleant particle density, and f is the fraction of particles
Fig. 15
Relative grain size versus solute content for a range of nucleant particle potencies, DTn as evaluated from Eq 27. The solute used is titanium added to pure aluminum.
that are active. From these considerations, Eq 28 becomes: a0 b0 Tn ffiffiffiffiffiffi þ D¼p 3 Q rf
(Eq 29)
By replotting results such as those in Fig. 15 in the form indicated by Eq 28 or 29, some useful insight is obtained on the separate factors controlling grain size. For example, in Fig. 16(a), the slope b (or b0 ) is related to the nucleant potency. A steeper slope corresponds to a lower potency. The intercept a (or a0 ) represents the number of nucleant particles that are active in grain initiation at infinite Q values, as shown in Fig. 16(b). One example is given in Fig. 17, where the measured grain size for several aluminum alloys is plotted against (1/Q) for different TiB2 addition levels. The zero level represents the background nucleant effect. With all TiB2 levels shown, the potency has increased to a constant amount, with further grain refinement related to the increased number of particles. The relationships shown in Fig. 16 have been applied to a wide variety of aluminum and magnesium alloys and have also been used to interpret other effects, such as the influence of melt superheating and poisoning solutes on grain size. As with the free-growth model, the constitutional undercooling model requires further development to account for fluid flow effects. Both the free-growth model and the constitutional undercooling model require input information on the particle sizes and number density, but the constitutional undercooling model also requires grain refinement results in order to evaluate the parameters in Eq 28 and 29. Once the model is calibrated, it can provide useful design guidance.
Fig. 16
(a) Effect of changing the nucleant particle potency on the gradient b with a constant number of particles. Particles that produce a slope b00 are the most potent particles. (b) Effect of adding more particles at constant potency on the value of the intercept a
Nucleation Kinetics and Grain Refinement / 285 1200
1000
Grain size, µm
800
0TiB2 0.005TiB2
600
0.01TiB2 0.02TiB2
400
0.05TiB2 200
0 0
0.1
0.2
0.3
0.4
1/Q, 1/K
Fig. 17
Measured grain size versus the reciprocal of the growth restriction factor, Q, for 0, 0.005, 0.01, 0.02, and 0.05% TiB2 additions in several aluminum alloys
Fig. 18
Micrograph of an Al-6Ti master alloy illustrating the range of size and shape of Al3Ti particles
Fig. 19
The Al3Ti structure and the arrangement of the (011) planes. Three successive (011) planes identified by dashed lines contain only aluminum atoms and offer a low disregistry to the atom arrangement on (012) planes of aluminum.
Master Alloy Processing Another important aspect of the grain refinement reaction that has not received much consideration relates to the influence of preparation conditions during master alloy processing and structural differences on the effectiveness of nucleant particles (Ref 5, 33). This feature has taken on new significance with the observation that Al3Ti particles may appear with different morphologies (Fig. 18) and may exhibit several possible crystallographic orientation relationships with aluminum crystals. The degree of refinement is observed to be dependent on the Al3Ti crystal morphology. Essentially, three types of Al3Ti crystals are observed to be present: a flake type exposing (001) planes to the melt that forms at high preparation temperatures (>760 C, or 1400 F), a blocky form that tends to form at low temperatures, and a petallike morphology. There is
some indication that each particle morphology can differ in nucleating potency. This has been examined in droplet studies where individual particles can be isolated into separate droplets (Ref 37, 38). The blocklike crystals expose (011) planes to the melt, as indicated in Fig. 19, for a [(011) Al3Tik (012) Al] relationship with aluminum. For flake-type crystals, the crystallographic axes of Al3Ti and aluminum are closely parallel, that is, (001)Al3Tik(001)Al. Although the latter relationship also gives a relatively small disregistry, the (011) Al3Ti surface as indicated in Fig. 19 is occupied by three successive layers of aluminum atoms, which can be viewed as somewhat distorted (012) aluminum lattice planes. The formation of aluminum crystals on such a surface merely requires a growth (that is, epitaxy) of this existing distorted lattice at little or no undercooling below the peritectic temperature Tp. The (001) Al3Ti surface, on
the other hand, contains titanium atoms so that a simple growth of the existing surface is not sufficient to generate an aluminum crystal, and a lower nucleating potency may be expected and is observed. In addition, the observation of different morphologies and twinning for Al3Ti crystals may account for the variety of orientation relationships that have been observed between Al3Ti and aluminum. Other particles, such as TiB2, TiC, TiN, ZrB2, and TaB2, have been examined, but the undercooling required to initiate solidification was minimized when Al3Ti particles were present. Usually, the grain-refining particles are added externally to the melt, but there is recent work that has demonstrated an in situ nucleant process that can operate on a commercial scale (Ref 66). From grain-refining tests with in situ nucleant particles, it was observed that the ascast grain size did not depend on the nucleant particle size, which is not consistent with the predictions of the free-growth model. Since the particle surface in contact with the melt is the critical point in grain initiation, it is possible that ex situ particles have size-dependent surface characteristics that are absent for in situ particles. Thermal Analysis Techniques. Grain-refining characteristics can also be conveniently assessed through thermal analysis techniques. In examining cooling curves after master alloy dilution, three main groups of characteristic curves are often observed, as shown in Fig. 20. The characteristic curve A in Fig. 20 is detected for a high degree of grain refinement and is related to a large number of aluminum crystals that form on surviving Al3Ti particles
286 / Principles of Solidification
Fig. 20
Representative types of cooling curves for an aluminum melt treated with different master alloys to yield varying levels of grain refinement. A, high degree of grain refinement coincides with Al3Ti crystals present in melt; B, poor grain refinement due to few Al3Ti crystals in melt; C, intermediate case in which there are a sufficient number of Al3Ti crystals but an insufficient local titanium concentration to promote grain refinement. The characteristic peritectic hump for all three curves lies between the growth temperature, TG, and the equilibrium liquidus temperature, TL. Source: Ref 31
at Tp. With continued cooling upon reaching the growth temperature, TG, the aluminum crystals can grow, consuming the liquid and raising the temperature to the equilibrium liquidus temperature, TL. This reaction creates a peritectic hump, which is an indication that Al3Ti crystals are present in the melt. The type A cooling curve in Fig. 20 was detected with blocky Al3Ti crystals following a short holding time but also with flake-type Al3Ti after a long contact time of 1 h or more. With a cooling curve of the type B configuration in Fig. 20, relatively few Al3Ti crystals remain; therefore, poor grain refinement occurs. A cooling curve of the type C configuration in Fig. 20 represents an intermediate case in which there are a sufficient number of Al3Ti crystals, but the local titanium concentration is not sufficient to allow for growth at high temperatures, a condition that yields poor grain refinement. Although the results of recent work represent a significant improvement in the description of grain refinement as well as in the improved characterization and predictive modeling of grain refinement agents and their effectiveness, the overall understanding of the grain refinement of aluminum or other metallic melts by compound particles remains incomplete. The success of the contemporary models reveals that while nucleant particle size is an important factor in grain refinement effectiveness, constitutional factors are also significant. Other chemical effects such as adsorption remain to be considered as well as the influence of fluid flow. The issue of the actual grain initiation mechanism is also open, and the resolution is likely to be related to the details of the nucleant
surface structure. However, the new models seem to provide the necessary foundation for further significant progress in the development of a real understanding of grain refinement.
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13. D. Turnbull, J. Appl. Phys., Vol 21, 1950, p 1022 14. F. Spaepen, Solid State Phys., Vol 47, 1994, p 1 15. J.J. Hoyt and M. Asta, Phys. Rev. B, Vol 65, 2002, p 214106 16. D.Y. Sun, M. Asta, J.J. Hoyt, M.I. Mendelev, and D.J. Srolovitz, Phys. Rev. B, Vol 69, 2004, p 020102 17. R.L. Davidchack, J.R. Morris, and B.B. Laird, J. Chem. Phys., Vol 125, 2006, p 094710 18. A. Inoue, Acta Mater., Vol 48, 2000, p 279 19. K.F. Kelton, A.L. Greer, and C.V. Thompson, J. Chem. Phys., Vol 79, 1983, p 6261 20. K.F. Kelton and A.L. Greer, J. Non-Cryst. Solids, Vol 79, 1986, p 295 21. “Standard Test Procedure for Aluminum Alloy Grain Refiners: TP-1,” The Aluminum Association, 1987 22. D. Turnbull and B. Vonnegut, Ind. Eng. Chem., Vol 44, 1952, p 1292 23. P.B. Crosley, A.W. Douglas, and L.F. Mondolfo, The Solidification of Metals, The Iron and Steel Institute, 1968, p 10 24. A. Hellawell, Solidification and Casting of Metals, The Metals Society, 1979, p 161 25. I. Maxwell and A. Hellawell, Acta Metall., Vol 23, 1975, p 895 26. I.G. Davies, J.M. Dennis, and A. Hellawell, Metall. Trans., Vol 1, 1970, p 275 27. J. Cisse, G.F. Bolling, and H.W. Kerr, J. Cryst. Growth, Vol B-14, 1972, p 777 28. A. Cibula, J. Inst. Met., Vol 76, 1949– 1950, p 321; Vol 80, 1951–1952, p 11 29. G.P. Jones and J. Pearson, Metall. Trans. B, Vol 7, 1976, p 223 30. F.A. Crossley and L.F. Mondolfo, Trans. AIME, Vol 141, 1951, p 1143 31. J.H. Perepezko and S.E. LeBeau, in Aluminum Transformation—Technology and Applications, American Society for Metals, 1982, p 309 32. J.A. Marcantonio and L.F. Mondolfo, J. Inst. Met., Vol 98, 1970, p 23 33. M.M. Guzowski, G.K. Sigworth, and D.A. Senter, Metall. Trans. A, Vol 18, 1987, p 603 34. L.F. Mondolfo, in Grain Refinement in Castings and Welds, G.J. Abbaschian and S.A. David, Ed., The Metallurgical Society, 1983, p 3 35. D.A. Granger, “Practical Aspects of Grain Refining Aluminum Alloy Melts,” Laboratory Report 11-1985-01, Aluminum Company of America, 1985 36. E. Nes and H. Billdal, Acta Metall., Vol 25, 1977, p 1031 37. M.K. Hoffmeyer and J.H. Perepezko, Scr. Metall., Vol 23, 1989, p 315 38. M.K. Hoffmeyer and J.H. Perepezko, Light Metals 1989, P.G. Campbell, Ed., TMS, 1989, p 913 39. M. Johnsson, L. Backerud, and G.K. Sigworth, Metall. Trans. A, Vol 24, 1993, p 481 40. P.S. Mohanty and J.E. Gruzleski, Acta Metall. Mater., Vol 43, 1995, p 2001
Nucleation Kinetics and Grain Refinement / 287 41. G.P. Jones, in Solidification Processing 1987, Institute of Metals, London, 1988, p 496 42. G.K. Sigworth, Scr. Mater., Vol 34, 1996, p 919 43. A.M. Bunn, P.V. Evans, D.J. Bristow, and A.L. Greer, in Light Metals 1998, B. Welch, Ed., The Minerals, Metals & Materials Society, 1998, p 963 44. A.L. Greer, A.M. Bunn, A. Tronche, P.V. Evans, and D.J. Bristow, Acta Mater., Vol 48, 2000, p 2823 45. A.L. Greer, Philos. Trans. R. Soc. (London) A, Vol 361, 2003, p 479 46. T.E. Quested and A.L. Greer, Acta Mater., Vol 52, 2004, p 3859 47. T.E. Quested and A.L. Greer, Acta Mater., Vol 53, 2005, p 2683, 4643 48. A.L. Greer and T.E. Quested, Philos. Mag., Vol 86, 2006, p 3665 49. S.A. Reavley and A.L. Greer, Philos. Mag., Vol 88, 2008, p 561
50. D. Turnbull, Acta Metall., Vol 1, 1953, p 8 51. P. Schumacher, A.L. Greer, J. Worth, P.V. Evans, M.A. Kearns, P. Fisher, and A.H. Green, Mater. Sci. Technol., Vol 14, 1998, p 394 52. W.T. Kim and B. Cantor, Acta Metall. Mater., Vol 42, 1994, p 3115 53. R. Gu¨nther, Ch. Hartig, and R. Bormann, Acta Mater., Vol 54, 2006, p 5591 54. W. Oldfield, Trans. ASM, Vol 59, 1966, p 945 55. J.D. Hunt, Mater. Sci. Eng., Vol 65, 1984, p 75 56. M. Rappaz and Ch.-A. Gandin, Acta Metall. Mater., Vol 41, 1993, p 345 57. M. Easton and D. St. John, Metall. Mater. Trans. A, Vol 30, 1999, p 1613 58. Y.C. Lee, A.K. Dahle, and D.H. St. John, Metall. Mater. Trans. A, Vol 31, 2000, p 2895 59. M.A. Easton and D.H. St. John, Acta Mater., Vol 49, 2001, p 1867
60. M.A. Easton and D.H. St. John, in Light Metals 2001, J.I. Anjier, Ed., The Minerals, Metals & Materials Society, 2001, p 927 61. M. Qian, D.H. St. John, and M.T. Frost, Scr. Mater., Vol 50, 2004, p 1115 62. D.H. St. John, M. Qian, M.A. Easton, P. Cao, and Z. Hildebrand, Metall. Mater. Trans. A, Vol 36, 2005, p 1669 63. M. Easton and D.H. St. John, Metall. Mater. Trans. A, Vol 36, 2005, p 1911 64. L. Lu, A.K. Dahle, and D.H. St. John, Scr. Mater., Vol 54, 2006, p 2197 65. P. Cao, M. Qian, and D.H. St. John, Scr. Mater., Vol 56, 2007, p 633 66. J.A. Megy, D.A. Granger, G.K. Sigworth, and C.R. Durst, in Light Metals 2001, J.I. Anjier, Ed., The Minerals, Metals & Materials Society, 2001, p 943
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 288-292 DOI: 10.1361/asmhba0005208
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Transport Phenomena During Solidification Jonathan A. Dantzig, University of Illinois at Urbana-Champaign
Fundamentals Some of the most important tools that materials scientists use regularly are equilibrium phase diagrams. These diagrams, derived from thermodynamics, define the phases present and their relative amounts in equilibrium as a function of overall alloy composition, C, and temperature, T. However, there is no information in the diagram about how a system evolves in space, x, and time, t. Transport phenomena is the generic name used to describe the dynamical and spatial aspects of the system behavior. There are more specific names. The study of energy transport, embodied in the temperature field T(x,t), is called heat transfer. The transport of solute, given by C(x,t), is called mass transfer. In fluid systems, it is also necessary to consider the velocity field v(x,t), the study of which is called momentum transport or, more commonly, fluid dynamics. Clearly, these subjects are much deeper and broader than the scope of a Handbook article. Therefore, the focus here is on an introduction to the topic, all given in the context of solidification. One begins with the balance equations for mass, energy, and solute and the necessary boundary conditions for solving problems of interest in casting and solidification. The phenomena cover a vast range of length and time scales—from atomic dimensions up to macroscopic casting size, and from nanoseconds for interface attachment kinetics to hours for casting solidification. It is essential to understand how to determine which phenomena are most important at the particular length and time scale of interest for the problem at hand. Methods for doing so are described in the section “Scaling” in this article. Finally, several examples are given of the application of transport phenomena in solidification, focusing in particular on microstructure formation.
Balance Equations The derivation of the general forms of the balance equations is given in Ref 1 and 2. Instead, a shorter form is presented here that begins with
the general balance equations and then specializes them using common models for materials behavior appropriate for metal systems. Since the subject of interest is solidification, it is necessary to examine more closely how these equations are applied when there is a moving solidification front. This should become more clear as each balance equation is developed. Mass Balance. The mass balance for a fluid of density r moving at velocity v is given by:
@r þ r ðrvÞ ¼ 0 @t
(Eq 1)
This is written in the so-called conservative form, which is quite useful if one decides to integrate over a control volume. Doing so, and choosing a fixed control volume V with surface S, yields: ð ð @ rdV þ rv ndS ¼ 0 @t V
S
(Eq 2)
where the divergence theorem has been applied to convert the second term to a surface integral. The surface normal vector is n. The interpretation of Eq 2 is clear: The time rate of change of total mass in the volume, represented by the first term, is balanced by the total mass leaving the volume through the surface. Now consider the control volume shown Fig. 1, which encloses a portion of the solid-liquid interface. The interface, solid, and liquid are moving at velocities v*, vs, and vL, respectively, and the interface normal vector is n*. (The subscripts “s” and “L” are used throughout to refer to the solid and liquid, respectively, and the superscript “*” indicates quantities on the liquid-solid interface.) Now evaluate Eq 2 as the thickness of the control volume goes to zero. The first term vanishes, and the second term then consists of four parts:
rs ðvs n v n Þ
where b is called the solidification shrinkage. Equation 4 states that there is a net flow of liquid into the interface in order to balance the density difference between the liquid and solid phases. This flow is the underlying cause of porosity and hot tearing in castings. It is less important for the other balances. Therefore, the shrinkage flow is omitted from the development of the momentum, energy, and solute balances that follow. Momentum Balance. The balance of linear momentum is the continuum form of Newton’s second law of motion, that is, that the sum of the forces on a body is balanced by the time rate of change of linear momentum. In conservative form:
where s is the stress tensor, and b is a body force such as gravity. The dyadic product vv is a tensor with components vivj. Integrating over a control volume as for the mass balance shows more clearly the relation to Newton’s second law: ð ð @ rvdV þ rvv ndS @t v S ð ð ¼ f ndS þ rbdV
S
(Eq 6)
v
where f = s n is the external traction force applied on the surface of the control volume. It is often convenient to separate the stress into vL n∗ v∗
(Eq 3) vs
The physical meaning of Eq 3 is made more clear by assuming that the solid is not moving (vs = 0) and by rearranging terms to obtain:
(Eq 5)
Liquid
þ rL ðv n vL n Þ ¼ 0
vL
@rv þ r ðrvvÞ ¼ r s þ rb @t
r rL n ¼ s v n ¼ bv n rL
(Eq 4)
Solid Interface
Fig. 1
Schematic of a control volume including the liquid-solid interface
Transport Phenomena During Solidification / 289 hydrostatic pressure, p, and the “extra” stress tensor, t, leading to the form s = pI + t. Following a procedure similar to the one followed to develop the interface form of the mass balance leads to balance equations for the normal and tangential components of the force at the interface: ðpL ps Þ ðts tL Þ : ðn n Þ ¼ 2g k
(Eq 7)
ðts tL Þ : ðn t Þ ¼ rsurf g t ¼ 0
(Eq 8)
where g is the surface energy of the liquid-solid interface, and k is the mean curvature. The notation “:” implies a double-dot product, and rsurf indicates the gradient within the surface. This term captures the so-called Marangoni effect, where gradients in temperature or composition along the surface create tangential forces that can drive a flow in the liquid. These forces are very important for welding, where there is a free surface, but less so for liquids and solids. Equation 7 provides the well-known relation that the surface tension is balanced by a pressure difference across the curved interface. Energy Balance. The energy balance is one of the most important balances for solidification problems. The most convenient thermodynamic variables are temperature, T, and pressure, p. The energy balance is written in terms of the enthalpy density, H, as:
@rH þ r ðrHvÞ @t ¼ r ðkrT Þ pr v þ t : D þ Q_
(Eq 9)
where D is the rate of deformation tensor, 2D = rv + (rv)T, and t:D is the work done by the shear stress, also called the viscous dissipation. For most solidification problems, the pressure effects, viscous dissipation, and internal heat generation, Q_ , usually can be neglected, leaving the form:
@rH þ r ðrHvÞ ¼ r ðKrT Þ @t
(Eq 13)
Often, the ratio Ks/KL 1. Solute Balance. The solute balance is developed in a similar way to the energy balance. For the sake of brevity, the treatment here is restricted to binary alloys (A and B) with no chemical reactions. In analogy with Fourier’s law for conduction, Fick’s law for diffusion is adopted, which gives the diffusive flux of species B as proportional to the composition gradient, J = Dr(rC). The composition is given as a mass fraction, that is, grams B divided by total grams, the units of the diffusivity D are m2/s, and thus, the flux has units grams B per unit area per unit time. Note that there are numerous subtleties associated with whether the mass appears inside or outside the gradient operator when the density is not constant, but these issues are beyond the scope of this article. The solute balance is given by:
@rC þ r ðrCvÞ ¼ r ðDrrCÞ @t
(Eq 14)
and the interfacial balance for solute is also analogous to the interfacial balance for energy:
ðrL CL rs Cs Þv n ¼ Ds rðrs Cs Þ n
DL rðrL CL Þ n (Eq 15)
As opposed to heat conduction, where Ks KL, usually Ds DL, and the first term on the right side of Eq 15 is often neglected. The left side is usually modeled by assuming local thermodynamic equilibrium at the interface, and therefore the compositions CS and CL are given by the equilibrium phase diagram. In that case, the segregation coefficient, k, is introduced, such that rsCs = krLCL. Introducing both of these simplifications in Eq 15 gives:
rL CL ð1 kÞv n ¼ DL rðrL CL Þ n
(Eq 16)
Scaling
ð ð ð @ rhdV þ rhv ndS ¼ KrT ndS @t V S S ð ¼ q ndS
(Eq 11)
where Fourier’s law for the conductive heat flux, q = KrT, has been applied. The interfacial balance, neglecting the shrinkage flow, is:
rs H ¼ Ks rTs n KL rTL n
(Eq 10)
Integrating over the control volume shows that the time rate of change of enthalpy in the control volume is balanced by that advected and conducted through the surface:
S
Inserting DH gives the Stefan condition:
rL HL rs Hs ¼ Ks rTs n KL rTL n
(Eq 12)
The left side of Eq 12 is the difference in enthalpy between the liquid and solid phases, the latent heat of fusion, DH; the right side is the net heat conduction away from the interface.
The governing equations, while necessary, are only the beginning. The geometry, boundary and initial conditions, and material properties also must be specified. When all of these data are gathered, one can, in principle, evolve the appropriate equations in time and find a solution to the problem at hand. This is usually done numerically, perhaps using one of the many available commercial codes. However, that is usually much more work than is either necessary or even possible, because the range of length and time scales involved with solidification processes is vast—from atomic-scale kinetic processes that take place over picoseconds to macroscopic heat flow at the size of the casting, extending over as much as an hour for large parts. Practical simulations can cover at most 3 to 4 orders of magnitude in both space and time, and this means that some phenomena must be omitted from the problem. But which ones should be kept and which ones
neglected (or are better to include via an analytical model)? Clearly, the important scales should be included, but how can one decide which those are? That is the subject of this section. The decision about which length and time scales are important is done by scaling the governing equations. This provides a systematic means for ensuring that the important terms are kept while the negligible terms are discarded. First, the procedure is described, then a simple example is given. Scaling begins by choosing characteristic scales for variables such as length, temperature, and so on that usually come from the problem geometry and boundary conditions. Some characteristic values are unknown and are defined by the governing equations. The characteristic values are used to define dimensionless variables that are order one. For example, suppose that the temperature, T, is known (e.g., from initial and boundary conditions) to lie in the interval T1 T T0. A dimensionless temperature is defined, y = (T T1)/(T0 T1) E [0,1]. Unknown characteristics are just given a symbol and carried along. When the dimensionless variables are defined, they are substituted into the governing equations. Each term in the equation then has a (dimensional) coefficient multiplying a scaled (dimensionless) term. The equation is then divided by the coefficient of a term that is expected to be important. This requires some engineering judgment, but it is expected that something about the problem at hand is known! The unknown characteristics are then evaluated by setting the value of the coefficient where they appear to 1. Now that all of the terms in the equation are order 1 or less, terms whose coefficients are small compared to 1 can be safely neglected and the remaining problem solved. A concrete example makes the scaling process seem much less mysterious. Consider a thin, circular rod of length L and radius R L, initially at temperature T0, quenched into a large bath at temperature T1. Some simplifying assumptions are made in order to focus on the scaling process. Assume that the bath temperature stays at T1, unaffected by the quench, and that the heat transfer from the surface of the rod to the bath can be characterized by a constant heat transfer coefficient, h. All of the material properties will be assumed to be constant, and the temperature distribution is axisymmetric. The problem is solved in polar coordinates. Finally, the rod is solid (v = 0), and there is no internal heat generation (Q_ ¼ 0). The energy equation for the rod (0 r R; L/2 z L/2) is then: @T 1@ @T @2T ¼a r þ 2 @t r @r @r @z
(Eq 17)
and the initial and boundary conditions are: T ¼ T0 K
@T ¼ hðT T0 Þ @r
t¼0
(Eq 18)
r¼R
(Eq 19)
290 / Principles of Solidification @T þK ¼ hðT T0 Þ @z
z ¼ L=2
(Eq 20)
where a = K/rcp is the thermal diffusivity. Now the problem specification is used to define the following dimensionless variables: r~ ¼
r ; R
z¼
z þ L=2 ; L
y¼
T T1 T 0 T1
(Eq 21)
Time must also be scaled, but the characteristic value tc is not yet known. Therefore, t~ ¼ t=tc is defined, and guidance is awaited. The scaled variables are now inserted into Eq 17, and the entire equation is divided by a(T0 T1)/R2, the coefficient of the term associated with conduction in the radial direction. The result is: R2 @y 1@ @y R2 @ 2 y ¼ r~ þ 2 2 atc @ t~ L @ r~ @~ r @~ r
(Eq 22)
Setting the dimensionless group on the left side to 1 gives the characteristic time tc = R2/a, and the ratio at/R2 is called the Fourier number. It can be seen that the geometry of the problem, where R L, indicates that axial conduction, associated with the term @ 2y/@z2, is negligible in comparison to radial conduction. This makes the problem effectively one-dimensional, a great simplification. The initial condition is simply y = 1 at t~ ¼ 0. The scaled form of the boundary conditions becomes: @y hR ¼ y @~ r K
r~ ¼ 1
(Eq 23)
@y hR L y þ ¼ @ K R
z ¼ 0;1
(Eq 24)
The dimensionless parameter hR/K is called the Biot number, Bi. Notice that information about the problem geometry is retained in the aspect ratio L/R, which now appears in the boundary condition. To review concepts of the scaling: The characteristic time for conduction is R2/a, conduction in the axial conduction can be neglected with respect to radial conduction for R/L, and the important parameter that determines the rate of heat loss to the environment is the Biot number, hR/K. The remaining one-dimensional problem can be solved, for example, using Fourier series
methods (Ref 3). The simplification to a onedimensional problem came at a price, however. Since the derivatives with respect to z were eliminated from the differential equation, the end conditions in Eq 24 cannot be satisfied. It is expected that there will be a small region near the boundary, called a boundary layer, where the solution is two-dimensional. In this region, the axial and radial conduction terms must be comparable, and scaling should make it obvious that this layer extends a few R from each end. Figure 2 shows temperature solutions for a problem where Bi = 10 and L/R = 20, at several values of t~. Notice that the temperature is nearly uniform at T1 at t~ = 0.7—this shows the significance of the characteristic time.
Transport and Microstructure Using the fundamentals presented in the preceding sections, solidification problems now can be discussed. Perhaps the most important thing to understand is that for most materials, Ds DL aL as, and these large differences produce a separation of characteristic time and length scales. Because the thermal diffusivity is much larger than the chemical diffusivities, the solidification rate is controlled by heat flow, and chemical diffusion tries to follow. This eventually leads to the formation of microstructure, such as that shown in Fig. 3 (Ref 4). Several aspects of the solidification process are now considered as well as how transport phenomena are expressed in the development of microstructure.
Planar Front Growth As discussed in other articles in this Handbook, the two processing variables that are used to correlate microstructures are the temperature gradient in the liquid at the interface, GL, and
the solidification front speed, vn*. For this reason, a common method for either assessing microstructural features experimentally or actually controlling them in practice is directional solidification. In a simplified version of the experiment, a thin layer of alloy is encapsulated between microscope slides that are then placed atop controlled hot and cold blocks. The blocks are maintained at predetermined temperatures and at a fixed distance, in order to establish a desired temperature gradient, GL. The slide assembly is then pulled at constant speed vn* through the gradient. The experiment is often done with a transparent organic alloy, and the solidification morphology is observed using a microscope. (Recently, similar experiments have been performed in metallic systems using synchrotron radiation to observe the microstructure (Ref 5). Next, a model of the system is built with the goal of finding out how the important parameters interact to produce the observed microstructural features shown in Fig. 3. It is observed that if the pulling speed is slow enough, the interface remains flat. In this case, after a brief transient at the beginning of the experiment, the process reaches a quasi-steady state where the microstructure is constant in time. It is convenient to analyze the process in a frame that is fixed on the interface (z = 0) and to assume that in this frame the transport is independent of time. The material moves through the interface at velocity vn*. It is assumed that the material properties are all constant (thus there is no shrinkage flow) and, for convenience, that the thermal diffusivities of the solid and liquid are equal. Finally, variations in all fields perpendicular to the pulling direction are ignored, making the problem one-dimensional. Note that all these assumptions are not essential to the analysis, and either everything could be included or the assumptions could be justified through a scaling analysis similar to that given for the quenched rod. For the sake of brevity, they are simply taken
0.8 t = 0.05
0.6 0.4
t = 0.25
0.2
t = 0.70
Fig. 2 Temperature distributions through the midplane of a quenched rod for several values of the dimensionless time
Fig. 3
Micrographs from directional solidification of a succinonitrile-acetone alloy. Left: initial planar interface with small perturbations; right: developing dendrites, with primary and secondary arms indicated. Source: Ref 4
Transport Phenomena During Solidification / 291 as given. The transport equations for heat and solute now reduce to: vn
@T @2T ¼a 2 @z @z
vn
@CL @ 2 CL ¼ DL @z @z2
vn
@Cs @ Cs ¼ Ds @z @z2
1
2
0z z0
(Eq 25) (Eq 26) (Eq 27)
Looking at the micrograph in Fig. 3, it seems that the system chooses a length scale l1, the primary dendrite arm spacing. These equations are partially scaled by defining z = z/l1, with the result: v l1 @T @2T n ¼ 2 a @z @z vn l1 @CL @ 2 CL ¼ DL @z @z2 v l1 @CS @ 2 CS ¼ n DS @z @z2
(Eq 28) (Eq 29) (Eq 30)
Each equation has as a prefactor a variation of a dimensionless group called the Pe´clet number, corresponding to the ratio of advective to diffusive transport of heat or solute, depending on the equation. Actual values can be taken from the experiments for the material properties, vn* and l1, to evaluate the Pe´clet numbers: vn l1 102 ; a
vn l1 1; DL
vn l1 104 Ds
(Eq 31)
Substituting these results into Eq 28 to 30 allows terms to be dropped that have coefficients that are much less than 1. It is also convenient to then redimensionalize the equations so that boundary conditions can be applied. The result is: @2T 0 @z2 @CL @ 2 CL ¼ DL @z @x2 @CS 0 @x
vn
(Eq 32) (Eq 33) (Eq 34)
The solution to Eq 32 is a simple linear profile. Applying the boundary conditions T = Tcold at z = zcold and T = Thot at z = zhot gives: T ¼ Tcold þ
Thot Tcold ðz zcold Þ zhot zcold
(Eq 35)
The solution clearly shows how the apparatus selects the temperature gradient. The result that the temperature field is linear and undisturbed by the moving interface is sometimes called the frozen temperature approximation, but it is really a result of scaling. The solution for the composition field in the liquid can be found after applying the boundary
λ2
condition given in Eq 16 at z = 0 and setting CL = C0 for z ! 1: CL ¼ C0 þ C0
1 k vn z=DL e k
(Eq 36)
Thus, there is a solute boundary layer ahead of the interface of thickness roughly 3DL/vn*. The solution to Eq 33 is just Cs is constant. Examination of Eq 36 shows that CL = C0/k at z = 0, and this permits evaluation of the constant, so that: Cs ¼ C0
L Z
X
(Eq 37)
This example, although rather simple, indicates how the transport phenomena naturally separate into different length scales due to the disparity in the rate of transport of energy and solute. This is very important for understanding the origin of microstructure during solidification. It is beyond the scope of this article to pursue the analysis much further, but it is important to understand that there is one other important aspect. On considering the stability of the planar front growth morphology, it is found that for small values of the ratio GL/vn*, the interface becomes unstable to perturbations and breaks down first into cells and then into dendrites. In the stability analysis, a length scale associated with the surface energy of the liquid-solid interface enters the problem through a material property called the capillary length, d0 = Gk/mLC0(1 k), where G is the Gibbs-Thomson coefficient, and mL is the slope of the liquidus line. The characteristic length selected by the dendritic arms can be correlated to the geometric mean of the two length scales, that is: l
pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi d0 ðDL =vn Þ
Microsegregation After the initially flat interface becomes unstable, it evolves into the dendritic pattern shown in the micrograph in Fig. 3. The structure is characterized by the primary (l1) and secondary (l2) dendrite arm spacings. In the previous analysis of transport, l1 is chosen as the length scale. At a finer scale, such as in the interdendritic region between secondary arms, l2 would be the appropriate length scale. Figure 4 illustrates the geometry. The fundamental concepts of scaling can be applied to this new problem. As illustrated, the array of secondary arms is approximately periodic, and a control volume can be chosen as shown, with dimensions l2 L and the expectation that l2 L. This time, there is no velocity, because the control volume is fixed in the dendrite array. Writing the governing equations for energy and solute within this volume yields: 2 @T @ T @T þ ¼a @t @x2 @z
(Eq 38)
Fig. 4
Schematic drawing of the local geometry of secondary dendrite arms
2 @CL @ CL @CL ¼ DL þ @t @x2 @z 2 @Cs @ Cs @Cs ¼ Ds þ 2 @t @x @z
(Eq 39) (Eq 40)
Next, some scaled variables are defined: x z ; z¼ ; l2 =2 L T Tsol y¼ Tliq Tsol
x¼
t¼
t ; tf
(Eq 41)
where tf is the local freezing time, that is, the time it takes to go from the liquidus temperature, Tliq, to the solidus, Tsol. Substituting these scaled variables into the governing equations and invoking the assumption that l2 L gives: @y 4atf @ 2 y ¼ 2 @t l2 @x2
(Eq 42)
@CL 4DL tf @ 2 CL ¼ @t l22 @x2
(Eq 43)
@Cs 4Ds tf @ 2 Cs ¼ @t l22 @x2
(Eq 44)
Metallurgists often use measured correlations between l2 and tf, which show that for aluminum alloys (Ref 6): 1=3 l2 105 m=s1=3 tf
(Eq 45)
For these same alloys, a 106 m2/s, DL 1010 m2/s, and Ds 1014 m2/s. Combining these data with Eq 45, one finds that for solidification times between approximately 10 s and 1 day, Eq 42 to 44 reduce to: @2y 0 @x2
(Eq 46)
@ 2 CL 0 @x2
(Eq 47)
@Cs 0 @t
(Eq 48)
The dendritic array is periodic, so the derivatives @y/@x and @CL/@x must be zero at both
292 / Principles of Solidification x = 0 and x = 1. From this, it is concluded that both the temperature and the composition in the liquid are constant, and that the composition in the solid does not change with time. However, the composition at the liquid-solid interface must adjust to maintain local equilibrium. Applying the solute balance at the interface from Eq 16 gives a relation between the composition of the solid, Cs, and the solid fraction, fs, known as the Scheil equation: Cs ¼ kC0 ð1 fs Þk1
(Eq 49)
(This is probably not completely obvious from what has been presented here—see Ref 1 or 7 for a more thorough derivation.) Equation 49 can be recast in terms of temperature using the phase diagram, with the result: fs ¼ 1
T Tf Tliq Tf
needed to homogenize the structure is simply l22 =Ds . A conservative approach would be to homogenize the structure for two to three times this value. This implies that the microstructure, in the form of l2, is known. From where does that come? It is correlated with the local solidification time, as in Eq 45, but how does one obtain tf? This comes from solving the transport equation for heat at the macroscopic scale, for example, using a commercial casting solidification analysis package. Note that there is also a connection between the microsegregation and heat flow in the code, because the user usually must specify the enthalpy as a function of temperature, which is obtained from Eq 50 from the definition: ðT HðT Þ ¼
1=ðk1Þ (Eq 50)
This equation is very important, because it provides a very good estimate for the microsegregation in cast alloys.
Homogenization The microsegregation described in the previous section is inevitable in normal casting practices. This is often undesirable, and many castings are heat treated to reduce the segregation. A detailed analysis of the problem could be done, but it really is not necessary to do so if one wishes only to estimate the proper homogenization time. It is known that the segregation is more or less periodic on the length scale of the secondary dendrite arm spacing. Based on scaling, then, the estimated time
cp dT þ Hð1 fs Þ
(Eq 51)
298
The integral begins at 298 K (77 F), where enthalpy is taken to be zero by convention.
Summary The governing equations for transport based on balance of mass, momentum, energy, and solute were written. Equations for these balances were developed for a moving solidification front. It was found that the differences in density, enthalpy, and composition between the liquid and solid are very important for the understanding of solidification processes. Considerable effort was spent to understand how material properties and geometry can be analyzed in the context of the governing equations, a process called scaling.
Several example problems were considered to show how the hierarchy of time and length scales associated with transport leads to important features of cast microstructures. An analysis of planar front growth revealed the important length scale associated with diffusion DL/vn*. On the scale of secondary dendrite arms, the limited transport of solute in the solid, combined with rapid transport in the liquid, leads to the Scheil equation for microsegregation. Finally, the principles of transport and scaling are applied to homogenization processes to ameliorate the undesirable segregation pattern. REFERENCES 1. J.A. Dantzig and C.L. Tucker III, Modeling in Materials Processing, Cambridge University Press, New York, 2001 2. R.B. Bird, W.E. Stewart, and E.N. Lightfoot, Transport Phenomena, 2nd ed., Wiley, New York, 2002 3. H.S. Carslaw and J.C. Jaeger, Conduction of Heat in Solids, 2nd ed., Oxford University Press, London, 1959 4. R. Trivedi and K. Somboonsuk, Constrained Dendritic Growth and Spacing, Mater. Sci. Eng., Vol 65, 1984, p 65–74 5. R. Mathiesen and L. Arnberg, X-Ray Monitoring of Solidification Phenomena in Al-Cu Alloys, Mater. Sci. Forum, Vol 508, 2006, p 69–75 6. T.F. Bower, H.D. Brody, and M.C. Flemings, Measurements of Solute Redistribution in Dendritic Solidification, Trans. AIME, Vol 236, 1966, p 624 7. W.F. Kurz and D.J. Fisher, Fundamentals of Solidification, Trans-Tech, Aedermannsdorf, 1984
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 293-298 DOI: 10.1361/asmhba0005209
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Plane Front Solidification Wilfried Kurz, Ecole Polytechnique Fe´de´rale de Lausanne, Switzerland
PLANE FRONT SOLIDIFICATION (PFS) is not a commonly used process since it is generally difficult to obtain the required conditions. It is only in cases where a crystal has to be highly homogeneous across the section that the corresponding processes are developed. One impressive example of PFS is the industrial production of large silicon single crystals, used mainly as substrates for integrated circuits. Other examples of PFS can be found in more basic work of solidification phenomena; the morphological stability of plane fronts at both small and large growth rates has generated substantial interest. This part is fundamental for the theoretical understanding of most solidification phenomena and the reason for examining the topic here. In this article, the PFS of a single phase is treated without taking convection into account. The sections in this article discuss the solute buildup at the solid-liquid interface forming transients and steady state, the morphological stability/instability and perturbation theory, and rapid solidification effects, including solute trapping and oscillatory instabilities. Finally, a microstructural selection map is presented that gives an overview of interface stability as a function of composition for a given alloy.
Plane Front Solidification Technology Silicon single-crystal growth is a technical application of PFS. The component that forms the substrate of a multitude of integrated
Fig. 1
Element distribution as a function of volume fraction solid, fs, in plane front Czochralski processing of silicon single crystals. Source: Ref 1
circuits with minute electronic elements must have reliable properties on a scale smaller than the smallest elements of the circuit. Silicon substrates must also have a high degree of chemical purity and a high degree of crystalline perfection. This can only be obtained by PFS in either the Czochralski (CZ) or the float zone (FZ) melting process. For PFS of silicon, very low growth rates, on the order of several millimeters per minute, are required. Figure 1 shows the distribution of some elements in silicon for CZ growth (Ref 1). The doping elements used, such as boron or
Table 1 Symbol
Cl Co Cs D Di Ds d fs G Gc J K k m P T* Tf Tl To Ts V Va Vc Vo y z z0 DCo DTo dC di dK dT e G f l o xc
Definitions and units of symbols Definition
Liquid composition Alloy composition Solid composition Diffusion coefficient in liquid Diffusion coefficient across interface Diffusion coefficient in solid Diameter Solid fraction Temperature gradient Concentration gradient Solute flux Curvature ðK1 1 þ K2 1 Þ Distribution coefficient (Cs/Cl) Liquidus slope Pe´clet number Interface temperature Melting temperature Liquidus temperature Temperature of equal free energy (k(V) = 1) Solidus temperature Interface velocity Velocity of absolute stability Velocity of constitutional undercooling Collision rate of atoms Coordinate parallel to interface Coordinate perpendicular to interface Coordinate of specimen Solidus-liquidus interval at Ts Solidus-liquidus interval (Tl Ts) Solute diffusion length (D/V) Interface diffusion length (Di/V) Capillary length (G/DTo) Thermal length (DTo/G) Amplitude of perturbation Gibbs-Thomson parameter Driving force for instability (G mGc) Wavelength Wave number (2p/l) Stability function
Units
% % % 2 m /s 2 m /s 2 m /s m ... K/m %/m % m/s 1 m ... K/% ... K K K K K m/s m/s m/s m/s m m m % K m m m m m Km K/m m 1 m ...
phosphorus, have distribution coefficients (k) close to unity, while many other elements have much smaller coefficients. (See Table 1 for the definitions and units of symbols used in this article.) Due to the different volumes of liquid and solid in the CZ and FZ processes, the composition profile of the single crystal produced is also different. In CZ, the originally large liquid volume produces a crystal where the concentration of the first half corresponds closely to kCo, leading in its second half to an important final transient (Fig. 1). In the FZ process, the liquid volume is small, and saturation of this liquid, at least for the elements with k values close to unity, is rapid. After a short distance, the composition profile reaches the mean concentration of the crystal, and the final transient is small. In the early years of silicon single-crystal growth, the FZ process was widely used; however, today (2008) large-diameter crystals are generally produced by the CZ process. The size of industrial single crystals has dramatically increased over the last 55 years (Fig. 2), and progress in crystal growth techniques has made substrates for integrated circuits progressively less expensive.
Plane Front Solidification Science An understanding of most solidification phenomena began essentially in the 1950s with the work on constitutional undercooling (CU).
Fig. 2
Plane front solidification processing. Evolution of the size of silicon single crystals over the years. Source: Ref 1
294 / Principles of Solidification Chalmers and his group initially produced bicrystals in order to study grain boundaries. Using the Bridgman method for PFS, they found various patterns at the decanted solidification front. These observations led to the most successful research on morphological instability: cell and dendrite growth (see the interesting article on the history of CU by Tiller in Ref 2). It was discovered by theory and experiment that a plane front is stable only if the growth rate is sufficiently small and the temperature gradient across the interface is positive and sufficiently large. The paper on CU by Tiller et al. in 1953 (Ref 3) was a major starting point for a half-century of solidification science (Ref 4). A decade later, Mullins and Sekerka (Ref 5, Ref 6) published their important papers on a rigorous linear stability analysis. They showed that, taking into account the interface energy, the wavelength of the morphological instabilities could be determined. Furthermore, a new limit of stability, absolute stability, was predicted for high growth rates. From perturbation theory, Langer and Mu¨ller-Krumbhaar (Ref 7) proposed marginal stability in 1978 as the operating point of dendrites (i.e., the selection of the tip radius for a given dendrite velocity), which substantially improved theoretical predictions and initiated much new work on dendrite theory.
at the composition kCo (Fig. 3(b)). The difference between the liquid and solid composition at the moving solid-liquid interface forms the source of the solute, J1, which dissipates into the liquid by diffusion, J2. Due to slow diffusion in the liquid and much slower diffusion in the solid, solute piles up at the solidification front. Neglecting the small flux into the solid, the fluxes at the interface may be written as: J1 ¼ V Cl ð1 kÞ J2 ¼ DGc
At low growth rates, interface stability requires a positive temperature gradient; that is, the temperature must increase from the solid into the melt. Such a situation is typical for directional solidification. Here, the heat flux is unidirectional to produce planar isotherms and interface shapes. For example, in the Bridgman process, a crucible that contains the alloy is moved through a heater and cooler, both being separated by a buffer that thermally isolates the central zone. For a theoretical analysis of the phenomena that are produced in such a process, the situation is often simplified, and the thermal transients at the beginning and end of solidification are not taken into account. It is assumed that the temperature gradient and the growth rate are instantaneously established and that they stay constant throughout the specimen of constant cross section.
(Eq 1)
J2 ¼ V ðCl Co Þ
1k k
(Eq 5)
and the liquid composition profile is: Cl ¼ Co þ Co expðVz=DÞ
(Eq 6)
(Eq 2)
where z is the distance into the melt from the solid- liquid interface, perpendicular to that interface. This concentration field is essential to understanding constitutional undercooling.
(Eq 3)
Final Transient
the flux into the melt becomes:
Initially, C*l is close to Co, and the sink term is smaller than the source term, that is, J2 < J1. Therefore, the concentration peak at the interface increases with distance and approaches the steady-state value Co/k (Fig. 3b). After a certain distance along the crucible, given by: z0 ¼ nD=V k
Transients and Steady State
Co ¼ Co
where C*l is the liquid composition at the solid-liquid interface. With the concentration gradient at the interface: Gc ¼ ðCo Cl ÞV =D
by the surface (A), is equal to the mass in the boundary layer (B) (Fig. 3(b)). In the steady state, the liquid composition far from the interface is equal to the solid composition, Co, and the interface grows at the solidus temperature of the alloy, Ts. In the steady state, the solute rejection at the interface is given by:
(Eq 4)
steady state is closely approached, that is, J2 ! J1. As a good approximation, n = 4. From Eq 4, it can be seen that for small k- values, the initial transient becomes very long, and steady state may never be reached. This, for example, is true for aluminum in silicon or salt in ice.
Steady State Mass balance considerations require that the missing mass in the solid, given by the difference between the mean alloy composition, Co, and the solid composition, C(z0 ) represented
When the boundary layer interacts with the end of the specimen, the dissipation of solute becomes limited due to the decreasing volume of liquid. Therefore, when k < 1, both the concentration of the liquid at the interface, C*l, and, correspondingly, of the solid, C*s = kC*l, increase. (In the case of k > 1 the final transient leads to a solid with a decreasing concentration.) The final transient becomes measurable when the interface reaches a multiple of the boundary layer thickness, D/V. It has been evaluated in terms of a series by Smith et al. (Ref 8). Unlike the initial transient, the final transient does not depend on k. Due to the very slow diffusion of substitutional elements in the solid, the final composition may reach very large values—at the limit, 100% of solute. The final concentration can be estimated by the Brody-Flemings equation modified by a function of the Fourier number (Ref 9, 10). For D ! 1 and Ds ! 0, the final transient is already present at the very beginning of
Initial Transient The phase diagram in Fig. 3(a) shows that the solubility of a solute in a solvent is different in the liquid and the solid phases. This reflects the thermodynamic requirement of equal chemical potentials. In most alloys, the melting point of the solvent is reduced by the solute and, according to the generally accepted definition, the distribution coefficient, k (= Cs/Cl), is smaller than unity. In this case, assuming that there is no nucleation barrier, the solid starts to grow
Fig. 3
Steady-state solute distribution. (a) Phase diagram and (b) concentration field in a directional solidification experiment
Plane Front Solidification / 295 solidification, and the concentration variation with solid fraction, fs, follows the GulliverScheil equation: Cs ¼ ð1 fs Þk1 kCo
(Eq 7)
The form of the solute distribution in this case resembles but, due to geometrical differences, is not exactly the same as the curves of Fig. 1.
Constitutional Undercooling Constitutional undercooling is a general phenomenon where heat diffusion is much more rapid than solute diffusion, thereby tending to produce steep concentration gradients (small concentration boundary layers) and low-temperature gradients (extensive temperature boundary layers). The concentration field ahead of the moving interface can be related through the liquidus slope, m = dTl/dC, to a corresponding liquidus temperature field. When this field is superimposed on the temperature field due to the heat flux (neglecting a very small undercooling from atom attachment kinetics), an undercooled liquid region is obtained under certain conditions. The concentration gradient for the steadystate concentration field (Eq 6) at the interface z = 0 is: (Eq 8)
This concentration gradient gives rise to a liquidus temperature gradient: dTl =dz ¼ ðdTl =dCÞðdC=dzÞ ¼ mGc
G ¼ mGc
(Eq 9)
(Eq 10)
or G=V ¼ To =D
Morphological Stability
Gc ¼ Co ðV =DÞ
Constitutional undercooling occurs when the heat flux gradient at the interface, G, is smaller than the liquidus temperature gradient. This case is shown in Fig. 4. The limiting condition of CU is obtained with the equality of the two gradients (Ref 3):
(Eq 11)
with DTo = mDCo. The left side of Eq 11 represents the process variables, and the right side represents the alloy properties. Stability corresponds to G/V > DTo/D. Equation 11 can also be understood as the equality of two lengths: the diffusion length, dC (= D/V), and the thermal length, dT (= DTo/G) (Ref 11).
correspond to this sequence of events. Eventually, the waves develop into cells and dendrites. Clearly, in order to form such an undulated interface structure, the system requires energy for creating the curvature that must come from interface undercooling. Not every structure is possible, and the strength of the CU defines the possible curvature (l-1). The details of the Mullins-Sekerka analysis can be found in the original papers (Ref 5, 6) or in Appendix 7 of Ref 10. Here, only the simplest possible interface stability analysis is developed, which, however, contains most of the essential physics (Chapter 3 of Ref 10). The starting point of a linear stability analysis is an infinitesimally small perturbation, that is, a wave of much smaller amplitude, e, than wavelength, l(e/l << 1). The interface temperature, T*, may be written as: T ¼ Tf þ mC1 K
Perturbation Analysis The simple criterion of CU depends only on transport phenomena. It does predict the possibility of morphological interface instability. However, since the interface energy is not considered, it does not say anything about the interface morphology, such as the wavelength, l, of a perturbation in the case of instability. The capillarity effect is taken into account in the perturbation theory, which was developed for the first time by Mullins and Sekerka (Ref 5, 6) for solidification. From the inclined interface of succinonitrile under the microscope (Fig. 5) (Ref 12), it can be seen that at the beginning of instability of a plane front, a wavy small-amplitude pattern develops across the solid-liquid interface. This happens first at the grain-boundary grooves, where solute piles up more strongly than at the plane front, then at the subgrain-boundary grooves and finally at the planar interface. The figure shows that the amplitudes
(Eq 12)
with the product of the Gibbs-Thomson parameter, G, and curvature, K, equal to the capillarity undercooling. With a sinusoidal deformation of the interface: z ¼ e sinðoyÞ
(Eq 13)
the temperature difference between the extremes of the perturbation (tip and depression) is given by: Tt Td ¼ mðCt Cd Þ ðKt Kd Þ
(Eq 14)
and: Kt ¼ Kd ¼ 4p2 e=l2
(Eq 15)
Due to the infinitesimal small amplitude of the perturbation and the much larger diffusion distances in metals, the temperature and concentration differences between the tips and the depressions can be obtained to a good approximation from the gradients existing at the unperturbed planar solid-liquid interface (small thermal and solutal Pe´clet number approximation, Pt < Pc < 1). The temperature and concentration differences are simply given by: Tt Td ¼ 2eG
(Eq 16)
Ct Cd ¼ 2eGc
(Eq 17)
This leads to a relationship for the perturbation wavelength of a marginally stable interface: li ¼ 2pð=fÞ0:5
Fig. 5 Fig. 4
Constitutional undercooling. (a) Interface concentration field with boundary layer dc. (b) Temperature fields at the solid-liquid interface
Image of a solid-liquid interface at the beginning of instability. For clarity, the interface of succinonitrile in a directional solidification experiment is inclined to the optical axis of the microscope. Grain boundaries and subgrain boundaries can be seen. The mean wavelength of the perturbations is 5 mm (200 min.). Source: Ref 12
(Eq 18)
where f = (mGc G) is the driving force for unstable growth. Equation 18 shows that the critical (marginally stable) perturbation wavelength, where neither the tip nor the depression will grow with respect to the mean interface location, depends on the ratio of the GibbsThomson parameter (resisting deformation)
296 / Principles of Solidification and the positive difference between the liquidus temperature gradient, mGc, and the temperature gradient, G (driving force of deformation). From the full Mullins-Sekerka analysis, a relationship is obtained between the amplitude growth and the wavelength, which is graphically shown in Fig. 6. The unstable modes of the perturbation lead to a range of wavelengths with li, the lower bond of possible instabilities. The wavelength range, li, is used in the Ivantsov-marginal stability (I-MS) dendrite theory (Ref 13) for calculating the tip radius. If it is assumed, for the sake of simplicity, that the thermal properties of liquid and solid are the same and that the temperature gradient across the interface is constant, that is, no latent heat is produced, then for the marginally stable state, the Mullins-Sekerka analysis gives:
o2 G þ mGc xc ¼ 0
(Eq 19)
The stability function, xc, has a value of unity for solute Pe´clet numbers up to P = 1 and drops to zero at higher P-values (Ref 14). It can be easily seen that, when xc equals unity, Eq 19 corresponds to Eq 18, and, in addition, when the interface energy is set to zero, the limit of CU is obtained (Eq 10).
Rapid Solidification Effects Absolute Stability When the interface velocity becomes large, the stabilizing capillarity effect dominates the
destabilizing solute diffusion effect. This phenomenon leads to a restabilization of the plane front when the growth velocity exceeds a critical value given by: Va ¼
DTo k
(Eq 20)
This phenomenon of absolute stability can be understood by comparing the variation of the controlling length scales with velocity: The diffusion distance decreases with V1, and the characteristic length of the perturbation decreases with V1/2. At high velocity, the diffusion distance becomes much smaller than the perturbation wavelength, and diffusion becomes essentially one-dimensional (localized). This phenomenon can also be qualitatively understood as the equality of two length scales: the diffusion length, dc (= D/V), and the capillarity length, dK (= G/DTo). According to this theory, morphological stability is therefore observed when the interface velocity is either below Vc (= DG/DTo) or above Va (= DDTo/kG).
Solute Trapping Up to now, it has been assumed that there is local equilibrium at the interface; that is, the compositions at the interface of solid and liquid correspond to the equilibrium phase diagram. This is a good assumption so long as the diffusion processes across the interface are much more rapid than the interface displacement. If, however, the interface velocity becomes large, the local equilibrium assumption is no longer applicable, and the distribution coefficient tends to unity. The relationship between k and V for a dilute solution can be written as (Ref 15):
Fig. 6
Amplitude growth (de/dt
1/e) as a function of the perturbation wavelength for three velocities (G = 10
4
kðV Þ ¼ K/m)
k þ ðdi V =Di Þ 1 þ ðdi V =Di Þ
(Eq 21)
where the parameter in parentheses represents the interface Pe´clet number (ratio of interface width, di, to diffusion distance across the interface, Di/V). A decreasing difference between the liquid and solid compositions (solute trapping) appears in the phase diagram where both liquidus and solidus lines converge to the To line (Ref 16–18), as shown in Fig. 7(a). When T1 = Ts = To, the distribution coefficient k = 1. In this limiting case, the Gibbs free energies of both phases are equal, but the chemical potentials are not. This can be expressed through the following equations for the interface temperature, T*, and the nonequilibrium liquidus slope, m(V) (Ref 18): T ¼ Tf þ mðV ÞCl þ
Fig. 7
Nonequilibrium solidification. (a) Equilibrium (cm) and nonequilibrium phase diagram mv. The interrupted line represents To, which corresponds to equal free energies, i.e., to k(V) = 1. (b) Interface temperature functions for plane front and cells/dendrites
mðV Þ ¼ m
m V 1 k Vo
1 kðV Þ þ kðV Þ lnðkðV Þ=kÞ 1k
(Eq 22)
(Eq 23)
Plane Front Solidification / 297 The last term of Eq 22 introduces the atom attachment kinetic undercooling, which depends on the collision rate of the atoms, Vo. In most alloys, absolute stability corresponds to the regime of local equilibrium loss, which reduces the limit, Va. A detailed analysis of this problem leads essentially to Eq 20, with DTo and k as velocity-dependent parameters (Ref 19, 20): Va ¼
DTo ðV Þ kðV Þ
(Eq 24)
where DTo(V) = Com(V)[k(V)1]/k(V). In Fig. 7(b), the interface temperatures of the various growth forms—plane front, cells, and dendrites—are shown as functions of V for a constant G (>0). The lower (partially interrupted) curve represents the plane front temperature, and the upper curve the cell/dendrite temperature. At two velocities, Vc and Va, the cell/dendrite tips have the same growth temperature as the plane front. The increase in the plane front temperature at growth velocities near to Va is due to solute trapping (k(V)), while the decrease of the plane front temperature at very high growth velocities is due to the increasing effect of atom attachment kinetics (Ref 21).
Oscillatory Instability Once the growth velocity exceeds absolute stability, the nonequilibrium conditions make the plane front unstable with time; that is, the planar interface develops oscillations. Oscillatory instabilities have been predicted by Corriel and Sekerka (Ref 22) and Davis and co-workers (Ref 19, 23, 24). Karma and Sarkissian (Ref 25) analyzed this phenomenon numerically and showed that the oscillatory instability induces large-amplitude relaxation oscillations at a critical velocity. This phenomenon can be qualitatively understood as follows: When, beyond Va, the isotherm displacement rate reaches the plane front branch with its positive T(V) slope, (Fig. 7(b)) the interface accelerates as k approaches unity, and less solute is rejected. When the accelerating plane front reaches the maximum in the T(V) curve, the solidification velocity must decrease in order to conform with the thermal field until it reaches the interface temperature minimum, where it may develop a cellular morphology. Either a banded, plane front structure, or bands of alternating plane front and cells is obtained (Ref 26, 27). Figure 8(a) shows the regimes of existence of the various instability phenomena according to the model of Huntley and Davis in a Co-V diagram (Ref 24). Below the concentration minimum at 5 103 at.% Fe, a plane front is stable at all growth velocities. At this point, CU and absolute stability have the same effect. The hatched region at the branch of absolute stability corresponds to the experimental observation of banded structures (Ref 26). The
schematic representation of Fig. 7(b) is shown in quantitative form for an Al-1at.%Fe alloy in Fig. 8(b) (Ref 29). This diagram has been calculated using the plane front and IMS dendrite model (Ref 13). For the parameters used in the figure, plane front growth is unstable with respect to cellular and dendritic growth between Vc = 5 105 and Va = 7 101 m/s. Beyond VM = 4 m/s, atom attachment kinetics dominate, and the plane front is stable in every respect. From this, it is evident that the results of both theories, stability theory (Ref 24) (Fig. 8a) and interface response (Ref 29, 30) (Fig. 8b), and of experiments are consistent and complementary.
microstructure formation phenomena and is therefore fundamental to solidification theory. Interface stability theory allows the limits of microstructure formation to be predicted. At low rates, it is the ratio G/V that controls the limit of CU, while at high rates, the limit of absolute stability depends on capillarity. Solute trapping reduces the limit of absolute stability. The wavelength of the marginally stable state, li, is useful for modeling dendritic growth with the I-MS theory. In the high V transition region from cellular-dendritic to absolutely stable plane front growth, banded microstructures appear that are indicative of oscillatory interface instability.
Solidification Microstructure Selection Maps
REFERENCES
Figure 8(a) represents a microstructural selection map for binary aluminum-iron alloys. It shows two V-shaped regimes of instability. The low-V branch of the continuous curve corresponds to constitutional undercooling, while the high-V branch of the continuous curve represents absolute morphological stability. At low-V (above the solid curve), morphological instability is predicted, and at high velocities (above the dashed curve), oscillatory instabilities develop. The V-range of instability increases with solute concentration. To the left and right of the curves, PFS is observed, and below the concentration minimum, PFS occurs for all interface velocities. Similar diagrams for various alloys, silvercopper, silicon-tin, aluminum-copper, and aluminum-tin, can be found in Ref 19 and 24. Figure 9 is a microstructural selection map for iron-nickel alloys (Ref 28, 30). If the coordinates of the axes are interchanged, Fig. 9 and Fig. 8 are similar diagrams for different alloys. In alloys such as iron-nickel, where two solid phases compete in the range of the peritectic composition, the microstructures that may form are more complicated. At concentrations below 3.8 at.% Ni, the d-phase (ferrite) is stable, while at high nickel concentrations, the g-phase (austenite) replaces the d-phase. Because kNi g > kNi d , the V-range of unstable g plane front growth is much narrower than that of d. In the low-V regime (high G/V values) at 4 at. %, two-phase growth is observed, either in the form of d/g bands or of peritectic coupled growth (Ref 31).
1. W. Zulehner, Mater. Sci. Eng. B, Vol 73, 2000, p 7–15 2. W.A. Tiller, Mater. Sci. Eng., Vol 65, 1984, p 3 3. W.A. Tiller, K.A. Jackson, J.W. Rutter, and B. Chalmers, Acta Metall., Vol 1, 1953, p 428
Fig. 8
Summary Plane front growth can be observed at both low and high velocities. It is a phenomenon that is rarely observed and seldom used in practice, except in single-crystal growth and in directional solidification experiments. However, the understanding of PFS is basic to most
Stability diagrams for an Al-1at.%Fe alloy as a function of growth rate, V, (G = 5 106 K/m). Source: Ref 29. (a) Composition dependence of morphological and oscillatory instabilities showing complete morphological stability below 5 103 at.%. Source: Ref 23. (b) Interface temperature (response) for plane front, Tp(V ), and dendritic front, Td(V ), compare with Fig. 7(b). Source: Ref 29. The characteristic interface velocities are Vc, limit of constitutional undercooling; Vm, velocity of minimum interface temperature; Va, limit of absolute stability; and VM, velocity of maximum interface temperature.
298 / Principles of Solidification
Fig. 9
Microstructure selection map (V- Co diagram) of iron-nickel alloys for G = 105 K/m showing regions of plane front solidification of cellular/dendritic structures of delta and of gamma, respectively. Source: Ref 28
4. W. Kurz, Mater. Sci. Forum, Vol 508, 2006, p 313–324 5. W.W. Mullins and R.F. Sekerka, JAP, Vol 34, 1963, p 323 6. W.W. Mullins and R.F. Sekerka, JAP, Vol 35, 1964, p 444 7. J.S. Langer and H. Mu¨ller-Krumbhaar, Acta Mater., Vol 26, 1978, p 1681 8. V.G. Smith, W.A. Tiller, and J.W. Rutter, Can. J. Phys., Vol 33, 1955, p 723
9. H.D. Brody and M.C. Flemings, Trans. Metall. Soc. AIME, Vol 236, 1966, p 615 10. W. Kurz and D.J. Fisher, Fundamentals of Solidification, 4th ed., Trans Tech, 1998 11. R. Trivedi and W. Kurz, Acta Metall. Mater., Vol 42, 1994, p 15–23 12. D.J. Fisher and W. Kurz, unpublished work, EPFL, 1977 13. W. Kurz, B. Giovanola, and R. Trivedi, Acta Metall., Vol 34, 1986, p 823
14. R. Trivedi and W. Kurz, Acta Metall., Vol 34, 1986, p 1663–1670 15. M.F. Aziz, JAP, Vol 53, 1982, p 1158 16. J.C. Baker and J.W. Cahn, in Solidification, American Society for Metals, 1971, p 23 17. W.J. Boettinger, S.R. Coriell and R.F. Sekerka, Mater. Sci. Eng., Vol 65, 1984, p 27 18. W.J. Boettinger and S.R. Coriell, in Science and Technology of the Undercooled Melt, P.R. Sahm, H. Jones, and C.M. Adam, Ed., Martinus Nijhoff, 1986, p 81 19. D.A. Huntley and S.H. Davis, Acta Metall. Mater., Vol 41, 1993, p 2025–2043 20. A. Ludwig and W. Kurz, Acta Mater., Vol 44, 1996, p 3643–3654 21. R. Trivedi and W. Kurz, Int. Mater. Rev., Vol 39, 1994, p 49–74 22. S.R. Corriel and R.F. Sekerka, J. Cryst. Growth, Vol 61, 1983, p 499 23. D.A. Huntley and S.H. Davis, Phys. Rev. B., Vol 53, 1996, p 3132 24. S.H. Davis, Chap. 6, Theory of Solidification, Cambridge Univ. Press, 2001 25. A. Karma and A. Sarkissian, Phys. Rev. E, Vol 47, 1993, p 513 26. M. Gremaud, M. Carrard, and W. Kurz, Acta Metall. Mater., Vol 39, 1991, p 1431 27. M. Carrard, M. Gremaud, M. Zimmermann, and W. Kurz, Acta Metall. Mater., Vol 40, 1992, p 983–996 28. M. Vandyoussefi, H.W. Kerr, and W. Kurz, in Solidification Processing 1997, J. Beech and H. Jones, Ed., University of Sheffield, 1997, p 564 29. W. Kurz and S. Dobler, Interfaces for the 21st century, M.K. Smith, M.J. Miksis, G.B. Mc Fadden, G.P. Neitzel, D.R. Canright, Ed, Imperial College Press, 2002 203–212 30. W. Kurz, Adv. Eng. Mater., Vol 3, 2001, p 443–452 31. S. Dobler, T.S. Lo, M. Plapp, A. Karma, and W. Kurz, Acta Mater., Vol 52, 2004, p 2795–2808
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 299-306 DOI: 10.1361/asmhba0005210
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Non-Plane Front Solidification Rohit Trivedi and Erin Sunseri, Ames Laboratory, USDOE and Iowa State University
DURING THE SOLIDIFICATION of alloys, a nonplanar interface is most commonly present under conditions of casting and welding, playing a crucial role in determining the properties of the solidified material. The most common nonplanar interfaces form cellular or dendritic microstructures, which lead to microsegregation and possible second-phase formation that control the properties and reliability of the solidified product. Since microstructures form a key link between properties and processing conditions, it is important to develop basic ideas on how processing conditions influence microstructures so that one can predict, modify, and control the microstructure of the material by designing appropriate composition and processing conditions.
Columnar and Equiaxed Microstructures The thermal field in a casting is very important in determining the microstructure of the cast alloy. Two distinctly different conditions may exist in a mold that can give rise to equiaxed or columnar crystals. First, consider the case of equiaxed crystals that form in an undercooled melt. As the solid nucleates and grows, latent heat and solute are rejected into the liquid so that negative temperature and concentration gradients are present in the liquid at the solid/liquid interface. As discussed by Kurz in the previous article, “Plane Front Solidification,” in this Volume, the stability of the interface at low velocity is controlled by constitutional undercooling and interface energy. Since the thermal and solute gradients in the liquid are negative, constitutional supercooling will always be present to cause the interface to become unstable. The interface energy will initially stabilize the spherical nucleus, but the effect of interface energy becomes small as the sphere radius becomes approximately an order of magnitude larger than the critical radius of nucleation. Consequently, instabilities will develop during the early stage of the growth. In most cubic metals, instabilities occur in six directions. These instabilities grow and form side branches, so that six dendrite branches form from a nucleus
and give an equiaxed shape to the solid, as shown in Fig. 1. Equiaxed structures generally form at the center of a casting, where a substantial part of the liquid becomes constitutionally undercooled. The nucleation density controls the final grain size, and the nucleation density is often increased by the use of inoculants. Columnar crystals form when the temperature gradients in the liquid and the solid are positive, such that the latent heat generated at the interface is dissipated through the solid. Such a temperature field gives rise to directional solidification and results in the columnar zone in a casting. The conditions for the stability of a planar front, discussed in the previous article, “Plane Front Solidification,” by Kurz, show that a planar front will be stable when V < Vc or V > Va, where, for clarity, the thermal gradient and composition are taken to be constant. These limiting velocities are given by Vc = G D/DT0 and Va = DDT0/kG, where D, k, and G are the diffusion coefficient in the liquid, the solute distribution coefficient, and the Gibbs-Thomson parameter, respectively, and DT0 = mC0(k 1)/k, in which m is the slope of the liquidus at the solidus temperature. Thus, within the velocity range Vc < V < Va, planar interfaces become unstable and reorganize into a periodic structure consisting of shallow or deep cells or dendrites (Ref 1), as shown in Fig. 2. If the velocity is increased further, dendrites transform to cells and then to a planar interface at the velocity Va.
Fig. 1
Equiaxed dendrites undercooled liquid
that
form
from
an
Microstructure and Processing Conditions The formation of planar, cellular, or dendritic microstructure for a single-phase growth is governed by the processing conditions of G and V and the composition of the alloy. The effect of processing conditions (G,V) on microstructure is illustrated in Fig. 3, in which regimes of planar, cellular, and dendritic microstructures are
Fig. 2
Columnar structures that form under directional solidification conditions. (a) Shallow cellular. (b) Deep cellular. (c) Dendritic interfaces
300 / Principles of Solidification
Fig. 3
Microstructure-processing diagram showing the regimes of planar, cellular, and dendritic microstructures as a function of the Gibbs-Thomson parameter (G) and the velocity (V ) that are present in different processing technologies
Fig. 4 shown for the directional solidification of Al4.0wt%Cu. Typical ranges of G and V, as well as the cooling rate ðT_ ¼ GV Þ, present in different solidification processing techniques are also shown in the figure. A planar front growth occurs at very low velocity, as in single-crystal casting, or at very high velocity, as in laser processing and atomization of very fine droplets. A cellular microstructure forms at low velocities when a planar front becomes unstable. It also forms at very high velocities before the planar interface becomes stable. Under normal casting conditions, dendritic structures are commonly observed (gray area) in most alloys used in casting. The scale of microstructure becomes finer as one goes from single-crystal growth to normal casting to continuous casting to strip casting to laser welding. The processing conditions (V and G) are critical not only in the selection of nonplanar microstructures, but they also govern microstructural scales that are related to properties.
Characteristic Lengths of Microstructures To relate properties of a material to the microstructure, it is first necessary to define microstructure in terms of some characteristic parameters that can be quantitatively measured and can be related to properties. For equiaxed dendritic structure, the important characteristic lengths of the microstructure are the dendrite tip radius, R, the secondary arm spacing, l2, and the average distance between the grains. For directional solidification, important lengths are the tip radius, R, the primary spacing, l, the secondary spacing, l2, and the length of the dendrite in the two-phase region, or the extent of the mushy zone, ‘. The formation of a nonplanar interface causes solute segregation patterns, since the solid grows with a smaller composition than the alloy composition. The rejected solute
(a) Array of cells showing variation in the shape and local spacing. (b) Rescaling of two cells with different local spacing gives a unique shape.
accumulates in the intercellular/interdendritic region for alloy systems with phase diagrams in which the solute segregation coefficient is less than 1 (decreasing liquidus temperature with concentration). From a practical viewpoint, the interarm spacings characterize the solute segregation pattern, and often, a secondphase distribution in the interdendritic region results, which also has a strong influence on the properties of the material. Theoretical models of microstructure formation in a casting are thus directed toward quantitatively understanding how the processing conditions, alloy composition, and thermophysical properties control the magnitudes of these spacings and of the tip temperature or tip composition. First, the basic model of cellular and dendrite growth is described, and then the conditions for the columnar-to-equiaxed transition are discussed.
Cellular Microstructures When growth conditions exceed the critical condition for planar interface stability, the thermal fluctuations cause perturbation in the shape of the interface. These perturbations grow with time and reorganize into a periodic array of cells. For velocities just above the critical velocity, shallow cells form in which the length of the cell is small and is of the same order of magnitude as the cell spacing (Fig. 2a). The tip region of the cell is broader, and the cell has a larger tip radius. As the velocity is increased, the shallow cells transform to deep cells in which the length of the cell is much larger than the cell spacing, as shown in Fig. 2 (b). As the velocity is increased, the tip region of the deep cell assumes a nearly parabolic shape, which is similar to the dendrite tip shape
so that the term cellular dendrite is often used to characterize this structure (Ref 2, 3).
Scaling Behavior of Cellular Microstructures The steady-state shape of the cell is rather complicated, since it is influenced by the solute as well as thermal field interactions between the neighboring cells. A detailed numerical model has been developed by Hunt and Lu (Ref 2) to characterize the shape, cell tip undercooling, and cell spacing as functions of G, V, and composition. A simple analytical expression for the cell tip undercooling or tip temperature (Tt), which is a good approximation in the limit of small k, has been developed by Bower, Brody, and Flemings (Ref 4) as: Tt ¼ TL ðGD=V Þ
(Eq 1)
where TL is the liquidus temperature of the alloy. The shape and the local cellular spacing are found not to be unique, but a narrow range of stable shapes and spacing has been observed for different cells in an array, as seen in Fig. 4(a). For spacing below this range, cell elimination occurs, while for spacing larger than this range, tip-splitting occurs to create a new cell. Although different cells in an array have different shapes and different local spacing, it is shown that the shapes are self-similar (Ref 5), and they scale with the local spacing; that is, all cell shapes in an array superpose when they are rescaled to the local spacing, as shown in Fig. 4(b) for two cells having minimum and maximum spacing value in the array. This scaling behavior implies that l/R is constant for all cells in an array. Although the shape and spacing vary for different cells, the position of the cell tips, or the tip temperature, is nearly constant.
Non-Plane Front Solidification / 301 Microsegregation The formation of cellular structures gives rise to solute segregation in the solid. The tip of the cell is at a higher temperature or lower composition for systems with k < 1, so that the solute is rejected in the intercellular region that will give rise to higher composition in the solid at the base of the cell. A periodic concentration profile will be present in the solid with a periodicity that is equal to the primary spacing. Since the liquid region between the primary cells is very small compared to the diffusion distance in the liquid, the concentration is nearly uniform in the lateral direction, that is, normal to the growth direction, so that under negligible diffusion in the solid, the solute profile in the solid is given by the Scheil-Gulliver equation (Ref 4, 6): Cs ¼ kC0 ½1 fs k1
(Eq 2)
where fs is the fraction of solid. When the diffusion in the solid is not negligible, the aforementioned equation is modified as: 0
k1 12a0 k
Cs ¼ kC0 ½1 fs ð1 2a kÞ
(Eq 3)
0
in which the parameter a is given empirically in terms of the dimensionless parameter a = Dstf/L2 as: a0 ¼ a½1 expð1=aÞ 0:5 expð1/2aÞ
(Eq 4)
Ds is the diffusion coefficient in the solid, and L is the distance between the cells. The modified equation ensures that the Scheil-Gulliver equation is obtained when Ds = 0, and lever rule is obtained when Ds = 1. For single-phase solidification in which the diffusion in the solid is important, microsegregation is characterized by the ratio of the maximum solid composition (at the cell base) to the minimum solid composition (at the cell tip), and this segregation ratio (SR) can be quantitatively related to properties. If ‘ is the length of the cell, GM the average temperature gradient in the two-phase region, and DT 0 is the cell tip undercooling with respect to the melting point isotherm of the solvent, then: SR ¼ 1 þ
‘ GM T 0
(Eq 5)
As the cellular structure becomes cellular dendritic or dendritic, ‘ increases sharply and DT 0 decreases under normal solidification conditions. Thus, the segregation ratio will increase significantly upon the formation of cellular dendrites or dendrites. The cellular spacing, l, is important since it characterizes the periodicity of the segregation pattern in the solid. Thus, if homogeneous composition is desired, the time required for the homogenization is related to the spacing as:
t ¼ 0:47 l2 =Ds
(Eq 6)
where t is the time required to homogenize to 1% of the original composition difference, and Ds is the diffusion coefficient in the solid. One of the important applications of the microsegregation equation is to predict the possible formation of different phases in the intercellular region. When the phase diagram shows a limited solubility in the solid, a new phase can form in the intercellular region when the local composition of the solid, predicted by the modified Scheil-Gulliver equation, reaches its solubility limit. In this case, the remaining liquid may form a new phase, such as a eutectic or peritectic phase, or a compound, depending on the phase diagram. As the composition in the liquid increases further, it is possible to form additional phases depending on the solubility limit for the new phase given by the phase diagram. The formation of a peritectic phase in the intercellular region of the primary phase in the tin-cadmium system is shown in Fig. 5. When the second phase forms, the properties of the material are related to the fraction of the second phases.
High-Velocity Cells The transition from dendrite to cell also occurs at high velocity, and these cells are generally referred to as high-velocity cells. Since the cellular structure is very fine at high velocities, it is sometimes referred to as microcellular structure. In most alloy systems, this transition occurs at very high velocity, which is near the absolute velocity for planar front solidification. Microstructural observations in very fine, atomized aluminum-silicon powders have characterized the transition from dendrites to cells at high velocity (Ref 7). The dendrite-to-cell transition also depends on the freezing range of the alloy, so that it can occur at relatively lower velocities in some systems. For example, experimental studies by Lee et al. (Ref 8) on Fe-18.0%Cr-5%Al have shown that the dendrite-to-cell transition occurs at approximately 0.050 mm/s (0.002 in./s), whereas the absolute velocity for planar front growth is approximately 50.0 mm/s (2 in./s). There are several similarities and differences between the low- and high-velocity cells: (1) The interface temperature in the low-velocity cellular regime increases with an increase in velocity, whereas it decreases with an increase in velocity in the high-velocity regime. (2) The primary spacing is found to decrease with velocity in both the regimes, and the spacing variation is found to be continuous between the two regimes. (3) The cells at low velocity grow in the direction of heat flow, while the growth of the high-velocity cells occurs along preferred crystallographic directions, which are the same as those for the dendrites.
Dendrite Growth Dendritic structures are classic examples of branching patterns in which the branched
Fig. 5
Cellular primary phase with the peritectic phase forming in the intercellular region
structure evolves with primary, secondary, tertiary, and eventually higher-order branches. Such a branched structure forms because it can effectively allow rapid dissipation of the latent heat of fusion as well as any solute redistribution required for growth due to the difference in compositions in the solid and the liquid. Dendritic structures most commonly form during the solidification of metallic systems. Three important aspects of dendritic structures that are examined here are dendrite growth direction, dendrite tip characteristics such as tip composition and tip temperature, and microsegregation that is controlled by the secondary arm spacing.
Dendrite Growth Direction Dendrites have been observed to grow in specific crystallographic direction. This crystallographic effect not only can form texture during the columnar growth, but it is also critical for the steady-state growth of dendrites and in the transition from chill to columnar crystals in a casting. In cubic crystals such as facecentered cubic and body-centered cubic structures, dendrites have been observed to grow in <100> directions. Chalmers (Ref 9) discussed that the dendrite growth direction will be along the axis of a pyramid whose sides are the most closely packed planes with which a pyramid
302 / Principles of Solidification can be formed. However, other growth directions have also been observed, particularly in aluminum-zine and aluminum-magnesium alloys, where the growth direction changes from <100> to <110> as the solute content is increased (Ref 10, 11). A more rigorous criterion for the dendrite growth direction has now been developed (Ref 12), showing that the growth direction at low velocities is controlled by the anisotropy in interfacial energy, while at very high velocities it is selected by the anisotropy in interface kinetics. For clarity, first consider a two-dimensional case in which the interface energy, g, in {100} planes, for materials with an underlying cubic symmetry, can be expressed as a function of the orientation as: gðyÞ ¼ g0 ð1 þ e4 cos 4yÞ
(Eq 7)
The anisotropy is characterized by only one parameter, e4, and g0 is the mean value of g. The anisotropy parameter, e4, can be measured experimentally from the equilibrium shape of a droplet in the solid; for the Al-4.0wt%Cu alloy, e4 = 0.0097 (Ref 13). The capillary undercooling for anisotropic interface energy is proportionalto the stiffness, S, which is defined as S ¼ g þ @ 2 g=@n21 . The growth of perturbations will be preferred at the orientations where the stiffness is at a minimum. In this case, the polar plot of 1/S with orientation will show protrusions at the orientations that have the minimum stiffness, as illustrated in Fig. 6 for a cubic system with preferred <001> growth directions. The aforementioned model assumes fourfold symmetry and (100) dendrite growth in cubic crystals. In order to predict orientations with other symmetries, a three-dimensional model must be considered. In three dimensions, two anisotropy parameters, e1 and e2, are required. From phase-field calculations, Haxhimali et al. (Ref 12) have shown that the term with a positive value of e1 favors a (100) orientation, whereas the n term with a negative value of e2 favors a (110) orientation. Combinations of these values allow for various crystallographic growth directions. These predictions have been used to explain different orientations that are observed in the aluminum-zinc and aluminummagnesium alloys (Ref 10, 11). The effect of anisotropy in interface energy on the dendrite tip radius selection and on the dendrite tip undercooling is presented in the next section.
complex casting (Ref 14). (2) Relationships for the dendrite tip composition and the dendrite tip temperature with composition and processing conditions are predicted. The dendrite tip composition is important in determining microsegregation patterns, and the dendrite tip temperature is critical in predicting the selection of dendritic structures over other (such as a eutectic) morphologies. (3) The concentration field ahead of the dendrite is required to evaluate the region of constitutional supercooling that is important in the determination of the columnar-to-equiaxed dendrites. Since the dendritic pattern evolves to effectively dissipate heat or solute, dendrite growth models have been based on the solutions of the thermal and mass transport equations. For most metals, the Schmidt number (the ratio of thermal to solute diffusivity) is large, so that the formation of columnar and equiaxed dendrite in alloys is largely dictated by the solute transport process and the interfacial free energy. There are two important aspects of the dendrite growth model based on the transport processes that must be addressed. First, the steady-state dendrite growth problem is a freeboundary problem in which the shape of the interface is not known a priori, but it must be determined in a self-consistent manner such that this shape is preserved during growth. Second, if only thermal and mass transport processes were important for the formation of dendrites, the dendrite tip will keep splitting to form finer dendrites or nondendritic structures to obtain more efficient heat and solute dissipation into the liquid. Thus, the thermal and solute diffusion processes are not sufficient to describe the steady-state dendrite growth, and a stabilizing effect such as that of interfacial free energy must be included in the dendrite growth model. A steady-state growth model was first developed by Ivantsov (Ref 15) under the constraint of constant interface temperature and
composition. He showed that a parabolic dendrite satisfies the steady-state shape-preserving constraint in that the entire interface will not change shape during growth. However, the solution gave a relationship between the growth rate, tip radius, and tip composition (or the tip undercooling). Thus, for a known V, the model predicts an infinite number of solutions that relate tip undercooling to the tip radius. Experimental studies, however, show that a unique tip radius is selected under given growth conditions so that another criterion is required for the selection of a specific value of the dendrite tip radius. Furthermore, if only diffusion is considered, the dendrite tip will be unstable, since it can split and dissipate solute more effectively. The stabilizing effect of interface energy must be considered to obtain the stable tip radius value. However, the presence of isotropic interface energy effect will cause the composition to vary along the interface, and it is shown that no steady-state solution of dendrite is possible (Ref 16). However, a steady-state solution is recovered if anisotropy in interface energy is introduced in the model. This stability condition introduces a dimensionless parameter, s*, that relates the dendrite tip radius to velocity. The parameter s* is referred to as the dendrite tip selection parameter and is a function of the anisotropy parameter, e4. A functional relationship between s* and e4 has been obtained numerically by Karma and Rappel (Ref 16). The Ivantsov solution, along with the stability condition, predicts the dendrite tip temperature, dendrite tip composition, and dendrite tip radius as a function of velocity and alloy composition (Ref 17). Although the interface energy and its anisotropy are critical in the stable, steady-state growth of dendrites, the solutions of the transport equations are not influenced by these effects. Thus, the interface energy effect on tip undercooling is negligible at low velocities, and the dendrite tip undercooling is close to the solutal undercooling at the tip. In addition, the dendrite tip radius and dendrite tip temperature do not depend on the thermal gradient. The variation in dendrite tip composition and dendrite tip temperature with velocity has been calculated for Al-3.0wt%Cu by using the Ivantsov model with the dendrite tip stability condition; the results are shown in Fig. 7. For casting conditions where the velocity (or the Peclet number, VR/2D) is small, the
Dendrite Growth Model Dendrite growth models have been developed to correlate the effect of processing variables, such as velocity and thermal gradient, on dendrite microstructures. The aim of the model is to predict the following important aspects of dendrite growth. (1) The relationship between the velocity and dendrite tip undercooling is critical in correlating microstructure with macroscopic heat flow profiles in a
Fig. 6
Schematic illustration of the polar plot of the inverse of stiffness showing bumps at easy growth directions
Fig. 7
Temperature and composition at the dendrite tip as a function of velocity in Al-3.0wt%Cu
Non-Plane Front Solidification / 303 dendrite growth model is often simplified by approximating the Ivantsov solution to give a much simpler relationship for the dendrite tip undercooling as (Ref 17):
Tt ffi K V
by the final secondary arm spacing near the base of the dendrite, which is given by (Ref 20, 21): 1=3
n
l2 ¼ const: tf
where K and n are constants whose values depend on the range of Peclet number that is relevant to the problem. The value of n is of the order of 0.3 to 0.5 for small Peclet number conditions. Note that this form of relationship is valid for columnar dendrites and also for equiaxed crystal if the ratio of thermal to solute diffusivity is large.
Microsegregation and Secondary Arm Spacing The growth of dendritic structures is associated with solute segregation since the solid rejects solute during growth (for k < 1), which accumulates in the surrounding liquid. In columnar growth, the secondary arm spacing controls the microsegregation profile in the solidified alloy. In equiaxed crystals, the solute accumulation in the liquid will also result in solute segregation at the boundaries of neighboring grains, so that the grain size is important. The primary dendrite arm spacing, l, becomes finer with increase in velocity or thermal gradient in the liquid, and a general relationship is given by (Ref 18, Ref 19): h pffiffiffi i1=2 l ¼ 4 2D=s ½V GR1=2
(Eq 10)
(Eq 8)
(Eq 9)
The variation in primary spacing with velocity for cellular and dendritic structure is shown in Fig. 8. There is an increase in primary spacing as the cells transform to dendrites. Recently, Lee et al. (Ref 8) have shown that a decrease in primary spacing is observed at the velocity where dendrites transform to high-velocity cells. For dendritic microstructures, secondary arm spacing is important since it governs the solute segregation pattern. The secondary arms originate close to the dendrite tip, and their initial spacing scales with the dendrite tip radius. However, the secondary spacing coarsens with time in the mushy zone. Initial coarsening occurs by the competition in the growth process among secondary arms. However, once the diffusion fields of the tips of the secondary branches interact with those of the neighboring dendrite, the growth of the secondary arms is reduced, and a coarsening process to reduce interfacial energy begins. The final secondary arm spacing near the dendrite base is significantly larger than that near the dendrite tip. This final secondary arm spacing controls the microsegregation profile in the solidified alloy. This microsegregation pattern is analogous to that discussed for the cellular structure, except that the periodicity of segregation is controlled
where tf is the local solidification time. For directional solidification, the local solidification time is inversely proportional to the cooling rate, which is a product of G and V, so that secondary arm spacing is generally related to the cooling rate, and this has been verified experimentally by Bower et al. (Ref 4) in the aluminum-copper system. The secondary spacing is the smallest length scale over which solute segregation occurs in dendritic structures, and it is correlated to the mechanical properties of single-phase alloys. In reality, these alloys are often not single phase but frequently contain small amounts of second phases, such as eutectic, precipitates, pores, and so on, that form due to microsegregation and control the properties.
Columnar-to-Equiaxed Transition Most practical metallurgical alloys solidify with a dendritic structure in a mold, and both columnar and equiaxed dendrites may be present in a casting (Fig. 9). As the liquid metal is poured into the mold, solid nuclei appear at the mold wall that give rise to a chilled zone. Some of these crystals are then favorably oriented for growth under directional heat flow conditions, and the selection of the grains occurs due to the anisotropy property of the interface energy. Initially, the chill crystals are oriented randomly. However, those crystals that are oriented with the preferred crystallographic direction in the heat flow direction will be selected. This is because the stiffness of the advancing front in the preferred growth direction is smaller, and thus, the capillary undercooling is smaller, so that they will grow ahead of the neighboring grains and crowd them out. These favorably oriented grains produce a columnar zone in a casting and exhibit [100] texture in most cubic metals. For alloy castings, an equiaxed zone can form ahead of the columnar crystal. This transition from the columnar to equiaxed zone is now examined.
Fig. 8
Effect of velocity on cellular, cellular dendrite, and primary dendrite spacings in the pivalic acid-ethanol system. Source: Ref 6
The basic concepts for the columnar-toequiaxed transition (CET) were first presented by Hunt (Ref 22). The transition from columnar to equiaxed grains occurs when enough large equiaxed grains form and grow ahead of the columnar front, hindering its movement. The size of the transition zone depends on the amount of competition between the columnar and equiaxed grains. To form an equiaxed zone, there are two requirements: the presence of nuclei ahead of the columnar front, and sufficiently fast growth of these nuclei to effectively block the columnar front. If either of these conditions are not met, then columnar grains will still predominate. One of the prerequisites for the formation and growth of nuclei is the presence of a constitutionally supercooled region ahead of the columnar front. Since the stable dendrite tip growth requires a balance between the constitutional supercooling and curvature effects, a constitutionally undercooled zone is always present ahead of the dendrite tip. Equiaxed growth can occur if the required nucleation undercooling is smaller than the undercooling present ahead of the columnar dendrites, as shown in Fig. 10. The length of the supercooled region, LT, ahead of the columnar front, shown in Fig. 10, is given by LT = DTt/G. This characteristic length provides the distance over which equiaxed nuclei can grow for the limiting case of zero nucleation undercooling. The second important parameter is the distance between the nuclei, LN, which forms the equiaxed zone. Once the nucleation undercooling is reached, nuclei begin to form, and the nucleation rate increases rapidly as the undercooling increases slightly. Thus, it is reasonable to assume that all effective nucleation sites are activated close to the nucleation undercooling. If N0 is the number of nuclei per unit volume, then 1=3 LN ¼ N0 . Once the nucleation undercooling is reached, the nuclei will form and grow as they approach
Fig. 9
Grain structures of an ingot casting showing the regions of chill, columnar, and equiaxed grains
304 / Principles of Solidification
Fig. 11
Fig. 10
Formation of equiaxed crystals ahead of the columnar front, and the length of the constitutionally supercooled liquid ahead of the columnar front
the columnar front. The increase in the radius of the grain with time (or undercooling) is obtained from Eq 8 for the undercooling changing from DTN to DTt. From the volume of the equiaxed grains at the columnar front and the number of nuclei per unit volume, the fraction of the equiaxed grains present at the columnar front can be determined. Hunt (Ref 22) proposed that equiaxed crystals will block the columnar front when the fraction of the equiaxed grains at the columnar front is 0.49. Since the nucleation sites are arranged randomly, a correction for the impingement of the equiaxed grains can be made by using the extended volume fraction concept, which requires that the extended volume fraction of the equiaxed grains be larger than 0.66 for the transition from columnar to equiaxed grains. Numerical calculations of CET can be obtained by calculating the columnar dendrite tip undercooling by using a more rigorous dendrite growth model (Ref 16). First, a simple picture of the condition for the CET is obtained, and then the results from the rigorous model are shown. The CET will occur if the time required for the nuclei to obtain required volume fraction to block the columnar front is less than the time available to the nuclei in the constitutionally supercooled zone. These times can be correlated with the characteristic distances. For a limiting case of very small nucleation undercooling, that is, DTN << DTi, CET occurs when the distance between the nuclei is less than the length of the constitutionally supercooled zone, that is, LN < LT. Substituting the
Numerical calculations of the columnar-to-equiaxed transition in Al-3.0wt%Cu alloy
values of these lengths, the CET condition 1=3 becomes G < N0 Tt . Thus, to promote CET, one needs a low thermal gradient, large number of nuclei, and a large dendrite tip undercooling that is obtained at higher velocity (Fig. 7). By using the general dendrite growth model to calculate the tip undercooling, Ga¨umann et al. (Ref 23) have numerically calculated the CET condition, shown on a G-V plot for different nucleation undercooling values, as shown in Fig. 11. Note that anisotropy in interface energy also influences CET, since a larger anisotropy parameter will sharpen the tip radius and reduce the tip undercooling at a given velocity. This will stabilize the columnar region, and the CET curve in Fig. 11 will move upward as the anisotropy parameter is increased. One of the ways to enhance the equiaxed zone is to increase the density of nuclei by the addition of effective grain refiners or inoculants. In casting and welding, the nuclei often originate from the detachment of the secondary arms close to the columnar dendrite tip region due to the fluctuating thermal field caused by convection. These fragments grow at the slightest undercooling; therefore, they are highly effective in producing CET. Fragmentation is often induced by stirring the liquid in the mold to obtain castings with equiaxed grains (Ref 24).
Nonplanar Eutectic Growth Eutectic alloys are important for casting because of their good fluidity caused by the small freezing range and low melting temperatures. The fine-scale composite microstructure also gives them superior mechanical properties. Thus, theoretical models have been developed to correlate eutectic spacing with velocity (Ref 25, 26), and very fine spacing of the order of 10 nm has been observed in aluminum-copper alloys at a velocity of 0.2 m/s (0.7 ft/s)
(Ref 27). A eutectic alloy in a binary system grows with a planar interface. However, it is observed that at higher velocities, a planar two-phase interface becomes unstable, and two-phase cellular and dendritic microstructures can form (Ref 28, 29). These nonplanar eutectic fronts in a model transparent system of carbon tetrabromide and hexachlorethane, observed by Han (Ref 29), are shown in Fig. 12, which shows the formation of planarto-cellular-to-dendritic two-phase eutectic as the velocity is increased. Two types of eutectic cells are observed: one in which the anisotropy effects are small (Fig. 12b), and the other (Fig. 12c), in which the anisotropy effects are significant. The eutectic cell is generally referred to as the colony formation. Detailed analytical (Ref 30) and numerical (Ref 31) models have been developed to examine the long wavelength instability of a planar eutectic front in binary alloys. They have clearly shown that a large-scale morphological instability of a two-phase planar eutectic front does not occur in binary alloys. The instability of a planar two-phase interface is attributed to the presence of a small amount of impurity (Ref 32). As in the case of a morphological instability of a planar single phase, the instability of a planar eutectic interface is due to the rejection of the third component by both of the solid phases that build up a boundary layer of the ternary impurity. When the solute buildup increases to the point where the concentration gradient of the third element produces a constitutionally supercooled region, the planar eutectic interface becomes unstable and forms two-phase cells, which then transform to two-phase dendrites as the velocity is increased. Since most binary alloys contain some impurity, a nonplanar eutectic will form when the velocity exceeds some critical value. Thus, colony formation is common in many rapidly solidified eutectics.
Non-Plane Front Solidification / 305 further increase in velocity, the eutectic spacing in the colony decreases. Furthermore, the spacing between the colonies is found to decrease with an increase in velocity, and the colony spacing is found to be approximately an order of magnitude larger than the eutectic spacing. However, as the planar interface becomes unstable, the interface temperature, the eutectic spacing at the center of the colonies, and the colony spacing are all found to increase initially, reach a maximum, and then decrease as the cellular structure transforms into a two-phase dendritic structure. Experimental results are analyzed by using appropriate theoretical models on cellular, dendritic, and eutectic growth.
Summary
Fig. 12
Typical microstructures observed in the carbon tetrabromide-hexachlorethane system of eutectic composition as a function of velocity. (a) At low velocity, below 1.0 mm/s, a macroscopically planar eutectic interface is observed. (b) Cellular two-phase interface forms when the velocity is increased. (c) Another type of cellular morphology forms depending on the orientation and the anisotropy in interface energy. (d) At higher velocities, a dendritic two-phase eutectic is observed. Source:Ref 29
G mc C0 ðkc 1Þ V kc D
Fig. 13
(a) Variation in colony tip temperature with velocity. (b) Variation in the lamellar spacing at the center of the colony with velocity. An increase in spacing is observed when the planar interface becomes cellular eutectic.
The model for the onset of instability has been based on the single-phase instability, which predicts that the eutectic interface will be unstable when:
(Eq 11)
where C0 is the concentration of the impurity, and mc is the slope of the liquidus with the impurity content. The solute distribution coefficient is taken to be the weighted average value from both the solid phases, that is, kc = b la kac + lb kbc c/l, in which l is the eutectic spacing, and la and lb are the width of the a- and b-lamella. The distribution coefficients of the solute for the a- and b-phases are kac and kbc, respectively. The variation of the interface temperature with velocity is predicted by using the analogy with the model for the single-phase cell and dendrite (Ref 33). For a planar eutectic interface, the interface temperature decreases with velocity until a critical velocity is reached where the interface becomes unstable. The interface temperature increases as 1/V in the cellular colony regime, and then decreases when the cellular colony transforms to a dendritic colony. Quantitative measurements on the variation of interface temperature, eutectic spacing, and colony spacing with velocity were carried out by Han (Ref 29) over the two-phase planar, cellular, and dendritic structures. The results are shown in Fig. 13(a). The eutectic spacing varies within a colony, and it increases near the boundary of the colony due to the curvature of the front, so that one characterizes the spacing at the center of the colony. The variation of eutectic spacing with velocity is shown in Fig. 13(b). It decreases with an increase in velocity in the planar front region, but then it increases discontinuously as the planar front becomes nonplanar. With
Nonplanar microstructures form most frequently during the solidification of alloys, and they play a crucial role in governing the properties of the solidified material. The formation of cellular and dendritic structures in one- and two-phase structures is presented with emphasis on the effect of processing conditions and composition on the selection of microstructure and microstructure scales. A general microstructure map is shown that correlates processing conditions of G and V, or the cooling rate, in different solidification technologies. The basic ideas and the processing conditions required to control the columnar and equiaxed microstructures are also emphasized. ACKNOWLEDGMENTS This work was carried out at Ames Laboratory, which is operated by Iowa State University for the Office of Basic Energy Science, Division of Materials Science, U.S. Department of Energy under Contract No. W7405-Eng-82. REFERENCES 1. B. Billia and R. Trivedi, Pattern Formation in Crystal Growth, Chap. 13, Handbook of Crystal Growth, Vol 1, D.T.J. Hurle, Ed., Pergamon Press, 1993, p 899 2. J.D. Hunt and S.-Z. Lu, Numerical Modeling of Cellular/Dendritic Array Growth: Spacing and Structure Predictions, Metall. Mater. Trans. A, Vol 27, 1996, p 611 3. R. Trivedi, Y.X. Shen, and S. Liu, Cell to Dendrite Transition in Directional Solidification of Alloys, Mater. Trans. A, Vol 34, 2003, p 395 4. T.F. Bower, H.D. Brody, and M.C. Flemings, Measurements of Solute Redistribution in Dendritic Solidification, Trans. TMS-AIME, Vol 236, 1966, p 624 5. R. Trivedi, Y.X. Shen, and S. Liu, A Unique Correlation among Cellular Microstructural Length Scales in Directionally Solidified Binary Systems, Proc. Int. Conf. on Advanced Materials and Processing,
306 / Principles of Solidification
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N. Chakraborty and U. Chatterjee, Ed., Tata McGraw-Hill, New Delhi, 2001, p 42 T.W. Clyne and W. Kurz, Solute Redistribution during Solidification with Rapid Solid State Diffusion, Metall. Trans. A, Vol 12, 1981, p 965 R. Trivedi, F. Jin, and I.E. Anderson, Dynamical Evolution of Microstructure in Finely Atomized Droplets of Al-Si Alloy, Acta Mater., Vol 51, 2003, p 289 J.H. Lee, H.C. Kim, C.Y. Jo, S.K. Kim, J.H. Shin, S. Liu, and R. Trivedi, Microstructure Evolution in Directionally Solidified Fe-18Cr Stainless Steels, Mater. Sci. Eng. A, Vol 413–414, 2005, p 30 B. Chalmers, Principles of Solidification, John Wiley and Sons, New York, 1964 F. Gonzales and M. Rappaz, Dendrite Growth Direction in Aluminum Zinc Alloys, Metall. Mater. Trans. A, Vol 37, 2006, p 2797 E. Sunseri, M.S. thesis, Iowa State University, 2008 T. Haxhimali, A. Karma, F. Gonzales, and M. Rappaz, Orientation Selection in Dendritic Evolution, Nature Mater., Vol 5, 2006 S. Liu, S.R. Napolitano, and R. Trivedi, Measurement of Anisotropy of CrystalMelt Interfacial Energy for Binary Al-Cu Alloy, Acta Mater., Vol 49, 2001, p 42, 71 M. Rappaz, Modeling of Microstructure Formation in Solidification Processes, Int. Mater. Rev., Vol 34, 1989, p 93 G.P. Ivantsov, Temperature Field Around Spherical, Cylindrical and Needleshaped
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Crystals Which Grow in Supercooled Melt, Dokl. Akad. Nauk SSSR, Vol 58, 1947, p 567 A. Karma and W.J. Rappel, Quantitative Phase-Field Modelling of Dendritic Growth in Two and Three Dimensions, Phys. Rev. E, Vol 57, 1998, p 4323 R. Trivedi and W. Kurz, Dendritic Growth, Int. Mater. Rev., Vol 34, 1990, p 93 J.D. Hunt, Primary Dendrite Spacing, Solidification and Casting of Metals, Book 192, The Metals Society, London, 1979, p 3 R. Trivedi, Interdendritic Spacing, Part II: A Comparison of Theory and Experiments, Metall. Trans. A, Vol 15, 1984, p 977 T.Z. Kattamis and M.C. Flemings, Dendrite Morphology, Microsegregation, and Homogenization of Low Alloy Steel, Trans. TMS-AIME, Vol 233, 1965, p 992 U. Feurer and R. Wunderlin, Einfluss der Zusammensetzung und der ErstarrungsBedingungen auf die Dendritenmorphologie Binarer Al-Legierungen, Fachbericht DGM, Oberursel, Germany, 1977 J.D. Hunt, Steady State Columnar and Equiaxed Growth of Dendrites and Eutectic, Mater. Sci. Eng., Vol 65, 1984, p 75 M. Ga¨umann, R. Trivedi, and W. Kurz, Nucleation Ahead of the Advancing Interface in Directional Solidification, Mater. Sci. Eng. A, Vol 226–228, 1997, p 763 S. Kobayashi, H. Tomono, and K. Toda, Ame´lioration de la Qualite´ des Brames de Coule´e Continue Graˆce au Brassage E´lectromagne´tique par Courant de Conduction,
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Rev. Me´tall., Cah. Int. Tech., Vol 80, 1983, p 887 K.A. Jackson and J.D. Hunt, Lamellar and Rod Eutectic Growth, Trans. TMS-AIME, Vol 236, 1966, p 1129 R. Trivedi and W. Kurz, Microstructure Selection in Eutectic Alloy Systems, Solidification Processing of Eutectic Alloys, D.M. Stefanescu, G.J. Abbaschian, and R. J. Bayuzik, Ed., Metallurgical Society of AIME, 1988, p 3–34 M. Zimmerman, M. Carrad, and W. Kurz, Rapid Solidification of Al-Cu Eutectic Alloy by Laser Remelting, Acta Metall., Vol 37, 1989, p 3305 H.W. Weart and D.J. Mack, Eutectic Solidification Structures, Trans. Metall. Soc. AIME, Vol 212, 1958, p 664 S.H. Han, “Stability of a Eutectic Interface during Directional Solidification,” Ph.D. thesis, Iowa State University, Ames, IA, 1995 V. Datye and J.S. Langer, Stability of Thin Film Lamellar Eutectic Growth, Phys. Rev. B, Vol 24, 1981, p 4155 A. Karma and A. Sarkissian, Morphological Instabilities of Lamellar Eutectics, Metall. Trans. A, Vol 27, 1996, p 635 W.A. Tiller, Liquid Metals and Solidification, American Society of Metals, 1958 D.G. McCartney, J.D. Hunt, and R.M. Jordan, The Structures Expected in a Simple Ternary Eutectic System, Part 1: Theory, Metall. Trans. A, Vol 11, 1980, p 1243
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 307-311 DOI: 10.1361/asmhba0005211
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Eutectic Solidification—An Introduction Sumanth Shankar, McMaster University
METALLIC EUTECTIC SYSTEMS have been used for several centuries to cast engineering components. The use of eutectic systems predates the understanding of the solidification of such systems. Systems wherein two pure metals can be mixed to specifically yield a low-melting alloy have enabled ease of casting near-net shape components of high strength and performance. Rapid developments in our understanding of the solidification of such systems over the second half of the 20th century have enabled improvements in the processing of eutectic metallic systems. Presently, commercially important metallic alloy systems used for casting engineering components are predominantly eutectic systems. Aluminumsilicon, aluminum-copper, and iron-carbon systems are typical examples.
Binary Eutectic Phase Diagram Figure 1 shows a schematic of a typical phase diagram of a binary eutectic system between two components, A and B. In Fig. 1, there is no solubility of B in A or vice versa. This type of eutectic system is typical of nonmetallic systems such as salts and organic alloys. The LiBr-LiF and NaF-RbF systems are typical examples of such systems. However, metallic eutectic systems seldom exhibit the behavior shown by Fig. 1. Figure 2 shows a schematic of a typical phase diagram of a binary metallic eutectic system. In Fig. 2, there is a solid solubility of component B in A (designated as a-phase) and a solid solubility of component A in B (designated
as b-phase). Although A and B have high melting points individually, the alloy composition denoted by point “E” in Fig. 2 has a considerably lower melting point than components A or B. Alloys with composition close to point “E” will have a short freezing range and a low melting point. Lead-tin, aluminum-silicon, aluminum-copper, iron-carbon (Fe-Fe3C), Silvergermanium, bismuth-lead, and Zn-Mg2Zn11 are typical examples of binary metallic eutectic systems exhibiting the behavior shown in Fig. 2. In Fig. 2, the lines CE and CF are the liquidus and solidus lines for the alloys, respectively, which solidify to form solid-solution a. Similarly, lines DE and DG are the liquidus and solidus lines for the alloys, respectively, which solidify to form the solid-solution b. Intermediately, consider the liquid of composition CE, which will cool until temperature TE, at which point the liquid will be a part of two liquidus lines, CE and DE, and be in equilibrium with two solid solutions of compositions Ca and Cb. Both of these solid solutions will form at temperature TE from the liquid composition CE. The combination of these solid solutions is known as the eutectic. The eutectic is a mixture of two solid solutions wherein the two solid phases intimately coexist. However, the distribution of these two solid phases in the eutectic structure will vary vastly among various binary metallic eutectic systems and will solely depend on the nature of the components that form the system. To further understand the various structures that evolve in a binary eutectic system during solidification, it is useful
to examine the various classifications of eutectic structures and relate them to principles of solidification.
Eutectic Structures For the solids with Ca and Cb to evolve from the liquid at temperature TE in Fig. 2, there must be a nucleation event and subsequent growth of the solid phases to form the final solidified eutectic structure. From studies of over 60 binary metallic eutectic systems containing solid solutions a and b (Ref 1–4), researchers have shown that if solid-solution a can be nucleated by b, then solid-solution b cannot be nucleated by a. For example, in the system shown in Fig. 2, assume that a is nucleated by b and b cannot be nucleated by a. For an alloy with a composition of B greater than CE (hypereutectic alloys), the eutectic b-phase will nucleate or grow epitaxially on the primary b-phase at the eutectic temperature, TE. Subsequently, the eutectic a-phase will nucleate on the b-phase with minimal undercooling to initiate the growth of the eutectic structure. However, in an alloy with a composition of B less than CE (hypoeutectic), the eutectic a-phase will initially nucleate or grow epitaxially on the primary a-phase; however, the eutectic b-phase cannot nucleate on the a-phase and hence will require a nucleating agent in the form of an impurity phase in the alloy or a substantially large undercooling to nucleate.
Classification of Eutectic Structures The most widely accepted classification of the eutectic structures was first proposed by Sheil (Ref 5, 6) and was later modified and elaborated by Jackson et al. (Ref 7–9). Jackson et al. (Ref 7–9) introduced a parameter, a, to explain the entropy of melting: a¼
Fig. 1
Schematic of a typical simple binary eutectic system with no solubility of component A in B or vice versa
Fig. 2
Schematic of phase diagram of a typical metallic binary eutectic system
SF R
(Eq 1)
where (DSF/R) is the entropy of melting in dimensionless units, and x is a crystallographic factor that is less than and almost unity. Solids with high entropy of melting will have a > 2, and these solids grow with crystalline facets (atomically
308 / Principles of Solidification smooth interface). Solids with low entropy of melting will have a < 2, and these grow isotropically with no facets (atomically rough interface). Eutectic structures can be broadly classified into predictable morphologies based on the value of a for the two solid solutions, respectively. Further, Scheil (Ref 5, 6) proposed that the volume fraction of solute phase in the eutectic structure will also influence the morphology of the eutectic structure. Hence, the morphology of the eutectic structure in a binary metallic system can be broadly classified as shown in Fig. 3. In Fig. 3, the regular or normal eutectic structure, such as the rodlike and the lamellar, will form when the two phases have low entropies of fusion (a < 2) (nonfaceted/ nonfaceted-type growth). The rodlike structure typically occurs when the volume fraction, VF, of the minor phase is less than approximately 30% (Fig. 3a), and the lamellar structure typically occurs when VF of the minor phase is greater than approximately 30% (Fig. 3b). The irregular or anomalous eutectic structure forms when one of the solid phases has high entropy of fusion (a > 2) (nonfaceted/faceted-type growth). In the nonfaceted/faceted-type growth, when VF < 30%, the faceted rodlike or spherical eutectic phases will typically occur (Fig. 3c), and when VF > 30%, the faceted branched (irregular) lamellar or the acicular structure will typically occur (Fig. 3d). Metallic systems very rarely exhibit a eutectic system wherein the two solids have high entropy of fusion (faceted/ faceted-type growth). These are typically found in nonmetallic systems. Unlike regular eutectic structures that are either rodlike or lamellar, the anomalous eutectic structures have a wide variety of morphology, depending on the system and the solidification conditions. Croker et al. (Ref 10) further refined the classification of the eutectic structures, as shown in Fig. 4. In Fig. 4, 5.5 cal mol1 K1 (23 J mol1 K1) was evaluated as the critical value to differentiate between normal eutectic structures, such as lamellar or rodlike
(two nonfaceted solids), and anomalous eutectic structures (nonfaceted/faceted). Moreover, Fig. 4 was drawn for various eutectic structures solidified at a growth rate of approximately 5 104 cm s1 (2 104 in s1). Figure 5 shows a comprehensive map of the location of the eutectic structures of various binary metallic alloy systems solidified at a growth rate of 5 104 cm s1 (2 104 in s1). In all of these alloys, the first eutectic solid forms by heterogeneous nucleation, and the second solid nucleates contiguously and heterogeneously but may not necessarily nucleate on the first solid phase. In Fig. 5, regions marked 1 and 2 are termed normal or regular eutectic structures and constitute structures formed with two solids growing without crystalline facets. At high volume fraction (VF) of over 32%, a lamellar eutectic structure dominates, as shown in region 1; however, at low VF (<32%), the rodlike eutectic structure dominates. Regions 3 to 6 represent anomalous eutectic structures obtained by one of the solids having a faceted structure and the other a nonfaceted structure. Region 3 occurs at VF 10%, wherein the
structure is that of broken lamellar. As the VF increases, the broken lamellar structure transitions to an anomalous irregular flake structure (region 4). As VF increases to greater than approximately 20%, the irregular structure transitions into a complex regular structure (region 5) that consists of large regions of regular eutectic structure that are cellular in nature. In between the complex cellular regular structure, the anomalous irregular structure exists. As VF increases to higher values, the VF of the irregular structure diminishes. Further increase in the value of VF leads to region 6, wherein the morphology of the eutectic structure is termed quasi-regular. Typically, fibrous morphology coupled with platelike morphology, with varying proportions of each, is characteristic of the quasi-regular structure. The solid phase with the high entropy value forms the matrix of this eutectic structure. The structure in this region appears to be more regular than the anomalous structures with lower values of VF because of the reduced tendency of overgrowth exhibited by the higher VF of the faceted phase. Eutectic structures in region 6 are very sensitive to
Fig. 3
Schematic of the four broad categories of regular eutectic structures. (a) Regular rodlike (volume fraction, VF, <30%; entropy of fusion, a, <2). (b) Regular lamellar (VF > 30% and a < 2). (c) Faceted rods or spherical (VF < 30% and a > 2). (d) Faceted branched lamellar or acicular (VF > 30% and a > 2). The growth direction in all of these schematics is perpendicular to the plane of the paper.
Fig. 4
General classification of eutectic structures as a function of the entropy of fusion (DSF) and the volume fraction (VF). There are six regions of classification based on the morphology of the eutectics. These are typical microstructures corresponding to the six regions: 1, Pb-AuPb2 eutectic; 2, Cd-Pb eutectic; 3, Ag-Bi eutectic; 4, CdSbCd eutectic; 5, Bi-Pb2Bi eutectic; and 6, Bi-Cd eutectic. The eutectic growth directions in all of these microstructures are perpendicular to the plane of the paper, and the growth rate is approximately 5 104 cm s1 (2 104 in s1). Source: Ref 10
Eutectic Solidification—An Introduction / 309
Fig. 5
Classification of eutectic structures into six distinct regions as a function of entropy of fusion (DSF) and volume fraction (VF). The structures obtained were for a growth rate of 5 104 cm s1 (2 104 in s1). Structures left of the vertical line at 5.5 cal mol1 K1 (23 J mol1 K1) are formed by nonfaceted/nonfaceted-type growth, and those to the right of this line are formed by nonfaceted/faceted-type growth. Region 1, predominantly regular lamellar (□); 2, regular rodlike (); 3, broken lamellar and a few fibrous (▲); 4, irregular morphologies (); 5, complex regular structures with contours showing the volume fraction regular cells ( ); and 6, quasi-regular structures ($). Source: Ref 10
▪
310 / Principles of Solidification growth conditions. Ledeburite (white iron eutectic of the Fe3C-Fe system) is a typical example of the quasi-regular structure. The classification of the eutectic structures mentioned in this system is intended to present a general and broad overview of the types of morphologies of various systems, depending on the entropy of melting and VF of the two solid phases. It must be noted that the variations in factors such as nucleation, solidification rate, levels of impurities, surface energies, and crystallographic relationships can significantly influence and alter the morphology of eutectic phases in a given binary alloy system. Nucleation plays a significant role in defining the morphology of eutectic structures. In nearly all cases of binary eutectic reactions, nucleation of the eutectic phases occurs by a heterogeneous mode, since the homogeneous mode is seldom viable in freezing alloys. Hence, the inoculant impurity levels, distribution in melt, and their potency are critical in defining the undercooling required for nucleation. The potency of nucleation by an impurity inoculant is mainly controlled by the surface energies (Ref 4). If the energy of the nucleus/liquid surface is higher than the energy of the nucleated phase/liquid surface, heterogeneous nucleation takes place. Further, if the second solid eutectic phase easily nucleates on the first phase, the resultant structure is typically that of a normal lamellar or rodlike eutectic. Alternatively, if the two eutectic solid phases nucleate independently on separate inoculant phases, the resultant structure is typically anomalous.
subsequent solidification of a- and b-phases together. The same mechanism occurs in alloy 3, wherein the a-phase separates out of the supersaturated liquid to freeze around the primary b-phase (forming a halo) to initiate simultaneous evolution of a- and b-phases from the eutectic liquid. In alloy 2, the b-phase cannot nucleate until point “B2” is reached, wherein the liquid is at a temperature in which the a-phase can nucleate. Subsequently, the evolution of b and a renders the liquid to eutectic composition for the simultaneous evolution of both the eutectic phases. In alloy 2, formation of a halo is unnecessary. The size of the halo is proportional to the distances ED. Large halos are observed in alloys (Fig. 7a) requiring high undercooling for nucleation and have a small slope of the metastable liquidus line; insignificant halos are observed in alloys (Fig. 7b) requiring small undercooling for nucleation and have a large slope for the metastable liquidus line (Ref 3). This discussion explaining the observed morphologies of eutectic structures in metallic binary systems is a general overview based on
thermodynamic considerations alone. There are several researchers (Ref 9, 11–15) who have successfully predicted the growth and structure of regular lamellar or rodlike eutectics, wherein the entropy of melting is less than 5.5 cal mol1 K1 (23 J mol1 K1) (Fig. 5). Constitutive relationships among growth rate, n, eutectic spacing, l, and degree of undercooling, DT, have been developed. The general relationships among them are:
2 nl ¼ Constant
(Eq 2)
and T pffiffiffi ¼ Constant n
(Eq 3)
Various theories have been proposed to explain these relationships, and several rigorous mathematical and phenomenological models have been developed to explain the various features of the regular lamellar and rodlike eutectic structures (Ref 9, 13–16). However, the
Formation of Halos In most eutectic structures, the minor solid phase typically solidifies around the primary solid phase, forming a halo. Halos are observed in eutectic microstructures of most alloy systems. In some, they are visually prominent, and in others they are insignificant. Sundquist et al. (Ref 3) explained the formation of these halos based on nucleation of the phases in a solidifying eutectic alloy. Figure 6 shows a schematic of a typical binary eutectic equilibrium phase diagram with the metastable line extensions of the two liquidus lines. Halo formation can be understood by examining the solidification of the three alloys marked 1, 2, and 3. In alloy 1, the nucleation of the primary phase, a, takes place at point “B1”, after sufficient undercooling from point “A1”. Immediately, the liquid composition moves to point “L1”, and further cooling eventually renders the liquid at point “E”. Nucleation of b-phase requires an undercooling as well and takes place at point “D1”. During cooling between points “E” and “D1,” additional growth of the a-phase occurs along the metastable liquidus line, EC8. The b-phase can be nucleated by a or some other impurity in the melt. The liquid at point “D1” is supersaturated with b-phase, and the excess b has to solidify from the liquid onto the primary a-phase (forming a halo of b on primary a) to render the remaining liquid back to the eutectic composition for
Fig. 6
Schematic of eutectic phase diagram showing the solidification sequence of alloys 1, 2, and 3. Points “A1,” “A2,” and “A3” are equilibrium freezing points for alloys 1, 2, and 3, respectively. Point “E” is the eutectic point. Point “B” is the actual temperature of initiation of solidification of primary a-phase. Line EC is the metastable liquidus line. Point “D” is the nucleation point of the second solid phase. Line DE is the excess of the secondary phase. The subscripts in all of the points denote the respective alloys 1, 2, and 3.
Fig. 7
Typical halos found in eutectic structures. (a) Ag-36 wt% Cu showing primary copper (black) with silver halo. (b) Al-35 wt% Cu showing eutectic radiating from primary CuAl2. Source: Ref 3
Eutectic Solidification—An Introduction / 311 predictability of the structure of eutectic growth produced by a pair of nonfaceted/faceted solid solutions is still an unsolved problem. There have been many theories proposed to explain the anomalous eutectic structures, but it has now come to be widely accepted that the influence of extraneous factors such as solidification rates, impurity levels, and crystallographic factors of solid phases play a significant role in determining these structures. It is best that these structures be explained on a case-by-case basis for each binary eutectic alloy and processing condition. Two of the most significant alloy systems for commercial engineering applications are the aluminum-silicon and the iron-carbon (Fe-Fe3C) systems. A detailed review of the evolution of eutectic structures in these systems is presented in subsequent articles in this Volume.
REFERENCES 1. E. Scheil and R. Zimmermann, Z. Metallkd., Vol 48, 1957, p 509 2. B.E. Sundquist and L.F. Moldolfo, Trans. Metall. Soc. AIME, Vol 221, 1961, p 157 3. B.E. Sundquist, R. Bruscato, and L.F. Moldolfo, J. Inst. Met., Vol 91, 1962–1963, p 204 4. L.F. Moldolfo, J. Aust. Inst. Met., Vol 10, 1965, p 169 5. E. Scheil, Z. Metallkd., Vol 37, 1946, p 1 6. E. Scheil, Z. Metallkd., Vol 45, 1954, p 298 7. K.A. Jackson, Liquid Metals and Solidification, American Society for Metals, 1958, p 174 8. K.A. Jackson, Growth and Perfection of Crystals, R.H. Doremus, D. Turnbull, and B.W. Roberts, Ed., John Wiley and Sons, New York, 1958, p 319
9. J.D. Hunt and K.A. Jackson, Trans. Metall. Soc. AIME, Vol 236, June 1966, p 843 10. M.N. Croker, R.S. Fidler, and R.W. Smith, Proc. R. Soc. (London) A, Math. Phys. Sci., Vol 335 (No. 1600, Oct 2, 1973, p 15 11. C. Zener, Trans. AIME, Vol 167, 1946, p 550 12. W.A. Tiller, Liquid Metals and Solidification, American Society for Metals, 1958 13. M.C. Flemings, Solidification Processing, McGraw-Hill, 1974 14. D.A. Porter and K.E. Easterling, Phase Transformation of Metals and Alloys, 2nd ed., CRC Press, Boca Raton, FL, 2004 15. W. Kurz and J.D. Fischer, Fundamentals of Solidification, Trans Tech Publications, Clausthol-Zellerfield, Germany, 1984 16. B. Chalmers, Principles of Solidification, John Wiley and Sons, New York, 1964
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 312-316 DOI: 10.1361/asmhba0005212
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Solidification of Eutectic Alloys— Aluminum-Silicon M.M. Makhlouf, Worcester Polytechnic Institute
Figure 1 shows the aluminum-silicon equilibrium phase diagram (Ref 5). It is a simple binary eutectic with limited solubility of aluminum in silicon and limited solubility of silicon in aluminum. The solubility of silicon in aluminum reaches a maximum 1.5 at.% at the eutectic temperature, and the solubility of silicon increases with temperature to 0.016 at.% at 1190 C (2175 F). There is only one invariant reaction in this diagram, namely Eq 1, where L is the liquid phase, a is predominantly aluminum, and z is predominantly silicon. It is now widely accepted that this eutectic reaction occurs at 577 C (1070 F) and at a silicon concentration of 12.6 wt%: L!þ
(Eq 1)
It is interesting to note that in 1926, Gwyer and Phillips (Ref 6) determined the composition and temperature of the eutectic reaction in the aluminum-silicon binary system to be 11.7 wt% Si and 577 C (1070 F), respectively. Although the eutectic temperature is in accord with the currently accepted value, the eutectic þ 0.1 composition was later changed to 12.2 at.% Si (Ref 7–9). The initial error in establishing the eutectic point is due to the fact that the temperature of formation of the primary constituents (aluminum and silicon), as well as the temperature of the aluminum-silicon eutectic plateau, is cooling-rate dependent. Primary silicon undercools more than primary aluminum;
0
10
20
30
40
Equilibrium Metastable extensions T0 curves
1400
1200
Depression of the eutectic temperature with increased cooling rate may also be explained on the basis of the coupled region effect (Ref 12). Coupled regions represent fields within the phase diagram where the two phases of the eutectic structure are organized in the solid in such a way that diffusion in the liquid occurs effectively at a duplex solid/liquid front (Ref 8, 12–15). Regions of coupled growth are shown schematically in Fig. 3 for both a symmetric and an asymmetric hypothetical phase diagram. In the aluminum-silicon system, silicon is a nonmetal with directed covalent bonds; therefore, it tends to grow anisotropically into faceted crystals. Hence, it requires more undercooling for its growth than the isotropic
Silicon, wt% 60 50
70
80
90
100 2500
L 2000
1000 1500
800
600 1000
Temperature, ⬚F
The Aluminum-Silicon Equilibrium Phase Diagram
hence, the eutectic structure forms at 10 to 12 C (18 to 22 F) below the eutectic temperature without appreciable recalescence (Ref 10). Consequently, at higher cooling rates, the system behaves as though the eutectic point is shifted to higher silicon contents, and the eutectic temperature is depressed. Figure 2 illustrates the effect of this apparent shift in the eutectic point on the microstructure of a typical aluminum-silicon alloy. Figure 2(a) depicts the microstructure of a eutectic aluminum-silicon alloy that is slowly cooled and shows no primary aluminum; on the other hand, Fig. 2(b) depicts the microstructure of the same alloy cooled at a significantly faster rate and shows primary aluminum dendrites and finer silicon particles.
Temperature, ⬚C
THE ALUMINUM-SILICON EUTECTIC REACTION has been the subject of intense interest ever since Pacz introduced aluminumsilicon alloys to the United States in 1921 (Ref 1). In the years that followed, there was an explosion of ideas concerning this technologically important transformation, and much of the current understanding of the reaction may be traced back to the work of Mondolfo (Ref 2) and to the more recent work of Lu and Hellawell(Ref 3, 4).
400 500 200
(Si)
0
0
-200 0 Al
Fig. 1
10
20
30
40
50 Silicon, at.%
60
70
80
90
100 Si
Equilibrium phase diagram for the aluminum-silicon system showing metastable extensions of the liquidus and solidus lines. Source: Ref 5
Solidification of Eutectic Alloys—Aluminum-Silicon / 313
Fig. 2
(a)
Fig. 3
Al-12.5wt%Si alloy (a) slowly cooled and (b) chill cast. Source: Ref 11
(b) Coupled zones in (a) symmetrical and (b) asymmetrical phase diagrams
aluminum phase. Consequently, the coupled region in the aluminum-silicon system is asymmetric.
Classification of the AluminumSilicon Eutectic Structure Typical eutectic structures of binary alloys form by the simultaneous growth of the two phases from the liquid; therefore, they may exhibit a variety of microstructures that can be classified according to two criteria: Lamellar versus fibrous morphology of the
individual phases: In general, when the two eutectic phases are present in approximately equal volume fractions, eutectics of binary alloys exhibit a lamellar structure. However, if one phase is present in a small volume fraction, this phase tends to be fibrous. As a rule of thumb, the eutectic structure obtained will tend to be fibrous when the volume fraction of the minor phase is less than 0.25; otherwise, it will tend to be lamellar (Ref 16). Regular versus irregular eutectic structure: If both phases in the eutectic structure are nonfaceted, the eutectic will exhibit a regular morphology. In this case, the microstructure is made up of either lamellae or fibers that have a high degree of regularity and periodicity. However, if one phase is faceted, the eutectic structure is often irregular.
The typical aluminum-silicon eutectic is closer to a lamellar structure than to a fibrous one, even though the volume fraction of silicon in the aluminum-silicon binary is less than 0.25. (The volume fraction of silicon in the aluminum-silicon eutectic is 0.143.) This structure is usually attributed to the strong anisotropy of growth of silicon and to the relatively low interfacial energy between silicon and aluminum (Ref 16). Moreover, the aluminum-silicon eutectic has always been designated an irregular eutectic; in fact, it is the most widely used example of the irregular eutectic structure.
Microstructure of the AluminumSilicon Eutectic The aluminum-silicon eutectic can exhibit either of two morphologies: an unmodified morphology and a modified morphology. The unmodified morphology, such as that shown in Fig. 2(a), is characterized by coarse and flaky silicon particles and is usually observed in slowly cooled foundry alloys (typically at cooling rates between 1 and 10 C/min, or 2 and 18 F/min) when no chemical modifiers are added. The silicon particles in an unmodified alloy seem to be disconnected when viewed at a low magnification without deep chemical etching. Consequently, it was earlier believed that each of the coarse silicon particles is an isolated crystal. Later, Gurtler (Ref 17) observed that radial silicon particles grew across the primary aluminum dendrites, as shown in Fig. 4, and concluded that such an occurrence is not possible unless the silicon particles extend in three dimensions. Crosley and Mondolfo (Ref 18) suggested that the needlelike silicon particles observed in unmodified alloys must be flakes or sheets. Now, with the availability of electron microscopy, it is confirmed that silicon in unmodified aluminum-silicon eutectics has a flakelike microstructure. At relatively fast cooling rates, such as those encountered in chill casting, the aluminum-silicon eutectic is much finer, and silicon assumes a fibrous morphology. Day and Hellawell (Ref 19) classified the morphology of eutectic silicon in directionally
Fig. 4
Continuity of eutectic silicon flakes around a primary aluminum dendrite. Source: Ref 18
solidified, unmodified eutectic aluminum-silicon alloys as a function of the temperature gradient (G), the growth rate (R), and the alloy composition, as shown in Fig. 5. The region marked “A” in Fig. 5 is characterized by massive silicon particles growing by long-range diffusion and by a solidification interface that is essentially planar and isothermal. Some of the silicon crystals in the eutectic are interconnected, and the individual silicon crystals are irregular, banded, and tend to contain numerous {111} growth twins. In the region marked “B,” silicon assumes a variety of morphologies depending on the temperature gradient (G). At higher temperature gradients, close to those of the region marked “A,” silicon appears as fibers with a round cross section, but as the temperature gradient decreases, silicon forms lateral branches with a characteristic fourfold symmetry (Ref 7, 21, 22). In the region marked “C,” the silicon phase displays a flakelike, interconnected morphology. This is the general class of irregular eutectic observed in chemically unmodified alloys. At high growth rates (R), the structure develops as shown in regions “C” and “D.” Figure 6 shows the microstructure of a chemically modified aluminum-silicon alloy. When observed at low magnification, the silicon particles in these alloys appear as isolated, more-orless spherical crystals, as shown in Fig. 6(a). Figure 6(b), which is a higher-magnification scanning electron microscope image of a deep-etched sample, reveals the fibrous morphology of this silicon phase. It should be noted, however, that in spite of the apparent similarities in the morphology of the silicon phase in the chemically modified alloy and the chill-cast alloy, the growth mechanism of silicon is different in the two cases.
Theories of Solidification of the Aluminum-Silicon Eutectic The flakelike, interconnected silicon phase characteristic of typical unmodified aluminumsilicon eutectics is believed to grow by twinning via a mechanism known as twin plane re-entrant edge (TPRE). The TPRE mechanism
314 / Principles of Solidification
A
G, K/mm
100
D
10
B
C
1 B+C
1
10
100
R, µm/s
(a)
(b)
(c)
(d)
(e)
Fig. 5
Plot of temperature gradient (G) versus growth rate (R) along with the microstructures in the different regions. (a) Region A, (b) Region B (R = 0.3 m/s, or 1 ft/s, and G = 20 C/mm, or 915 F/in.). (c) Region B (R = 0.3 m/s, or 1 ft/s, and G = 3 C/mm, or 135 F/in.). (d) Region C. (e) Region D. Source: Ref 19, 20
Fig. 6
Microstructure of an aluminum-silicon alloy modified with strontium. (a) At low magnification. (b) At higher magnification and deeply etched. Source: Ref 20
was first introduced by Hamilton and Seidensticker (Ref 23) to explain the growth of germanium dendrites and was later extended to the growth of silicon. The equilibrium habit of silicon is an octahedron bound by eight {111} planes. A twin crystal is half of the equilibrium crystal reflected across the remaining solid along the twin composition plane. Consequently, the outline of the twin silicon crystal consists of six edges of the intersection of pairs of {111} planes, as shown in Fig. 7. The
external angles between these bounding planes are 141 and 219 . The bounding planes that make a 141 external angle form a re-entrant corner, while those planes that make a 219 external angle form a ridge. Because of the more favorable bonding for an atom joining a re-entrant corner than one joining a ridge, re-entrant corners are more favorable nucleation sites than ridges. Thus, the presence of a re-entrant corner leads to rapid growth along the [211] direction. This rapid growth on the
re-entrant corner stops when a trigonal solid that is bounded entirely by ridges is formed, as shown in Fig. 7(b). However, if the crystal has two twins instead of one, as shown in Fig. 7(c), it will have six re-entrant corners located along the h211i direction. In this case, growth on the re-entrant corners generates more reentrant corners, as shown in Fig. 7(d), and the newly generated re-entrant corners relieve the blockage of nucleation sites that is caused by the formation of ridges. Figure 7(e) shows a solid with several steps that are growing simultaneously. Although Billings and Holmes (Ref 24) experimentally verified that the TPRE mechanism is responsible for the growth of germanium dendrites and observed that all germanium dendrites invariably contain two or more twins and never a single twin, it has not been shown that the TPRE mechanism is the mechanism responsible for growth of silicon dendrites or the flakelike silicon. Moreover, Kitamura et al. (Ref 25) questioned the adoption of the TPRE mechanism to explain growth in real crystals, because it was originally developed for perfect crystals. Kitamura et al. argue that screw dislocations, which are prevalent in real crystals, should be accounted for when developing a growth mechanism for faceted crystals such as silicon. Sunagawa and Yasuda (Ref 26) proved by means of x-ray topography that growth in quartz crystals occurs at the re-entrant corner of a twin junction due to the concentration of screw dislocations. Extension of the Sunagawa and Yasuda hypothesis to eutectic flakelike silicon crystals would invariably cast doubt that the TPRE is the operating mechanism in these crystals. The contemporary understanding of the formation of the aluminum-silicon eutectic (Ref 27, 28) has its origins in the nucleation rather than the growth aspects of the eutectic reaction and is based on the fact that aluminum-silicon alloys invariably contain iron. For example, alloys that are prepared from very pure materials, that is, 99.999% purity aluminum and 99.999% purity silicon, could have up to 50 ppm Fe (Ref 29–32). Although this level of iron is normally considered a trace-level impurity of little consequence, it has been established that it plays a significant role in the solidification of the aluminum-silicon eutectic (Ref 27). It is maintained that iron, except when present in exceedingly low amounts (0.0015%), results in the formation of an iron-containing b-(Al, Si,Fe) phase that plays an important role in the nucleation of the aluminum-silicon eutectic. Therefore, it is more appropriate to think of this system when discussing the eutectic reaction as essentially an Al-Si-Fe ternary system with the eutectic phases being Aleut + Sieut + b-(Al,Si, Fe), rather than a binary aluminum-silicon system. In hypoeutectic alloys of this system, nucleation of the eutectic phases proceeds as illustrated schematically in Fig. 8. During solidification, the primary aluminum phase forms as dendrites at the liquidus temperature of the
Solidification of Eutectic Alloys—Aluminum-Silicon / 315
219⬚ 219⬚ ––
[211] 219⬚ (b) Twin planes
II 219⬚ 141⬚
(a)
I ––
––
[211]
I
[211]
––
[211] 141⬚
141⬚
219⬚
Twin planes
Twin planes
(c)
(d)
Fig. 8
Sequence of events during nucleation of the eutectic phases in aluminum-silicon hypoeutectic alloys. Source: Ref 27
Twin planes (e)
Fig. 7
Schematic illustrating the twin plane re-entrant edge mechanism. (a) Crystal with a single twin. (b) Closure of twins due to ridge formation. (c) Crystal with two twins. (d) Creation of extra re-entrant corners I and II. (e) Propagation of crystal due to re-entrant corners. Source: Ref 23
10 µm (a)
Fig. 9
50 µm (b)
Microstructure of an aluminum-silicon alloy. (a) Unmodified. (b) Modified with strontium. Source: Ref 33
alloy. This is followed by the evolution of a secondary b-(Al,Si,Fe) phase at some temperature between the liquidus temperature and the eutectic temperature of the alloy, depending on the concentration of iron in the alloy. At the eutectic temperature, and at an undercooling of 0.4 to 0.8 C (0.7 to 1.4 F), eutectic silicon (Sieut) nucleates on the secondary b-(Al,Si,Fe) phase in the solute field ahead of the growing aluminum dendrites. Once nucleated, the eutectic silicon grows as flakes into the eutectic liquid. The liquid surrounding the eutectic silicon flakes becomes enriched with aluminum as it is being depleted of silicon; consequently, eutectic aluminum (Aleut) nucleates and grows on the edges and tips of the eutectic silicon flakes. Finally, the aluminum dendrites stop growing as they impinge on the growing eutectic aluminum grains.
Theories of Chemical Modification of the Aluminum-Silicon Eutectic Chemical modification refers to the morphological transformation that occurs in the eutectic phases of aluminum-silicon alloys upon addition of trace quantities of certain elements, such as strontium or sodium, and converts the otherwise platelike (flake) silicon phase, shown in Fig. 9(a), to the fine, fibrous (corallike) structure of Fig. 9 (b). This morphological transformation significantly enhances the mechanical properties and overall performance of components that are cast from these alloys (Ref 32). Consequently, much of the fundamental research in the aluminum-silicon alloy system since Pacz introduced it into the United States in 1921 (Ref 1) has been directed toward understanding the mechanism
behind the modification of the morphology of the eutectic silicon phase by trace element additions (Ref 33, 34). However, and in spite of many hypotheses proposed to explain the modification of the eutectic structure in aluminum-silicon alloys (Ref 34), the genesis of this technologically important morphological transformation remains uncertain. This is due mainly to the lack of conclusive evidence provided by experimentation in support of the proposed hypotheses. The main assumption underlying many popular hypotheses on the mechanism of chemical modification of the silicon phase morphology is that the eutectic silicon phase nucleates on the primary aluminum dendrites during solidification of the hypoeutectic alloys, and that the modifying element inhibits the growth of the eutectic silicon phase, thus transforming its morphology from platelike to fibrous (Ref 18, 35, 36). In 1987, Lu and Hellawell (Ref 3) introduced the concept of impurity-induced twinning to explain the mechanism by which chemical modifiers such as sodium and strontium affect the morphology of eutectic silicon. Soon after, this became the accepted eutectic modification theory, in spite of its inability to explain several observed phenomena associated with chemical modification of the aluminum-silicon eutectic. According to the impurity-induced twinning theory, chemical modifiers are impurities that poison already growing atomic silicon layers by becoming adsorbed onto surface steps and kinks, thus preventing the attachment of silicon atoms to the crystal. Furthermore, the adsorbed impurity atoms induce twinning by altering the stacking sequence of atomic layers as the newly added layers seek to grow around the adsorbed impurity atom. Lu and Hellawell assumed a face-centered cubic structure and calculated the ratio of impurity atom radius (ri) to matrix atom radius (r) that is required for impurity-induced twinning. They found it to be ri/r = 1.6457. Analyses for sodium in modified aluminumsilicon alloys using Auger electron spectroscopy (Ref 3) and electron microprobe measurements (Ref 37) show that sodium segregates in the
316 / Principles of Solidification silicon fibers and in the aluminum matrix, which seems to validate the impurity-induced twinning theory. However, careful examination shows that the impurity-induced twinning theory cannot explain many observed phenomena that are associated with chemical modification of the aluminum-silicon eutectic. Particularly, it cannot explain the relatively large undercooling during solidification that is observed with the evolution of the eutectic phases when modifying elements are present. It also cannot explain the occurrence of eutectic modification without chemical additives (i.e., eutectic modification due to an increased superheat and/or a relatively fast cooling rate). It is interesting to note that sodium, which has less than the ideal radius ratio for eutectic modification as predicted by the impurity-induced twinning theory, is a better modifier than ytterbium and calcium, which are very close to the ideal radius ratio. Moreover, lithium, which is much smaller than the ideal radius ratio, modifies the eutectic silicon morphology when added in large concentrations to aluminum-silicon alloys (Ref 38). Because of the inability of the impurity-induced twinning theory to account for these facts, research efforts that aim to develop a better understanding of chemical modification and to rationalize these well-established facts into a consistent theory are prevalent. The most promising of these efforts are those that maintain that the morphology of the aluminum-silicon eutectic is decided during its nucleation in the liquid phase (Ref 28). REFERENCES 1. A. Pacz, U.S. Patent 1,387,900, 1921 2. L. Mondolfo, J. Aust. Inst. Met., Vol 10, 1965, p 169
3. S. Lu and A. Hellawell, Metall. Trans. A, Vol 18, 1987, p 1721 4. S. Lu, and A. Hellawell, J. Cryst. Growth, Vol 73, 1985, p 316 5. J.L. Murray and A.J. McAlister, Bull. Alloy Phase Diagrams, Vol 5, 1984, p 74 6. A.G.C. Gwyer and H.W.L. Phillips, J. Inst. Met., Vol 36, 1926, p 283 7. J.A. Bell and W.C. Winegard, J. Inst. Met., Vol 93, 1965, p 318 8. H.W. Kerr, J.A. Bell, and W.C. Winegard, J. Aust. Inst. Met., Vol 10, 1965, p 64 9. R.A. Meussner, U.S. Naval Research Laboratory Report, NRL 5331, U.S. Naval Research Laboratory, 1959 10. M.L.V. Gayler, J. Inst. Met., Vol 39, 1927, p 157 11. A. Hellawell, Prog. Mater. Sci., Vol 15, 1970, p 1–78 12. G. Tamman and A.A. Botschwar, Z. Anorg. Chem., Vol 157, 1926, p 26 13. G. Chadwick, Prog. Mater. Sci., Vol 12, 1963, p 208 14. V.D.L. Davies and J.M. West, J. Inst. Met., Vol 92, 1963–1964, p 208 15. E. Scheil, Z. Metallkd., Vol 37, 1946, p 1 16. W. Kurz and D.J. Fisher, Fundamentals of Solidification, Trans. Tech. Publications, 1988 17. G. Gurtler, Z. Metallkd., Vol 44, 1953, p 503 18. P.B. Crosley and L.F. Mondolfo, Mod. Cast., Vol 46, 1966, p 89 19. M.G. Day and A. Hellawell, Proc. R. Soc. (London) A, Vol 305, 1968, p 473 20. H.V. Guthy, M.S. thesis, Worcester Polytechnic Institute, 2002 21. F.N. Rhines and W.F.B. Timpe, Z. Metallkd., Vol 48, 1957, p 109
22. Z. Obinata and N. Komatsu, Sci. Rep. Tohoku Univ. A, Vol 9, 1957, p 107 23. D.R. Hamilton and R.G. Seidensticker, J. Appl. Phys., Vol 31, 1960, p 1165 24. E. Billings and P.J. Holmes, Acta Cryst., Vol 8, 1955, p 353 25. M. Kitamura, S. Hosoya, and I. Sunagawa, J. Cryst. Growth, Vol 47, 1979, p 93 26. I. Sunagawa and T. Yasuda, J. Cryst. Growth, Vol 65, 1983, p 43 27. S. Shankar, Y.W. Riddle, and M.M. Makhlouf, Acta Mater., Vol 52, 2004, p 4447 28. M.M. Makhlouf, S. Shankar, and Y.W. Riddle, AFS Trans. Paper 05-088(02), p 2005 29. C.R. Ho and B. Cantor, Acta Mater., Vol 43, 1995, p 3231 30. D.L. Zhang and B. Cantor, Metall. Trans. A, Vol 24, 1993, p 1195 31. C.R. Ho and B. Cantor, J. Mater. Sci. Eng. A, Vol 173, 1993, p 173 32. B. Cantor, J. Mater. Sci. Eng., Vol 151, 1997, p 226 33. M.M. Makhlouf and H.V. Guthy, J. Light Met., Vol 1, 2001, p 199 34. F. Yilmaz and R.J. Eliott, Mater. Sci., Vol 24, 1989, p 2065 35. A.K. Dahle, K. Nogita, J.W. Zindel, S.D. McDonald, and L.M. Hogan, Metall. Mater. Trans. A, Vol 32, 2001, p 949 36. G. Heiberg, K. Nogita, A.K. Dahle, and L. Arnberg, Acta Mater., Vol 50, 2002, p 2537 37. O.A. Atasoy, Z. Metallkd. Vol 78, 1987, p 177 38. L. Clapham, Ph.D. thesis, Queens University, 1987
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 317-329 DOI: 10.1361/asmhba0005213
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Solidification of Eutectic Alloys— Cast Iron* Andreas Bu¨hrig-Polackzek and Daniel dos Santos, Gießerei-Institut
Structure of Liquid Iron-Carbon Alloys Observations from x-ray, neutron diffraction, and sound velocity measurements on liquid binary iron-carbon alloys at temperatures approximately 20 C (35 F) above the liquidus indicate that for up to 1.8% C, the distance between nearest iron neighbors, rI, as well as the number of nearest neighbors, NI, in the first coordination sphere increases (Fig. 1). Above 1.8% C, the distance remains constant, while the number of nearest neighbors continues to grow. Above 3.5% and up to 5.5% C, both the distance and the number of neighbors remain constant. Above 3.5% C, short-range order regions rich in carbon exist in the melt. This means that the melt becomes more dense with the addition of carbon. A maximum packing
0.27 rI
11
10
0.26 NI
9
8
0.25 0
1
2 Carbon, wt%
3
4
Fig. 1
Variation of distance between nearest neighbors (rI) and the number of nearest neighbors (NI) as a function of the percentage of carbon in the ironcarbon alloy. Source: Ref 1
10
Nucleation of Eutectic in Cast Iron Structure of Graphite in the AusteniteGraphite Eutectic. The graphite phase is a faceted crystal bounded by low index planes. For graphite crystallizing from an iron-carbon melt, the normally observed bounding planes are (0001) and f1010g, as shown in Fig. 3(a). The crystallographic structure of graphite and the possible growth directions, A and C, are shown in Fig. 3(b). Because unstable growth occurs on
Distance between nearest neighbours (rl), nm
Depending on composition, cooling rate, and liquid treatment, it is also possible to produce a mixed white-gray eutectic called mottled structure. The two basic types of eutectic are very different, with mechanical properties such as strength, ductility, and hardness varying over very large intervals as a function of the type and the amount of eutectic formed. To understand the mechanism of the solidification of cast iron, it is necessary first to discuss the structure of liquid iron-carbon alloys.
12
10
9
rI ↑, NI ↑
8
rI const. NI ↑
rI, NI const.
8
7
7 Short range ordered regions
6 5 4 0
9
6 5
0.5
1.0 1.5 2.0 2.5 3.0 3.5 Carbon concentration, wt%
4.0
Fig. 2
4 4.5
Viscosity of iron-carbon alloys, cP
stable diagram, Fe-Fe3C, the white eutectic or austenitic (g), iron carbide (Fe3C) forms. If solidification follows the stable diagram iron-graphite (a significant amount of silicon is required for this to occur), the gray eutectic, austenite-graphite (Gr), results.
f1010g planes, the edges of the platelike graphite crystals are not well defined. Graphite growing out of liquid iron-carbon alloys has a layer-type structure, with strong covalent bonds (4.19 105 to 5 105 J/mol, or 3.09 105 to 3.70 105 ft lbf/mol) between atoms in the same layer. There is a trielectronic bond of each atom with its neighbors, while the fourth electron is common for the layer giving the metallic Number of nearest neighbours (Nl), atoms
If solidification occurs according to the meta-
density is reached at 3% C, and it remains constant at higher carbon concentrations. The excess carbon forms carbon-rich regions (nonhomogeneities) in the melt. Viscosity measurements (Fig. 2) show a correlation between viscosity and percentage of carbon. This correlation can be further explained in terms of increased viscosity as the interatomic distance becomes smaller. Liquid iron-carbon alloys with low carbon content (<3.5% C, that is, steels and cast irons poor in carbon) are microscopically homogeneous. Liquid iron-carbon alloys with high carbon (>3.5% C, that is, cast irons rich in carbon) are colloidally dispersed systems with microgroups of carbon in liquid solution. The nature of these microgroups is not clear. It is hypothesized that they are either Fe3C clusters (Ref 3, 4) or Cn clusters (Ref 1, 5). Because the nucleation energy for Fe3C is smaller than that for graphite, it is thermodynamically possible for the carbon-rich regions to exist as Fe3C clusters. Other investigators consider the Cn clusters to be stable (Ref 3). Their size is supposed to be in the range of 1 to 20 mm (40 to 800 min.), and it increases with the carbon equivalent, lower silicon content, and lower holding time and temperature. It is to be expected that the carbon-rich configurations existing in molten iron-carbon alloys are in dynamic equilibrium and that they diffuse within the melt.
Viscosity of iron-carbon alloys, Pa · s ⫻ 10−3
CAST IRON is a binary iron-carbon or a multicomponent Fe-C-X alloy that is rich in carbon and exhibits a considerable amount of eutectic in the solid state. Two such possible eutectics may result, as follows:
Viscosity of iron-carbon alloys as a function of carbon concentration at a temperature 20 above liquidus. Source: Ref 2
*This article has been updated and revised from D.M. Stefanescu, Solidification of Eutectic Alloys: Cast Iron, Casting, Vol 15, ASM Handbook, ASM International, 1988, p 168–181.
318 / Principles of Solidification properties of graphite. Weak molecular forces exist between layers (4.19 103 to 8.37 103 J/mol, or 3.09 103 to 6.17 103 ft lbf/mol). The prism plane is a high-energy plane at which impurities absorb preferentially. Strength and hardness are higher in the C direction of the graphite crystal. Complete destruction of the graphite structure occurs only at approximately 4000 C (7230 F). This explains the presence of some graphite aggregates in molten iron even at temperatures considerably higher than the liquidus temperature. Depending on the chemical composition and the temperature gradient/growth rate ratio, G/R, and/or cooling rate, G R, a variety of graphite shapes can solidify as part of the austenitegraphite eutectic. Basically, they are as follows: Flake (actually plate) graphite (FG) Compacted/vermicular graphite (CG) Coral graphite Spheroidal (nodular) graphite (SG)
Crystalline structure of graphite. (a) Crystal of graphite bounded by (0001) and ð1010Þ type planes; the hexagonal arrangement of the atoms within the (0001) plane is shown relative to the bounding ð1010Þ faces. (b) Hexagonal structure of graphite showing the unit cell (heavy lines). Source: Ref 6
Fig. 3
Schematic of graphite types occurring in the austenite-graphite eutectic. (a) Flake graphite. (b) Compacted/ vermicular graphite. (c) Coral graphite. (d) Spheroidal graphite
2.0
212
752
932 108
50
90
40
72
30
54
20
36
10
18
0
0
Fig. 5
100
200 300 Superheating, ⬚C
18
Undercooling (ΔT ), ⬚F 36 54
72
90 50
Undercooling by rapid freezing
400
0 500
Relationship between superheating and maximum undercooling in flake graphite cast iron. Source: Ref 7
1.6 Eutectic cell count per mm
32 60
Superheating, ⬚F 392 572
0
Undercooling by superheating
40
1.2
30
0.8
20
0.4
10
0 0
Fig. 6
10
20 30 Undercooling (ΔT ), ⬚C
40
0 50
Influence of undercooling on eutectic cell count in flake graphite cast iron. Source: Ref 8
Eutectic cell count per in.
Fig. 4
Undercooling (ΔT ), ⬚F
A schematic of these graphite types, showing the traces of the f1010g planes, is given in Fig. 4. Nucleation of the Austenite-Flake Graphite Eutectic. Experimental evidence indicates that various types of nuclei become effective at various temperatures. Indeed, if only one type of nuclei would be active in a cast iron melt, an undercooling-superheating curve will show a single arrest for the temperature region over which the nuclei become effective. However, as shown in Fig. 5, a number of steps are observed in the relationship, suggesting that various foreign nuclei become effective as superheating is increased. Increasing the superheating apparently destroys the effective nuclei. Consequently, the number of eutectic grains (eutectic cells) decreases as the superheating increases, and as a result undercooling increases. However, when undercooling, DT, is increased at constant superheat by increasing the cooling rate or the growth rate, R, the eutectic cell count will increase (Fig. 6). Mathematically, this is shown by r* / 1/DT, where r* is the critical size radius. As DT increases, r* decreases, yielding an increasing number of critical size nuclei as well as more eutectic cells. Homogeneous nucleation is improbable in cast iron because typical undercoolings are much smaller than required by the theory of homogeneous nucleation (1 to 10 C, or 2 to 20 F, as opposed to 230 C, or 415 F). Nevertheless, some Cn clusters (as discussed earlier in this section) or undissolved graphite may act as heterogeneous nuclei for graphite formation. When two separate melt charges of similar composition containing white iron and gray iron are compared, the resulting structure provides support for the idea that some type of residual graphite serves as graphite nuclei during solidification (Ref 9). Further, additions of graphite to the melt enhance nucleation because the eutectic cell count is increased while the chilling tendency decreases (Ref 10).
Undercooling (ΔT ), ⬚C
Solidification of Eutectic Alloys—Cast Iron / 319 A rather wide variety of compounds have been claimed to serve as nuclei for FG cast iron, including oxides (for example, silicon dioxide) or silicates, sulfides, nitrides (boron nitride), carbides (for example, Al4C3), and intermetallic compounds (Ref 9). In addition, a number of metals, such as sodium, potassium, calcium, strontium, barium, yttrium, and the lanthanides, act as inoculants in FG iron and can therefore be considered to play a significant role in the heterogeneous nucleation of the austenite-FG eutectic (Ref 9, 11–13). High silicon concentrations in the melt can also contribute to the heterogeneous nucleation of graphite. The heterogeneous nucleation of FG iron occurs mainly on silicon dioxide particles (Ref 14). This assertion is based on experimental results as well as on the following theoretical considerations. If the oxygen dissolved in the iron is in equilibrium with carbon in an iron-acid lining-slag-atmosphere system, Eq 1 can be written: ðSiO2 Þ þ 2½C Ð ½Si þ 2fCOg
metals that have been proved to act as nuclei are as follows: Group I
Group II
Group III
NaHC2 KHC2
CaC2 SrC2 BaC2
YC2 LaC2
These carbides have a number of properties that support the hypothesis of their role as nuclei: Due to the strong ionic bond between carbon
and the metal, they will not dissolve in the
melt but will precipitate as insoluble particles. They are regions of concentration in the melt as compared with the bulk composition of the liquid and can serve as heterogeneous nuclei. They have carbon planes in their structures that can serve as a basis for the epitaxial growth of graphite from the melt (Fig. 7). A schematic of the mechanism of inoculation by metals forming saltlike carbides is shown in Fig. 8. A metal from the I, II, or III group in the periodic table is added to the melt and reacts with carbon to form a saltlike carbide, MzC.
(Eq 1)
The equilibrium temperature for Eq 1 can be calculated as: log
½Si ½C2
¼
27; 486 þ 15:47 T
(Eq 2)
where T is the temperature in degrees K. Below the equilibrium temperature, solid silica is formed by precipitation deoxidation; above the equilibrium temperature, silica is decomposed. Silica serves as nuclei for graphite growth. The formation of silica particles in turn requires nucleation, which can be promoted by inoculation. Therefore, the inoculation of FG iron can be described as a deoxidation process with the following requirements:
Fig. 7
Epitaxial growth of graphite on a CaC2 crystal. Source: Ref 11
Fig. 8
Schematic of the inoculation of cast iron by metals to form salt-like carbides. Source: Ref 11
Enough oxygen must be available in the
melt.
There must be the addition of elements with
high affinity to oxygen, such as calcium, aluminum, and barium, to form nuclei for silicon dioxide precipitation (heterogeneous catalysis). Inoculation temperature must be a maximum of 50 C (90 F) above the equilibrium temperature to avoid dissolution of the silicon dioxide particles. The cooling rate after inoculation must be low enough to allow sufficient time for silicon dioxide formation, but high enough to prevent coalescence and segregation of silicon dioxide. A different theory of heterogeneous nucleation in FG iron has also been proposed (Ref 11). Based on the study of the influence of the inoculation of pure iron-carbon and Fe-C-Si alloys with pure metals, it was concluded that the addition of elements capable of forming saltlike carbides increases the number of eutectic cells and therefore the number of nuclei in the cast iron. Saltlike carbides are carbides containing the ion C2 2 . Typical saltlike carbides of
320 / Principles of Solidification Graphite then grows heterogeneously from the melt on the carbon planes of the carbide. Nucleation of the Austenite-Spheroidal Graphite Eutectic. Because most of the inoculants used in FG iron treatment are also effective for SG irons, it is reasonable to assume that similar mechanisms and substrates are active in both types of irons. Extensive transmission electron microscopy and scanning electron microscopy studies have been done to identify the composition of graphite nuclei in SG irons, with a rather wide variety of results. Numerous types of compounds have been found in the middle of graphite spheroids; therefore, it was hypothesized that they could act as nuclei. Some of these compounds are as follows: 3MgO2SiO22H2O (cristobalite) (Ref 15), xMgOyAl2O3zSiO2, xMgOySiO2, xMgOySiO2zMgS (Ref 16), MgS (Ref 17), Te + Mn + S (Ref 18), and lanthanide sulfides (Ref 19). A rather extensive work on what appeared to be substrates for spheroidal graphite in chilled ductile iron concluded that the substrates are duplex sulfide-oxide inclusions with a diameter of 1 mm (40 min.). The core is probably made of calcium-magnesium or Ca-Mg-Sr sulfides, while the outer shell is made of complex MgAl-Si-Ti oxides with a spinel structure (Ref 20). The orientation relationships were established as follows. For the nucleus core/nucleus shell: ð110Þsulfide k ð111Þoxide ½1010sulfide k ð211Þoxide For the nucleus shell/graphite: ð111Þoxide k ð0001Þ graphite ½110oxide k ½1010graphite The x-ray diffraction data showed that the first few graphite layers, adjacent to the (111) oxide, ˚, had a dilated lattice (0.264 nm, or 2.64 A ˚ ). It was suginstead of 0.246 nm, or 2.46 A gested that the spacing within the graphite layers decreases away from the oxide until unconstrained spacing is reached. Dislocations were frequently observed in the matrix, and it was suggested that these were generated to relieve some of the elastic strain in the graphite layers adjacent to the oxide. These observations allow for a theory of nucleation of spheroidal graphite similar to the catalytic theory of nucleation of flake graphite on silicon dioxide, as discussed previously. It can be assumed that the nucleation process begins with the formation of complex sulfides that serve as nuclei for complex oxides, which in turn serve as nuclei for spheroidal graphite. Recent experimentations (Ref 21) suggest that even adding small amounts of sulfur (less than 0.0059%) to hypereutectic postinoculated ductile irons results in higher nodule counts than in comparable melts without any later sulfur additions. This strengthens the theory that sulfur plays an important role in the formation of substrates for the graphite. Structure of Iron Carbide in the AusteniteIron Carbide Eutectic. The iron carbide
(cementite) consists of an orthorhombic unit cell with 12 iron atoms and 4 carbon atoms per cell and therefore has a carbon content of 6.7 wt%. Its density is 7.6 g/cm3. Nucleation of the Austenite-Iron Carbide Eutectic. Very little information is available on the nature of the nuclei of the white eutectic. Nevertheless, it is accepted that the nucleation of Fe3C occurs at lower undercooling than that of graphite (Ref 8). This is supported by the slopes of the solubility lines in the phase diagram, which shows that the supersaturation of cementite increases faster than that of graphite, resulting in an increased probability for the nucleation of cementite as the temperature is lowered. Cooling curve data also show that the solidification of white iron begins at a lower temperature than that of gray iron. There is some evidence that silicon dioxide and aluminum oxide can serve as substrates for the growth of the Fe-Fe3C eutectic and that the nature of the substrate can influence the morphology of the eutectic (Ref 22).
Nucleation of Primary Phases in Cast Iron Depending on the carbon equivalent (or saturation degree), the primary phase can be either austenite for hypoeutectic cast iron or graphite for hypereutectic cast iron. Primary Austenite. The formation and growth characteristics of austenite dendrites have received considerably less attention from investigators than eutectic cell count and morphology simply because dendrites are not readily discernible in the structure. Nevertheless, it has been demonstrated that inoculants that are effective in increasing the cell count (such as calcium-titanium, strontium, and 75% FeSi) have little effect on the solidification pattern of primary austenite dendrites (Ref 23). A number of elements, such as titanium, vanadium, and aluminum, were proved to increase the number of dendrites and their length when compared with the base iron. For titanium and vanadium, this was attributed to the formation of carbides, nitrides, and carbonitrides, which then act as substrates for austenite solidification. The influence of aluminum was not explained. Other elements, such as cerium and boron, although being instrumental in increasing the number and length of dendrites, were not considered to be nucleants for austenite. It was suggested that these elements restrict the growth of the eutectic cells, which subsequently results in larger undercooling for eutectic solidification. This means that more time (a larger temperature interval) is available for the nucleation and growth of austenite (Ref 23). Primary Graphite. It is reasonable to assume that nuclei that are active in the solidification of eutectic graphite also serve as nuclei for hypereutectic primary graphite. Nevertheless, it must be noted that inoculation does not seem to increase
significantly the number of eutectic cells in hypereutectic FG iron. On the other hand, inoculation is quite effective in hypereutectic SG iron.
Growth of Eutectic in Cast Iron Coupled Zone in Cast Iron The degree and type of eutectic growth that occurs in cast iron can be determined by using tools such as growth rate curves to locate coupled zone regions, isothermal time-temperature diagrams to gage susceptibility to carbide formation, and growth rate-composition plots to ascertain parameters that affect both directional and multidirectional solidification. In cast irons, the coupled growth region of the eutectic is asymmetric. It is possible to construct a theoretical coupled zone for gray iron from the condition of equal growth rate of the austenite (g) and graphite (Gr) phases (Ref 24). First, one must consider the growth rate curves for g and Gr in the iron-carbon system (Fig. 9). For FG iron, the growth rate of austenite, Rg, and that of graphite along the ½1010 direction, RGr ½1010, intersect; therefore, a coupled zone
Fig. 9
Suggested growth rate curves for austenite (g) and graphite (Gr) in the Fe-C-Si alloys. (a) Austenite-flake graphite eutectic. (b) Austenitespheroidal graphite eutectic. (c) Different graphite growth rates in lamellar growth. RGr ½1010 determines the lamellar habits of graphite flakes, while RGr [0001] gives the rate of thickening for the flake. Source: Ref 6
Solidification of Eutectic Alloys—Cast Iron / 321 Gr γ
ΔTGr⬘ ΔTγ⬘
γ
Gr
ΔTγ/Gr ΔTGr⬙
L
2
E ΔTγ/Gr ΔTγ ⬙
Growth rate, R
1
P Halo (shell)
ΔTγ /Gr
ΔTγ⬘
Undercooling, ΔT
Graphite leading phase
Q
ΔTGr⬘ ΔTGr⬙
R
ΔTγ ⬙
(a)
RGr⬙ = R g (Co line)
(b)
Fig. 10
Constructions of coupled growth line in Fe-C-Si alloys. (a) Austenite (g) and graphite (Gr) growth curves. (b) Construction of the equal growth rate curve Rg = RGr. Source: Ref 24
20 400 10 5 Au + Eu1 2
Temperature, ⬚C
Velocity (v), μm/s
Eu1
Gr + Eu1
Velocity (v), μin./s
4 ⫻ 103
D
L+A
1976
1080 L + A + Gr + C
1940
1060 L+A
N
1040
1904
F
1020 1000 Z 980 0.1
40
2012
I 0.5
1
2
1
1868 1832
L+A+C
K
Temperature, ⬚F
L 1100
100 R = 7K/mm
2048
1120
Gray eutectic
50
R
B
L + A + Gr
White eutectic Gray eutectic + γ- dendrite
500 200
1140
4 ⫻ 104
103
5 10 20 Time, s
50 100
(a) 0.5
1140 0.2
B
O
R
1120 1100
2120
1160
1155
1145
2102
Gr + Eu1
Au + Eu1
Ce + Eu2
Au + Eu2
2012 N
D
Liq 1080
1976 1940
1060 Liq + A 1040
1904
L Liq + A + Gr + C
2084
1140
1020
Temperature, ⬚F
Temperature, ⬚C
Eu1 1150
2048 Liq + A + Gr Temperature, ⬚F
(a)
3.75 4.0 4.25 4.5 4.75 Concentration (C∞), wt% carbon
4 5.0
Temperature, ⬚C
0.1 3.5
1868 F
1000 980 0.1
I
Z 0.5
1
K 2
1832 Liq + A + C
5 10 20
50 100
Time, s (b)
1135
1130 3.75
Eu2
4.0 4.25 4.5 4.75 Concentration (C∞), wt% carbon
Fig. 12 2066 5.0 5.25
Time-temperature transformation diagrams for isothermal solidification of eutectic cast iron. (a) Flake graphite cast iron. (b) Spheroidal graphite cast iron. L, liquid; A, austenite; C, cementite; Gr, graphite. Source: Ref 8
(b)
Fig. 11
and are equal (DTg/Gr); at points 1 and 2 on the curves, the growth rates are the same, but the undercooling differs. Figure 10(b) shows the phase diagram in the vicinity of the eutectic point, with DTg0 and DTg00 drawn for growth conditions represented by the points of Fig. 10 (a). Point “Q” in Fig. 10(b) corresponds to DTg/Gr, where Rg = RGr. At the left of the RGr = Rg line, primary g dendrites will form, while at the right of this line, graphite will be the leading phase. Coupled eutectic growth can only occur for combinations of undercoolings and compositions on the dotted line RGr = Rg. The transition from a fully eutectic to a eutectic plus dendrite structure in pure iron-graphite alloys of eutectic composition has been determined, and the g-iron and graphite-iron eutectic boundaries have been calculated (Fig. 11).
Coupled zone of directionally grown ironcarbon eutectic. (a) Growth rate-composition diagram showing calculated curves and experimental points. (b) Temperature-composition diagram with superimposed phase diagrams, stable (Eu1) and metastable (Eu2). Au, austenite dendrites; Gr, primary graphite plates; Ce, primary cementite plates. Source: Ref 25
can be constructed. For SG iron, where the predominant growth direction is along [0001], the two rates, Rg and RGr[0001], do not intersect, which means that coupled growth is impossible. Figure 10(a) shows the point at which the graphite and g growth curves for gray iron cross
Isothermal Solidification Solidification studies are usually performed athermally because of the high rate of the liquid-solid transformation. Nevertheless, useful information can be extracted from the isothermal time-temperature-transformation diagrams shown in Fig. 12. It is apparent that SG iron is more susceptible to carbide formation than FG iron. Graphite precipitates earlier in SG iron than in FG iron at all undercoolings, although the time interval for complete gray solidification is smaller in FG irons.
Growth in Directional Solidification It is quite obvious from the previous discussion that undercooling and composition must be carefully selected to achieve coupled growth of the eutectic in cast iron. The g-FG eutectic is a coupled irregular eutectic of the faceted (Gr)/nonfaceted (g) type. The g-SG eutectic is a divorced eutectic. The basic parameters affecting the morphology of the eutectic are the G/R ratio and composition.
322 / Principles of Solidification As shown in Fig. 13, it is possible to achieve a variety of structures in cast iron when varying G/R and/or the level of impurities (for example, cerium). The gray eutectic can solidify with a planar interface. The theoretical relationship l2R = constant (Jackson and Hunt), where l is the spacing, is not obeyed in the growth of the g-FG eutectic. Different experimental results are compared with the theoretical behavior in Fig. 14. The departure of experimental values from the theoretical line is a clear indication of irregular rather than regular cooperative growth in FG cast iron. For G/R ratios selected to produce a planar interface during the directional solidification of FG iron, it has been measured that the growth constant in the R = m DT2 equation is m = 8.7 108 m s1 K2 (Ref 25). In the case of SG iron, even for hypereutectic irons, the graphite spheroids do not grow in independent austenite envelopes but rather are associated with austenite dendrites (Fig. 15). The eutectic austenite is dendritic and cannot be distinguished from the primary austenite. Spheroidal graphite precipitates directly from the melt, becomes enveloped in an austenite shell that is very soon incorporated into the dendrites, and then grows together with the dendrites (Ref 27, 28). Directional solidification experiments (Ref 29) demonstrated that nucleation of SG in eutectic and hypereutectic irons occurs over the entire eutectic solidification temperature interval.
Growth in Multidirectional Solidification Flake (Lamellar) Graphite Eutectic. The austenite-FG eutectic solidifies with the formation of eutectic colonies (cells) that are more or less spherical in shape. It is generally thought that each eutectic cell is the product of a nucleation event. The eutectic cell is made of interconnected graphite plates surrounded by austenite. The degree of ramification of graphite within the cell depends on undercooling, with higher undercooling resulting in more graphite branching (Fig. 16). The leading phase during the eutectic growth is the graphite. Graphite spacing is determined by the same parameters as for regular eutectics, with branching occurring as a response to interface instability. In turn, interface instability is determined by localized changes in composition, convection currents, crystallographic orientation different from the heat extraction direction, and a change in temperature gradient. Figure 17 shows a scanning electron microscopy (SEM) micrograph of graphite eutectic cell and prior austenitic dendrite structures in gray cast iron. The variations in graphite structures have been classified, together with the length of the flakes, by standards that have been used for many years. Flake graphite in gray cast iron can be designated as: Type
A: uniform orientation
distribution,
random
Type B: rosette grouping, random orientation Type C: superimposed flake sizes, random
orientation
Type D: interdendritic segregation, random
orientation
Type E: interdendritic segregation, preferred
orientation The formation of the eutectic flake graphite (types A, B, C, and D) is greatly influenced by the amount by which the iron melt cools below the equilibrium temperature for the austenite-graphite eutectic before appreciable solidification occurs. Gray iron with type A graphite undergoes only small amounts of undercooling, whereas that with type D graphite undercools significantly below this equilibrium temperature. The undercooling that occurs with type B graphite is intermediate between the two, producing fine graphite flakes, like type D, in the center of the eutectic cells or rosette and a coarser type, like type A, at the outer cell boundaries. Type E graphite occurs in strongly hypoeutectic gray irons with carbon equivalents well below 4.3%. Figure 18(a) shows an SEM image of flake graphite and respective three-dimensional reconstruction (Fig. 18b) using focused ion beam (FIB) nanotomography. In the represented volume, it is possible to see that almost all graphite particles are connected, and in the cropped-out magnified regions (Fig. 18c) some inclusions can be found. A chemical analysis (energy-dispersive x-ray) of these inclusions determined them to be MnS precipitates (Ref 31). Calculations using data from continuous cooling experiments given in Ref 32 allowed estimation of the growth constant m in the range of 7.25 108 to 9.5 108 m s1 K2, which is of the same magnitude as the values given previously for the case of directional solidification. The solidification of off-eutectic FG iron will start with the nucleation and growth of a primary phase, which is austenite dendrites for hypoeutectic iron or primary graphite plates for hypereutectic iron. It is not unreasonable to assume that the primary phases may play a role in the nucleation of the eutectic, with further growth of the eutectic occurring, as experimental evidence shows, on the primary phases. Spheroidal Graphite Eutectic. Growth of the austenite-SG eutectic is more complicated and less understood than that of the g-FG eutectic, although a good number of theories have been proposed. As discussed previously in this section, the g-SG eutectic is a divorced eutectic. It has been rather widely accepted that the growth of this eutectic begins with nucleation and the growth of graphite in the liquid, followed by early encapsulation of these graphite spheroids in austenite shells (envelopes). A schematic of the process is shown in Fig. 19. Graphite nucleation and growth deplete the melt of carbon in the vicinity of the graphite; this creates conditions for austenite nucleation
and growth around the graphite spheroid. Once the austenite shell is formed, further growth of graphite can occur only by solid diffusion of carbon from the liquid through the austenite. Calculations of the diffusion-controlled growth of graphite through the austenite shell were originally made based on Zener’s growth equation for an isolated spherical particle in a matrix of low supersaturation (Ref 33). Equation 3 was derived: RGr ¼
VmGr g rg X g=L X g=Gr D Vmg c rGr ðrg rGr Þ X Gr X g=Gr
(Eq 3)
where RGr is rate of growth of graphite; VmGI and Vmg are the molar volumes of graphite and austenite, respectively; Dgc is the diffusivity of carbon in austenite; rGr and rg are the radii of graphite and austenite, respectively; Xg/L and Xg/Gr are the molar fractions of austenite at the austenite/liquid and austenite/graphite boundaries, respectively; and XGr is the molar fraction of carbon in graphite. It has also been shown that rg = 2.4 rGr. A slightly different approach, based on a steady-state diffusion model of carbon through the austenite shell (Ref 34), also allowed derivations for RGr and rg. However, recent research has shown that the solidification mechanism of SG iron is more complicated and that austenite dendrites play a significant role in eutectic solidification (Ref 27, 28). The eutectic austenite is dendritic and can scarcely be distinguished from primary austenite dendrites. The sequence of solidification is as follows: At the eutectic temperature, austenite den
drites and graphite spheroids nucleate independently in the liquid. Limited growth of spheroidal graphite occurs in contact with the liquid. Flotation or convection then determines the collision of spheroidal graphite with the austenite dendrites. Graphite encapsulation in austenite can occur before or immediately after the contact between graphite and austenite dendrites. Further growth of graphite occurs by carbon diffusion through the austenite shell.
This sequence is illustrated schematically in Fig. 20. Graphite shape in SG iron is not perfectly spheroidal. Depending on cooling rate and chemical composition, significant deviations from the true spheroidal shape can be obtained (various nodularities) (Fig. 21). It remains now to explain the reasons for the change of the habits of graphite from lamellar to spheroidal. Many theories have been proposed (see, for example, review works in Ref 6 and 9), but because of space limitations, only a few are discussed in this section. Many theories capitalize on the observation that the graphite/liquid surface energy is higher
Solidification of Eutectic Alloys—Cast Iron / 323 0.4
Spacing (λ), mm
0.2
C
0.1 0.06 0.04 D 0.02
0.01 10−4
B A
10−3 Growth rate (R ), mm · s−1
0.01
l-R relationships in flake graphite cast iron. A, Lakeland l = 3.8 105 R0.5 cm; B, Nieswaag and Zuithoff l = 0.56 105 R0.78 cm (0.004% S); C, Nieswaag and Zuithoff l = 7.1 105 R0.57 cm (>0.02% S); D, Jackson and Hunt l = 1.15 105 R0.5 cm (theoretical). Source: Ref 6
Fig. 14
Fig. 13
Schematic showing the influence of composition (% cerium), G, and R over the eutectic morphology of Fe-C-Si alloys. SG, spheroidal graphite; CG, compacted graphite; FG, flake graphite. Source: Ref 26
Fig. 15
Microstructure of the solid/liquid interface for a directionally solidified hypereutectic cerium-treated spheroidal graphite cast iron. Source: Ref 27
in SG iron than in FG iron. These theories explain SG formation by either simply implying that a sphere will have less free surface energy than a lamella with the same volume above a certain critical interface energy (Ref 35) or by suggesting that the high interface energy will curve the growing crystal in order to decrease
the energy/volume ratio, resulting in spheroidal rather than lamellar graphite (Ref 36). The defect growth of graphite theory explains SG formation based on the possible growth mechanisms of faceted crystals such as graphite (Ref 37). Three growth mechanisms are considered: two-dimensional nucleation
(Fig. 22a), step of a defect (twisted) boundary (Fig. 22b), and screw dislocation. The first two mechanisms are governed by exponential laws and apply to the ð1010Þ surface, but the third is governed by a parabolic law and applies to the (0001) surface of the graphite crystal.
324 / Principles of Solidification
Fig. 16
Schematic of solidification of flake graphite. (a) Typical eutectic colonies (cells). (b) Growth sequence for a eutectic colony. g, austenite; Gr, graphite
Fig. 17
Fig. 18
Fig. 19
SEM micrograph showing graphite, eutectic cell, and prior dendrite structure in gray cast iron. Original magnification: 200. Source: Ref 30, courtesy of Gary F. Ruff
(a) SEM micrograph of flake graphite, deep etched. (b) Reconstructed three-dimensional structure. (c) Cropped-out magnified region with inclusions and flake graphite. Source: Ref 31
Isothermal growth of a graphite spheroid within an austenite shell and growth of the shell with a smooth interface. (a) Growth of spheroidal graphite in contact with melt. (b) Envelopment by austenite. (c) Growth of spheroidal graphite within the austenite shell. Source: Ref 24
Solidification of Eutectic Alloys—Cast Iron / 325 When weak, reactive impurities such as sulfur are present in the melt, a contaminated environment occurs. These elements change the edge energy of steps, resulting in a relative position change of the growth rates involved, as shown in Fig. 23(a). The curve for growth on the step of a defect boundary, Rstep, is at a lower undercooling than those for growth by two-dimensional nucleation, R2D, or by screw dislocation, Rscrew. In a pure environment such as an Fe-C-Si alloy with no sulfur contamination, the growth rate curves are displaced to higher undercooling (Fig. 23b). In a melt of sufficient purity, or when increasing cooling rate, the higher degree of undercooling may allow growth with Rscrew so that graphite spheroids can form. This has been achieved experimentally for pure nickelcarbon alloys by increasing the cooling rate of the melt, or for ultrapure iron-carbon alloys by cooling slowly in a vacuum. In an environment with reactive impurities (for example, magnesium), the impurity will react with the surface, and the growth at a step of a twist boundary will be neutralized. Only the curves for R2D and Rscrew are left, and they are displaced to greater undercoolings (Fig. 23c). Another theory relates graphite shape in cast iron with undercooling (kinetic plus constitutional) during solidification (Ref 6). As shown in Fig. 24, each graphite form has its own temperature for growth, which is achieved by a specific cooling rate and composition. Steps on surfaces can change graphite morphology from plate to rod. With an increase in undercooling, pyramidal instabilities will occur on the faces of the graphite crystal. At undercoolings of 29 to 35 C (50 to 65 F), instabilities occur on the ð1011Þ faces of the pyramid, and it is suggested that graphite spheroids form at these undercoolings. Finally, at large undercoolings of 40 C (70 F), the growth form noted is a pyramidal one, bounded by ð1011Þ faces. These pyramidal crystals are part of the series of imperfect forms observed particularly in thick-wall SG iron castings. Imperfect SG shapes, such as those in Fig. 21(b) and (c), are due to incorrect kinetic undercooling and an imbalance of impurities influencing constitutional undercooling. An insufficient amount of impurities influencing kinetic undercooling (for example, magnesium) will favor the formation of intermediate graphite shapes, such as compacted/vermicular, while an insufficient amount of impurities influencing constitutional undercooling (for example, lead) will perturb the surface of the spheroid and promote the formation of protuberances. Surface Adsorption Theory. This somewhat older theory of graphite growth postulates that the change from lamellar to spheroidal graphite occurs because of the change in the ratio between growth on the ð1010Þ face and growth on the (0001) face of graphite (Ref 38). For
equilibrium conditions, the Gibbs-Curie-Wulf law states that the crystalline phase with the higher interface energy has a slow rate of growth in the normal direction. Bravais’s rule stipulates that the growth rate in the direction normal to a plane is inversely proportional to the density of atoms located on the plane. Accordingly, it follows that under equilibrium conditions, the crystallographic plane with the highest density of atoms has the lowest interface energy and the minimum growth rate in a direction perpendicular to the plane. Nevertheless, under the nonequilibrium conditions prevailing during the solidification of cast iron, kinetic considerations become important. Assuming growth by two-dimensional nucleation, the highest rate of growth will be experienced by the face with the higher density of atoms, where the probability for nucleation is higher. Therefore, in a pure environment, the highest growth rate will be in the (0001) direction of the graphite crystal (Fig. 25), resulting in the formation of unbranched single crystals (coral graphite). In a contaminated environment, surface-active elements such as sulfur or oxygen are absorbed on the high-energy plane ð1010Þ, which has fewer satisfied bonds. Subsequently, the ð1010Þ plane face achieves a lower surface energy than the (0001) face, and growth becomes predominant in the ð1010Þ direction, resulting in lamellar (plate) graphite. Finally, the reactive impurities (such as magnesium, cerium, and lanthanum) in an environment scavenge the melt of surface-active elements (sulfur, diatomic oxygen, lead, antimony, titanium, and so on), after which they also block growth on the ð1010Þ prism face. A polycrystalline SG results. Compacted/Vermicular Graphite Eutectic. The sequence of growth of compacted/vermicular graphite during the eutectic transformation is shown schematically in Fig. 26, based on experimental data from Ref 39 on rapidly quenched samples from successive stages during the solidification process. It can be seen that at the beginning, graphite precipitates as spheroids, which then degenerate during growth and subsequently develop into compacted graphite. Compacted graphite develops as interconnected segments within an austenitic matrix. Typical compacted graphite is shown in Fig. 27(a). The reconstructed three-
(a) 10 μm
(b) 10 μm
(c) 1 μm
Fig. 21
Fig. 20
Schematic illustrating the progression of growth in austenite-spheroidal graphite eutectic
SEM micrographs of deep-etched spheroidal graphite samples showing a fractured graphite spheroid (a). Nodularity decreases from (a) through (c).
326 / Principles of Solidification dimensional form (Fig. 27b) was obtained using FIB nanotomography (Ref 31). It was proposed that the growth of compacted graphite occurs by the twin/tilt of boundaries (Ref 40). The initiation of a twin/tilt is related to the unstable growth of the ð1010Þ interface, which can be induced by the presence of some reactive elements. The formation of a ð1012Þ twin/tilt plane is shown in Fig. 28. It was further hypothesized that when insufficient reactive element is present, the tilt orientation of twin boundaries can alternate, resulting in typical compacted graphite (Fig. 29a). On the other hand, when there is enough impurity in the melt, the tilt orientation is singular, and spheroidal graphite may occur (Fig. 29b). Based on this growth mechanism, compacted graphite can be considered to grow either by transition from flake graphite (Fig. 30a) or by degeneration of spheroidal graphite (Fig. 30b). The degree of interconnection of the compacted graphite within the eutectic cell is also apparent from Fig. 30. Based on cooling curve configuration, it is reasonable to assume that the growth rate of compacted graphite is similar to that of spheroidal graphite, which is not surprising because the growth direction is supposedly in the [0001] direction of the graphite crystal. Austenite-Iron Carbide Eutectic. Growth of the austenite-iron carbide (ledeburite) eutectic begins with the development of a cementite plate on which an austenite dendrite nucleates and grows (Fig. 31). This destabilizes the Fe3C, which then grows through the austenite. As a result, two types of eutectic structure develop: a lamellar eutectic with Fe3C as the
leading phase in the edgewise direction, a, and a rodlike eutectic in the sidewise direction, c. Cooling rate has a significant influence on the morphology of the g-Fe3C eutectic. At moderate undercoolings, ledeburite structure is expected. High cooling rates, as obtained in quenching experiments, produce a degenerated eutectic structure dominated by Fe3C plates. A coarse mixture of Fe3C and g fills the spaces
Suggested DT-R correlation for ð1010Þ and (0001) crystal faces of graphite growing in various environments. (a) Contaminated environment (for example, sulfur-containing Fe-C-Si alloy). (b) Pure environment (for example, Fe-C-Si alloy). (c) Environment with reactive impurities (for example, magnesium-containing Fe-C-Si alloy). Three growth mechanisms are discussed: A, on the step of the defect boundary (Rstep); B, two-dimensional nucleation (R2D); C, screw dislocation (Rscrew). Source: Ref 37
between the Fe3C plates. This structure obviously does not result from cooperative growth (Ref 8). A platelike carbide structure associated with equiaxed eutectic grains can be obtained by increasing the undercooling by superheating, by decreasing the silicon content, or by adding chromium or magnesium. The gray or white solidification mode of cast iron apparently depends on the relative nucleation probability and the growth rates of the graphite and Fe3C phases. In turn, this will be a function of the cooling rate and chemistry of the alloy. As shown in Fig. 32, only the graphite eutectic can nucleate and grow between the eutectic temperature for the gray eutectic (1153 C, or 2107 F) and that for the white eutectic (1148 C, or 2098 F). Below 1148 C (2098 F), both eutectics can occur. The growth rate of the g-Fe3C eutectic rapidly exceeds that of the g-Gr eutectic, and at a temperature of approximately 1140 C (2085 F), the gray-towhite transition occurs (Ref 6, 43).
Fig. 23
Growth of graphite in the ð1010Þ direction. (a) Growth by two-dimensional nucleation on ð1010Þ faces illustrates that steps on (0001) surfaces will advance only as far as bounding crystal edges. (b) Growth from step to twist boundary illustrates that the (1010) faces grow by nucleation of planes at the step. Source: Ref 37
Fig. 22
Fig. 24
Correlation among the different types of instability observed in graphite growth and growth morphologies with increasing undercooling, DT. (a) DT = 4 C (7 F). (b) DT = 9 C (16 F). (c) DT = 30 C (54 F). (d) DT = 40 C (72 F). Source: Ref 6
Solidification of Eutectic Alloys—Cast Iron / 327
Cooling Curve Analysis Cooling curves have been extensively used in the experimental study and production control of cast irons. The history of solidification can be traced on the cooling curve. A detailed description of the principles involved in this analysis and its application in cast iron technology can be found in other articles in this Volume. REFERENCES 1. S. Steeb and U. Maier, Structure of Molten Fe-C Alloys by Means of X-Ray and Neutron Wide Angle Diffraction as well as Sound Velocity Measurements, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 1 2. W. Krieger and H. Trenkler, Arch. Eisenhu¨ttenwes., Vol 42 (No. 3), 1971, p 175 3. A.A. Vertman et al., Liteinoe Proizvod., No. 10, 1964 4. J. Luo, Q. Zhai, and P. Zhau, Microstructures of Liquid Pure Iron and Fe-C Alloy near the Melting Point, Acta Metall. Sin. (Chinese Lett.), Vol 39, 2003, p 5–8 5. L.S. Darken, “Equilibria in Liquid Iron with Carbon and Silicon,” Technical Publication 1163, American Institute of Mining, Metallurgical, and Petroleum Engineers, 1940, p 1 6. I. Minkoff, The Physical Metallurgy of Cast Iron, John Wiley & Sons, 1983 7. W. Patterson and D. Ammann, Solidification of Flake Iron-Graphite Eutectic in Gray Iron, Giesserei, Jan 1959, p 1247
8. H.D. Merchant, Solidification of Cast Iron—A Review of Literature, Recent Research on Cast Iron, H.D. Merchant, Ed., Gordon and Breach, 1968, p 1 9. J.F. Wallace, Effects of Minor Elements on the Structure of Cast Iron, Trans. AFS, Vol 83, 1975, p 363 10. S.L. Liu, C.R. Loper, Jr., and S. Shirvani, Inoculation of Gray Cast Irons with Carbonaceous Materials, Trans. AFS, Vol 93, 1985, p 501 11. B. Lux, Nucleation of Graphite in Fe-C-Si Alloys, Recent Research on Cast Iron, H.D. Merchant, Ed., Gordon and Breach, 1968, p 241 12. J.V. Dawson and S. Maitra, Recent Research on the Inoculation of Cast Iron, Br. Foundryman, April 1967, p 117 13. D.M. Stefanescu, Inoculation of Gray Iron with Sodium and Barium, Giesserei-Prax., No. 24, 1972, p 429 14. W. Weiss, The Importance of Deoxidation in the Crystallization of Cast Iron, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 69 15. M.B. Zeedijk, Identification of the Nuclei in Graphite Spheroids by Electron Microscopy, J. Iron Steel Inst., July 1965, p 737 16. P. Poyet and J. Ponchon, Contributions to the Study of Graphite Nucleation in Cast Iron for Ingot Molds, Fonderia, No. 277, 1969, p 183 17. J.C. Mercier et al., Inclusions in Graphite Spheroids, Fonderia, No. 277, 1969, p 191 18. H. Nieswaag and A.J. Zuithoff, “The Occurrence of Nodules in Cast Iron Containing Small Amounts of Tellurium,” Paper
19. 20.
21.
22. 23.
24.
25.
26.
27. 28.
29.
Fig. 25
Schematic of the change in the growth rate of graphite due to the absorption of foreign atoms in spheroidal graphite eutectic. Three variations of an Fe-C-Si cast iron are as follows. (a) With nodularizer added as reactive impurity environment. (b) Pure environment. (c) Contaminated environment in which surface-active elements such as oxygen and sulfur are absorbed into system. For (a) and (b), density in the basal plane, VB, is greater than the density in the prism face, VP, and either branch polycrystalline or unbranched single crystals result. For (c), VB < VP. Sulfur adsorption makes prism faces the most densely packed, and graphite flakes are subsequently formed.
presented at the 37th International Foundry Congress (Brighton, England), Sept 1970 R.J. Warrick, Spheroidal Graphite Nuclei in Rare Earth and Magnesium Inoculated Irons, Trans. AFS, Vol 76, 1968, p 722 M.M. Jacobs, T.J. Law, D.A. Melford, and M.J. Stowell, Basic Processes Controlling the Nucleation of Graphite Nodules in Chill Cast Iron, Met. Technol., Vol 1, Part II, No. 1974, p 490 O.M. Suarez, R.D. Kendrick, and C.R. Loper, Jr., Late Sulfur Additions to Postinoculated High-CE Ductile Iron Melts, Trans. Am. Foundrymens Soc., Vol 108, 2000, p 63–69 J. Rickard and I.C.H. Hughes, Eutectic Structure in White Cast Iron, BCIRA J., Vol 9 (No. 1), 1961, p 11 G.F. Ruff and J.F. Wallace, Control of Graphite Structure and Its Effect on Mechanical Properties of Gray Iron, Trans. AFS, Vol 84, 1976, p 705 B. Lux, F. Mollard, and I. Minkoff, On the Formation of Envelopes Around Graphite in Cast Iron, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 371 H. Jones and W. Kurz, Growth Temperatures and the Limits of Coupled Growth in Unidirectional Solidification of Fe-C Eutectic Alloys, Metall. Trans. A, Vol 11, 1980, p 1265 D.M. Stefanescu, P.A. Curreri, and M.R. Fiske, Microstructural Variations Induced by Gravity Level During Directional Solidification of Near-Eutectic IronCarbon Type Alloys, Metall. Trans. A, Vol 17, 1986, p 1121 S.K. Biswal, D.K. Bandyopadhyay, and D.M. Stefanescu, unpublished research A. Rickert and S. Engler, Solidification Morphology of Cast Irons, The Physical Metallurgy of Cast Iron, Vol 34, H. Fredriksson and M. Hillert, Ed., Proceedings of the Materials Research Society, North Holland, 1985, p 165 D.K. Banerjee and D.M. Stefanescu, Structural Transitions and Solidification Kinetics of SG Cast Iron during Directional Solidification Experiments, Trans. Am.
Fig. 26
Schematic of the sequence of development of compacted/vermicular graphite: (a) small spheroids; (b) and (c), some spheroids have tails; (d) compacted graphite plus spheroidal graphite; and (e) compacted graphite
328 / Principles of Solidification
Fig. 29
Change of direction of the tilt of the boundary due to the amount of reactive impurity. (a) Insufficient spheroidization. (b) Sufficient spheroidization. Source: Ref 40
(a)
(b)
Fig. 27
(a) SEM micrograph of deep-etched compacted graphite iron sample. (b) Reconstructed three-dimensional form. Source: Ref 31
Fig. 28
Diagram of sequences (a) through (c) involved in the formation of a twin/tilt plane for graphite growing in the [1010] direction. Source: Ref 40
Foundrymens Soc., Vol 99, 1991, p 747– 759 30. D.M. Stefanescu, Cast Iron, Casting, Vol 15, ASM Handbook, ASM International, 1988 31. A. Velichko, C. Holzapfel, and F. Mu¨cklich, 3-D Characterization of Graphite Morphologies in Cast Iron, Adv. Eng. Mater., Vol 9 (No.1–2), 2007, p 39–45 32. H. Fredriksson and S.E. Wetterfall, A Study of Transition from Undercooled to Flake Graphite in Cast Iron, The
Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 227 33. S.E. Wetterfall, H. Fredriksson, and M. Hillert, Solidification Process of Nodular Cast Iron, J. Iron Steel Inst., May 1972, p 323 34. K.C. Su, I. Ohnaka, I. Yamauchi, and T. Fukusako, Computer Simulation of Solidification of Nodular Cast Iron, The Physical Metallurgy of Cast Iron, Vol 34, H. Fredriksson and M. Hillert, Ed., Proceedings of the Materials Research Society, North Holland, 1985, p 181
35. H. Geilenberg, A Critical Review of the Crystallization of Graphite from Metallic Solutions After the “Surface Tension Theory,” Recent Research on Cast Iron, H.D. Merchant, Ed., Gordon and Breach, 1968, p 195 36. J.P. Sadocha and J.E. Gruzleski, The Mechanism of Graphite Spheroid Formation in Pure Fe-C-Si Alloys, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 443 37. I. Minkoff and B. Lux, Graphite Growth from Metallic Solution, The Metallurgy of Cast Iron, B. Lux et al., Ed., Georgi Publishing, 1975, p 473 38. K. Herfurth, Investigations into the Influence of Various Additions on the Surface Tension of Liquid Cast Iron with the Aim of Finding Relationships Between the Surface Tension and the Occurrence of Various Forms of Graphite, Freib. Forschungsh., Vol 105, 1965, p 267 39. E.N. Pan, K. Ogi, and C.R. Loper, Jr., Analysis of the Solidification Process of Compacted/Vermicular Graphite Cast Iron, Trans. AFS, Vol 90, 1982, p 509 40. P. Zhu, R. Sha, and Y. Li, Effect of Twin/ Tilt on the Growth of Graphite, The Physical Metallurgy of Cast Iron, Vol 34, H. Fredriksson and M. Hillert, Ed., Proceedings of the Materials Research Society, North Holland, 1985, p 3 41. X. Den, P. Zhu, and Q. Liu, Structure and Formation of Vermicular Graphite, The Physical Metallurgy of Cast Iron, Vol 34, H. Fredriksson and M. Hillert, Ed., Proceedings of the Materials Research Society, North Holland, 1985, p 141 42. M. Hillert and H. Steinhauser, The Structure of White Iron Eutectic, Jernkontorets Ann., Vol 144, 1960, p 520 43. M. Hillert and V.V. Subba Rao, “The Solidification of Metals,” Publication 110, The Iron and Steel Institute, 1968
Solidification of Eutectic Alloys—Cast Iron / 329
Fig. 30
Transition from (a) flake to compacted graphite and from (b) spheroidal to compacted graphite based on the twin/tilt of boundaries growth mechanism. Source: Ref 41
1170
1153 ⬚C
Spontaneous change from white to gray mode of solidification
B
1150
2120
2102
1148 ⬚C 2084
1140 Spontaneous change from gray to white mode of solidification
A
1130 0.01
Temperature, ⬚F
Temperature, ⬚C
1160
2066 10−3
10−4
10−5
10−6
Growth rate (R ), cm · s−1
Fig. 32
Gray-to-white eutectic transition as a function of undercooling and growth rate. A, white structure; B, gray structure. Source: Ref 43
Fig. 31
Schematic illustrating growth of ledeburite (austenite-iron carbide) eutectic. (a) Lamellar eutectic with cementite as the leading phase in the
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 330-337 DOI: 10.1361/asmhba0005214
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Peritectic Solidification THE TERM PERITECTIC, in the science of heterogeneous equilibria, may be used to define all reactions in which two or more phases (gas, liquid, solid) react at a defined temperature, Tp, to form a new phase that is stable below Tp. Usually, the term peritectic refers to reactions in which a liquid phase reacts with at least one solid phase to form one new solid phase. This reaction can be written as a + liquid ! b. Furthermore, the term peritectoid denotes the special case of an equilibrium phase in which two or more solid phases (which are stable above the temperature Tp) react at Tp to form a new solid phase. This reaction can be written as a + b ! g. The phases formed during a peritectic or peritectoid reaction are a solid solution of one of the components, an allotropic phase of one of the components, or an intermetallic base. Schematics of peritectic phase diagrams are shown in Fig. 1. Many interesting alloys undergo these types of reactions, for example, iron-carbon- and iron-nickel-base alloys as well as copper-tin and copper-zinc alloys. Controlled peritectic reactions and transformations are rarely used to optimize the microstructure of engineered materials. Usually, peritectic structures are avoided because of their detrimental effect on material properties. However, an understanding of peritectic solidification structure development is necessary in order to concisely interpret microstructures. Also, there may be some technical importance attributed to the possibility of producing metallic alloys with aligned fibers or plates (Ref 2–5) and of achieving grain refinement by means of peritectic decay of primary crystals (Ref 6, 7).
Phase Diagrams Phase diagrams are very instructive when describing peritectic phase transformations. Figure 2 shows a phase diagram with a peritectic reaction. This diagram shows that, under equilibrium conditions, all alloys to the left of I will solidify to a-crystals. Similarly, all alloys to the right of III will solidify to b-crystals. Alloys between II and III first solidify to a-crystals and then transform to stable b-crystals. Alloys between I and II also solidify to a-crystals, but they are partially transformed to b-crystals later. The volume fraction of each phase will be given by the lever rule if the alloy solidifies under equilibrium conditions (the lever rule is defined in the article “Thermodynamics and Phase Diagrams” in this Volume). In practice, the lever rule usually will not give the volume fraction of the different phases, because the nucleation and growth kinetics, determined by the diffusion rate, in the solid phases determine the time required to reach equilibrium.
Peritectic Solidification Peritectic structures form by means of at least three mechanisms: Peritectic reaction, where all three phases (a, b,
Direct precipitation of b from the melt, when
enough undercooling below the peritectic temperature (Tp) occurs. Direct precipitation of b from the melt can occur when a peritectic reaction or peritectic transformation is sluggish, as is often the case. Peritectic reactions can proceed only as long as a-phase and liquid are in contact. The b-solid nucleates at the a/liquid interface, readily forming a layer that isolates a from the liquid. This mechanism occurs through short-range diffusion, as shown in Fig. 3(a). In contrast, peritectic transformations occur through a long-range mechanism, where A atoms and B atoms migrate through the a-layer to form solid b at the a/b and the b/liquid interfaces, respectively (Fig. 3b). These distinctions between peritectic reaction and peritectic transformation were introduced by Kerr (Ref 8). The mechanism by which b has formed is not apparent from the final microstructure. In any case, all three of the mechanisms require some undercooling, because the driving force is zero at the peritectic temperature. The time dependence is pronounced for the peritectic transformation. Therefore, the amount of b formed will depend on the cooling rate or on the holding time if isothermal conditions are established. Localized nucleation of b and shape changes of b by solution reprecipitation, which is driven
and liquid) are in contact with each other
Peritectic transformation, where the liquid
and the a-solid solution are isolated by the b-phase. The transformation takes place by long-range diffusion through the secondary b-phase.
Typical peritectic phase diagrams. (a) Peritectic reaction a + liquid ! b and peritectoid reaction a + b ! g. (b) Peritectic formation of intermetallic phases from a high-melting intermetallic. (c) Peritectic cascade between high- and low-melting components. Source: Ref 1
Fig. 1
Fig. 2
Phase diagram with a peritectic reaction
Peritectic Solidification / 331
Mechanisms of peritectic reaction and transformation. (a) Lateral growth of a b-layer along the a/liquid interface during peritectic reaction by liquid diffusion. (b) Thickening of a b-layer by solid-state diffusion during peritectic transformation. The solid arrows indicate growth direction of b; dashed arrows show the diffusion direction of the atomic species. Source: Ref 8
Fig. 3
Fig. 5
Start of the peritectic reaction in a directionally solidified Cu-20Sn alloy. Primary a-dendrites (white) are covered by peritectically formed b-layer (gray) shortly after the temperature reaches Tp. Matrix (dark) is a mixture of tin-rich phases. Mechanically polished, etched in HNO3. Original magnification: 40. Source: Ref 12
Types of peritectic systems. (a) Type A system where the b/a + b solvus and the b-solidus have slopes of the same sign. (b) Type B system where the slopes have opposite signs. (c) Type C system where the b-phase has a limited composition. Source: Ref 1
Fig. 4
by surface energy through diffusion in the liquid, also influence the microstructure of peritectic alloys. The tendency for final solidification microstructures to be far from equilibrium is another factor limiting their widespread commercial use. St. John and Hogan (Ref 9) ascertained that the degree of completion may be approximated based on certain characteristics of the phase diagram. From their work on the aluminumtitanium system and the work of Titchener and Spittle (Ref 10), three classes of peritectic phase diagram (Fig. 4) have been identified based on the shape of the peritectic (b) solid-solution region (Ref 11). Additional information is available in the article “Invariant Transformation Structures” in Metallography and Microstructures, Volume 9 of ASM Handbook.
Peritectic Reaction In the classical description of peritectic reactions, heterogeneous nucleation of b occurs at the a/liquid interface at the peritectic equilibrium temperature, Tp. If nucleation is limited to a few locations, and lateral growth of the b-nuclei does not readily occur, no continuous layer of the peritectic phase is formed. Two types of the peritectic reactions can occur: Nucleation and growth of the b-crystals in
the liquid without contact with the a-crystals
Nucleation and growth of the b-crystals in
contact with the primary a-phase
In the first case, the secondary phase is nucleated in the liquid and does not contact the primary phase because of surface tension conditions. Following nucleation, the secondary phase grows freely in the liquid. At the same time, the primary phase will dissolve. The secondary phase will not develop a morphology similar to a precipitated primary phase. This type of peritectic reaction has been observed for the reaction g + L ! b in the aluminummanganese system (Ref 12). The second phase tends to grow around the primary phase at increasing cooling rates. Similar reactions have been observed in nickel-zinc (Ref 13) and aluminum-uranium (Ref 14) systems. In the second type of reaction, which is the most common, nucleation of the secondary b-phase occurs at the interface between the primary a-phase and the liquid. A lateral growth of the b-phase around the a-phase then takes place. In an ideal peritectic reaction, undercooling is rather low (up to a few degrees Kelvin), and a plateau is observed in the cooling curve, as in the aluminum-titanium system (Ref 8). Envelopes of the peritectic phase around the primary phase form by direct reaction in some systems—copper-tin and silver-tin, for example— through interrupted directional solidification experiments (Ref 12). Figures 5 and 6 depict microstructures of Cu-20Sn and Cu-70Sn alloys
Fig. 6
Start of the peritectic reaction in a directionally solidified Cu-70Sn alloy. The primary e-phase (dark) is covered by the peritectically formed Z-layer (white), which thickens with increasing undercooling below Tp. The matrix is the Sn-Z eutectic. Mechanically polished, etched in HNO3. Original magnification: 100. Source: Ref 12
that demonstrate the onset of the peritectic reactions a + liquid ! b and e + liquid ! Z, respectively. Figure 7 shows, at a higher magnification, the homogeneous thickness of the blayer around the a-dendrites. The copper-tin phase diagram is shown in Fig. 8. Remelting is normally much faster than solidification, and there are no forces hindering dissolution. Therefore, dissolution will occur at the same rate as precipitation of the secondary phase. Thus, the peritectic reaction can proceed very rapidly by liquid diffusion over a very short distance in the lateral direction, as shown
332 / Principles of Solidification
Fig. 7
Start of the peritectic transformation in the same directionally solidified Cu-20Sn alloy shown in Fig. 5, but at higher magnification. Note the homogeneous thickness of the b-layers (gray) around the primary a (white). The matrix (dark) is a mixture of tinrich phases. Mechanically polished, etched in HNO3. Original magnification: 160. Source: Ref 12
Fig. 8 in Fig. 3(a). The thickness of the layers has been calculated with fairly good agreement to experimental results for copper-tin and silvertin alloys on the basis of maximum growth rate or minimum undercooling from the laws derived for solidification at low undercoolings (Ref 12). The thickness depends to some extent on the cooling rate but more strongly on the interfacial energies, s, with s (liquid/b) + s (a/b) s (liquid/a) as the determining factor. Nucleation of the secondary phase on the primary phase is sometimes favored. When this occurs, a number of crystals will form around the primary phase, as illustrated in Fig. 9 (Ref 13).
Peritectic Transformation and Direct Precipitation The precipitation of b directly from the liquid and the solid depends on the shape of the phase diagram and the cooling rate. The diffusion process through the b-layer depends on the diffusion rate, the shape of the phase diagram, and the cooling rate. The thickness of the b-layer will normally increase during subsequent cooling. There are three reasons for this: Diffusion through the b-layer Precipitation of b directly from the liquid Precipitation of b directly from the a-phase
The thickness of the b-phase envelope surrounding the a-phase is determined by the peritectic reaction followed by an increase in thickness due to a precipitation directly from the liquid. The rate of the peritectic transformation is influenced by the diffusion rate and the extension of the b-phase region in the phase
Copper-tin phase diagram. Source: Ref 15
diagram. If the diffusion rate is small, the peritectic transformation will be negligible compared to the peritectic reaction. During continuous cooling, this diffusion-controlled growth is affected by precipitation from the liquid and from the primary a or by dissolution of b, according to the slopes of the solubility limits in the phase diagram. Under simplifying conditions, the growth rates have been calculated numerically and found to be in reasonable agreement with the experimental findings in the copper-tin and silver-tin systems (Ref 12). The kinetics of peritectic transformations can more easily be studied under isothermal conditions. Then, b is formed exclusively by diffusion of the two atomic species in the b-layer and at the a/b and b/liquid interfaces. From calculations described in Ref 16 and 17, one can assume that the growth of the b-layer is controlled by the diffusion through it at a temperature just below the peritectic temperature. The diffusion process and the concentration profile and its relation to the phase diagram are illustrated in Fig. 10. A mass balance gives the following: a ! b : Db ¼
ðxb=L xb=a Þ ‘b
d‘b=a ðxb=a xa=b Þ dt
(Eq 1)
and L ! b : Db ¼
b=L
d‘ dt
ðxb=L xb=a Þ ‘b
ðxL=b xb=L Þ
(Eq 2)
Fig. 9
Microstructure in a Cd-10Cu sample that has passed a peritectic reaction. The primary Cu5Cd8 crystals are white, the dark matrix is cadmium, and the peritectically formed CuCd3 is gray.
where ‘b is the thickness of the b-phase, t is the time, and Db is the diffusion coefficient in the b-phase. All other terms are concentrations that are defined in Fig. 10: d‘b d‘b=a d‘b=L Db ¼ þ ¼ b ðxb=L xb=a Þ dt dt ‘ dt 1 1 þ L=b b=a a=b b=a ðx x Þ ðx x Þ
(Eq 3)
Peritectic Solidification / 333 dissolved alloying elements, the diffusion rates are much higher than for substitutionally dissolved elements in face-centered cubic metals. In this case, the diffusion process has a much larger influence on the precipitation process.
Iron-Carbon System A peritectic transformation occurs in the important iron-carbon system The iron-carbon phase diagram (Fig. 11) is very similar to the diagram illustrated in Fig. 10(a). A detailed numerical calculation of the transformation for one alloy has been performed (Ref 18). However, a simpler analytical model is used in this discussion (Ref 19). The model is the same as the isothermal model described earlier but is used with the assumption that the boundary conditions change during cooling. Equation 3 can therefore be used. In Eq 3, d‘b/dt can be expanded to (d‘b/dT)(dT/dt), where dT/dt is the cooling rate of the sample. Using the phase diagram, the concentrations can be transferred to a temperature. If the slope of the lines in the phase diagrams and the cooling rate are both assumed to be constant, Eq 3 can be simplified in the following way: ‘b d‘b ¼ A
Fig. 10
The peritectic transformation in a system with a high diffusion rate in the b-phase. (a) The phase diagram. (b) Concentration distribution. The dashed lines in the concentration profile are for a system with a low diffusion rate in the b-phase where the volume fraction of b increases with decreasing temperature. See Eq 6 and corresponding text. See also Fig. 13(b).
Integrating Eq 3 results in: ð‘b Þ2 ¼Db ðxb=L xb=a Þ 1 1 þ ðxb=a xa=b Þ ðxL=b xb=L Þ
(Eq 4)
Equations 3 and 4 show that the growth rate increases with increasing undercooling. For example, at the peritectic temperature, the expression (xb/L xb/a) is 0 and increases with increasing undercooling. Equations 3 and 4 also show that the growth rate is dependent on the diffusion coefficient. For substitutionally dissolved alloying elements in face-centered cubic metals, the diffusion coefficient near the melting point is of the order of 1013 m2/s. In such a case, the growth rate will be very low, and the time for the peritectic transformation will be unrealistically large. In a normal casting process, the reaction rate will be so low that the amount of b-phase formed by the peritectic reaction will be negligible in comparison with the precipitation of b from the liquid. For bodycentered cubic metals and for interstitially
dt
dT dT
(Eq 5)
where A is a constant. By integrating Eq 5 under the assumption of a constant cooling, one can obtain a relation that can be used to calculate the temperature interval under which the peritectic transformation takes place. This has been done in Ref 19, and the result of those calculations is shown in Fig. 12. Figure 12 shows that the peritectic reaction in iron-carbon alloys is very rapid and is finished at a maximum of 6 to 10 K below the peritectic temperature. The shape of the curve is a result of the change in the volume fraction of ferrite with the carbon content. The curve is in agreement with the results of other experiments (Ref 18, 20). As mentioned previously, the effect of the phase diagram must also be considered (Fig. 4). This was not taken into consideration when deriving Eq 3 and in the previous discussion for iron-carbon alloys. The effect of the phase diagram is illustrated in Fig. 10(a) and 13(a). Figures 10(a) and 13(a) show two different types of phase diagrams, in which the slopes of the a/b regions are different. In the first case (Fig. 10a), the b-layer will increase in thickness both at the interface a/b and at the interface b/L. In the second case (Fig. 13a), the b-layer will increase only at the side against the liquid and will decrease at the side against the a-phase. These two cases are somewhat more difficult to treat theoretically than the isothermal example or the case involving a high diffusion rate. The concentration profiles for the transformation are given in Fig. 10(b), as shown by the dashed lines, and Fig. 13(b). Equation 3 will now be changed in the following way:
b dx dy y ¼ ‘a=b b Db dx þ L=b x xb=L dy y ¼ ‘b=L b Da dx þ a=b x xb=a dy y ¼ ‘a=b
d‘b Db ¼ dt ðxa=b xb=a Þ
þ
‘L dxL xL=b xb=L
(Eq 6)
where the first and second terms on the right side describe the increase in thickness due to diffusion into the b-phase and a-phase, respectively, from the boundary a/b. The third term describes the increase in thickness due to diffusion into the b-phase from the boundary b/L. The last term is the increase in the b-phase due to the changing of the composition in the liquid dxL with a decrease in the temperature dT according to the phase diagram. Solving Eq 6 is a difficult task that can require the use of numerical methods. The equation can be simplified by assuming that the diffusion distance corresponds to the thickness of the b-phase and that there are no concentration gradients in the a-phase. Figure 14 compares simulated and experimental determinations of the thickness of the secondary phase layer for a Cu-20Sn alloy (Ref 12), showing close agreement between the experiments and the theory. Equation 6 has been solved in Ref 19 by assuming that diffusion into the b-phase could be described in the same manner as in Ref 21 for back diffusion during segregation at a primary precipitation. The model was used to calculate the concentration distribution of austenite for iron-nickel alloys, as shown in Fig. 15. These calculations are in agreement with experiments reported in Ref 22 and 23.
Primary Metastable Precipitation of Beta In many systems, the secondary b-phase has been observed to form as a primary phase for alloys with a composition left of point III in Fig. 2. This is especially true for iron-base alloys (Ref 22, 23). The formation of a metastable b-phase directly from the liquid under different cooling rates can be demonstrated by examining the binary iron-nickel phase diagram. In Fig. 16, the metastable extensions of the equilibrium between liquid and ferrite and between liquid and austenite are represented by dashed lines. The melting point of pure iron as austenite is 11 C (20 F) below the melting point of pure iron as ferrite. It can be seen that there is a difference in the partition coefficient between ferrite and liquid on one hand and austenite and liquid on the other. The growth rates were calculated for needles of ferrite and austenite. The alloy composition chosen is to the left of the peritectic point, and the calculations of the growth rates were carried out for different undercoolings.
334 / Principles of Solidification
Fig. 11
Carbon-iron phase diagram. Source: Ref 15
Fig. 12
Temperature range of peritectic reaction in iron-carbon alloys as a function of carbon content and the solidification rate. The temperature gradient, G, is 6000 K/m.
Only ferrite can form at low undercoolings. Austenite can also form if the temperature is chosen just below the extension of its liquidus
line, but ferrite grows faster and should therefore dominate in the solidification process. However, below a critical line, austenite will have the highest growth rate and may dominate, although the temperature is still above the peritectic temperature. This kinetic advantage of austenite is due to its partition coefficient being smaller than that for ferrite. The critical line is indicated by a dashed line in Fig. 16. The partition coefficient of an element can be closer to unity for ferrite than for austenite in ternary alloys, and ferrite will then have the kinetic advantage and will be favored by a high cooling rate. It has also been reported that an aligned structure can be formed in peritectic systems (Ref 2). However, it has been argued that this type of structure is formed when a- and b-primary phases are growing at the same rate (Ref 12). In addition, the composition of the liquid must be chosen so the sum of the volume fraction of the two solid phases will be unity.
Peritectic Cascades The theoretical analysis shows that the rate of the peritectic transformation is influenced by the diffusion rate and the extension of the b-phase region in the phase diagram. If the diffusion rate is small, the peritectic transformation will be negligible compared to the peritectic reaction. The thickness of the b-phase envelope surrounding the a-phase is determined by the peritectic reaction followed by an increase in thickness due to a precipitation directly from the liquid. In many systems, one peritectic reaction at a high temperature is followed by one or more peritectic reactions at lower temperatures. If the diffusion rate is low in the initially formed peritectic layer, a second peritectic layer can be formed when the second peritectic temperature is reached. This type of series of peritectic reactions is referred to as a cascade (Ref 13). A characteristic microstructure formed by cascading peritectic reactions in a Sn-17Co alloy is shown in Fig. 17.
Peritectic Solidification / 335
Fig. 15
Nickel distribution after peritectic reaction in a steel containing 4 wt% Ni. The temperature gradient was 60 K/cm. Calculations were made at different solidification rates. The dotted line shows the nickel distribution at the start of the peritectic reaction. d is primary ferrite, g is austenite. Source: Ref 22
Fig.
13
The peritectic transformation during continuous cooling in a system with a low diffusion rate and where the volume fraction of b increases with decreasing temperature. (a) The phase diagram. (b) Concentration distribution
Fig. 17
Microstructure of a Sn-17Co alloy that was cooled to 1000 C (1830 F) and held 2 h, then cooled to 225 C (435 F) and held 22 h. The primary g-phase has completely transformed and dissolved into relatively small CoSn crystals that form the dark centers surrounded peritectically by CoSn2 layers (gray), followed by a peritectoid envelope of the CoSn3 phase. The temperatures of peritectic formation of CoSn and CoSn2 are approximately 910 and 540 C (1670 and 1005 F), respectively; CoSn3 forms peritectoidally at approximately 235 C (455 F). Scanning electron micrograph (backscattered electron image). Original magnification: 200
Fig. 16
The binary iron-nickel system. The broken line shows the undercooling at which ferrite, d, and austenite, g, grow at the same rate. Source: Ref 24
Peritectic Transformations in Multicomponent Systems
Fig. 14
Thickness of the secondary phase layer as a function of temperature below the peritectic temperature in the copper-tin system. The solidification rate was 100 mm/min (4 in./min). The diffusion units are given in cm2/s. The volume fraction is defined as ‘b/l, where l corresponds to the dendrite arm space.
Alloys often consist of more than two alloying elements. Investigations of iron-base alloys have shown that peritectic reactions are very common in stainless steels (Ref 22, 25). The peritectic reaction in these alloys gives the same type of nickel distribution as shown in Fig. 15. In stainless steels, the peritectic reaction will become a eutectic reaction if the chromium content is increased to 20% or more. This transition is also influenced by the molybdenum content, as shown in Fig. 18.
Fig. 18
The transition from a peritectic to a eutectic reaction as a function of chromium and molybdenum content in a stainless steel containing 11.9% Ni
Both chromium and nickel are substitutionally dissolved elements. Iron-base alloys often consist of carbon with some other elements. Carbon is interstitially dissolved and has a very high diffusion rate. The other alloying elements are primarily substitutionally dissolved with very low diffusion rates. This gives rise to
336 / Principles of Solidification
Fig. 19
Three stages of a peritectic reaction in a unidirectionally solidified high-speed steel. (a) First-stage structure. Dark gray is austenite, white is ferrite. The mottled structure is quenched liquid. (b) Subsequent peritectic transformation of (a). (c) Further peritectic transformation of (a) and (b). Dark gray in the middle of the white ferrite is newly formed liquid. Source: Ref 26
transformations that are determined by the movement of the substitutional elements, and carbon is distributed according to equilibrium conditions. As a result, a normal peritectic transformation does not occur. To fulfill the criterion that carbon should follow the equilibrium conditions, liquid must be formed at the border between ferrite and austenite. This reaction is illustrated in Fig. 19. This type of reaction has been both experimentally and theoretically analyzed in Ref 26. ACKNOWLEDGMENTS This article has been adapted from H. Fredriksson, “Solidification of Peritectics,” Casting, Vol 15, Metals Handbook, 9th ed., ASM International, 1988, p 125–129, and H.E. Exner et al., “Peritectic and Peritectoid Structures,” Metallography and Microstructures, Vol 9, ASM Handbook, ASM International, 2004, p 156–164. REFERENCES 1. H.E. Exner and G. Petzow, Peritectic Structures, Metallography and Microstructures, Vol 9, ASM Handbook, American Society for Metals, 1985, p 675–680 2. W.J. Boettinger, The Structure of Directionally Solidified Two-Phase Sn-Cd Peritectic Alloys, Metall. Trans. A, Vol 5, 1974, p 2023–2031 3. H.D. Brody and S.A. David, Controlled Solidification of Peritectic Alloys, Solidification and Castings of Metals, The Metals Society, London, 1979, p 144–151 4. A.P. Tichener and J.A. Spille, The Microstructure of Directionally Solidified Alloys That Undergo a Peritectic Transformation, Acta Metall., Vol 23, 1975, p 497–502 5. A. Ostrowski and E.W. Langer, Unidirectional Solidification of Peritectic Alloy,
6. 7.
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17. 18.
Solidification and Castings of Metals, The Metals Society, London, 1979, p 139–143 J.A. Sartell and D.J. Mack, The Mechanism of Peritectic Reactions, J. Inst. Met., Vol 93, 1964, p 19–24 I. Maxwell and A. Hellawell, An Analysis of the Peritectic Reaction with Particular Reference to Al-Ti Alloys, Acta Metall., Vol 23, 1975, p 901–909 H.W. Kerr, J. Cisse, and G.F. Bolling, On Equilibrium and Non-Equilibrium Peritectic Transformation, Acta Metall., Vol 22, 1974, p 677 D.H. St. John and L.M. Hogan, Acta Metall., Vol 35, 1987, p 171–174 A.P. Titchener and J.A. Spittle, Met. Sci., Vol 8, 1974, p 112–116 J.W. Rutter et al., Mater. Sci. Technol., Vol 14, 1998, p 182–186 H. Fredriksson and T. Nyle´n, Mechanism of Peritectic Reactions and Transformations, Met. Sci., Vol 16, 1982, p 283–294 G. Petzow and H.E. Exner, Zur Kenntnis Peritektischer Umwandlungen, Radex Rundsch., 3/4, 1967, p 534–539 S. Uchida, “Systematik and Kinetik Peritektischer Umwandlund,” Ph.D. thesis, Technical University, 1980 Alloy Phase Diagrams, Vol 3, ASM Handbook, ASM International, 1992, p 2.110 M. Hillert, Eutectic and Peritectic Solidification, Solidification and Casting of Metals, The Metals Society, London, 1979, p 81–87 D.H. St. John and L.M. Hogan, The Peritectic Transformation, Acta Metall., Vol 25, 1977, p 77–81 Y.K. Chuang, D. Reinisch, and K. Schwerdtfeger, Kinetics of the Diffusion Controlled Peritectic Reaction During Solidification of Iron-Carbon Alloys, Metall. Trans. A, Vol 6, 1975, p 235–238
19. H. Fredriksson and J. Stjerndahl, Solidification of Iron-Base Alloys, Met. Sci., Vol 16, 1982, p 575–585 20. J. Stjerndahl, “The Solidification Process of Iron Base Alloys,” thesis, Department of Casting of Metals, Royal Institute of Technology, 1978 21. H.D. Brody and M.C. Flemings, Solute Redistribution in Dendritic Solidification, Trans. Met. Soc. AIME, Vol 236, 1966, p 615–624 22. H. Fredriksson, The Solidification Sequence in an 18-8 Stainless Steel, Metall. Trans., Vol 3, 1972, p 2989–2997 23. H. Fredriksson and J. Stjerndahl, On the Formation of a Liquid Phase During Cooling of Steel, Metall. Trans. B, Vol 6, 1975, p 661 24. H. Fredriksson, Segregation Phenomena in Iron-Base Alloys, Scand. J. Metall., Vol 5, 1976, p 27–32 25. H. Fredriksson, The Mechanism of the Peritectic Reaction in Iron-Base Alloys, Met. Sci., Vol 11, 1976, p 77–86 26. H. Fredriksson, Transition from Peritectic to Eutectic Reaction in Iron-Base Alloys, Solidification and Casting of Metals, Publication 192, The Metals Society, London 1977, p 131–136 SELECTED REFERENCES I.V. Belyaev, M. Kursa, and Y. Drapala,
Directed Crystallization of Peritectic Transformed Alloys, METAL 2007: 16th International Metallurgical and Materials Conference, 2007 C. Bernhard and G. Xia, Influence of Alloying Elements on the Thermal Contraction of Peritectic Steels during Initial Solidification, Ironmaking Steelmaking, Vol 33 (No. 1), 2006, p 52–56
Peritectic Solidification / 337 B.K. Dhindaw et al., Characterization of the
Peritectic Reaction in Medium-Alloy Steel Through Microsegregation and Heat-ofTransformation Studies, Metall. Mater. Trans. A, Vol 35 (No. 9), 2004, p 2869–2879 B.K. Dhindaw and H. Fredriksson, Characterisation of the Peritectic Reaction in Medium Alloy Steel Through Directional Solidification and Micro-Segregation Studies, Trans. Indian Inst. Met., Vol 56 (No. 5), 2003, p A5 R.J. Dippenaar, High-Temperature Phase Transitions and Mechanical Properties in Steel of Peritectic Composition, Mater. Sci. Forum, Vol 539–543, 2007, p 4243–4248 S. Dobler et al., Peritectic Coupled Growth, Acta Mater., Vol 52 (No. 9), 2004, p 2795– 2808 A. Karma and M. Plapp, New Insights into the Morphological Stability of Eutectic and Peritectic Coupled Growth, JOM, Vol 56 (No. 4), 2004, p 28–32 E. Kozeschnik, W. Rindler, and B. Buchmayr, Scheil-Gulliver Simulation with Partial Redistribution of Fast Diffusers and Simultaneous Solid-Solid Phase Transformations, Int. J. Mater. Res. (Z. Metallkd.,) Vol 98 (No. 9), 2007, p 826–831 Y. Liu et al., Primary Cellular/Dendritic Spacing Selection in Rapidly Solidified Peritectic Alloy, Mater. Lett., Vol 59 (No. 23), 2005, p 2915–2919 T.A. Lograsso, B.C. Fuh, and R. Trivedi, Phase Selection during Directional Solidification of Peritectic Alloys, Metall. Mater. Trans. A, Vol 36 (No. 5), 2005, p 1287–1300 D. Ma et al., On Secondary Dendrite Arm Coarsening in Peritectic Solidification,
Mater. Sci. Eng. A, Vol 390 (No. 1–2), 2005, p 52–62 N.J. McDonald and S. Seetharaman, “Peritectic Transition in the Hypoperitectic IronNickel System,” Ferrous Alloys, April 27–30, 2003 (Indianapolis, IN) N.J. McDonald and S. Sridhar, Peritectic Reaction and Solidification in Iron-Nickel Alloys, Metall. Mater. Trans. A, Vol 34 (No. 9), 2003, p 1931–1940 N. McDonald and S. Sridhar, Observations and Analysis of the Peritectic Reaction for Fe-Co Alloys, Z. Metallkd., Vol 96 (No. 3), 2005, p 304–310 J. Miettinen, Thermodynamic-Kinetic Model for the Simulation of Solidification in Binary Copper Alloys and Calculation of Thermophysical Properties, Comput. Mater. Sci., Vol 36 (No. 4), 2006, p 367–380 J.H. Perepezko, Nucleation Controlled Phase Selection during Solidification, Mater. Sci. Eng. A, Vol 413–414, 2005, p 389–397 D. Phelan, M. Reid, and R. Dippenaar, Experimental and Modelling Studies into High Temperature Phase Transformations, Comput. Mater. Sci., Vol 34 (No. 3), 2005, p 282–289 D. Phelan, M. Reid, and R. Dippenaar, Kinetics of the Peritectic Phase Transformation: In-Situ Measurements and Phase Field Modeling, Metall. Mater. Trans. A, Vol 37 (No. 3) 2006, p 985–994 R. Pierer and C. Bernhard, High Temperature Behavior during Solidification of Peritectic Steels Under Continuous Casting Conditions, MS&T ’06: Materials Science and Technology 2006 Conference and Exhibition, 2007
S. Raghavan, M.J.M. Krane, and D. Johnson,
“Different Rule Sets for Cellular Automata Modeling of Peritectic Dendritic Growth,” Modeling of Casting, Welding and Advanced Solidification Processes—X, May 25–30, 2003 (Destin, FL) L. Snugovsky et al., Phase Equilibria in SnRich Corner of Cu-Ni-Sn System, Mater. Sci. Technol., Vol 22 (No. 8), 2006, p 899–902 Y. Su et al., Microstructure Evolution of Ti-Al Peritectic System during the Initiated Stage of Directional Solidification, Mater. Sci. Forum, Vol 546–549, 2007, p 1447–1450 Y. Su et al., Microstructure Selection during the Directionally Peritectic Solidification of Ti-Al Binary System, Intermetallics, Vol 13 (No. 3–4), 2005, p 267–274 Y. Su et al., Numerical Model for Nucleation of Peritectic Alloy during Unidirectional Solidification, J. Mater. Sci., Vol 42 (No. 14), 2007, p 5360–5368 M. Sumida, Growth Analysis of Banded Microstructure Formed during Directional Solidification of Peritectic Alloys, Mater. Trans., Vol 45 (No. 2), 2004, p 388–391 R. Trivedi, The Role of Heterogeneous Nucleation on Microstructure Evolution in Peritectic Systems, Scr. Mater., Vol 53 (No. 1), 2005, p 47–52 R. Trivedi and J.H. Shin, Modelling of Microstructure Evolution in Peritectic Systems, Mater. Sci. Eng. A, Vol 413–414, 2005, p 288–295 M.F. Zhu and C.P. Hong, Modeling of Microstructure Evolution in Eutectic and Peritectic Solidification, Modeling of Casting, 2003
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 338-347 DOI: 10.1361/asmhba0005215
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Microsegregation Harold D. Brody, University of Connecticut
MICROSEGREGATION is the nonuniform distribution of alloying elements within a volume characteristic of the solidification microstructure, that is, chemical nonuniformity that can be observed and measured on the scale of the dendrite spacing and/or the grain size. The cored dendritic structure of a cast Ti-15%Mo alloy is shown, for example, in Fig. 1. Microsegregation results when the partitioning and redistribution of solute between liquid and solid phases that accompany the solidification of alloys does not reach equilibrium, and uniform chemical potential is not attained within all phases of the as-cast alloy. Diffusion, typically diffusion of solute in the primary solid solution, is the rate-limiting step in the redistribution of solute during solidification and postsolidification cooling and is the dominant factor in determining the extent of microsegregation in as-cast alloys. In contrast, solute redistribution processes that do not preserve the local solute concentration result in macrosegregation. Bulk convective flow, typically of solute-enriched (or solute-depleted) liquid in the mushy zone, is the dominant factor in determining the extent of macrosegregation in as-cast alloys. The segregation ratio, S, the local maximum solute content divided by the local minimum, and the fraction of nonequilibrium constituents, f nsq, are two, among several, measures of the extent of microsegregation. Because solute redistribution processes continuously alter the apparent microstructure during solidification and postsolidification cooling, interrupted solidification and in situ imaging are used to follow the kinetics of processes that lead to microsegregation. Microsegregation and other manifestations of nonequilibrium solidification limit castability, further processability, and the final properties of the alloy. Microsegregation and its effects can be eliminated or reduced by homogenization and solution heat treatments, which, in the case of wrought alloys, are combined with mechanical deformation processes to reduce the scale of the solidification microstructure.
Solidification Models In general, alloys freeze over a temperature range, and the solute is continually redistributed,
or partitioned, among solid and liquid phases in an attempt to maintain chemical equilibrium. Microsegregation results when overall equilibrium is not attained. For detailed modeling of solute redistribution and microsegregation, see the article “Modeling of Microsegregation and Macrosegregation” in this Volume. Herein, the process and material parameters that control microsegregation are discussed in relation to the manifestations of microsegregation in simple and then increasingly complex alloy systems. First, the two extremes of solute redistribution are discussed, equilibrium solidification and nonequilibrium Gulliver-Scheil solidification, for which solid redistribution of solute within the primary solid phase is the distinguishing parameter. C The partition ratio, ki ¼ CLiSi (or distribution coefficient), describes the difference in compo sition for solute i of the liquid, CLi , and solid, CSi , phases that would be at equilibrium at a temperature within the freezing range as given by the equilibrium tie line. (See the article “Thermodynamics and Phase Diagrams” in this Volume for phase diagram fundamentals.) For ki < 1, the liquid is enriched in solute i during solidification. For ki > 1, the liquid is depleted in solute i during solidification. The further ki is from unity, the greater the potential for microsegregation. If the liquidus and solidus for a terminal solid solution are straight, the partition ratio is a constant.
can be used to compute the progress of solidification: fS ¼
CL CO T TL CL CS ðT TM Þð1 kÞ
(Eq 1)
where fS is the fraction solid, TM is the equilibrium melting temperature, TL is the liquidus temperature for the alloy, and the latter approximate ratio assumes straight liquidus and solidus phase boundaries. Equilibrium is one extreme of solidification behavior (Ref 2–4).
Equilibrium Solidification For solidification to proceed through the freezing range, the equilibrium lever law requires that there must be (1) negligible undercooling before the nucleation of the solid phases; (2) negligible undercooling at the solid-liquid interface due to curvature, interface kinetics, or a solute boundary layer (local equilibrium at the solid-liquid interface); (3) sufficient mass transport in the remaining liquid phase for the composition of the liquid to remain uniform as it enriches (or depletes) in solute; (4) constant local average solute concentration, that is, negligible macrosegregation; and (5) sufficient mass transport in the forming solid phase(s) to maintain uniform chemical potential. The equilibrium lever rule
Fig. 1
As-cast microstructure showing dendritic coring in Ti-15%Mo alloy, and electron microprobe trace showing variation of molybdenum content across dendrite structure. Source: Ref 1
Microsegregation / 339 Nonequilibrium Solidification, The Gulliver-Scheil Relation In most commercial casting and welding practices, and for many alloys, equilibrium freezing would not be attained. The freezing range would be extended. The primary dendritic phase would be cored. Nonequilibrium phases will be formed. The nonequilibrium lever rule proposed by Gulliver and Scheil describes another extreme of solute redistribution during solidification. Conditions 1 through 4 listed previously are assumed by the Gulliver-Scheil relation, which may be used to compute the progress of solidification for condition 5, negligible redistribution of solute within solid phases by diffusion or other processes: fS ¼ 1
1=1k CO TL TM 1=1k 1 (Eq 2) CL T TM
The composition of the primary phase forming at the solid-liquid interface is given by another form of the Gulliver-Scheil relation: CS ¼ ð1 fS Þk1
(Eq 3)
The composition of the remaining liquid phase changes continuously and uniformly from the initial melt composition, CO, and the composition of the primary solid phase changes from kCO. Solidification is not complete until an invariant reaction, such as the melting point or a eutectic, is reached. The Gulliver-Scheil conditions lead to the maximum extent of microsegregation (Ref 2–7).
Nonequilibrium Solidification with Back Diffusion The actual solidification path and the extent of microsegregation for a cast alloy will be somewhere between the two extremes. The conditions 1 through 4 represent the solidification of most alloys over a wide range of welding and casting processes. The kinetics of solute redistribution in the solid phase(s), condition 5, to the first approximation, depend on solute diffusion within the local characteristic volume. The mass transport Fourier number, a, is a useful predictor of the extent of microsegregation: a¼
DS tdiff DSðTL Þ yf das 2 L2diff
(Eq 4)
2
where the characteristic time for diffusion, tdiff, is approximated by the local solidification time, yf, and the characteristic diffusion distance, Ldiff, is approximated by one-half the final dendrite arm spacing (das) for the local region of the casting under consideration (Ref 1–4, 8–22). For a highly grain-refined casting, onehalf the grain size could be used to approximate Ldiff (Ref 20). The solid diffusion coefficient, DS, changes exponentially with temperature
during solidification and during postsolidification cooling. For simplicity, the diffusivity around the liquidus temperature, DSðTL Þ , can be selected to estimate a. The transition from the extremes of Gulliver-Scheil to equilibrium behavior occurs as 0.01 < a < 100. For low-melting-temperature alloys, microsegregation approaches the Gulliver-Scheil condition. For high-melting-temperature alloys, and especially for interstitial solute elements, solute redistribution tends to equilibrium. The measurement and kinetics of microsegregation are discussed for the binary isomorphous system, titanium-molybdenum; for binary eutectic systems, aluminum-copper and aluminum-silicon; for a binary peritectic system, copper-Zinc; for a multicomponent eutectic system, Al-Si-Cu-Mg; and, finally, for systems with both eutectic and peritectic reactions, Fe-C-Cr and nickel-base superalloy.
Binary Isomorphous Systems, k > 1: Titanium-Molybdenum For many titanium-base alloys, as-cast microsegregation is close to the minimum— approaching equilibrium—and the observed segregation ratios are close to unity. In those titanium-base alloy systems that freeze over a narrow temperature range close to or above the melting point, the partition ratios are close to 1, and solute diffusivities in b-titanium are relatively high. Solute redistribution is extensive in b-titanium dendrites during solidification. Microsegregation present at the end of freezing is reduced by diffusion as the alloy cools in the mold from the high solidus temperatures. The relatively high diffusivities of alloying elements in b-titanium are countered by relatively large as-cast dendrite arm spacing (when compared, for example, to aluminum alloys with local solidification times), presumably the result of relatively high solute diffusivities in the liquid. Solute redistribution is less extensive and microsegregation more noticeable when alloying elements that lower the freezing range, such as iron, are added to titanium (Ref 1, 21, 22). Titanium-molybdenum is an isomorphous system. The liquidus and solidus boundaries
slope upward from the melting temperature of titanium to the melting point of molybdenum. For the compositions 5 < wt% Mo < 30%, k 1.4. As titanium-molybdenum alloys freeze, the liquid is depleted of molybdenum and enriched in titanium. The equilibrium freezing ranges, DTf, for these binary alloys (Table 1) are relatively narrow. The liquidus temperature, TL, for a Ti-5%Mo alloy would be approximately 1700 C (3100 F) and for a Ti-30%Mo alloy would be approximately 1870 C (3400 F) (Table 1) (Ref 1). At equilibrium, titanium-molybdenum alloys would freeze as uniform single-phase b-titanium. According to the Gulliver-Scheil assumptions, all the alloys would be cored b-titanium with G-S , at the maximum molybdenum content, Cmax the dendrite centers and pure titanium in G -S the interdendritic boundaries ðCmin ¼ 0%MoÞ. The effective solidus temperature, according to the Gulliver-Scheil conditions, would be the melting point of pure titanium for all titanium-molybdenum alloys. Figure 1 shows the microstructure of an as-cast Ti-15%Mo alloy polished and etched to reveal the coring. Above the micrograph is a plot of the molybdenum content measured by an electron microprobe line scan. As expected, the maximum molybdenum content is at the dendrite centers, and the minimum molybdenum content is at the interdendritic meas boundaries. Table 1 shows the values of Cmax meas and Cmin measured after cooling to room temperature for Al-5, 10, 15, and 30% Mo alloys arc melted and drop cast into copper chill molds. Table 1 includes the segregation ratio computed from the listed data and a residual segregation index (SI): SI ¼
meas meas Cmax Cmin G -S G S Cmax Lmin
(Eq 5)
where the superscript “meas” indicates a measured composition, and the superscript “G-S” indicates a composition computed from the Gulliver-Scheil relation. Both measures of the extent of segregation in titanium-molybdenum alloys are compared to the estimated value of a, the mass transport Fourier number, which is based on measured solidification times, and
Table 1 Solidification and microsegregation data for isomorphous titanium-molybdenum alloy castings DTf
TL Mo, wt%
C
F
C
F
-S CG max , wt%
5
1701 3094
10 18
7
10
1734 3153
19 34
14
15
1767 3213
29 52
21
30
1866 3391
57 103
52
Source: Ref 1
Cmeas max , wt%
Cmeas min , wt%
S, Cmax Cmin
meas meas SI, Cmax Cmin G-S -S CG max Cmin
uf, s
das, mm
5.4 5.3 11.5 11.1 17.4 16.8 35.6 34.6
4.0 4.2 8.5 8.8 12 12.5 22.8 24.6
1.35 1.26 1.36 1.27 1.45 1.34 1.56 1.41
0.20 0.16 0.21 0.17 0.26 0.20 0.30 0.24
0.9 29 1 31 1.1 33 7.6 131
28 130 33 140 26 128 32 128
a,
4Ds uf das2
0.06 0.09 0.06 0.09 0.03 0.04 0.03 0.035
340 / Principles of Solidification dendrite arm spacing and an estimate of DS at the liquidus temperature of the alloy. The qualitative trend is excellent. Both indices show a decrease in extent of microsegregation with increasing values of a. The trend is slightly stronger with the segregation index (Ref 1).
Binary Eutectic Systems, k < 1: Hypoeutectic Aluminum-Copper Alloys
the composition of the solid changes continuously to be equal to the interface composition (indicated by the dotted curve in Fig. 3), and the final distribution of solute (between the equilibrium solidus and equilibrium solvus temperatures) is uniformly CO, the same as the initial melt composition. For the Gulliver-Scheil conditions, there is no redistribution of solute behind the solid-liquid interface, and the
composition that forms in the solid at the interface remains unchanged as the region cools to the eutectic temperature, where solidification is complete, and continues unchanged as the region cools to the ambient temperature. Nonequilibrium Solidification with Back Diffusion—Directional Solidification. A mass transport Fourier number of a = 0.5 represents well the measured behavior for a
660 1200
1150
Eq
620
600 1100 a = 0.5
Temperature, ⬚F
Temperature, ⬚C
640
580 G-S 1050 560
540 0.00
0.10
0.20
0.30
0.40
0.50
0.60
0.70
0.80
0.90
1000 1.00
Fraction solid
Fig. 2
Solidification curves (temperature vs. weight fraction solid) for three cases of Al-4.5%Cu alloy: equilibrium (Eq: a ! 1), Gulliver-Scheil (G-S; a ! 0), and directional solidification with a = 0.5
6.00
5.00
Cu concentration, wt%
Two convenient ways to represent the extent of microsegregation are the solidification curve, the progress of solidification with temperature in the freezing range, and the solute distribution curve, the fraction of solid within the characteristic volume enclosed by a range of isoconcentration surfaces. These curves are illustrated in Fig. 2 and 3 for a model binary eutectic system, aluminum-rich aluminum-copper alloys, in which much of the initial research on alloy solidification has been reported (Ref 3, 8–11). The melting temperature of aluminum is 660.2 C (1220.4 F); the eutectic temperature is 547.8 C (1018.8 F) for the invariant reaction: liquid (33.2% Cu) ! a-aluminum (5.65% Cu) + y-phase (CuAl2); and the a-phase solvus recedes by approximately 1% Cu per 50 C (90 F) below the eutectic. Over the freezing range for hypoeutectic aluminum-copper alloys, 0.14 < k < 0.17. In practice, an Al-4.5%Cu alloy would not solidify completely until it cools to its eutectic temperature. The eutectic microconstituent, a-aluminum + y-CuAl2, would form in the last regions to freeze locally. The primary solidification phase would be cored—a gradient in copper content will be apparent on polished and etched surfaces. Solidification Curve. The progress of solidification with temperature in the freezing range, the solidification curve, for an Al-4.5%Cu alloy is compared for the two extremes in Fig. 2. For equilibrium freezing, the alloy is 100% solid at the equilibrium solidus. For the Gulliver-Scheil conditions, the Al-4.5%Cu alloy is 88% solid (by weight) at the equilibrium solidus temperature and 91% solid at the eutectic temperature. The liquid remaining at the eutectic temperature transforms to the eutectic microconstituent, 9% by weight. The freezing range is extended (and the incipient melting point is lowered) by 25 C (45 F). The solidification curve is shallow near the liquidus temperature and becomes progressively steeper as it approaches the eutectic temperature. Solute Distribution Curve. The solute distribution curve (solute profile) is a plot of the solute content versus the fraction of solid within a characteristic volume that would be enclosed by the isoconcentration surface at the composition indicated on the ordinate. Figure 3 compares, for the two extremes, the copper concentration in the solid phase(s) at the solid/ liquid interface as a function of the cumulative fraction of solid. For equilibrium freezing,
Eq-final
4.00
G-S
3.00
a = 0.5
2.00
Eq-intf
1.00
0.00 0.00
0.10
0.20
0.30
0.40
0.50
0.60
0.70
0.80
0.90
1.00
Fraction solid
Fig. 3
Solute distribution curve (weight percent copper vs. weight fraction solid contained) for three cases of Al-4.5% Cu alloy: equilibrium (Eq: a ! 1), Gulliver-Scheil (G-S; a ! 0), and directional solidification with a = 0.5. For equilibrium, the final as-cast solute distribution (Eq-final) and the solid composition at each stage of solidification (Eq-intf) are shown.
Microsegregation / 341 directionally solidified experimental casting of Al-4.5%Cu alloy (Ref 9,10). The solidification curve (Fig. 2) and as-cast copper distribution (Fig. 3) indicate that the freezing range still extends from the equilibrium liquidus to the eutectic temperature, where approximately 8 wt% of eutectic microconstituent will form. The minimum copper content is approximately 1.35% Cu, and the copper profile increases slowly across the first 40% of the primary phase. The maximum copper content in the primary phase is still 5.65%. As a result of back diffusion, the average copper content of the primary a-phase is 2% as compared to 1.65% for the Gulliver-Scheil conditions. Most of the copper segregates to the last region to freeze locally, the eutectic microconstituent, which has an average composition of 33.2% Cu. The segregation ratio, based on the primary phase, nsq will be S 5:65=1:35 ¼ 4:2 and fsut 8% for a = 0.5, as opposed to the extremes: nsq S 5:65=0:69 ¼ 8:2 and fsut 9% for Gullinsq ¼ 0% ver-Scheil conditions and S = 1 and fsut for equilibrium. Note that in all cases, the rate of solidification with temperature is very high near the liquidus temperature; that is, the initial slope of the solidification curve, @T =@fS , is shallow. For the Al-4.5%Cu alloy, 50% of the solid forms within the first 10 of the freezing range. Near the end of primary phase formation, the solidification curve is steep; the fraction solid changes slowly as the remaining liquid and now coherent solid coexist over the majority of the temperature range of freezing. For nonequilibrium solidification, the steep portion of the solidification curve leads to the isothermal transformation of the eutectic liquid to the eutectic mixture at the eutectic temperature. Concomitantly, the initial slope of the solute distribution curve, @C=@fS , is shallow, becoming progressively steeper as the eutectic composition is approached. Influence of Local Cooling Rate. The mass transport Fourier number, a, and the extent of back diffusion depend on the local solidification time and the reciprocal of the square of the characteristic length. The as-cast dendrite arm spacing has been found to depend on the local solidification time to a power, n, between 0.3 and 0.5. Using n = 0.39, the exponent found 1=5 for aluminum-copper alloys is a / yf . The extent of microsegregation decreases, but very slowly, as the local solidification time increases, and cooling rate decreases. Table 2 illustrates
Table 2
the decrease in segregation ratio and fraction eutectic as the local solidification time varies from a local solidification time and dendrite spacing characteristic of a weld (a = 0.1), to permanent mold and sand castings, to extremely low cooling rates. Over this wide range of processing conditions, the variation in microsegregation is minor. Slow cooling and long solidification times are not effective in reducing microsegregation. Conversely, short solidification times, to obtain a fine microstructure, followed by slow cooling in the mold or by solution heat treatment are more effective in removing microsegregation. Dendrite Coarsening. Dendrite (and grain) coarsening that involves remelting can be another source for solute redistribution in solid phases during solidification. The apparent dendrite arm spacing increases during the time solid and liquid phases coexist. Several mechanisms for dendrite coarsening have been proposed and observed (Ref 15, 16, 18, 23–26). Those coarsening mechanisms that involve remelting of solid phases redistribute the solute. As an example, for Al-4.5%Cu consider that an existing thin dendrite arm with high curvature melts, and a neighboring thicker dendrite arm with lower curvature thickens (as a result of diffusion of copper through the liquid from the thicker to the thinner arm). The solid being deposited on the thickening arm would have a higher copper content than the disappearing dendrite arm (that formed at higher temperatures). Consequently, the average copper content of the primary phase will increase, and the fraction of nonequilibrium eutectic will be reduced. For Al-4.5%Cu alloys, the extent of solute redistribution by solid diffusion or by coarsening is low during the initial stages of freezing. The copper content builds up slowly from kCO (note Fig. 3). The gradient for diffusion is small, as would be the difference in copper content between melting and solidifying dendrite arms. During the later stages of primary phase solidification, the copper gradient is higher. Solute redistribution by solid diffusion becomes appreciable. Coarsening involving melting of existing, cored dendrite arms would reduce microsegregation appreciably. On the other hand, studies of dendrite coarsening mechanisms indicate remelting mechanisms are more likely in the early stages of solidification (Ref 26). Certainly, in low-melting systems (e.g., tin and lead alloys), diffusion in the solid is
Variation in microsegregation with change in solidification time for Al-4. 5%Cu
Process
Weld Chill Permanent mold Sand mold Directional
a
0.1 0.15 0.2 0.3 0.5
uf, s
1.2 7.4 27 172 1756
das, mm
Cmin, wt% Cu
S
feut, wt%
8 16 27 56 138
0.93 1.00 1.05 1.15 1.29
6.1 5.7 5.4 4.9 4.4
8.7 8.5 8.3 8.0 7.5
negligible, and deviation from Gulliver-Scheil behavior is due to dendrite coarsening. For higher-melting alloy systems, higher solute diffusivities and the mass transport Fourier numbers indicate solute redistribution is predominantly by back diffusion. Influence of Microsegregation and Partitioning on Density and Solidification Shrinkage. As the compositions of the solid and liquid phases change as a result of partitioning, the relative densities of the phases can change significantly. The alloy Al-4.5%Cu provides an example where substantial density changes due to solute partitioning strongly influence castability. As an Al-4.5%Cu alloy cools through the freezing range, the copper content of the liquid increases from 4.5 to 33.2%, while the average composition of the primary solid phase only increases from 0.69 to 2% Cu. The first solid to form is denser than the initial liquid. After only a few percent of solidification, the copper-enriched liquid is denser than the primary solid a-phase. Equiaxed dendrites and dendrite fragments float in the copper-enriched liquid. Solidification expansion accompanies a large fraction of the formation of the primary a-phase. Importantly, though, the final solidification of eutectic liquid at 33.2% Cu to the eutectic mixture (a + y) is accompanied by a 10% decrease in specific volume. A large fraction of the overall solidification shrinkage of the Al-4.5%Cu alloy occurs at the last-stage of solidification. This tendency for late-stage solidification shrinkage makes Al-4.5%Cu alloys susceptible to microshrinkage porosity and to hot tearing. Interaction of Microsegregation and Macrosegregation. Macrosegregation is the difference in the local average solute content from the initial alloy after solidification, C, composition, CO. Macrosegregation is evident in the as-cast microstructure (and in properties) by the way it changes microsegregation. Using Al-4.5%Cu alloy as an example, the main manifestation of macrosegregation in the as-cast microstructure will be the amount of nonequilibrium eutectic constituent (Ref 27). For positive macrosegregation, C > CO , there will be slightly less primary phase formation and noticeably more nonequilibrium eutectic. The increased eutectic will likely be accompanied by increased microporosity. Standard solution treatment will leave more undissolved y-phase. Without corrective treatment, regions of Al-4.5%Cu alloys with positive macrosegregation can be expected to have reduced ductility. For negative macrosegregation, C < Co , there will be slightly more primary phase formation and noticeably less nonequilibrium eutectic. Regions of Al-4.5%Cu alloys with negative macrosegregation can be expected to have reduced strength. Microsegregation influences the relative density of solid and liquid phases, the driving force for interdendritic fluid flow, and the extent of macrosegregation. Interdendritic flow of solute
342 / Principles of Solidification (Figure 1 is an example for k > 1.) Then, a second, perpendicular trace, crossing at the first minimum, can be made to check for a lower minimum; another perpendicular trace can be made until the extreme minimum in that plane of polish is located (Ref 10). Quantitative metallography can be used to measure the volume fraction of secondary phases. That measurement is converted to a weight fraction, and the difference between the measured value and the equilibrium value is reported as f nsq. For example, y-phase is a nonequilibrium constituent in Al-4.5%Cu. All measured y-phase can be reported as fnsq , or the measured value can be divided by the mass fraction of y-phase in the eutectic microconstineq (Ref 8). tuent to report fsut An array of electron probe point counts can be used to statistically determine solute distribution curves (such as the copper profile in Fig. 3) (Ref 32–36). Composition is measured across a rectangular grid of n m = P points, taking care that the spacing between points is not a multiple of the dendrite arm spacing. For a binary alloy with k < 1, such as Al4.5%Cu, the measured points can be sorted and indexed (j) in order of ascending solute content. Measured solute content, Cij, is plotted versus fS ¼ j=P , the ratio of the index of the sorted point, j, to the total number of points, P. The reproducibility and accuracy of this statistical analysis can be tested by comparing the average composition of the measured points to CO and by separating the points at random into two groups, sorting, plotting the two groups, and comparing the solute profiles obtained from one-half of the points and the profile using all the points. If agreement is not satisfactory, more points should be added to the grid. Uncertainty in the composition measurements around
Cmin and Cmax may distort the shape of the solute profile. Also, noting that the measured volume may contain more than one phase, composition measurements that fall outside of the stability range of the phases can be prorated. Generally, secondary phases will have a fixed composition, such as y-phase in hypoeutectic aluminum-copper alloys. Figure 4 presents, as an example, solute (zinc) distribution curves measured for the binary alloy Cu-33%Zn (Ref 33). Figure 5 is an example of solute (copper and silicon) distribution curves for multicomponent and multiphase Al-Si-Cu-Mg, where the solute concentrations in the primary a-phase dendrites were computed by assuming the silicon-phase and y-phase have fixed compositions (Ref 36). Fig. 4 and Fig. 5 are discussed in later sections. Observation and Measurement of Kinetic Processes Leading to Microsegregation. The kinetics of solute redistribution processes can be examined more effectively by interrupted solidification and/or in situ solidification. Because of dendrite (or grain) coarsening, diffusion in the solid, and bulk flow of solid and liquid, the solidification morphology and solute distribution in the solid phases change continuously during solidification and can continue to change during postsolidification cooling. Measurements and images made after continuous cooling through the solidification range and to room temperature do not provide unambiguous information on the dynamic evolution of microstructure and solute distribution through the solidification range. Interrupted solidification refers to a variety of techniques to discontinuously increase the cooling rate of an alloy that has cooled only part way through the solidification range. In principal, the microstructure present before
35 34.5 34 Zn concentration, wt%
enriched (or depleted liquid) in the mushy zone that results from the requirement to feed solidification shrinkage and from thermal and solutal convection due to density gradients is the dominant factor in determining the extent of macrosegregation. For an Al-4.5%Cu alloy, the density of the liquid in the mushy zone increases as the local temperature decreases, the fraction of solid increases, and the liquid enriches in copper. The densest liquid is at the eutectic isotherm. Copper-enriched interdendritic liquid tends to settle, and copper-depleted solid grains and fragments tend to float. In contrast, for the aluminum-silicon alloys discussed subsequently, the density of the liquid in the mushy zone decreases as the liquid enriches in silicon. Solutal convection will be the opposite as for aluminum-copper alloys. In both cases, horizontal temperature gradients will give rise to solutal convection and the potential for macrosegregation. One method to minimize the extent of macrosegregation is to balance the alloy composition to minimize density changes as the solute is partitioned between solid and liquid phases during solidification. Measurement of As-Cast Microsegregation. Qualitatively, the extent and distribution of microsegregation will be evident in cut and polished surfaces. For optical and scanning electron microscope images, an etchant may be used to bring out the microstructure. Coring in the primary phase will be revealed by a gradation in color or texture after etching. The likely positions of relative composition maxima and minima can be inferred from the microstructures. Differential heat treatment can be used to emphasize composition differences. As an example, consider hypoeutectic aluminumcopper alloys. The alloy can be solutionized at a temperature below the equilibrium solvus and then quenched. The dendrite cores will appear single phase (a), and the regions above the solvus composition for the solutionizing temperature will precipitate y-phase. The boundary between single and two-phase regions will mark an isoconcentration surface. Changing the solutionizing temperature will indicate the position of a range of isoconcentration surfaces. Serial sectioning can be used to determine the shape of isoconcentration surfaces and as-cast dendrite morphology in three dimensions (Ref 28–30). Differential heat treatment is an especially useful technique to bring out microsegregation patterns in systems that undergo polymorphic transformations that mask microsegregation, for example, steels and titanium alloys (Ref 21, 31). The segregation ratio, S, the segregation index, SI, and the fraction of nonequilibrium microconstituent, f nsq, have already been mentioned as quantitative indicators of the extent of as-cast microsegregation. Electron microprobe traces can be used for quantitative analysis of Cmin and Cmax. For example, an electron beam trace (line scan) can be made across a dendrite arm, crossing an area where a minimum in solute concentration is expected.
Columnar
33.5
Equiaxed
33 32.5 32 31.5 31 30.5 30
0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1
Cumulative fraction of measured points
Fig. 4
Measured solute distribution curves for Cu-33%Zn alloy, for two directionally solidified ingots with the same thermal history but different grain structures. Adapted from Ref 33
Microsegregation / 343 Measured Cu, Si distribution of alloy 1 (as-cast, 0.10 ⬚C/s or 32 ⬚F/s) 4 % Cumin = 0.8 Cu and Si in a-phase, wt%
3.5
gprimary = 78% gbinary = 12%
3
geutectic = 10% ga = 92.7%
2.5
Cu
gSi = 5.6%
2
Fe
gq = 1.7%
1.5 Si 1 0.5
Mg
0 0
Fig. 5
0.2
0.6 0.4 Cumulative fraction of measured points
0.8
1
Silicon and copper solute distribution curves measured by electron microprobe grid in a five-component aluminum alloy, Al-6.6%Si-3.4%Cu-0.2%Mg-0.5%Fe. Source: Ref 36
Fig. 6
Al-7%Si-3.5%Cu-0.3%Mg-0.5%Fe alloy slowly cooled to two temperatures in the solidification range and then quenched. (a) Aluminum-rich dendrites and enriched liquid are present at 571 C (1060 F). (b) Al5FeSi intermetallic plates start to form at 564 C (1047 F). Colonies of aluminum and silicon eutectic will form at a slightly lower temperature, followed by Q-phase and y-phase. Source: Ref 37
solidification is interrupted will be coarse and easily contrasted from the fine microstructure that will form in regions that were liquid immediately before the quench. In one variation of interrupted solidification, several small crucibles containing molten alloy of the same composition are continuously cooled at a controlled, relatively slow rate; each crucible is extracted from the furnace at a different temperature (and fraction solid) within the freezing range, and the partly solidified alloy is quenched. When the samples are cut, polished, and etched, the regions that were solid before the quench will have a relatively coarse microstructure and be in contrast with the
regions that were liquid at the time the samples were quenched (Ref 10). The general morphology of the solid at the time and temperature of the quench, the phases present, and the composition profile within the preexisting solid will be apparent and measureable. For example, the microstructure of a multicomponent aluminum alloy slowly cooled and then quenched from a temperature within the freezing range is presented in Fig. 6 (which is discussed in a later section) (Ref 10, 37). Care must be exercised in analyzing the microstructure in interrupted solidification samples. There will be some delay after the quench before the solid growing at the preexisting
interface breaks down into fine dendrites. The fraction solid at the temperature of the quench will be overestimated. Further, the rapid change in solidification rate will lead to inverse macrosegregation in the fine structure at the preexisting solid-liquid interface and give the false appearance of a solute boundary layer. On the other hand, measurements of the solute content in the coarse solid phase will reliably represent the solute concentration in the solid at the time of the quench. For example, electron beam microprobe measurements of the minimum solute content in the a-aluminum dendrites of Al-4.5%Cu alloy samples quenched at different stages of solidification increase from kCO 0.65 to 1.35% Cu. Thermal analyses and theoretical estimates of the solute buildup at dendrite tips indicate undercooling less than 1 C (1.8 F), which would increase the copper content of the initial solid by only 0.05%. The reduction in as-cast microsegregation from the GulliverScheil extreme can be attributed primarily to back diffusion (Ref 9, 10). In situ solidification refers to techniques to image and measure solidification morphology and solute distribution in alloys as they evolve during solidification. Microradiography makes use of solute partitioning to provide contrast and make optically opaque alloys sensitive to dynamic processes. The microsegregation in thin sections of as-cast alloys, such as Al-4.5%Cu, is readily apparent in radiographic images wherein the wavelength of incident x-rays is selected to maximize the difference in aluminum and copper absorption coefficients. If thin sections are solidified in front of an x-ray beam and sequential images captured, the liquid and solid phases will be distinguished as the liquid progressively is enriched in copper, the more absorbing element (Ref 26, 38–41). Unlike imaging of as-cast samples, for dynamic images, exposure time must be short to obtain a “stop-action snapshot.” Wavelength, coherency, and intensity of the incident beam must be optimized to obtain images with good contrast and brief exposure times. Figure 7 is an example of an in situ image of equiaxed dendritic grains in a thin section (0.1 mm, or 0.004 in.) of an Al-30%Cu alloy solidifying under a shallow temperature gradient (Ref 39). In successive images, these dendrites are seen to float on the copper-enriched liquid, rotate, and impinge, hence isolating liquid pools at what will become grain-boundary junctions. Figure 8, by contrast, is an in situ image of columnar dendrites of Sn-13%Bi solidifying under a steep temperature gradient (Ref 42). The composition of the interdendritic liquid changes from CO = 13% Bi at the dendrite tips (top) to CE = 52% Bi at the dendrite roots (bottom). For this low-temperature alloy, the secondary dendrite arms appear to separate the alloy into isolated characteristic volumes where solidification proceeds in accordance with the Gulliver-Scheil model.
344 / Principles of Solidification
Fig. 7
Synchrotron microradiograph of Al-25%Cu alloy solidifying under shallow temperature gradient. Equiaxed aluminum-rich dendrites (light) are floating in copper-enriched liquid (dark). Source: Ref 39
Fig.
8
Microradiograph of Sn-13%Bi alloy directionally solidifying in a steep temperature gradient. Liquid (dark) at top is at the initial melt composition (13% Bi), and remaining liquid at bottom is at eutectic composition (51% Bi). Composition of the columnar dendrites varies from 3.5 to 21% Bi. Source Ref 42
Thermal analysis of aluminum-copper alloys indicates the freezing range is consistent with nonequilibrium solidification, and the cooling curve resembles the solidification curves in Fig. 2. Conversely, the solidification curve, if known, can be used to compute the rate of heat evolution through the freezing range. Just as the solidification curves for Gulliver-Scheil and nonequilibrium freezing with back diffusion are very similar, the enthalpyversus-temperature curves for the two solidification paths will not differ appreciably.
Binary Eutectic Systems: Hypoeutectic Aluminum-Silicon Alloys Microsegregation is discussed in terms of the behavior of another model binary eutectic system, aluminum-rich aluminum-silicon alloys.
The eutectic temperature is 577 C (1071 F) for the invariant reaction: liquid (12.2% Si) ! a-aluminum (1.65% Si, a-phase) + siliconphase (essentially pure silicon); the a-phase solvus recedes by approximately 1% Si per 50 C (90 F) below the eutectic temperature. The partition ratio for aluminum-silicon alloys is approximately the same as for aluminum-copper alloys. The diffusivity of silicon in a-aluminum is approximately 2½ times greater than the diffusivity of copper when compared at the same temperatures. Silicon is an important component of most aluminum-base casting alloys, such as Al-356, because it improves castability by increasing fluidity and resistance to hot tearing and decreasing the tendency for shrinkage microporosity. In commercial casting alloys, specified silicon content exceeds the solubility limit at the eutectic temperature, and the a-aluminum + silicon eutectic is an equilibrium microconstituent of the as-cast microstructure. Consider, for example, an Al-7%Si alloy. Approximately 50% of the a-aluminum primary phase solidifies above the eutectic temperature, and the remaining 50% of solidification is the eutectic microconstituent (secondary a-aluminum + silicon). The secondary a-aluminum forms at the solubility limit, 1.65% Si. The silicon content at the center of the dendrite arms, Cmin, increases by back diffusion from approximately 0.9% Si at the liquidus temperature to 1.15% Si at the eutectic temperature. At the eutectic temperansq 2%, and ture, S 1:65=1:15 ¼ 1:4; fsut nsq fSi 0:2%. Anomalous Microsegregation. As an Al7%Si alloy cools below the eutectic temperature, the solubility of silicon in the a-phase decreases rapidly. Within the interdendritic regions, silicon diffuses out of the a-phase and increases the volume fraction of the siliconphase particles. The direction of the silicon composition gradient reverses, and silicon diffuses away from the dendrite cores toward the interdendritic regions. The slower the postsolidification cooling rate, relative to the solidification rate, the more the a-phase will be depleted in silicon. The maximum silicon content in the a-phase will be in the center of the dendrite, and the minimum silicon content will be in the interdendritic a-phase. The silicon microsegregation profile will be the opposite of that observed in aluminum-copper alloys. This solute distribution has been called anomalous microsegregation (Ref 43).
Binary Peritectic Systems: Copper-Zinc a-Brass Alpha-brass, Cu-30–33%Zn, is an example of a hypoperitectic system in which solute redistribution and as-cast microsegregation are between the extremes of Gulliver-Scheil and equilibrium behavior. The melting temperature of copper is 1083 C (1981 F); the peritectic
temperature is 903 C (1657 F) for the invariant reaction: liquid (38% Zn) + a-brass (32.5% Zn) ! b-brass (36.9% Zn); the solubility limit for a-phase increases from 32.5% Zn at the peritectic temperature to 39% Zn at room temperature. A Cu-33%Zn alloy will begin to freeze as a-brass at the liquidus temperature with an initial composition of kCo 28.3% Zn. For equilibrium solidification, the alloy would be 91% a-brass and 9% liquid at the peritectic temperature and transform to 89% a-brass + 11% b-brass right below the peritectic temperature. As the alloy continues to cool toward room temperature, the b-brass will transform to a. For equilibrium solidification, the segregation ratio is S = 1 and fbnsq ¼ 0 in the as-cast microstructure. At the other extreme, Gulliver-Scheil solidification, the alloy would be 62% a-brass and 38% liquid at the peritectic temperature. Right below the peritectic temperature, b-brass would nucleate and surround the a-brass. The remaining liquid would transform to increasingly zinc-rich b-brass that would remain in the interdendritic region of the as-cast microstructure. For Gulliver-Scheil solidification, the segregation ratio is S 57=28:3 ¼ 2 and fbneq ¼ 38%. The diffusivity of zinc in a-brass near the Cu-33%Zn liquidus temperature is approximately 7 times the diffusivity of copper in a-aluminum around the liquidus temperature for Al-4.5%Cu. Back diffusion is prevalent in a-brass. Right above the peritectic temperature, approximately 22 wt% liquid will remain. As the alloy cools below the peritectic temperature, the b-phase will nucleate and surround the a-brass. The peritectic reaction (L + a ! b) occurs in three steps: (1) L ! b at the b/L interface, (2) zinc diffuses through the b layer from the b/L to the a/b interface, and (3) a ! b at the a/b interface. Diffusion through the b layer (step 2) is the rate-controlling step for the peritectic reaction. Eventually, at the nonequilibrium solidus, the final zinc-enriched liquid transforms to zinc-rich b-brass. For a local solidification time of 1500 s, the segregation ratio is S 44=31:3 ¼ 1:42 and fbnsq 8% (Ref 33). The influence of grain structure on microsegregation has been investigated for a-brass. Figure 4 compares for columnar and equiaxed grain structures the zinc distributions measured by point counting on 13 by 13 grids and plotting the points in ascending order of measured zinc content. The columnar grain structure was obtained by directional solidification. The equiaxed structure was obtained by adding 0.75% Fe to the liquid before directional solidfication. The equiaxed grain size is approximately the same as the secondary dendrite arm spacing in the columnar-grained casting. The locations compared in Fig. 4 have essentially the same thermal history. The difference in zinc profile for the equiaxed and columnar structures is within the precision of the measurement technique. The dependence in a-brass of the extent of microsegregation on grain structure is weak (Ref 33).
Microsegregation / 345
Multicomponent, Multiphase Eutectic Systems: Hypoeutectic Al-Si-Cu-Mg Alloys Common heat treatable aluminum casting alloys are strengthened by precipitation of the precursor phases to Mg2Si (b0 ) and CuAl2 (y0 and y00 ). Al-6.5%Si-3.5%Cu-0.5%Mg-0.5%Fe is representative of Al-319 alloys, which are used for automotive structural applications that require good castability, good heat resistance, good machinability, and tolerate low ductility. The solidification sequence is L ! a-aluminum (from liquidus 600 C, or 1110 F), L ! aaluminum + silicon-phase binary eutectic (from 566 C, or 1050 F), L ! a-aluminum + silicon-phase + Q-phase ternary eutectic (from 524 C or 975 F), L ! a-aluminum + silicon-phase + Q-phase + y-phase quaternary eutectic (at 505 C, or 940 F). Figure 6(b) is a micrograph of the microstructure of an Al-6.5%Si-3.5%Cu-0.5%Mg-0.5%Fe alloy that was slowly cooled to 564 C (1047 F) and then water quenched. The coarse a-dendrites and Al5FeSi platelets formed before the quench. The fine-structured quenched liquid is enriched in copper, silicon, magnesium, and iron. Divorced eutectic colonies of a and silicon, Qphase and y-phase will form at lower temperatures. Analyses of the interrupted solidification microstructures and of the compositions of the quenched solid provide an understanding of the kinetics of the solidification processes (Ref 36, 37, 43, 44). Note that the liquidus and eutectic temperatures of the multicomponent alloy are considerably lower than for the binary alloys, which tends to lower the diffusivity of the alloying elements in the solid phases. On the other hand, the presence of less than ½% Mg more than doubles the diffusivity of both copper and silicon in the primary a-phase. Also, the diffusion of copper in the binary eutectic a-aluminum + silicon-phase region is increased 5 to 20-fold over the diffusivity in the single-phase region. During primary phase solidification, the partition ratios for silicon, copper, and magnesium are less than 1 (Ref 44, 45). Copper redistribution during primary phase solidification approaches the Gulliver-Scheil conditions, and silicon redistribution approaches equilibrium. When the silicon nucleates, the copper content of the remaining liquid and of the forming a-aluminum continues to increase. The silicon concentration of the liquid and of the forming a-phase decreases slightly. Silicon begins to diffuse from the primary phase at the center of the dendrites into the two-phase region. After ternary and quaternary eutectic solidification, and as the alloy cools below the eutectic composition, the solubilities of copper and silicon in the a-phase decrease sharply. Copper and silicon diffuse away from the a-aluminum + silicon-phase toward the interdendritic region. The interdendritic y and silicon-phase particles act as sinks
for diffusing copper and silicon. Figure 5 presents the copper and silicon profiles measured by electron microprobe in the a-phase of a sample from a sand casting, allowed to cool in the mold. The silicon profile within the a-phase is at a maximum at the center of the dendrite and at a minimum in the interdendritic regions, that is, anomalous microsegregation. The copper profile within the a-phase is a minimum at the center of the dendrites and a maximum near, but not at, the interdendritic region. Knowledge of the as-cast microsegregation in this alloy can be used to optimize the solution heat treatment prior to quenching and aging. A brief solution treatment will return the silicon distribution in the a-phase to the level present at the eutectic temperature, over 1%. Magnesium and copper in the core of the dendrites will be at a low level. The predominant precipitate on aging will be b0 . After moderate solution treatment, magnesium will diffuse into the dendrite cores, and b00 precipitates will be prevalent. For longer solutionizing times, copper will diffuse into the dendrite cores, and y0 will contribute to strengthening. Brief solutionizing times, or rapid cooling from just below the eutectic temperature, can be used to achieve adequate hardness and strength and save energy. Extended solution treatment can be used to substantially increase strength and ductility if required (Ref 36).
Multicomponent Multiphase Steel: Fe-C-Cr Alloying elements in steel are both substitutional and interstitial solutes. Interstitial elements, such as carbon, are highly mobile, especially in ferrite, d-iron. Redistribution of carbon during solidification can be expected to attain equilibrium. Substitutional elements, such as chromium, are much less mobile, and their diffusivities depend strongly on the temperature range of solidification. In the ironchromium binary system, the freezing range is very narrow; the liquidus temperature is close to the melting temperature of iron, and k approaches 1. In both iron-carbon and ironchromium binary alloys, microsegregation will be negligible, and solidification will follow equilibrium conditions: S = 1 and fbneq ¼ 0. Fe-1.5%Cr-1%C is the nominal composition for 52100, a wear-resistant, hardenable bearing steel. The liquidus temperature is reduced by 100 C (180 F) in the ternary alloy, and the freezing range is extended. Diffusivity of chromium in the ternary alloy is reduced considerably from the binary alloy. In the ternary alloy, redistribution of carbon attains equilibrium (uniform activity but not uniform composition), and redistribution of chromium by diffusion is negligible. The dendrites are cored with respect to both carbon and chromium; the freezing range is extended to the line of twofold saturation where carbides form. Carbon diffuses
to attain uniform chemical potential. The presence of chromium alters the activity of carbon, and the nonuniform distribution of chromium causes the carbon to diffuse to a nonuniform distribution. In directionally solidified columnar grain ingot, SCr 4; in small, slowly cooled, equiaxed ingot, SCr 3; and in commercial wrought 52100 alloy, SCr 2 (Ref 31). Solidification sequence and residual microsegregation in austenitic stainless steels depend, sensitively, on composition and solidification rate. Retention of a small fraction of ferrite is considered to reduce the susceptibility of stainless steel castings and weldments to the ratio of nickel (and other austenite stabilizers) and chromium (and other ferrite stabilizers) to hot tearing and crater cracking. Figure 9 presents detailed measurements of the isoconcentration lines in an Fe-25%Cr-20%Ni alloy that illustrate the as-cast dendritic morphology (Ref 30, 46). Measurements such as these are used to select geometry for simulation models of solute redistribution during dendritic solidification (Ref 19).
Multicomponent, Multiphase Nickel-Base Superalloy Nickel-base superalloys, such as Mar-M-200, contain several alloying elements, some, such as titanium and aluminum, with k < 1 and some, such as molybdenum, with k > 1. Alloying elements such as titanium and aluminum tend to decrease the density of the liquid during
Fig. 9
Isoconcentration profile in an Fe-25%Cr-20%Ni columnar dendrite. Source: Ref 45, 46
346 / Principles of Solidification solidification, and others, such as tungsten, tend to increase the liquid density. Hafnium is added to directionally solidified columnar-grained Mar-M-200 airfoil castings to modify the ductility of transverse grain boundaries during elevated-temperature service. For hafnium, k 1, and its addition extends the freezing range and significantly increases eutectic fraction. The increased eutectic liquid amount enhances the castability of columnar-structured Mar-M-200 + hafnium by decreasing the tendency to hot tearing (in the last stages of solidification). On the other hand, the increase in the volume and density of the eutectic liquid increases the susceptibility of the hafnium-modified alloy to macrosegregation (Ref 47). REFERENCES 1. J.I. Nurminen, “Solidification Structure and Segregation in Binary Titanium Base Alloys,” Ph.D. dissertation, University of Pittsburgh, 1972. 2. B. Chalmers, Principles of Solidification, Wiley, New York, 1964 3. M.C. Flemings, Solidification Processing, McGraw-Hill Book Co., New York, 1974 4. W. Kurz and D.J. Fisher, Fundamentals of Solidification, Trans Tech Publications Ltd., 1986 5. G.H. Gulliver, The Quantitative Effect of Rapid Cooling upon the Constitution of Binary Alloys, J. Inst. Met., Vol 9, 1913, p 120–157 6. E. Scheil, Bemerkungen Zur Schictkristallbildung, Z. Metallk., Vol 34, 1942, p 70–72 7. M.E. Glicksman and R.N. Hills, Origins of Non-Equilibrium Segregation Theory, Proc. M.C. Flemings Symp. on Solidification and Materials Proc., R. Abbaschian, H.D. Brody, and A. Mortensen, Ed., TMS, 2001, p 81–85 8. A.B. Michael and M.B. Bever, Solidification of Aluminum Rich Aluminum-Copper Alloys, Trans. Met. Soc.-AIME, Vol 200, 1956, p 47 9. H.D. Brody and M.C. Flemings, Solute Redistribution in Dendritic Solidification, Trans. Met. Soc.-AIME, Vol 236, 1966, p 615–624 10. T.F. Bower, H.D. Brody, and M.C. Flemings, Measurements of Solute Redistribution in Dendritic Solidification, Trans. Met. Soc.-AIME, Vol 236, 1966, p 624–634 11. D.H. Kirkwood, Microsegregation, Mater. Sci. Eng., Vol 9 (No. 6), 1984, p 101–109 12. T.W. Clyne and W. Kurz, Solute Redistribution during Solidification with Rapid Solid State Diffusion, Metall. Mater. Trans., Vol 12 (No. 6), 1981, p 965–971 13. J.A. Warren and W.J. Boettinger, Prediction of Dendritic Growth and Microsegregation Patterns in a Binary Alloy Using the Phase-Field Method, Acta Metall. Mater., Vol 43 (No. 2), 1995, p 689–703
14. V.R. Voller and S. Sundarraj, Modeling of Microsegregation, Mater. Sci. Technol., Vol 9 (No. 6), 1993, p 474–481 15. A. Roosz, E. Halder, and H.E. Exner, Numerical Calculation of Microsegregation in Coarsened Dendritic Microstructures, Mater. Sci. Technol., Vol 2 (No. 11), 1986, p 1149–1155 16. T. Kraft, M. Rettenmayr, and A. Roosz, Undercooling Effects in Microsegregation Modeling, Scr. Mater., Vol 35 (No. 1), 1996, p 77–82 17. T. Kraft and Y.A. Chang, Predicting Microstructure and Microsegregation in Multicomponent Alloys, JOM, Vol 49 (No. 12), 1997, p 20–28 18. V.R. Voller and C. Beckermann, A Unified Model of Microsegregation and Coarsening, Metall. Mater. Trans. A, Vol 30 (No. 8), 1999, p 2183–2189 19. T. Umeda, Microsegregation, Proc. M.C. Flemings Symp. on Solidification and Materials Proc., R. Abbaschian, H.D. Brody, and A. Mortensen, Ed., TMS, 2001, p 161–170 20. V. Mathier, A. Jacot, and M. Rappaz, Coalescence of Equiaxed Grains during Solidification, Mater. Sci. Eng., Vol 12, 2004, p 479–490 21. H.D. Brody and S.A. David, Application of Solidification Theory to Titanium Alloys, The Science and Technology of Titanium, R. Jaffee and N. Promisel, Ed., Pergamon Press, Oxford, 1970, p 21–34 22. J.I. Nurminen and H.D. Brody, Dendrite Morphology and Microsegregation in Titanium Base Alloys, Titanium Science and Technology, Vol 3, R.T. Jaffee and H.M. Burte, Ed., Plenum Publishing Corp., New York, 1973, p 1893–1914 23. T.Z. Kattamis, J.C. Coughlin, and M.C. Flemings, Influence of Coarsening on Dendrite Arm Spacing of Aluminum-Copper Alloys, Trans. Met. Soc.-AIME, Vol 239, 1967, p 1504–1511 24. P.W. Voorhees, Coarsening in Binary Solid-Liquid Mixtures, Metall. Trans. A, Vol 21, 1990, p 27–37 25. T.Z. Kattamis and P.W. Voorhees, Coarsening of Solid-Liquid Mixtures: A Review, Proc. M.C. Flemings Symp. on Solidification and Materials Proc., R. Abbaschian, H.D. Brody, and A. Mortensen, Ed., TMS, 2001, p 119–170 26. B. Li, H.D. Brody, and A. Kazimirov, Real-Time Observation of Dendrite Coarsening in Sn-13%Bi Alloy by Synchrotron Microradiography, Metall. Mater. Trans. A, Vol 38 (No. 3), 2007, p 599–605 27. G.E. Nereo and M.C. Flemings, Macrosegregation, I, Trans. Met. Soc.-AIME, Vol 239, 1967, p 1449–1461 28. B.P. Bardes and M.C. Flemings, Dendrite Arm Spacing and Solidification Time in Cast Aluminum-Copper Alloy, Trans. AFS, Vol 74, 1966, p 405–412
29. J. Alkemper and P.W. Voorhees, Three Dimensional Characterization of Dendritic Microstructures, Acta Mater., Vol 49 (No. 5), 2001, p 897–902 30. I. Ohnaka, Microsegregation and Macrosegregation, Casting, ASM Handbook, Vol 15, ASM International, 1988 31. M.C. Flemings, R.V. Barone, D.R. Poirier, and H.D. Brody, Segregation in Iron Base Alloys, JISI, April 1970, p 371–381 32. P.J. Ahearn and M.C. Flemings, Dendrite Morphology of a Tin-Bismuth Alloy, Trans. Met. Soc.-AIME, Vol 239, 1967, p 1590–1593 33. J.S. Lee, “Dendrite Structure and Solute Redistribution in Copper-33%Zinc Alloy,” M.S. thesis, University of Pittsburgh, 1972 34. M.N. Gungor, A Statistically Significant Experimental Technique for Investigating Microsegregation in Cast Alloys, Metall. Trans. A, Vol 20 (No. 11), 1989, p 2529– 2533 35. M. Ganesan, D. Dye, and P.D. Lee, A Technique for Characterizing Microsegregation in Multicomponent Alloys and Its Application to Single Crystal Superalloy Castings, Metall. Mater. Trans. A, Vol 36 (No. 8), 2205, p 2194–2204 36. Y. Ma, “Integrated Study of the Influence of Casting and Heat Treatment Parameters on Microstructures of Multiphase Multicomponent Aluminum Alloys,” Ph.D. dissertation, University of Connecticut, 2008 37. C. Lin, “Sequence of Phase Appearance and Microstructure Development during the Solidification of Al-Si-Cu-Mg-Fe Alloys,” M.S. thesis, University of Connecticut, 2003 38. B. Li, H.D. Brody, and A. Kazimirov, Real Time Observation of Dendrite Coarsening in Sn-13%Bi Alloy by Synchrotron Microradiography, Phys. Rev. E, Vol 70, 2004 39. B. Li, H.D. Brody, D.R. Black, H. Burdette, and C. Rau, A Compact Design of a Temperature Gradient Furnace for Synchrotron Microradiography, Meas. Sci. Technol., Vol 17, 2006, p 1883–1887 40. B. Li, H.D. Brody, and A. Kazimirov, Synchrotron Microradiography of Thermal Gradient Zone Melting in Directional Solidification, Metall. Mater. Trans. A, Vol 37, March 2006, p 1039–1044 41. H. Nguyen Thi, G. Reinhart, A. Buffet, T. Schenk, N. Mangelinck-Noe¨la, H. Jung, N. Bergeon, B. Billia, J. Ha¨rtwig, and J. Baruchel, In Situ and Real-Time Analysis of TGZM Phenomena by Synchrotron X-Ray Radiography, J. Cryst. Growth, Vol 310 (No. 11), 2008, p 2906–2914 42. L.H. Cadoff, “Effect of Dendrite Coarsening on Solute Redistribution in Cast Alloys,” unpublished research, University of Pittsburgh, 1970 43. M. Qian, F. Yi, D. Zhang, X. Pan, H.D. Brody, and J.E. Morral, Solute Redistribution and Phase Appearance in As-Cast
Microsegregation / 347 Al-Si-Cu-Mg-Fe Alloys, Metallurgical Modeling for Aluminum Alloys, M. Tiryakioglu and L.A. Lalli, Ed., ASM International, 2003, p 69–78 44. X.M. Pan, C. Lin, J.E. Morral, and H.D. Brody, An Assessment of Thermodynamic Data for the Liquid Phase in the Al-Rich Corner of the Al-Cu-Si System and Its Application to the Solidification of 319
Alloy, J. Phase Equilibria Diffusion, Vol 26 (No. 3), June 2005, p 225–233 45. D. Zhang, J.E. Morral, and H.D. Brody, Measurements for Cu and Si Diffusivities in Al-Si-Cu Alloys by Diffusion Couples, Mater. Sci. Eng. A, 2007, p 217–221 46. T. Umeda, J. Matsuyama, H. Murayama, and M. Sugiyama, Tetsu-to-Hagane´ (J. Iron Steel Inst. Jpn.), Vol 63, 1977, p 441
47. R. Sellamuthu, H.D. Brody, and A.F. Giamei, Effect of Fluid Flow and Hafnium Content on Macrosegregation in the Directional Solidification of Nickel Base Superalloys, Metall. Trans. B, Vol 17, 1986, p 347–356
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 348-352 DOI: 10.1361/asmhba0005216
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Macrosegregation Christoph Beckermann, University of Iowa
MACROSEGREGATION refers to spatial variations in composition that occur in metal alloy castings and range in scale from several millimeters to centimeters or even meters. These compositional variations negatively impact the subsequent processing behavior and properties of cast materials and can lead to rejection of cast components or processed products. It is present in virtually all casting processes, including continuous, ingot, and shape casting of steel, cast iron, aluminum and copper alloys, casting of single-crystal superalloys, semisolid casting, and growth of semiconductor crystals. Macrosegregation is especially important in large castings and ingots. Because of the low diffusivity of the solutes in the solid state and the large distances involved, macrosegregation cannot be mitigated through processing of the casting after solidification is complete. The cause of macrosegregation is relative movement or flow of segregated liquid and solid during solidification. Most alloy elements have a lower solubility in the solid than in the liquid phase, as shown by the partial equilibrium phase diagram in Fig. 1. During solidification, the solutes are therefore rejected into the liquid phase, leading to a continual enrichment
k=
C0
Tm
of the liquid and lower solute concentrations in the primary solid. This segregation occurs on the scale of the microstructure that is forming, which often consists of dendrites having arm spacings of the order of 10 to 100 mm. It is therefore termed microsegregation and results in a nonuniform, cored solute distribution in the dendrite arms. Consider now a small volume element that contains several dendrite arms and the liquid between them, that is, an element inside the liquid-solid (mushy) zone. In the absence of any solute transport into or out of the volume element, the average composition of the mixture inside of the volume element will remain constant at the nominal alloy composition, C0. However, if liquid or solid having a solute concentration different from the one of the liquid or solid inside of the volume element flows into the volume element, the average composition of the mixture in the volume element will change away from the nominal alloy composition. Since such flow generally occurs over large distances, the casting becomes macroscopically segregated. Positive (negative) macrosegregation refers to compositions above (below) the nominal alloy composition. All macrosegregation must average out to zero over the entire casting. There are numerous causes of liquid flow and solid movement in casting processes: Flow that feeds the solidification shrinkage
CS CL
Temperature
m CL
CS
Liquidus
Solidus
A
C0
B
Alloying element B, mass%
Fig. 1
Schematic of a partial equilibrium phase diagram of a binary alloy. See text for an explanation of parameters.
and the contractions of the liquid and solid during cooling Buoyancy-induced flows due to thermal and solutal gradients in the liquid. The thermal and solutal buoyancy forces can either aid or oppose each other, depending on the direction of the thermal gradient and whether the rejected solutes cause an increase or a decrease in the density of the liquid. Flows due to capillary forces at liquid-gas interfaces Residual flows from pouring Flows induced by gas bubbles Forced flows due to applied electromagnetic fields, stirring, rotation, vibration, and so on Movement of small (equiaxed) grains or solid fragments that have heterogeneously nucleated in the melt, separated from a mold wall or free surface, or melted off dendrites.
The solid can either float or settle, depending on its density relative to the liquid. Deformation of the solid network in the mushy zone due to thermal and shrinkage stresses, head pressures, or external forces on the solid shell Efforts to prevent macrosegregation are aimed at controlling liquid flow and movement of solid. Examples include adjustments to the alloy composition or thermal gradients to induce a stable density stratification in the liquid; application of nozzles, baffles, porous materials, or electromagnetic fields to redistribute the flow; controlled deformation such as soft reduction during continuous casting of steel to squeeze highly enriched liquid away from the centerline of slabs; changing the cooling rate or the alloy composition to decrease the dendrite arm spacing and hence increase the resistance to flow through the solid network; and increasing the prevalence of equiaxed grains in a casting, for example, by inoculation or stirring, in order to obtain a more uniform and dense solid distribution.
Macrosegregation Induced by Flow of the Liquid Liquid flow through the mushy zone, where the developing solid network can be assumed to be rigid and fixed, is the most common cause of macrosegregation. Such macrosegregation can be best understood by considering the local solute redistribution equation (LSRE) derived by Flemings and coworkers (Ref 1–3). This equation is based on a solute balance on a small-volume element inside the mushy zone, as illustrated in Fig. 2. As in a standard Scheil-type analysis (see other articles in this section of the Volume), local solute diffusion in the solid phase is neglected, and the liquid inside the volume element is assumed to be well mixed and in equilibrium. Then, the solute concentration in the liquid, CL, is given as a function of temperature by the liquidus line of an equilibrium phase diagram (Fig. 1) according to CL = (T Tm)/m (where T is the temperature, Tm is the melting point of the pure solvent, and m is the slope of the liquidus line).
Macrosegregation / 349 Liquid flow
velocity in the direction normal to the isotherms, vT is the isotherm velocity, and: b¼
Solid
rS rL rS
is the solidification shrinkage, where rS and rL are the densities of the solid and liquid, respectively. In one dimension, mass conservation yields that the liquid velocity needed to feed the solidification shrinkage is given by:
Liquid
r b un jshrink ¼ vT 1 S ¼ vT rL ð1 bÞ
Fig. 2
Schematic of liquid flowing through a small volume inside the mushy zone. The gray ovals represent cuts through dendrite arms.
Mushy zone
Temperature, T
Liquid solute concentration, CL
Isotherms
Isotherm Normal liquid velocity, un velocity, nT
Schematic of liquid flowing through a mushy zone
Distance from chill, in. 0
0.5
1.0
1.5
2.0
2.5
3.0
– Copper concentration, (C/mass)%
4.6 Measured
(Eq 3)
Note that shrinkage-driven flow is in the direction opposite to the isotherm velocity, that is, in the direction of decreasing temperature toward regions of higher solid fraction. Only if the alloy expands upon solidification, b < 0, and unjshrink has the same sign as vT. Figure 3 illustrates the meaning of some of the quantities appearing in the previous equations. Since the liquid solute concentration is uniform along isotherms, liquid flowing in or out of the volume element in the direction of the isotherms has the same solute concentration as the liquid inside of the volume element. Therefore, liquid flow parallel to the isotherms in the mushy zone does not induce macrosegregation. Defining a flow factor, x, as: un x ¼ ð1 bÞ 1 vT
Liquid flow
Fig. 3
(Eq 2)
(Eq 4)
Equation 1 can be integrated to yield a modified Scheil equation: CL ¼ ð1 fS Þðk1Þ=x C0
(Eq 5)
where x and k are assumed to be constant. No macrosegregation occurs if the flow factor is equal to unity and the standard Scheil equation is recovered. There are two obvious cases for which x = 1 (no macrosegregation): No shrinkage (b = 0) and no liquid flow nor-
4.4
mal to the isotherms (un = 0)
Simulated
No liquid flow other than the one required
to feed the solidification shrinkage in one dimension: un = vT b/(1 b) or un = un jshrink
4.2 Co = 4.4% 4.0
3.8
0
25
50
75
Distance from chill, mm
Fig. 4
Inverse segregation in an A1-4.1 Cu ingot with unidirectional solidification
Accounting for advection of solute by the liquid flow in or out of the volume element, and for the different densities of the solid and liquid phases, the following LSRE results: dfS ð1 fS Þ ð1 bÞ un ¼ 1 dCL vT CL ð1 kÞ
(Eq 1)
where fS is the solid volume fraction, k is the partition coefficient, un is the liquid flow
For x 6¼ 1, macrosegregation occurs, since the liquid concentration, CL, varies differently from that predicted by the standard Scheil equation. Based on Eq 5, three different macrosegregation mechanisms and outcomes can be identified (assuming, for the purpose of this discussion, that k < 1): 1. Negative macrosegregation: x > 1 or un/vT < b/(1 b), or equivalently, un < un jshrink. At the same solid fraction, CL is lower than for x = 1, indicating that negative macrosegregation will result. For the case of b 0, this condition is met if the liquid flows in the direction of decreasing temperature toward regions of higher solid fraction (un < 0) and with a speed (absolute value of velocity) that is greater than jun jshrinkj.
2. Positive macrosegregation: 0 < x < 1 or 1 > un/vT > b/(1 b), or equivalently, vT > un > un jshrink. At the same solid fraction, CL is higher than for x = 1, indicating that positive macrosegregation will result. For the limiting case of b = 0, this condition is met if the liquid flows in the direction of increasing temperature toward regions of lower solid fraction (un > 0) but with a velocity that is lower than the isotherm velocity (un < vT). For b > 0, this condition can be met even if the liquid flows in the direction of decreasing temperature toward regions of higher solid fraction (un < 0), since unjshrink is negative. 3. Remelting: x < 0 or un > vT. This condition is met if the liquid flows in the direction of increasing temperature toward regions of lower solid fraction and with a velocity that is higher than the isotherm velocity. Based on Eq 1, such flow results in the solid fraction decreasing, rather than increasing, with decreasing temperature. This means that remelting occurs and that open (solid-free) channels can form inside the mushy zone. Once the casting is solidified, the macrosegregation inside of these channels is positive. The direction and magnitude of the liquid flow velocities in the mushy zone depend on numerous factors. For a given driving force, the permeability of the mush is the single most important parameter that limits the flow (Ref 4). In general, the permeability decreases with increasing solid fraction. A smaller dendrite arm spacing also reduces the permeability. The following examples illustrate in more detail the aforementioned macrosegregation mechanisms. Inverse Segregation. Near a chill face, the positive macrosegregation shown in Fig. 4 is often observed (Ref 5). This so-called inverse segregation (Ref 6) is a consequence of the second macrosegregation mechanism discussed previously. At an impenetrable chill face, the liquid velocity vanishes (un = 0). Furthermore, for a finite shrinkage, unjshrink is negative. Therefore, the condition un > unjshrink is always satisfied at the chill face, and positive macrosegregation results. If a gap is present between the mold and the casting, the solute-rich interdendritic liquid can be pushed into the gap by metallostatic pressure. This is termed exudation and results in even higher positive macrosegregation at the surface. In direct chill casting of aluminum ingots, this macrosegregation can be so severe that the exuded surface layer must later be removed by scalping the ingot. Macrosegregation due to Nonuniform Shrinkage Flow. The first macrosegregation mechanism listed previously has been demonstrated for the casting geometry shown in Fig. 5 (Ref 5). The casting is chilled from the bottom, such that the isotherm velocity is generally in an upward direction. The downward liquid flow is caused solely by shrinkage. However, due to the reduction in the cross section of
350 / Principles of Solidification
0.26
0.27 Section B–B C
0.24 0.22
Section A–A
0.20
∂T = 0 ∂n
20
B
B ∂T = 0 ∂n
Z
– Copper concentration, (C/mass)%
Simulated 0.23
Measured (0.25C–0.25Si)
0.25% C
0.21
23
0.025 mm/s ∂T = 0 ∂n
0
0.25
0.19 −15
−10
−5
0
5
10
Position (Z), mm A
A
Fig. 6 25
Macrosegregation observed along Z-direction (cross-sectional mean value) for an Fe-0.25C specimen
y ∂T =0 ∂n
y=H
–Z fs = 1
50 mm
Riser
∂T = 0 ∂n
Simulated fluid flow at 50 s after cooling and macrosegregation in an Fe-0.25C specimen
the casting near midheight, the liquid must flow at a higher downward velocity in this region. Hence, the condition un < un jshrink is satisfied in this region (since both velocities are negative), and the negative macrosegregation seen in Fig. 5 and 6 forms (Ref 5). Macrosegregation due to Buoyancy-Driven Flow. The first and second macrosegregation mechanisms can be observed simultaneously if the liquid flow pattern in the mushy zone is of a recirculating nature. This has been demonstrated for the horizontally solidified aluminum-copper casting shown in Fig. 7 (Ref 5). Here, a counterclockwise recirculating flow pattern is induced in the liquid by thermal and solutal buoyancy forces. Shrinkage-induced flow is negligibly small in this case. Figure 7 shows that the flow in the bulk liquid region penetrates into the mushy zone. In the upper part of the mushy zone, the flow is in the direction of decreasing temperature (un < 0). According to the first macrosegregation mechanism, such flow causes negative macrosegregation. Negative macrosegregation was indeed measured in the upper part of the casting, as shown in Fig. 8. Conversely, in the lower part of the mushy zone, the interdendritic liquid flows in the direction of increasing temperature
S
S+L
Cooled
Fig. 5
1 mm/s
y=0
12 mm diam Cooled 29 mm diam
0.1 mm/s
H = 100 mm
Carbon, %
0.28
250 mm
Fig. 7
∂T =0 ∂n
Simulated fluid flow at 400 s after cooling in a horizontally solidified Al-4.4Cu casting
(un > 0), back into the bulk liquid region. This causes positive macrosegregation according to the second mechanism. This was again verified by the measurements (Fig. 8). Note that there is no macrosegregation at the midheight of the casting, because there the liquid flow is primarily parallel to the isotherms. Buoyancy-driven flow is also responsible for the positive macrosegregation seen in the riser of large steel castings and ingots, as shown in Fig. 9. In steel casting, the interdendritic liquid is often lighter than the bulk liquid and flows up. This up-flow is in the direction of increasing temperature (un > 0), since the riser is the last part of the casting to solidify. Consequently, positive macrosegregation results according to the second mechanism.
Channel Segregation. The third macrosegregation mechanism leads to the formation of so-called channel segregates. One such channel segregation defect is commonly referred to as freckles, an example of which is shown in Fig. 10. Freckles sometimes occur in upward, directional solidification of single-crystal nickel-base superalloy castings. During solidification of nickel-base superalloys, a number of light elements (such as aluminum and titanium) are rejected into the liquid, and some heavy elements, such as tungsten, are preferentially incorporated into the solid, leading to strong, upward solutal buoyancy forces in the mushy zone. Despite the presence of a stabilizing thermal gradient, these buoyancy forces can trigger convection cells. As the upward liquid
Macrosegregation / 351 1.0 100 mm from chill
Fractional height, y/H
0.8 Simulated 0.6
0.4
0.2
50 mm from chill
0 4.0
Fig. 8
4.5 5.0 4.0 4.5 – Carbon concentration, (C/mass)%
5.0
Solute distribution in the Al-4.4Cu ingot shown in Fig. 7
Riser A-segregation in branched columnar dendrite zone
Positive segregation in riser
Fig. 10 V-segregation in equiaxed grain zone Negative cone Columnar dendrite zone
Fig. 9
Chill zone
Typical macrosegregation observed in steel ingots
velocities in the mush exceed the isotherm velocity (un > vT), the solid fraction decreases with decreasing temperature, according to the third macrosegregation mechanism. Eventually, this leads to open channels in the mush through which solute-rich liquid streams upward into the bulk liquid region. The numerical simulation results shown in Fig. 11 illustrate these flow patterns (Ref 7). The rodlike channels are no larger than a few millimeters in diameter and have a spacing of, at most, a few centimeters. Later, the channels are filled with dendrite fragments, which are then observed as frecklelike chains of equiaxed crystals in the otherwise single-crystal columnar structure (Fig. 10). The resulting positive macrosegregation in the channels can also be seen in Fig. 11. Channel segregates are also observed in large steel ingots. Here, they are termed A-segregates and occur in the columnar zone of the ingot (Fig. 9). Strong buoyancy forces can cause the liquid to flow upward in the mush with a velocity normal to the isotherms that is greater than the isotherm velocity (un > vT), resulting in remelting. The channels are usually inclined toward the center of the ingot and thus appear as the A-shape illustrated in Fig. 9.
(a) Freckles in a single-crystal nickel-base superalloy prototype blade. (b) Closeup of a single freckle. The freckles are approximately 2 mm (0.1 in.) in diameter.
Macrosegregation Induced by Movement of the Solid The movement or flow of solid during solidification can also induce macrosegregation in castings. Unfortunately, no simple theory, such as the LSRE discussed in the previous section, exists to explain the macrosegregation patterns resulting from solid movement (Ref 8). Therefore, only a few examples are provided to illustrate some of the complex phenomena that have been observed. The most common form of solid movement is the settling or floating up of small solid pieces formed early in the solidification process. These solid pieces may be dendrite fragments that separated from an existing solid structure or equiaxed grains that nucleated in the bulk liquid. They settle or float depending on their density relative to the liquid. The solid pieces generally have a composition different from the nominal alloy composition, and their movement to different parts of the casting thus induces macrosegregation. One example is provided by the cone of negative macrosegregation that is often observed in the bottom third of steel ingots, as illustrated in Fig. 9. This macrosegregation pattern is confined to the equiaxed zone and is caused by the settling of dendrites that are relatively poor in solute (Ref 9, 10). Buoyancy-driven flow of the liquid also induces negative macrosegregation in the bottom half of steel ingots, but the degree is not as severe as that due to settling of solid (Ref 11). Similarly, negative centerline macrosegregation in direct chill casting of aluminum ingots is sometimes
attributed to the settling of unattached solute-poor solid, although shrinkage-driven flow has been found to have a similar effect (Ref 12). On the other hand, the floating of spheroidal graphite (kish) can cause strong positive macrosegregation in the upper part of iron castings. Another example of solid-movement-induced macrosegregation is given by the V-segregates in the equiaxed region of tall steel ingots. This macrosegregation pattern is also illustrated in Fig. 9 and consists of rodlike solute-rich streaks that appear periodically along the centerline. The V-segregates arise from equiaxed crystals settling in the core and forming a loosely connected network that can easily rupture due to metallostatic head and liquid being drawn down to feed solidification shrinkage. Fissures then open up along shear planes oriented in a V-pattern and are filled with enriched liquid (Ref 9, 10). Macrosegregation can also be induced by deformations of the fully and partially solid regions during casting. One important example is centerline segregation in continuously cast steel slabs (Ref 10). Close to the bottom of the liquid pool, inadequate roll containment can cause the already solidified steel shell to bulge. The bulging draws solute-rich liquid into the center of the slab, where it freezes. This results in a rather sharp and thin line of positive macrosegregation along the center of the slab, as can be seen from the sulfur print in Fig. 12. Deformation of the solid can also be induced by thermal stresses. This is a particularly common phenomenon in shape casting, due to the restraint that the mold and/or the cores can offer. However, deformation of the solid can also
352 / Principles of Solidification ACKNOWLEDGMENT Figures 4 to 9 are Fig. 8 to 12 and 7, respectively, in I. Ohnaha, Microsegregation and Macrosegregation, Casting, Vol. 15, Metals Handbook, ASM International, 1988, p136–141 REFERENCES
Fig. 11
Simulation of freckle formation during upward directional solidification of a single-crystal nickel-base superalloy in a 5 cm (2 in.) wide by 15 cm (6 in.) tall enclosure. (a) Melt velocity vectors in the liquid region and mushy zone (longest vector corresponds to a velocity of approximately 5 mm/s, or 0.2 in./s) and solid fraction contours (in five equal intervals of 20%, starting with 0% for the uppermost contour at approximately midheight). (b) Macrosegregation pattern. The gray shades indicate the titanium concentration, with the darkest shade inside the open channels in the mush corresponding to the largest positive macrosegregation.
Fig. 12
Sulfur print showing centerline macrosegregation in a continuously cast steel slab
occur in continuous casting due to uneven cooling. The stresses are able to propagate into the high solid-fraction portion of the mushy zone, where the solid network is already coherent. If the solid network is under compression, interdendritic liquid will be forced to “drain” out of the mush. Conversely, if this part of the mush is in tension, liquid is drawn into it. In either case, complex macrosegregation patterns can result. Under sufficiently severe tension, the solid network will rupture, and the fissure may fill with solute-rich liquid. The resulting cracklike positive macrosegregation in a casting is often referred to as a healed hot tear. If the flow of the interdendritic liquid is insufficient to feed the fissure, an open hot tear will be present in the casting.
1. M.C. Flemings and G.E. Nereo, Macrosegregation: Part 1, Trans. Met. Soc. AIME, Vol 239, 1967, p 1449–1461 2. R. Mehrabian, M. Keane, and M.C. Flemings, Interdendritic Fluid Flow and Macrosegregation; Influence of Gravity, Metall. Trans., Vol 1, 1979, p 1209–1220 3. M.C. Flemings, Solidification Processing, McGraw-Hill, New York, 1974 4. D.R. Poirier, Permeability for Flow of Interdendritic Liquid in Columnar-Dendritic Alloys, Metall. Trans. B, Vol 18, 1987, p 245–255 5. I. Ohnaka and M. Matsumoto, Computer Simulation of Macrosegregation in Ingots, Tetsu-to-Hagane (J. Iron Steel Inst. Jpn.), Vol 73, 1987, p 1698–1705 6. J.S. Kirkaldy and W.V. Youdelis, Contribution to the Theory of Inverse Segregation, Trans. Met. Soc. AIME, Vol 212, 1958, p 833–840 7. M.C. Schneider, J.P. Gu, C. Beckermann, W.J. Boettinger, and U.R. Kattner, Modeling of Micro- and Macrosegregation and Freckle Formation in Single-Crystal Nickel-Base Superalloy Directional Solidification, Metall. Mater. Trans. A, Vol 28, 1997, p 1517–1531 8. C. Beckermann, Modelling of Macrosegregation: Applications and Future Needs, Int. Mater. Rev., Vol 47, 2002, p 243–261 9. M.C. Flemings, Principles of Control of Soundness and Homogeneity of Large Ingots, Scand. J. Metall., Vol 5, 1976, p 1–15 10. G. Lesoult, Macrosegregation in Steel Strands and Ingots: Characterization, Formation and Consequences, Mater. Sci. Eng. A, Vol 413–414, 2005, p 19–29 11. J.P. Gu and C. Beckermann, Simulation of Convection and Macrosegregation in a Large Steel Ingot, Metall. Mater. Trans. A, Vol 30, 1999, p 1357–1366 12. A.V. Reddy and C. Beckermann, Modeling of Macrosegregation due to Thermosolutal Convection and Contraction Driven Flow in Direct Chill Continuous Casting of an Al-Cu Round Ingot, Metall. Mater. Trans. B, Vol 28, 1997, p 479–489
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 353-356 DOI: 10.1361/asmhba0005217
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Interpretation and Use of Cooling Curves— Thermal Analysis H. Fredriksson, The Royal Institute of Technology, Sweden
THERMAL ANALYSIS is a classical method of determining phase diagrams. By melting and cooling an alloy of known composition and registering the temperature-time curves, the liquidus temperature for the respective alloy can be determined. Figure 1 shows an idealized view of the relationship between the cooling curves and the phase diagram for an alloy A-B. The starting point for solidification is given by a change in the slope of the cooling curve due to the evolution of the heat of solidification. This starting point is normally used to define the liquidus temperature. The method is described more fully in Ref 1. The idealized behavior illustrated in Fig. 1 is not seen in experimentally determined temperature-time curves. For example, Fig. 2, an experimentally determined cooling curve for an ironcarbon alloy with 3.2% C, shows that the formation of austenite and the eutectic reaction occur at a temperature well below the liquidus and the eutectic temperatures, respectively. The reason for this depends on the nucleation and growth of the crystals in the liquid. In this article, this effect is analyzed, and the use of thermal analysis in industrial processes and in research is discussed.
Quantitative Thermal Analysis Cooling curves describe a balance between the evolution of heat in the sample and the heat transport away from the sample. It is often easiest to describe this in a mathematical form. However, heat transport away from the liquid is often very complex and can be described in many different ways because of the experimental conditions. Currently, heat transport is often described by numerical methods. In the following discussion, only the expression dQ/dt is used to define the heat transport away from the experimental setup. This heat transport must be equal to the evolution of heat in the sample. Before and after the solidification process, this is balanced by a loss of heat capacity. Before the start of the solidification process, this can be written in analytical form as:
dQ dT ¼ V r Cp dt dt
(Eq 1)
where V is the volume of the sample, r is the density of the liquid, Cp is the heat capacity, and dT/dt is cooling rate of the liquid. From a cooling curve similar to that shown in Fig. 2, dT/dt is the slope of the curve before
solidification begins. By evaluating the slope as a function of time in an experiment, dQ/dt can be determined if r and Cp are known. This can later be used in the analysis of the temperaturetime curve. When the liquid cools to the liquidus temperature, crystals are nucleated and begin to grow. The heat balance from Eq 1 is now: dQ ¼ dt
Relationship between a cooling curve and phase diagram for alloy A-B
df dT
dT dt
(Eq 2)
where DH is the heat of solidification per unit volume, and df/dT is the volume fraction of solid formed at a changing temperature. The change in shape of the temperature-time curve now depends on the volume fraction of solid formed. The larger the rate of increase of volume fraction of solid, the smaller the slope (dT/dt) of the temperature-time curve. The volume fraction of solid formed is related to the number of crystals and their growth rate. Crystals are normally nucleated when the temperature of the liquid reaches the liquidus temperature as well as the eutectic temperature. In spite of this, the temperature of the liquid will decrease further, and initially the curve
Fig. 2
Fig. 1
V r Cp þ rH
Cooling curve for an iron-carbon alloy with 3.2% C. The solidification starts with a primary precipitation of austenite, followed by a eutectic reaction.
354 / Principles of Solidification can be extrapolated to show where no crystals are growing. After a certain time, the slope of the curve begins to decrease, slowly at first and then more rapidly, forming a small plateau or even a small temperature increase near the liquidus, as seen in Fig. 2. In the early stages of crystal formation, the growing crystals are small (as is heat development) and therefore have no effect on the bulk liquid temperature. The growth rate is also low because of the low supercooling close to the liquidus temperature. The supercooling increases when the temperature decreases; therefore, the growth rate will increase as the area of the crystal increases. The result is that more heat is given off and the slope of the temperature-time curve begins to decrease. The change in the slope will begin earlier and will proceed faster. The more crystals there are in the liquid, the earlier the change in slope starts, and the faster it will proceed. When the temperature increases, the growth rate of the free crystals decreases again, but they are so large by this time that the volume solidifying per unit of time will also be large. As such, the temperature will continue to increase until the growth rate and the growth area of the crystals offset the heat extraction. The larger the area of the growing crystals, the lower the growth rate and the closer the temperature of the bulk liquid to the liquidus or eutectic temperature.
Deviation from Equilibrium The solidification process illustrated by the cooling curve in Fig. 2 begins with a primary precipitation of austenite, followed by a eutectic reaction. In this case, the liquid was cast into a sand mold. The heat extraction from the liquid can be approximated by Chvorinov’s law and follows the expression: dQ K ¼ pffiffi dt t
Fig. 3
(Eq 3)
where t is the time and K the so-called modulus constant. The value K can be determined by using Eq 1 and the slope of the cooling curve before solidification begins. The fraction of solid can be determined by using Eq 2 and 3 and the data given in Fig. 2. The results of these calculations are shown in Fig. 3(a). This type of analysis has been performed and is discussed in 2 to 4; it can be used to analyze the deviation from solidification under equilibrium conditions. The Fe-3.2C alloy discussed earlier in this article can be used as an example to illustrate this analysis. According to the iron-carbon phase diagram, the liquidus temperature for this alloy is 1268 C (2314 F). At this temperature, dendritic crystals of austenite are formed. Because carbon has a very high diffusivity in austenite, the lever rule can be used to determine the fraction of solid as a function of temperature. The lever rule can be applied to any two-phase field of a phase diagram to determine the amounts of the different phases present at a given temperature in a given alloy. A horizontal line, referred to as a tie line, represents the lever, and the alloy composition represents its fulcrum. The intersection of the tie line with the boundaries of the two-phase field fixes the compositions of the coexisting phases, and the amounts of the phases are proportional to the segments of the tie line between the alloy and the phase compositions. The rule is derived with respect to temperature, and the derivative df/dT is obtained. This derivative is inserted into Eq 2. The temperature-time curve is then easily calculated. The result is shown in Fig. 3(b) as a dashed line. The deviation from the real curve and that calculated under equilibrium conditions is greatest at the beginning of the precipitation of austenite. After the initial stage during the cooling down to the eutectic temperature, the lever rule will describe the fraction of solid formed. The eutectic reaction for the equilibrium case occurs without any undercooling. This deviation from equilibrium is interpreted as follows.
The fraction of solid phases as a function of time for the alloy with a cooling curve shown in Fig. 2. (b) Deviation of the cooling shown in Fig. 2 resulting from equilibrium solidification condition (see dashed line)
The austenite dendrites can be assumed to nucleate randomly in the liquid. They grow in leading tips in a starlike fashion. The secondary dendrite arms form behind the tips. The growth of the dendrite tips is dependent on undercooling. This undercooling is determined by the growth rate, and it increases with increasing growth rate. The growth rate of the tips is determined by the heat extraction and by the number of growing crystals. When the dendrite crystals have collided, a network of dendrite arms is formed. The solidification now proceeds by a thickening of the dendrite arms. Similarly, eutectic colonies are randomly nucleated in the liquid; they grow spherically and collide with each other. When the collision occurs, the growth area decreases, and the temperature of the liquid decreases. Theoretical models that combine growth equations, statistical laws for the collision of crystals, and heat extraction laws have been developed in recent years to describe the previous discussion (Ref 2–6). In the future, these models will probably be used to predict the structure formed during the solidification process and the resultant properties of the material.
Differential Thermal Analysis Thermal analysis has a long history of use in research. To increase accuracy, the experiments are normally carried out by so-called differential thermal analysis. Simplified Thermal Analysis. Before discussing the theoretical basis of differential thermal analysis, a simplified experimental setup is considered. This setup consists of a crucible with a melt placed inside a furnace. A thermoelement is placed in the center of the melt to measure the temperature of the sample, Ts. Another thermocouple is placed at the wall of the furnace; the purpose of this thermocouple is to regulate the temperature of the furnace, Tf, as a function of time. The furnace is assumed to cool at a constant rate in degrees Kelvin per minute. Figure 4 shows the temperature-time cycle for an aluminum-copper alloy that contains 5% Cu. The furnace has cooled during the experiment at a rate of 10 K/min (285 C/ min, or 510 F/min). Figure 4 shows that the temperature of the sample is somewhat higher than that of the furnace. This divergence increases slightly with time. A stationary condition, that is, the same cooling rate for the sample and the furnace, never occurs; the cooling rate of the sample is always somewhat lower than that of the furnace. The melt begins to solidify at the point designated tstart. Heat is emitted by the sample, and the temperature difference between the sample and the furnace increases. A hump sometimes appears in the temperature-time curve of the sample. The maximum point, or the deviation point, on the cooling curve is usually used to establish the liquidus temperature.
Interpretation and Use of Cooling Curves—Thermal Analysis / 355
Fig. 4
Thermal analysis of an Al-5Cu alloy cooled at a rate of 10 K/min (285 C/min, or 510 F/min). Source: Ref 6
the solidus temperature, tsol, during the experiment. To facilitate evaluation of the liquidus and solidus temperatures, the temperature-time curve of the sample is differentiated in most cases. In general, the derivative of the temperature-time curve for the sample shows a discontinuity when the heat of solidification ceases to develop. The solidus temperature is generally defined by this discontinuity. This sharp change becomes less pronounced in cases in which the solidification interval is large and the reaction rate is low toward the end of solidification. The reaction rate is normally evaluated by a simple mathematical analysis. This is described by: dTs df r Csp Vs þ H r Vs dt dt
(Eq 4)
¼ ZðTs Tf Þ
Fig. 5
Schematic of principal setup for differential thermal analysis equipment. Tn is the standard temperature, Ts is the sample temperature, and Tf is the furnace temperature.
The solidification rate, or the reaction rate, decreases after the initial period, and the cooling rate of the sample increases. Toward the end of solidification, or when everything has solidified, the cooling rate of the sample becomes greater than that of the furnace. This is a result of the large temperature difference that exists between the sample and the furnace when a considerable amount of solidification heat is given off. Afterward, the cooling rate of the sample approaches that of the furnace. It is often very difficult to determine the end of the solidification process and to determine
where dTs/dt is the cooling rate of the sample, CSP is the specific heat of the sample, Vs is the volume of the sample, DH is the heat of solidification, df/dt is the volume fraction of solid formed per unit time, Ts is the sample temperature, Tf is the furnace temperature, and Z is the heat transfer coefficient. The second term on the right-hand side in Eq 4 is 0 before solidification begins. The heat transfer coefficient Z can be evaluated at this time by measuring dTs/dt, Ts, and Tf. The reaction rate df/dt can be determined by using that value and by remeasuring the temperature and its first-time derivative during solidification. Differential Thermal Analysis. Figure 5 shows a schematic setup for differential thermal analysis. The sample to be analyzed and the standard sample, with a thermocouple in each, are placed in a furnace. The temperature of the furnace and the cooling rate are controlled and measured with a separate thermocouple. The standard sample and the analysis sample must have the same temperature at the beginning of the experiment. During the cooling cycle, the samples will undergo one or more phase transformations that give off heat. This emission of heat implies that, during the time of reaction, the temperature of the analysis sample differs from that of the standard sample. The temperature difference between the analysis sample and the standard sample is recorded at the same time the furnace temperature is recorded. Using the experimental setup shown in Fig. 5, one normally registers the difference in temperature between the standard sample and the analysis sample, not the difference between the sample and the furnace. Therefore, only (Ts Tn) and Ts are known, where Tn is the temperature of the standard sample. In order to be able to calculate the reaction velocity, the term (Ts Tf) in Eq 4 is replaced with the term (Ts Tn) + (Tn Tf). For the standard sample, it is true that Z (Tn Tf) is equal to:
r Csp Vn
dTn dt
(Eq 5)
where dTn/dt is the cooling rate of the standard sample. Replacing this in Eq 4 gives:
dTs df þ ðHÞ r Vs ¼ ZðTs Tn Þ dt dt (Eq 6) dTn þ r CnP Vn dt
r CsP Vs
If the heat of solidification, DH, is known, the reaction rates can be calculated by measuring the temperatures and the cooling rates. If the heat transfer coefficient is known, the heat of solidification can also be determined by integrating Eq 6 over the entire solidification interval (from f = 0 to f = 1). Simplifications of Eq 6 are often found in the literature when the reaction rate and the heat of solidification are determined. It is assumed that the first term on the left-hand side is equal to the last term on the right-hand side, which leads to the following approximate relationship: df Z ðTs Tn Þ ¼ dt ðHÞVs
(Eq 7)
Equation 7 shows a direct proportionality between the measured temperature difference and the reaction rate. In Eq 7, Z/DHVs can be considered constant. The difference in temperature between the sample and the normal (Ts Tn) is measured and is thus known as a function of time. The constant can be determined by integrating Eq 6 from f = 0 to f = 1. This integration is often done graphically. Equation 7 is also often used in interpreting differential thermal analysis data. However, one should keep in mind that doing so introduces many faults. This can be demonstrated by a comparison between Eq 6 and 7. For example, it is assumed that CsP , Vs, Vn, and CnP are equal. This is often not the case. To avoid this problem, the sample itself can be used as an imaginary standard sample. The temperature difference can then be determined by a simple computer analysis (Ref 6). To determine the heat of solidification, the heat transfer coefficient Z must be determined by a series of reference experiments such as that described for the simple type of experiment in Eq 4.
Determining Liquidus and Solidus Temperature Cooling curves can be used in many different ways and are often used to determine the liquidus and solidus temperatures. In such applications, the small plateau temperature shown in Fig. 2 is often used. However, as indicated earlier, this temperature does not correspond to the liquidus temperature, because of the kinetics during crystal growth. To increase the accuracy, the cooling curve is often differentiated. In this
356 / Principles of Solidification
Fig. 6
Liquidus temperature determined for an aluminum- copper alloy. Source: Ref 7
Fig. 7 Temperature-time curve for an iron-base alloy with 1.01% C, 0.25% Si, and 0.46% Mn. The dashed line represents the cooling rate dT/dt = 0. Source: Ref 8
Fig. 8
Thermal analysis used for determining the chemical composition of cast irons
case, the temperature at which the derivative changes is used as the liquidus temperature. As indicated earlier, this will not provide the correct liquidus temperature. The only way to obtain the correct liquidus temperature is to perform a series of experiments at different
cooling rates. The plateau temperature can then be plotted as a function of the cooling rate. The liquidus temperature is determined by extrapolating the line to zero cooling rates. An example of this method is shown in Fig. 6. It is more difficult to determine the solidus temperature than the liquidus temperature. Toward the end of the solidification process, the derivative of the cooling curve often shows a large change (Fig. 7). The solidus temperature is defined in general by this large change. However, the temperature evaluated in this way varies with the cooling rate much more than the liquidus temperature does. This is due to the microsegregation that occurs during solidification. This microsegregation is very much influenced by the cooling rate (Ref 9). Therefore, thermal analysis by cooling is nearly useless for determining the equilibrium solidus temperature. This temperature can be determined only by analyzing the melting process on heating curves. Thermal analysis is often used in industrial processes. One of the most common applications is determination of the chemical composition of cast iron. In this case, a linear relationship has been found between the liquidus temperature and the carbon equivalent (CE), which is defined as: CE ¼ %C þ
%Si %P þ 3 3
(Eq 8)
An example of the use of thermal analysis or determining the chemical composition of cast irons is shown in Fig. 8. Given the present knowledge of thermal analysis, it should be noted that these diagrams can be used with accuracy only if the same conditions are prevalent every time. This is due to the influence of the cooling rate and the nucleation frequency on the plateau temperature. Casting temperature, for example, can influence the results. It is well known that the superheat of the liquid influences the number of crystals. The higher the temperature, the lower the number of growing crystals. The result will be a lower liquidus temperature. However, in most cases, a foundry will use the same melting procedure consistently and can therefore determine chemical composition easily and accurately by thermal analysis, provided a curve similar to the one shown in Fig. 8 is developed for their specific conditions.
3.
4.
5.
6. 7. 8. 9.
SELECTED REFERENCES X. Chen, H. Geng, and Y. Li, Study on the
REFERENCES 1. K.W. Andrews, Physical Metallurgy, Vol 1, William Clowes & Sons, 1973 2. H. Fredriksson and S.E. Wetterfall, The Metallurgy of Cast Iron, Proceedings of the
Second International Symposium on the Metallurgy of Cast Iron (Geneva), May 1974, Georgi Publishing, 1974 H. Fredriksson and I.L. Svensson, Computer Simulation of the Structure Formed During Solidification of Cast Iron, Proceedings of the Third International Symposium on the Metallurgy of Cast Iron, North Holland, 1984 D.M. Stefanescu and C. Kanetkar, Computer Modelling of the Solidification of Eutectic Alloys: The Case of Cast Iron, Proceedings of the Society for Metals and Materials Science, Computer Simulation Division, The Metallurgical Society (Toronto, Canada), 1985 H. Fredriksson and A. Olsson, Mechanism of Transition from Columnar to Equiaxed Zones in Ingots, Mater. Sci. Tech., Vol 2, 1986, p 508–516 H. Fredriksson and B. Rogberg, Met. Sci., Vol 13, 1979, p 685–690 L. Ba¨ckerud and B. Chalmers, Trans. Met. Soc. AIME, Vol 245, 1969, p 309–318 A Guide to the Solidification of Steels, Jernkontorets, Stockholm, 1977 M.C. Flemings, Solidification Processing, McGraw-Hill, 1974
Eutectic Modification Level of Al-7Si Alloy by Computer Aided Recognition of Thermal Analysis Cooling Curves, Mater. Sci. Eng. A, Vol 419 (No. 1–2), March 15, 2006, p 283–289 D. Emadi, L.V. Whiting, S. Nafisi, and R. Ghomashchi, Applications of Thermal Analysis in Quality Control of Solidification Processes, J. Therm. Anal. Calorimetry, Vol 81 (No. 1), 2005, p 235–241 J. Medved and P. Mrvar, Thermal Analysis of the Mg-Al Alloys, Mater. Sci. Forum, Vol 508, 2006, p 603–608 J. Sertucha, R. Suarez, J. Izaga, L.A. Hurtado, and J. Legazpi, Prediction of Solid State Structure Based on Eutectic and Eutectoid Transformation Parameters in Spheroidal Graphite Irons, Int. J. Cast Met. Res., Vol 19 (No. 6), Dec 2006, p 315–322 D.J. Varga L., Examination of the Eutectic Crystallization of Cast Iron by Thermal Analysis, Mater. Sci. Forum, Vol 508, 2006, p 87–92 S. Vogelgesang, G.-S. Leo, and B. Tonn, Thermal Analysis of Grain Refined Brass, Livarski Vestn., Vol 51 (No. 3), 2004, p 133–141
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 357-361 DOI: 10.1361/asmhba0005218
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
X-Ray Imaging of Solidification Processes and Microstructure Evolution Ragnvald H. Mathiesen and Lars Arnberg, Norwegian University of Science and Technology
IN THE PURSUIT for improved insight into the interplay between different transport phenomena that control alloy solidification processes, the inherent difficulties with providing direct observation of the microstructure evolution have been a serious drawback. Typically, experimental observations of metals and alloys have been limited to metallographic and electron microscopy investigations of either quenched or as-solidified material. On the other hand, in situ studies of microstructure evolution and several important process phenomena have been accessible through optical video microscopy studies with a selection of transparent, organic model systems available for primaryphase solidification, eutectics, and peritectics displaying morphologies and growth patterns analogous to those found in metals (Ref 1). Such studies have been extensive and contributed to advance theory (Ref 2–4) as well as provided benchmark data used for modeling (Ref 5, 6). Despite their impact, however, these model systems are far from complete as analogs to metals and alloys; several of their physical properties tend to differ distinctively from those of metals, for example, solute/solvent mass ratios, heat conductivities and capacities. Perhaps the most important difference concerns their Prandtl numbers, the ratio between viscous and diffusive properties, which are typically 3 to 4 orders of magnitude larger for the model systems. Therefore, solidification processes in metals may often exhibit a greater complexity than in the transparent organics, since buoyancy-driven flow is much more prominent in the former category. Accordingly, it is of significant interest to find and develop techniques that allow for direct in situ observations of real metallic alloys.
X-Ray Imaging Metal transparency and interaction with x-rays have long been recognized as obvious candidate principles from which methods for in situ monitoring of solidification processes could be developed. Yet, the limited x-ray
source brightnesses and detection efficiencies available have hampered the practical impact of x-rays as a diagnostic tool for imaging studies at the characteristic time and length scales of microstructure evolution. A general description of contrast formation by x-ray transmission in matter can be given by a wave-optical approach. Let Ei(x, y, z) be an incident monochromatic plane-wave propagating along the x-axis of a Cartesian laboratory coordinate system. The transverse cross sections of the incident wave-field, Ei (y, z), and the field transmitted through the sample, Et(y, z), are related via a complex sample transmission function, F(y,z): Et ðy; zÞ ¼ Fðy; zÞEi ðy; zÞ
(Eq 1)
with Fðy; zÞ ¼ Aðy; zÞ expðijðy; zÞÞ
(Eq 2)
The amplitude, A(y,z), and phase, j(y,z), relate to the x-ray absorption and refractive properties of the material, respectively, and can be expressed as: ð 1 Aðy; zÞ ¼ exp mðx; y; zÞdx 2 2p ðnðx; y; zÞ 1Þdx l
(Eq 3)
In-House Laboratory X-Ray Imaging of Solidification Processes
ð
jðy; zÞ ¼
105 or less. However, provided that phase contrast could be used, it would help to resolve phases with similar atomic compositions but adequate differences in density. Current state-of-the-art x-ray imaging detectors operate with limiting spatial resolutions similar to optical microscopes. Hence, for any phase contrast to be resolvable by the camera system, an extremely high degree of transversal coherence of the incident wave-field is required, corresponding to at least 105 l. Since a highphoton flux also must be maintained at the sample to allow for real-time studies, conditions for phase contrast can only be delivered with modern high-brilliance synchrotron radiation (Ref 7). Finally, it should be noted that the quantity measured in x-ray imaging experiments is the intensity Iðy; zÞ / Et ðy; zÞ Et ðy; zÞ ¼ jEt ðy; zÞj2 , with Et as the complex conjugate Et in Eq 1. Since the actual phase field j(y,z) is not retained in the measurements, true phase contrast imaging requires some form of mathematical reconstruction. Nevertheless, effects of phase contrast may be directly visible in I(y,z) from interference between different exit wavefield components.
(Eq 4)
where m is the linear absorption coefficient, n is the real part of the refractive index, and l is the x-ray photon wavelength. With Z as the atomic number and ra as the atomic density, m can be shown to vary roughly as raZ 4 l3, except at energies corresponding to photoelectric excitations, where it is discontinuous. Thus, absorption contrast is mainly available between phases with a distinct difference in atomic composition. The (n 1) refractive properties follow approximately a raZ l2 dependence in between the discontinuities occurring at photoelectric excitations and remain extremely weak with all materials, deviating from unity by only
The first in situ x-ray investigations of solidification were based on conventional in-house source radiography and used for studies of solute redistribution and boundary layer propagation (Ref 8, 9). The spatial resolutions available, Dr 100 mm, prevented detailed studies of solid-liquid interface morphologies, but the resolution in x-ray absorption contrast was adequate to verify proximity to Scheil-Gulliver conditions as well as to demonstrate the influence of buoyancydriven flow on the solute boundary layer by comparative measurements varying the growth direction relative to gravity (Ref 9). Later, the resolution of in-house-based radiography setups have been improved significantly by the introduction of microfocus sources. Such setups have
358 / Principles of Solidification been used to study topics like convection and microstructure evolution in gallium-indium eutectics (Ref 10, 11), hydrogen porosity formation in aluminum-copper (Ref 12), and liquid droplet formations, striations, and engulfment in aluminum-base monotectics (Ref 13, 14) reporting Dr 25 to 30 mm. In a recent real-time in-house study of metal foaming, Dr 5 mm was reported (Ref 15). As the nominal spatial resolution is improved, the temporal resolution and the source brilliance become more critical factors. Temporal resolution here refers to either: 1. the inverse of the frame-grabbing rate, that is, 1/(Dtexp + Dtr.o.), where Dtr.o. is the read-out dead-time of the camera system, when each frame is subjected to analysis; or 2. the accumulation of the former when several frames must be integrated to provide image statistics sufficient for analysis. The temporal evolution of the system under investigation will lead to blurring whenever vobj(r) Dtexp > Dr, where vobj(r) is the motion of a contrast object in local position r in the sample, relative to the x-ray camera system, and Dtexp is the exposure time. Hence, higher spatial resolutions can be obtained only by reducing the exposure time or by slowing down the dynamics and kinetics of the system. The potential for reduction of Dtexp is determined by the x-ray source brilliance, the photon-detection quantum efficiency and signal-to-noise of the detector system, and finally the transmission and contrast provided for by the sample under investigation. The latter is shown by Eq 1 to 4 to be a function of both the x-ray source and the energy in use, as well as the intrinsic x-ray properties of the different sample constituents. Decent spatial resolutions are possible to realize in laboratory studies of metal foaming and hydrogen porosity, only because the ratio between the x-ray attenuation lengths of H2 gas and metal are adequately large, 106 or more at relevant energies. The x-ray attenuation length of a medium, at a particular monochromatic x-ray photon energy, is defined as the length of propagation where photoelectric absorption by the medium has reduced the xray beam amplitude by a factor of e1, with e as the natural exponential base. In comparison, with E varying from 10 to 50 keV, the largest ratio found between x-ray attenuation lengths of any two base metals (magnesium and heavier) is 100. Consequently, high-resolution in situ studies with fully metallic samples usually require longer exposure times or improved performance of the x-ray sources and detectors. Nevertheless, in a recent study of directional solidification in gallium-indium, employing a state-of-the-art microfocus source and detector, a spatial resolution of Dr 10 mm and a temporal resolution of Dt = (Dtexp + Dtr.o.) 0.5 s was used to investigate dendritic
growth and analyze the buoyancy-driven flow of rejected solute (Ref 16).
X-Ray Imaging Studies of Solidification Processes with Synchrotron Radiation The conditions for in situ x-ray imaging are vastly improved by the eminent source brilliance and x-ray energy selectivity offered at modern synchrotron radiation (SR) sources. These facilities have opened for the use of x-ray imagingbased techniques to investigate interface morphology evolution, solute transport, and various process phenomena at spatiotemporal resolutions gradually approaching those of light optical video microscopy. In general, there are three distinct viable imaging techniques made available by SR for the real-time investigation of solidification microstructures in alloys: two-dimensional x-ray topography, two-dimensional x-ray radiography, and ultrafast three-dimensional x-ray tomography.
X-Ray Topography X-ray topography is a diffraction-based imaging technique, in principle, the x-ray variant of dark-field electron microscopy (for a review of the technique, see Ref 17). Synchrotron whitebeam x-ray transmission topography was first employed in solidification science for in situ studies of single-crystal melting-solidification processes in pure tin (Ref 18) and aluminum (Ref 19). Topography may be carried out in reflection or transmission geometry and with a nonmonochromatic (white beam) or monochromatic radiation, where the combination of a white beam and transmission geometry normally would be preferred for in situ solidification studies. Other applications include studies of planarcellular and cellular-dendritic morphology transitions during directional solidification in aluminum-copper alloys (Ref 20), including the role of strain to the onset of interface perturbations (Ref 21, 22) (Fig. 1). The main advantage of topography over the other x-ray imaging techniques is that it is unique in providing in situ information on microstructure stresses induced during solidification. Also, contrast is available at solid-liquid interfaces even for dilute alloys and pure substances, since only the solid gives rise to diffraction signals. On the other hand, in situ topography is severely limited by the output signal strength. In the x-ray energy regime below 100 keV, the elastic x-ray scattering cross sections of any element are 103 or less of their absorption cross-sectional counterparts. Typical obtainable spatiotemporal resolutions for in situ studies of metals, employing state-of-the-art live-capturing x-ray cameras and a white-beam transmission mode, are of the order Dr 15 mm and Dt 2 s.
Fig. 1
(200) reflection x-ray topograms showing local cell rotation during directional solidification of an Al-5.7wt%Ni alloy. (a) t0. (b) t0 + 360 s. Source: Ref 22
X-Ray Radiography Currently, synchrotron x-ray radiography is the technique most frequently used for in situ studies of alloy solidification microstructures and phenomena. It offers superior spatiotemporal resolutions, down to Dr 1.5 mm and Dt 0.25 s, and is reasonably flexible, provided that the alloy system under investigation contains constituent phases that can be resolved by some form of transmission contrast. With third-generation SR sources, a high degree of transversal beam coherence can be obtained, facilitating the use of x-ray phase contrast in addition to conventional absorption contrast (Ref 7). Although actual phase-contrast objects have to be reconstructed, by employing sample-camera distances within the so-called near-field of the first Fresnel region, low-order interference fringes between waves refracted in slightly different directions across phase boundaries can appear directly in the radiograms as sharp contrast features (Ref 7, 23). Thus, direct radiographic contrast enhancement can be obtained locally at interfaces and can resolve boundaries between aggregates of similar atomic compositions. Provision of additional interference contrast at phase boundaries is particularly advantageous to studies of curved fronts, because any cell or dendrite will exhibit a loss in transmission contrast at the interface due to its three-dimensional curvatures. Unlike radiation from conventional sources, synchrotron radiation provides a high-photon flux over a continuous energy band. In most cases, monochromatization is affordable, permitting selection of an x-ray energy that yields an optimum compromise between sample transmission and contrast. The first in situ SR radiography experiments in solidification science were directional solidification investigations of various tin-lead alloys to study evolution of columnar and equiaxed dendritic microstructures (Ref 24). In addition to the characteristics of the source, access to a high-resolution, fast-readout, low-noise charge-
X-Ray Imaging of Solidification Processes and Microstructure Evolution / 359 coupled device camera specially designed for time-resolved studies (Ref 25) was critical to the outcome of the experiments. Successive studies with the same setup on various aluminum-copper alloys 1. microstructure evolution during primaryphase columnar and equiaxed dendritic growth, as well as stable and destabilized eutectic growth (Ref 25); 2. destabilization of columnar dendritic fronts by local, gravity-induced settlements of rejected solute, including spatiotemporally resolved analysis of the solid-liquid interface morphology and chemical analysis of the liquid (Ref 26), shown in Fig. 2; and 3. dendrite detachment initiated by various mechanisms in different regions of the mushy zone (Ref 27), with particular emphasis on the role of solute pile-up and buoyancy-driven flow of fragments to promote mesoscale segregation, columnar front blocking, and the initial stage of columnarto-equiaxed transition (Ref 27, 28). In situ synchrotron radiography has also been used to study fragmentation (Ref 30), coarsening (Ref 31), and temperature gradient-zone melting (Ref 32) during directional solidification of tinbismuth alloys; columnar-to-equiaxed transitions and equiaxed dendritic growth in grainrefined and nonrefined aluminum-nickel (Ref 33, 34); solidification dynamics during faceted growth of AlPdMn quasi-crystals (Ref 34, 35); coarsening during semisolid processing of grain-refined aluminum-germanium (Ref 36); equiaxed dendritic growth in grain-refined aluminum-copper (Ref 29); and finally monotectic solidification in aluminum-indium (Ref 37) and aluminum-bismuth (Ref 29) demonstrating how hydrodynamics, short-range droplet-droplet interactions, and coagulation mechanisms apparent in the two-phase immiscible liquid zone can influence the monotectic reaction and perturb the primary-phase morphology.
X-Ray Tomography Ultrafast x-ray tomography is the most exclusive novelty offered with modern synchrotron facilities. Radiography and topography experiments provide two-dimensional transmission projected images of the sample. In a tomography experiment, a sequence of two-dimensional radiograms are collected while the sample is pivoted by a few hundred increments over a 180 rotation about an axis perpendicular to the propagation direction of the x-ray beam. Each radiogram provides the transmission contrast (amplitude and/or phase) corresponding to a particular sample orientation. Thus, by applying a suitable reconstruction algorithm, the three-dimensional structure of the sample can be rendered. With modern synchrotron sources, complete sets of projection data can be acquired within 10 s, providing a full threedimensional set of approximately 109 pixel
Fig. 2
X-ray radiograms showing planar eutectic and columnar dendritic growth during directional solidification parallel with gravity in an Al-30wt%Cu alloy. Left column: every fifth image from a part of the sequence representing the microstructure evolution at times t0, 2.25 and 4.5 s. The images also display the local liquid concentration, Cl(Cu), in a relative manner (as indicated by the horizontal contrast bar). Right column: examples of quantitative information extracted by image processing. (a) Liquid constitution contour map at 4.5 s, starting at Cl = 33 wt%Cu (black) and with increments of 0.5 wt%Cu. (b) Velocities for the eutectic front and columnar tips 1 to 3, enumerated from left to right, normalized against the sample pulling velocity, vsp 22.5 mm/s. (c) The time evolution of the two-dimensional solid-liquid interface projections over the full video sequence. Source: Ref 26, 29
360 / Principles of Solidification values, with an attainable three-dimensional spatial resolution of 2 mm (Ref 38). However, in order to achieve such spatial resolution in studies of evolving systems, the contrast object shapes must remain semisteady within the 2 mm resolution limit over the 10 s time required to collect a complete tomogram. Another complication is that, while topography and radiography with state-of-the-art x-ray cameras allow for direct on-line observation of the systems under study, tomography visualization is indirect because reconstruction is required. Hence, it is less flexible with respect to adjustment and fine-tuning of experimental parameters; and on-line adjustments are practically unavailable. The first demonstration of ultrafast in situ tomography in solidification science was a study of the solidification of inoculated Al-4wt%Cu alloys (Ref 39), employing a white beam to maximize the incident flux, and reporting a spatial resolution of 3 mm with a 10 s tomogram acquisition time. Within these limits, any interface velocities beyond 0.3 mm/s will cause temporal blurring, and accordingly, a slow cooling rate of 0.1 K/s was used. Image processing allowed for extraction of quantities such as three-dimensional solid-liquid interface morphologies, solid fraction versus temperatures, solid-phase shrinkage upon cooling, and copper concentration in both liquid and solid, although the latter is likely to be somewhat hampered by the fact that a white beam was used. When using a white beam, chemical analysis requires corrections for both the x-ray energy shift of the sample and the nonlinear energy
Fig. 3
response of the detector. Also, quantitative analysis had to be limited to the later stages of solidification ( fs > 0.5). In general, early stages of solidification are difficult to assess with tomography, since these tend to involve growth velocities beyond the blurring limit, and any gravityinduced relative solid-liquid flow prior to the onset of a coherent crystal network could lead to blurring by motion of the grains. More recent studies have focused on coarsening effects during partial remelting in Al-15.8wt%Cu (Ref 40) and solidification experiments in nonrefined Al-7wt%Si-10wt% Cu alloys (Ref 41), shown in Fig. 3. In these studies, cooling rates up to 3 K/s were used, while the freezing range suitable for analysis was limited between 0.18 > fs > 0.28. The lower boundary is set by the limiting contrast required for image processing, extended here as compared to Ref 39 by the heavier copper alloying. The upper limit corresponds to the point where eutectic silicon starts to grow on the aluminum dendrites, at which time further analysis is impaired since hardly any contrast can be obtained between aluminum and silicon. Coarsening phenomena have also been studied during semisolid processing of aluminum-germanium alloys, combining in situ real-time radiography and semistatic tomography (Ref 36), thus yielding information on the temporal evolution of the system by radiography and snapshots of the three-dimensional microstructure at discrete processing times by tomography. Other efforts have been made to combine different in situ x-ray imaging
Ultrafast tomography. Top row shows two-dimensional regional slices taken from the three-dimensional tomograms (size: 1.4 by 1.4 by 1.4 mm3, or 0.00009 by 0.00009 by 0.00009 in.3) at solidification times of 60, 660, and 1260 s. Bottom row shows three-dimensional reconstructions of a dendrite at solidification times of (a) 660, (b) 840, and (c) 1020 s, localized in the region of interest indicated in the 660 s slice. The grayscale in the tomograms displays the solid curvature variations; a region with prominent coarsening is highlighted. Source: Ref 41
techniques to obtain complementary information, for example, simultaneous radiography and topography (Ref 42); however, the spatiotemporal resolutions obtainable with the combined experiment are limited since the x-ray beam requirements set by topography are nonoptimal with respect to in situ radiography.
Further Reading The reader is advised to consult the original works cited for more details on the particular qualitative and quantitative results obtained in the various studies, many of which provide extensive comparisons of their results to governing theories. The original works also contain a vast amount of details on the explicit experimental conditions used for the various studies, that is, x-ray characteristics, furnace systems, detectors, and other information regarding equipment or experimental procedures. Finally, a few video demonstrations of real in situ observations of solidification microstructure evolution and various process phenomena can be accessed in the on-line version of Ref 29, distributed via the TMS website (http://www. tms.org/TMSHome.html).
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H. Klein, J. Ha¨rtwig, B. Grushko, B. Billia, and J. Baruchel, Application of Synchrotron X-Ray Imaging to the Study of Directional Solidification of Aluminium-Based Alloys, J. Cryst. Growth, Vol 275, 2005, p 201–208 H.N. Thi, J. Gastaldi, T. Schenk, G. Reinhart, N. Mangelinck-Noe¨l, V. Cristiglio, B. Billia, B. Grushko, J. Ha¨rtwig, H. Klein, and J. Baruchel, In-Situ and Real-Time Probing of Quasi-Crystal Solidification Dynamics by Synchrotron Imaging, Phys. Rev. E, Vol 74, 2006, article 031605 S. Zabler, A. Rueda, A. Rack, H. Riesemeier, P. Zaslansky, I. Manke, F. GarciaMoreno, and J. Banhart, Coarsening of Grain-Refined Semi-Solid Al–Ge32 Alloy: X-Ray Microtomography and In Situ Radiography, Acta Mater., Vol 55, 2007, p 5045–5055 H. Yasuda, T. Nagira, M. Yoshia, A. Sugiyama, I. Ohnaka, K. Kawasaki, K. Umetani, and K. Uesugi, TimeResolved X-Ray Imaging of Dendritic and Monotectic Solidification of Sn-, Al- and Cu-Based Alloys, Proc. Int. Conf. Solidification and Casting of Metals (Sheffield, U.K.), July 2007, p 306–310 M. DiMichel, J.M. Merino, D. FernandezCarreiras, T. Buslaps, V. Honkimaki, P. Falus, T. Martins, and O. Svensson, Fast Microtomography Using High Energy Synchrotron Radiation, Rev. Sci. Instrum., Vol 76, 2005, article 043702 O. Ludwig, M. DiMichel, L. Salvo, M. Sue´ry, and P. Falus, In-Situ ThreeDimensional Microstructural Investigation of Solidification of an Al-Cu Alloy by Ultrafast X-Ray Microtomography, Metall. Mater. Trans. A, Vol 36, 2005, p 1515– 1523 N. Limodin, L. Salvo, M. Sue´ry, and M. DiMichiel, In Situ Investigation by XRay Tomography of the Overall and Local Microstructural Changes Occurring During Partial Remelting of an Al-15.8 wt.% Cu Alloy, Acta Mater., Vol 55, 2007, p 3177–3191 N. Limodin, E. Boller, L. Salvo, M. Sue´ry, and M. DiMichel, In-Situ Fast X-Ray Tomography Investigation of Microstructural Changes Occurring During Partial Remelting and Solidification of Al-Cu Alloys, Proc. Int. Conf. Solidification and Casting of Metals (Sheffield, U.K.), July 2007, p 316–320 H.N. Thi, H. Jamgotchian, J. Gastaldi, J. Hartwig, T. Schenk, H. Klein, B. Billia, J. Baruchel, and Y. Dabo, Preliminary In Situ and Real-Time Study of Directional Solidification of Metallic Alloys by X-Ray Imaging Techniques, J. Phys. D, Appl. Phys., Vol 36, 2003, p A83–A86
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 362-369 DOI: 10.1361/asmhba0005220
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
The Oxide Film Defect—The Bifilm John Campbell, University of Birmingham England
MOST METALS start their lives in the liquid state. Once melted, they are usually subjected to various transfers between furnaces or ladles, involving pouring or other types of surface turbulence. These actions entrain (fold in) the surface film to create entrainment defects. These are principally: Bubbles (that in turn create bubble trails) Bifilms, which are doubled-over surface
films that act as cracklike defects
Various entrained debris, traditionally known
as exogenous or external inclusions The bifilm defect is the subject of this article. It appears to be a common but serious and almost overlooked casting defect. Large populations of bifilms are usually introduced into metals at an early stage of their production. In general, their presence has been unsuspected because, even though they can have a large area, they often are only nanometers thick and not easily detected by conventional nondestructive techniques. They can be serious defects because they act as cracks in the liquid, and the cracks become frozen into the casting. The presence of cracks in suspension in liquid metal explains many otherwise inexplicable features of cast products, such as porosity, hot tearing, the morphologies of second phases, and impaired reliability of mechanical properties. The fundamental difference between such entrained defects (associated with a macroscopic unbonded interface) and defects and inclusions grown in the melt is of central significance for the failures of metals by mechanical or corrosion-type mechanisms. Analysis of bifilms provides a simple, powerful, and elegant concept to explain many features of the metallurgy of castings. Since bifilms occur mainly during turbulence involved in the pouring of liquid metals, they similarly can be avoided by avoiding turbulence. Although the effects of bifilms can be reduced (but often not eliminated) by expensive postcasting operations such as hot isostatic pressing or working, the major future potential lies in techniques for their avoidance. Some casting operations are already taking the first steps in such new technology and benefiting technically and commercially.
Bifilms The oxide film on the surface of a liquid can become entrained in the bulk liquid. The entrainment action is an enfolding of the surface, leading automatically to the folding-in and submersion of its surface film (Fig. 1). Such surface turbulence leads to entrainment events that are usually rapid, requiring only milliseconds. The creation of a new surface and the time available for the formation of a new oxide film on this fresh surface is so limited that the entrained film can be very thin, of the order of nanometers (Ref 1). Most films on liquid metals are oxides (although nitrides and carbon films are also common), and most are solid at the temperature of the melt. When entrained, the dry surfaces of the films fold inward, necessarily facing each other and making a double film, hence the name bifilm. The opposing halves of the film cannot bond, being a ceramic-to-ceramic interface; thus, all entrained oxides effectively constitute a crack (feature “A-B” in Fig. 1) in the liquid. Most pouring actions involve such severe surface turbulence that they fill the liquid with bifilm cracks. The bifilms are usually invisible and often undetectable by most techniques because the surface oxides are usually only a few molecules thick, approximately 20 nm, even though they may be square millimeters or square centimeters in extent. Occasionally, bifilms that are fractions of a meter (3 ft) across are seen in large alloy steel and nickel-base alloy castings. If they extend across the wall of the casting, they can form efficient leak paths and are the main reason for the rejection of castings for leakage faults. If bifilms are so common and their effects are so serious, how is it that they have been universally overlooked until now? The fact is that they are difficult to detect with standard metallurgical techniques. Despite decades of experience with the scanning electron microscope (SEM) in examining fractures, no one during the early years was aware of the presence of bifilms, with the predictable consequence that none were observed. Now that the features of bifilms are known, it is possible to recognize these features
and find bifilms in profusion on nearly every metal fracture surface that is examined. Furthermore, the bifilms are sometimes so extensive that they cover 100% of the fracture surface. (An example in Fig. 7 is described later.) They can remain largely invisible even in the SEM because they are so thin that the structure of the metal is clearly seen through the film. The presence of a film is seen only at the highest magnifications, where the characteristic creases and folds point to the presence of a defect that can be identified as a film originally generated on the liquid surface. Its near-mirror image resides on the opposite fracture surface as the “butterfly wing” pattern, which can often be seen if the fracture surfaces are placed together. Its characteristic fluorescence in the electron beam and its chemical composition are further identifying characteristics. Bifilms are easily opened and observed when the metal is in the liquid state. Fox and Campbell (Ref 2) assumed that there would be a thin layer of air (or other residual gas) trapped between the halves of the bifilm. The gas would be retained in both the minute surface roughness and the larger irregular folds and creases. Therefore, they predicted that if the pressure above the melt were reduced by a factor of 10, the gas entrained between the halves of the bifilm would be expanded by a factor of 10. Fox reduced the pressure to 0.01 atm
Fig. 1
Schematic representation of surface turbulence causing the entrainment of bifilms and bubbles. Small entrained bubbles form pores that decorate the bifilm crack “A-B.” Large bubbles, “C” and “D,” are buoyant, creating bubble trails.
The Oxide Film Defect—The Bifilm / 363 (1 kPa) and thus expanded the bifilms by 100 times, sufficient to make them observable in a radiographic study (Fig. 2). In subsequent work, the factor-of-10 reduction in pressure (down to 0.1 atm, or 10 kPa, residual pressure) has been found to be adequate, allowing the reduced pressure test (RPT) to be a powerful technique for observing and quantifying bifilm size and number. Bifilms, when searched for, can usually be found in large numbers. In aluminum alloys, a concentration of 109 m 3 (1.0 per mm3, or 16,000 per in.3) appears average, and even 1012 m 3 (1000 per mm3, or 16 million per in.3) is relatively common. A bifilm concentration of 1015 m 3 (1.0 million per mm3 or 16 million per in.3) would be necessary if a separate bifilm were required to be in place to initiate the fracture of the microscopic eutectic silicon particles that are commonly seen on fracture surfaces of aluminum-silicon alloys. As an alternative quantifying measure, the total length of bifilms in a standard sample can be measured. A convenient standard sample is that produced by the RPT, weighing approximately 100 g (0.2 lb). When sectioned vertically through its center and polished, its area of approximately 35 by 35 mm (1.4 by 1.4 in.) can be scanned using an automated metallographic system. In this way, the number can be counted and the average maximum length of the defects can be determined. Such inspections may reveal the total length of bifilms on the sampled surface to be only 0 to 1 mm (0 to 0.04 in.) in a high-quality melt, whereas a much worse length of 50 mm (2 in.) is more common. However, the length sometimes approaches 500 mm
Fig. 2
(20 in.) (Ref 3). The latter figure corresponds approximately to a sobering content of 5 mm2 (0.008 in.2) of bifilm per cubic millimeter. This work illustrates not only the wide spectrum of numbers and sizes but also that populations of very different sizes of bifilms occur because of different melt conditions. In addition, different populations accumulated from different stages of manufacture often coexist in the same sample. It is important to keep in mind that bifilms are not always oxides. Sometimes, they are nitrides, sulfides, amorphous carbon, graphite, or other materials. They are not always solid but can be partly solid or even liquid and, in some cases, not present at all. A bifilm would not be present when casting pure gold or platinum, since these noble metals do not possess surface oxides. Alternatively, a metal may contain no oxide films if it has been filtered or handled extremely carefully so as to avoid entraining its surface. When bifilms are present, as is usually the case, the quantity can be controlled to some extent by careful handling of the liquid metal. Entire commercial casting operations have been revolutionized by attention to the metal handling and pouring practices. The production of scrap castings has been reduced while simultaneously improving properties, usually with a reduction in costs.
Entrainment The minimum criterion for entrainment is a critical Weber number (We > 1), in which the inertial forces in the liquid that are tending to
Comparison of x-ray radiographs taken from the same melt of Al-7Si-0.4Mg solidified at the same time under (a) 1 atm (100 kPa), showing diffuse images of furled bifilms, and (b) 0.01 atm (1 kPa), showing unfurled bifilms. Source: Ref 2
perturb the surface exceed the surface tension forces that tend to restrain the surface, holding it in a compact form. In practice, the condition works out to be equivalent to a critical velocity, above which the surface can be propelled upward, achieving a height sufficient that it may enfold its surface as it falls back under gravity. This critical velocity is in the region of 0.5 m/s (1.5 ft/s) for nearly all liquids and is a useful parameter that is now being built into many filling system designs for castings, with excellent results if carried out correctly (Ref 4). At the present time, practically every pouring event violates the critical velocity criterion and introduces oxides to the melt. This is simply the result of the acceleration under gravity causing the critical velocity of 0.5 m/s (1.5 ft/s) to be exceeded if the fall of the liquid exceeds approximately 10 mm (0.4 in.). This is the critical fall height. This small distance is the reason that castings must be bottom gated at all times. Gating at the joint, if part-way down the casting, or anywhere above the floor of the casting risks the entrainment of defects. A second mechanism for entrainment of bifilms is the contraction of the surface, which happens, for instance, as the surface expands and contracts because of random perturbations (Ref 1). Such disturbances occur continuously in many melting operations. Each time the surface expands, more film area is created as its area stretches. However, each time it contracts, the newly created oxide raft cannot contract and therefore must fold, crumpling concertina-fashion. With continuing disturbances of the surface, the folds lengthen irreversibly in a kind of natural ratcheting mechanism. Eventually, the folds become sufficiently long to be dragged as cracklike defects into the bulk liquid. Turbulent processes such as semilevitation melting, and possibly even conventional induction melting, are likely to introduce large densities of defects. Clearly, the casting industry needs techniques that avoid the introduction of such damage. At the least, pouring transfers must be reduced, and, if necessary, the damage introduced must be reduced or reversed by additional treatment. However, there are now casting systems that eliminate pouring and other turbulence in the whole melting and casting operation, with significant benefits to the properties and reproducibility of products (Ref 5). One of the central features of surface entrainment is that the events are cumulative. For instance, for aluminum alloys, the first turbulence suffered by the material often occurs at the aluminum primary smelter. This first addition of bifilms is supplemented by a second population introduced by the prealloying producer. Several more populations follow in the foundry by melting, pouring into a transfer ladle, and finally pouring into the mold. Naturally, some of the original defects are lost by sedimentation, flotation, or by attachment to the walls of containers. Some may be lost by
364 / Principles of Solidification filtration, although it is known that filters have only a limited ability to trap inclusions (Ref 1). The remelting of foundry returns, and especially machining chips, adds to the total bifilm concentration. However, in this way, aluminum alloys are special. They have a surface oxide film that is marginally more dense than the liquid, but the entrained air in the folds of the bifilm means that many bifilms are near-neutral buoyancy. This feature, combined with the nearly zero Stokes velocity because of the high drag factor provided by its film morphology, means that separation by floating or sinking is especially slow, taking hours or days in conditions where the melt is gently convecting or, in the case of induction melting, strongly stirring. In addition to the problem of the slow separation by buoyancy when suspended freely in a liquid, the situation becomes much worse when the density of defects is such that they form an interconnected latticework throughout the melt. Such a situation is seen in an aluminum alloy in Fig. 2. These frameworks in the liquid are expected to be rather stable with time, since the mutual contacts mean that flotation or sedimentation is now difficult. Most liquid metals and alloys (containing normal populations of defects of interest in this account) pour similar to water because the viscosity of most engineering metals is not significantly different from that of water. Unfortunately, the sight of a melt pouring like porridge, yet at a temperature well above its melting point, is not that uncommon. It is a sign of the presence of an extremely high density of entrained films. (The mechanical properties of such cast material have not been reported in the literature but, in the experience of the author, are disastrously poor.) Some stainless steels, particularly the superduplex stainless types, have especially strong and tough oxide films that create macroscopic problems when pouring; occasionally, bifilms are found to form a linked network completely through cast walls over 100 mm (4 in.) thick. These are first indicated by radiography as areas of apparent shrinkage or sponge porosity. However, when the material is excavated by grinding and tested by dye penetrant, it reveals the appearance of a spider’s web of cracks. If the casting is to be salvaged, such material must be completely dug out and refilled with weld metal. Most other metals are not so disadvantaged. For instance, in the case of carbon/manganese steels, entrained defects are very different in density, thus providing a significant driving force for separation by flotation. However, when generous quantities of aluminum are used for deoxidation, the entrained alumina films still have a high drag because of their aspect ratio, and so separate slowly. If the steel is deoxidized with calcium after the aluminum addition, the surface film of solid alumina is converted into a low-melting-point eutectic of mixed alumina and calcia. The folding-in of this surface film is then simply the wetting
of two liquids that immediately mutually assimilate, forming spherical droplets that float out with maximum speed. As a result, calciumtreated steels are much cleaner than steels simply deoxidized with aluminum (Ref 6). In addition, those residual oxides that do not manage to detrain constitute the favored spherical-shaped inclusions that yield improved workability (Ref 7). Other steels, particularly the common carbon and low-alloy steels, have surface oxides that are partly liquid because of their compositions and the high casting temperatures. An entrainment event now results in the surface film adhering to itself as it is furled up by the internal turbulence in the liquid. The entrained film becomes a compact, sticky ball. This bifilm is glued shut and cannot unfurl again. Having a compact form, it can float out quickly, forming the large drossy or slaglike defects sometimes known as reoxidation defects, which are often centimeters in diameter. These decorate the cope surfaces of large steel castings and must be laboriously dug out and back-filled with weld metal (Ref 8). The sharp difference in behavior between cast aluminum alloys and cast steels is probably highly significant. The rapid detrainment of bifilms in most types of steels explains their generally high levels of cleanliness and their high ductilities, usually between 25 and 50%. Cast aluminum alloys frequently have high levels of bifilms because they are so sluggish to separate, resulting in ductilities often measured in low single figures, or even 0%. Only in the last few years, by using extremely careful handling in a specially designed plant, have higher ductilities in cast aluminum alloys been reproducibly produced in the range of 10 to 20%. The prospects of aluminum alloys matching the ductilities of steels are beginning to seem achievable.
Furling The ravelling (or furling) of the entrained bifilm into a compact mass is almost certainly a standard feature of the history of each bifilm. When a bifilm is entrained, the internal turbulence beneath the surface of the liquid will be strong. Because the bifilms are thinner than tissue paper, they have little rigidity of their own, so they are easily scrambled up by the turbulent eddies in the heavy flowing liquid. As they tumble along channels into a mold, they are expected to continuously ravel and unravel in different ways (Ref 1). In this compacted, convoluted form, a large bifilm may be able to pass through a relatively fine filter, explaining to some extent the partial effectiveness of filters. Convoluted, tangled bifilms are commonly seen in micrographs of cast light alloys. Radiographs of such compacted bifilms are seen as faint, diffuse images in Fig. 2. Also, while the bifilm is furled in this way, it is still a crack but nothing like a familiar stress crack. Its convoluted form means that its
deleterious effect on mechanical properties of the solidified cast metal is minimized, its effective crack length is minimized, and the effect of the crack is reduced because of the mechanical interlocking of its complex form. The bifilm only develops its maximum damaging effect later, as described in the next section.
Unfurling Once the pouring phase of a transfer is finished, or a casting operation is complete so that the melt has arrived in the mold, internal turbulence subsides, and a period of quiescence follows. During this quiet period, the bifilm may experience a number of influences that encourage the bifilm to unravel, opening fold by fold until it forms a substantially planar cracklike defect close to its original size when first entrained. This unfurling of the compacted defect is a critically important stage. This process defines the properties of many cast materials. In the simplest case, if freezing occurs quickly after the filling of the mold, the bifilms are frozenin as convoluted cracks of overall small, compact shape before they have a chance to unfurl, rendering them relatively harmless. Mechanical properties are therefore usually good. Everyone in a light-alloy foundry knows that the properties of chilled castings are better than those of slowly cooled castings. However, if freezing occurs slowly, the bifilm has time to unfurl, expanding to take up its full size as a planar crack and thus regaining its potential to seriously reduce properties. There are a number of reasons why bifilms unfurl, tending to return to their starting size. An understanding of the unfurling mechanisms is key to the understanding of casting properties. The mechanisms are discussed subsequently.
Gas Precipitation from Solution (Formation of Gas Microporosity) It is worth emphasizing that porosity in liquid metals cannot occur as a result of any classical nucleation processes, such as those involving the rearrangement of atoms (alternatively stated, the condensation of vacancies) to create an embryonic bubble. The impossibility of homogeneous nucleation is, of course, widely accepted. What may be less evident is that classical heterogeneous nucleation, although clearly significantly easier, is still not sufficiently easy to make pore nucleation possible in cast products (Ref 1). Thus, nucleation of porosity by any conceivable process is almost certainly impossible. In other words, interfaces in atomic contact are probably impossible to separate in any practical circumstance. This is because, despite any degree of wetting (i.e., assessed macroscopically by such concepts as contact angle), the bonding between atoms on an atomic scale is sufficiently strong to be inseparable in any conceivable practical situation (Ref 9). It is worth reflecting that all solid phases
The Oxide Film Defect—The Bifilm / 365 that have precipitated and grown in the liquid have grown atom by atom and thus are necessarily in perfect atomic contact with the matrix. None of these can be favorable substrates for pore formation or decohesion. Conversely, if an interface is not in atomic contact with the liquid, the separation of this interface would be expected to be an easy initiation site for a pore or a decohesion event. However, such an interface cannot have been grown in the liquid; therefore, it must have been an entrained interface. As such, it will be separated from the liquid by an unbonded oxide film, together with a separating layer of residual entrained gas. Figure 3 illustrates how entrainment always introduces nonbonded interfaces, making such inclusions fundamentally different from those grown in the melt or in the solid. All such unbonded, entrained interfaces will constitute excellent initiation sites for pores. Therefore, bifilms are excellent initiation sites for porosity formation. It can be quickly appreciated that the separation of these unbonded films does not resemble nucleation in the classical sense. No nucleation barrier exists, as in the case of true nucleation phenomena; no critical embryo exists; and no surface tension or surface energy considerations apply, because no liquid surfaces are involved. The presence of the bifilm means that the pore simply and gradually inflates (i.e., the collapsed, folded balloon slowly becomes more like a sphere), as illustrated in Fig. 4. If the continued expansion of the bifilm occurs at a late stage of solidification, its expansion is constrained by the presence of surrounding dendrites, thus forming interdendritic porosity (Fig. 4d). It can be immediately appreciated that the conventional division of round pores being attributed to gas, and interdendritic pores to shrinkage, is not correct. The different morphologies merely result from the round pore opening early and the interdendritic pore opening later during the progress of freezing. One clear example of the variation in bifilm unfurling rates is seen in the fascinating morphology of subsurface porosity in castings. This
is a common condition in which an array of pores lies just 1 or 2 mm (0.04 or 0.08 in.) under the surface of the casting. They are grown by the flux of hydrogen diffusing into the casting from reaction between the metal and its mold environment as a result of the thermal degradation of its organic binder (Fig. 5a). The pores grow in a rather uniform environment of hydrogen flux and are therefore gas pores. The 1 to 2 mm (0.04 to 0.08 in.) depth is almost certainly the result of the increasing concentration of gas ahead of the advancing solidification front, similar to snow gradually piling ahead of the snowplow, until a critical concentration is reached at which the pore can form (Ref 1). Curiously, the pores do not grow to a uniform shape and size. Some display a round morphology like a bubble (Fig. 5b), while others have a dendritic morphology that may be assumed to be a shrinkage pore (Fig. 5c). Clearly, the round pores have nucleated early while still surrounded by liquid, but the dendritic pores have formed late, after the dendrites have formed. These differences in the time of pore initiation cannot be the result of differences in growth conditions, since the pores often lie adjacent to each other, and the hydrogen flux is uniform as a result of the high rate of hydrogen diffusion in aluminum. Therefore, it seems that the different morphologies result from variations in the ease of initiation, with some bifilms opening easily and some only with difficulty. This is in agreement with expectations based on the randomness of turbulent bifilm compaction leading to a random difficulty of unfurling. In general, it seems that gas pores form only from the opening of bifilms or other entrained
Fig. 4 Fig. 3
Inclusions, entrained via the surface, carrying their coating of surface film and remnant of oxide bifilm trail
Unfurling and inflating of bifilms. (a) Simply folded bifilm. (b) Tangled, convoluted bifilm. (c) Overinflation of a bifilm, resulting in a spherical pore. (d) Overinflation of a bifilm later in solidification, resulting in an interdendritic pore
unbonded interfaces. In the absence of suitable substrates such as bifilms, the gas cannot precipitate and thus remains in supersaturated solution. For many years, it has been a well-known fact in casting operations that filtration can
Fig. 5
SEM images of subsurface porosity in an Al-7Si0.4Mg alloy around a sand core bonded with a phenolic urethane binder. (a) general view. Closeups of (b) spherical pore and (c) dendritic pore. Source: Ref 1
366 / Principles of Solidification improve the soundness of castings (Ref 10). (Occasionally this has been interpreted as the ability of a filter to filter out hydrogen!)
Shrinkage due to Solidification Contraction Clearly, shrinkage on solidification, if occurring in an enclosed volume, will reduce the local pressure. The combined atmospheric and metallostatic pressures may be reduced to zero and possibly even further, becoming negative as in a tensile stress. Such a reduction in the hydrostatic pressure will open the bifilm (exactly as in the RPT example described previously). In a trapped volume of liquid, the accompanying shrinkage during solidification will simply expand the bifilms to conserve volume. If the poor feeding condition is not severe, the opening may be so slight as to be overlooked, and the metal will appear perfectly sound but will exhibit reduced mechanical properties. This property loss for no apparent reason is widely experienced. Further opening will lead to the typical scattering of microshrinkage pores and a further reduction in properties. If the shrinkage is severe, macroshrinkage porosity can result, and the strength may be zero. The circumstances are the mirror image of those in which gas drives the unfurling; an increase of pressure on the inside is physically precisely equivalent to a reduction of pressure on the outside of the bifilm. Thus, it is logical that both gas and shrinkage can act in combination; one pushes from the inside, and the other pulls from the outside, leading to the unfurling of the bifilm to make a planar crack and possibly further inflation to make a pore.
straightened between the advancing dendrites, similar to a carpet unrolled across a floor. Such flat cracks are common in many fields of casting and appear to be the cause of hot tears. An example of dendrite straightening is seen on the fracture surface of an Al-4.5Cu alloy in Fig. 6(a). In this image, the large bifilms straightened by dendrites were observed by video radiography to be entrained during the filling of the mold. No metal-to-metal contact could be observed on this fracture surface even under close examination by the SEM, explaining the lamentable elongation to failure of only 0.3%. The whole surface was found to be covered with a transparently thin oxide film (Fig. 6b). The main dendrite rafts in Fig. 6(a) are seen to progress from each side of the casting, pushing the bifilms ahead and finally meeting in the middle. The light, slightly fluorescing regions making the north/south ridge across the diameter of the image were found to be heaps of spare oxide film, pushed ahead of the dendrites. In contrast, in a different part of the same casting that had been observed in the same radiographic study to enjoy a nicely laminar fill profile, the excised test bar produced 3.0% elongation to failure. The fracture surface (Fig. 6c) was poor, exhibiting many shrinkage cavities and implying the presence of a high
population of smaller bifilms. However, there were some metal-to-metal contact areas on this fracture surface, as evidenced by the patches of ductile dimples that were observed (Fig. 6d). Examples of dendrite straightening in other alloy systems have been reported elsewhere (Ref 1). These include spheroidal graphite cast irons, creating the so-called plate fracture effect (Ref 12). The latter is particularly disturbing because so-called ductile iron, marketed especially for its ductility, can unexpectedly fracture in a brittle manner, with the crack appearing to propagate along large planar grain boundaries. In industrial practice, but unfortunately not in the scientific literature, there is now much evidence for this mechanism. Hot tears in castings appear to be universally caused by the presence of bifilms, which invariably can be eliminated by using a well-designed filling system (Ref 4). (See also the article “Filling and Feeding System Concepts” in this Volume.) The use of poorly designed filling systems also explains the wide scatter in results that has long troubled the scientific study of hot tearing phenomena. Little attention has been devoted to the filling system design for cast testpieces for such research. Thus, the random entrainment effects from the turbulent fill have introduced serious quantities of bifilms that have clouded results with “noise.”
Linear Contraction in the Solid State (Hot Tearing) If the cast alloy is stretched, as often happens to parts of a casting during freezing, it is clear that the bifilms can easily be opened. The phenomenon of hot tearing during the solidification of castings has been shown to be eliminated by the removal of bifilms in the liquid. In the absence of bifilms, the casting simply stretches as a result of its low strength and high ductility at elevated temperature. This is achieved by improving the quality of the melt and providing a filling system design for the casting that does not reintroduce bifilms (Ref 1, 4).
Dendrite Pushing The advancing solidifying dendrites cannot grow through a bifilm because of its layer of air. Consequently, the bifilm is pushed by the advancing dendrites. This is why the majority of bifilms appear concentrated in grain boundaries or in interdendritic regions. If a bifilm is overtaken, becoming trapped inside an advancing array of dendrites like a growing dendrite forest, and one end remains attached to the wall of the casting (the forest floor), the bifilm is
Fig. 6
SEM images of the fracture surfaces from a single Al-4.5Cu alloy casting showing (a) the part that suffered a turbulent back-wave and consequent entrained bifilm flattened by dendrite growth and yielding 0.3% elongation. The nearly invisible oxide film covering the whole surface is seen in (b) in a folded area (arrow). (c) View of a more distant point of the same casting, which enjoyed a more tranquil fill, thus exhibiting a fracture surface influenced only by the small but more numerous original bifilms in the rather poor melt. (d) Area of ductile dimples confirming the true metal-to-metal contact in the tensile fracture from this second part of the casting, which yielded 3% elongation. Source: Ref 11
The Oxide Film Defect—The Bifilm / 367 Nucleation and Growth of Intermetallics It seems that bifilms, at least the oxide film variety that has been mainly observed so far, are favored sites for the nucleation and growth of a wide variety of intermetallics in a wide variety of matrices. In fact, the author has found that of the intermetallics precipitating from the liquid thus far studied, each one has formed on an oxide bifilm substrate. Such precipitation occurs on one or both of the wetted outside surfaces of the bifilm that originally grew on the melt surface. The growth of the original film occurred atom by atom, resulting in the external interface of the bifilm (the interface between the melt and the oxide) being in perfect atomic contact (effectively bonded to the melt), in contrast to its dry inner, unbonded surfaces. Cao and Campbell (Ref 11) were the first to report the association between iron-rich intermetallics in aluminum alloys and bifilms as substrates. An excellent example of an intermetallic grown on an oxide substrate is the beta-iron precipitate Al5FeSi seen in Fig. 7. Particles of silicon eutectic in this image also appear to have formed on bifilms, explaining their central cracks. Strontium-rich intermetallics have been observed growing on the underside of surface oxide films (Ref 14), as have carbides in vacuum-cast nickel-base superalloys (Ref 15). The concept of the presence of bifilms in aluminum melts has allowed a new explanation to be put forward for the mechanism of silicon modification. It is based on the finding that the elements sodium and strontium appear to deactivate bifilms as favored substrates for the nucleation and growth of silicon. As a result, the silicon phase no longer forms as large (apparently) brittle plates by easy nucleation and growth on bifilms in suspension in the liquid but is now forced to grow at the lower temperature of the eutectic, with its spacing dictated by diffusion and therefore on a significantly finer scale. The detailed account is outlined elsewhere (Ref 16).
Fig. 7
SEM image of intermetallics (eutectic silicon and b-iron platelets) nucleated and grown on bifilms, causing the bifilms to straighten and create long central cracks. Source: Ref 13
The precipitation of a rigid intermetallic crystal on a rather flimsy film of substrate has interesting consequences. The progressive growth of the crystal can force straightening of the film. However, if the film is thicker or unable to unfurl because its folds are glued by, for instance, low-melting-point patches or perhaps effective spot welded by high-temperature treatment of the liquid (Ref 17), the film sometimes dictates that the crystal grows in a compact and/or irregular form (Ref 16). Thus, the influence of the substrate on the precipitate or vice versa depends on the relative rigidities of the two, and this can be highly variable and mixed, even within one sample. Interestingly, there is evidence that the growing intermetallics can be broken up, fragmenting into smaller sections in a form of the particle refining process (Ref 18). This appears to be the result of the simultaneous precipitation of hydrogen into the bifilm, causing it to inflate and unfold, while the intermetallic is attempting to grow on this mechanically unstable substrate. This refinement mechanism can be easily visualized. The hydrogen is in solution in the liquid metal. Thus, it will attempt to equilibrate with any void, filling the void until the internal pressure in the void leads to re-solution in the melt at the same rate as hydrogen evolves into the void. In this case, the void is the minute air gap between the folded-over films. Hydrogen will diffuse through the films and into the air gap, gradually increasing the pressure of hydrogen in the gap. The gap will widen, inflating the bifilm, first changing from a practically closed crack to an open crack; eventually, if there is enough hydrogen, the bifilm takes on the appearance of a balloon. Now, it has effectively become a hydrogen pore. During the early part of this inflation, the bifilm starts as a convoluted, compact defect. As hydrogen is taken up, it unfurls, opening outward fold by fold (a large, limp balloon does exactly this). Any crystal attempting to grow on the outer surfaces of the bifilm will find the ground moving beneath its feet. The crystal is likely to fragment, appearing to grow as many separate small crystals instead of a single large crystal. Southin and Romeyn (Ref 19) suggest that an analogous process is happening in the case of dendrite solidification on a smoke-deposited, and therefore fragile, carbon surface. (This effect is frequently used by aluminum casters to increase the apparent fluidity of the melt, because the normal hindrance to flow provided by the dendrite array growing from the surface of the casting is prevented; the dendrites formed on the weak substrate continuously break away from the walls of the mold, leaving the channel clear for continued flow of the liquid.) Sometimes, the central bifilm is not seen in the intermetallic particle, which has caused some investigators to assume it is not present. However, conditions during freezing are not always conducive to opening the bifilm to render it visible. These conditions are easily
arranged, but most casting practices naturally go to some lengths to avoid such conditions by using well-degassed metal, low iron contents (for aluminum alloys), and a good feeding system, keeping the melt pressurized during freezing to maintain tightly closed bifilms. Summarizing the unfurling mechanisms, it is interesting to note that they are already generally well known to casting metallurgists as deleterious to ductility. These deleterious factors are otherwise known as gas porosity, shrinkage porosity, tensile prestrain, large grain size, and (for aluminum alloys) high iron content.
Reliability and Reproducibility of Mechanical Properties The simple fact that castings can vary erratically in their properties has never been satisfactorily explained, particularly when all steps to eliminate variation have apparently been taken. The variation from casting to casting, day to day, and week to week in a closely monitored production plant has been nothing short of baffling. The presence or absence of bifilms in the starting material and the random variations in the pouring of the castings are features that have remained unsuspected. One of the first clear demonstrations of property variation due to pouring was the experiment carried out by Green and Campbell (Ref 20) in which the manner of pouring was varied. The turbulence of the different systems clearly had a major impact on the reliability of the casting. In this work, the use of Weibull statistics was introduced to quantify the reliability. Different reliabilities were found for essentially identical alloys poured at identical temperatures, the only difference being the degree of turbulence during the final transfer of metal and the final content of bifilms. (It is noted that other damaging features that would have influenced the results, such as porosity, the precipitation of beta-iron phase, shrinkage, and so on, are consequences of the presence of bifilms and thus are secondary influences.) Interestingly, Green and Campbell’s approximation of the Weibull distribution in two-parameter form has recently been challenged (Ref 21). It seems that the full three-parameter form of the Weibull analysis is more valid and yields additional key information about the potential safety implications of the scatter. Green and colleagues (Ref 22) went on to show that the fatigue properties of Al-7Si0.4Mg alloy were controlled by its bifilm content. In fact, in 98% of all fatigue fractures closely observed by SEM in this work, bifilms were the initiators of cracks. The remaining 2% exhibited up to 100 times greater fatigue lives. These were originally thought to fail from fractures initiated by slip planes. However, on a recent re-examination of these results, it seems that the slip planes are more likely to have been planar fractures resulting from the straightening of oxide bifilms by dendrite pushing. If this new
368 / Principles of Solidification interpretation is correct, it is not clear at this time how far fatigue life can be extended if bifilms are eliminated. The 100-times benefit seems an underestimate. There has been much discussion to claim that pores are more important than bifilms in initiating failure, particularly fatigue failure (Ref 23, 24). In a sense, the differentiation is, in practice, almost irrelevant since both are entrainment defects, the pore being simply a rather more open type of bifilm. In addition, the pore definitely has zero cross-sectional tensile strength. In contrast, the bifilm has mixed strength, on average perhaps 50% of the strength of the matrix (Ref 1). This is the result of mechanical interlocking of folds and the interference of “jogs” in the film when subjected to shear, as well as variations in alignment. A planar bifilm aligned with the tensile direction will cause no observable reduction in strength, whereas when perpendicular to the tensile axis, it could behave as having zero strength. There are genuine reasons why pores may be more important defects than bifilms. However, the films are usually vastly greater in area (for instance, in Fig. 1, the pores are seen merely to “decorate” the bifilm), have a sharp aspect ratio, and are so thin that they are routinely overlooked by observers. When an observer claims to have identified a pore as the origin of a fatigue crack, he has probably failed to notice that the pore is almost certainly a part of a much larger-area bifilm, whose presence is difficult to detect because of its thinness and consequent transparency. Planar fracture regions are often observed in association with pores. These are often interpreted as cleavage surfaces, grain-boundary surfaces, or slip planes. Close examination often reveals that these features are dendrite-straightened oxide films. Careful experiments on a nickel-base superalloy, vacuum cast and subjected to hot isostatic pressing (HIP), demonstrated that these features only occurred in those castings that were poured turbulently (Ref 25). Open or partially open bifilms have practically zero strength. The application of pressure to the casting can therefore produce significant benefit by closing bifilms, even if the stability of the oxide interfaces resists bonding. Thus, pressure applied immediately after casting, prior to solidification, has been repeatedly demonstrated over the years to be beneficial to mechanical properties (Ref 26). Similarly, in the solid state, the mechanism of HIP can be easily understood to increase mechanical properties by closing bifilms, even if the defects are not closed by bonding (i.e., welded) (Ref 27).
Central Role of Bifilm Defects in Fracture Brittle failure appears to be accurately described using some variant of the Griffiths
crack concept, in which the elastic energy built up in the material exceeds the energy to propagate a pre-existing crack. Where there is no preexisting defect, however, fracture is probably not possible. The arguments for this are not presented here; see the review article in Ref 26. Essentially, the arguments emphasize the fact that metals have theoretical strengths in the region of 10 to 20 GPa (1450 to 2900 ksi), which is so high that such bonds cannot normally be broken. As correctly conceptualized by the Griffiths crack model, it follows that only a solid or liquid phase containing a defect can fracture. For instance, an inclusion or second-phase particle precipitated in a solid matrix that has nucleated and grown by diffusion of solutes will effectively be in perfect atomic contact with the matrix and therefore incapable of creating a crack. Either the inclusion, the matrix, or both will yield plastically, but a cavity or crack cannot open. Interestingly, silicon particles in aluminum-silicon alloys would not be expected to fracture during tensile testing. However, it is widely known that solids often appear to fail by decoherence from inclusions, second-phase particles, or grain boundaries. There are many examples, including the opening of grain boundaries in creep in such materials as nickel-base superalloys and superplastic forming in aluminum-zinc alloys. From these arguments, it follows that the metallurgical features, such as the grain boundaries alone, cannot be the origin of the failures. Such failures most likely originate from bifilms lying in the grain boundaries, where they will often occur as the result of dendrite pushing or grain growth that is expected to be arrested at a bifilm. As usual, the bifilms will be invisible to casual inspection, so that the incorrect identification of the cause of failure is common. In summary, failure is not expected to occur from any inclusion or other feature that has grown within the metal, whether in the liquid or solid phase, because its boundaries are expected to be in more-or-less perfect atomic contact with the matrix, giving such interfaces great tensile strength. Conversely, failure is expected from features entrained in those liquid metals that have a solid surface film. These features are necessarily associated within the “paper bag” of oxide film dragged into the matrix from the liquid surface (Fig. 3), creating an unbonded interface around the entrained particle. The practical outcome of the quest for highquality metal, low in bifilm content, has been the development of exciting new filling system designs for castings, with great benefits to integrity, soundness, enhanced properties, and the simultaneous reduction of costs on the industrial scale (Ref 5). The improvement in reliability and repeatability has been most significant. These benefits are highly prized targets in the marketplace and represent the new vision for future generations of casting production.
REFERENCES 1. J. Campbell, Castings, 2nd ed., Butterworth Heinemann, Oxford, 2003 2. S. Fox and J. Campbell, Scr. Mater., Vol 43 (No. 10), 2000, p 881–886; Int. J. Cast Met. Res., Vol 14 (No. 6), 2002, p 335–340 3. D. Dispinar and J. Campbell, Int. J. Cast Met. Res., Vol 17 (No. 5), 2004, p 280–294 4. J. Campbell, Casting Practice, Butterworth Heinemann, Oxford, 2004 5. J. Campbell, Eng., Vol 221 (No. 3), 1981, p 185–188 6. J. Finardi, Trans. AFS, Vol 108, 2000, p 409–417 7. R. Kiessling, Non-Metallic Inclusions in Steel, The Metals Society, London, 1978, p 42–86 8. J.M. Svoboda, R.W. Monroe, C.E. Bates, and J. Griffin, Trans. AFS, Vol 95, 1987, p 187–202 9. J. Campbell, Mater. Sci. Technol., Vol 22, 2006, p 127–145, 999–1008 10. K.J. Brondyke and P.D. Hess, Trans. Metall. Soc. AIME, Vol 230 (No. 7), 1964, p 1542–1546 11. X. Cao and J. Campbell, Metall. Mater. Trans. A, Vol 34, 2003, p 1409–1420 12. S.I. Karsay, Ductile Iron Production: The State of the Art 1992, QIT Fer et Titane Inc., Canada, 1993 and 2000, p 196–198 13. J. Mi, R.A. Harding, and J. Campbell, Int. J. Cast Met. Res., Vol 14 (No. 6), 2002, p 325–334 14. S.M. Miresmaeili, J. Campbell, S.G. Shabestari, and S.M.A. Boutorabi, Metall. Mater. Trans. A, Vol 36 (No. 9), 2005, p 2341–2349 15. A.K.M.B. Rashid and J. Campbell, Metall. Mater. Trans. A, Vol 35, 2004, p 2063– 2071 16. J. Campbell, Shape Casting: The John Campbell Symposium (San Francisco, CA), M. Tiryakioglu and P.N. Creapeau, Ed., TMS, 2005, p 3–12 17. R. Ahmad and R.I. Marshal, Int. J. Cast Met. Res., Vol 15 (No. 5), 2003, p 497–504 18. H. Sina, M. Emamy, M. Saremi, A. Keyvani, M. Mahta, and J. Campbell, The Influence of Ti and Zr on Electrochemical Properties of Al Sacrificial Anodes, Mater. Sci. Eng. A, Vol 431, 2006, p 263–276 19. R.T. Southin and A. Romeyn, 1980 Warwick Conference, Solidification Technology in the Foundry and Cast House, Metals Society, 1983, p 355–358 20. N. Green and J. Campbell, Trans. AFS, Vol 102, 1994, p 341–347 21. H. Zahedi, M. Emamy, A. Razaghian, M. Mahta, J. Campbell, and M. Tiryakiog˘lu, Metall. Mater. Trans. A, Vol 38, March 2007, p 659–670 22. C. Nyahumwa, N.R. Green, and J. Campbell, Trans. AFS, Vol 106, 1998, p 215– 223; J. Mech. Behavior Mater., Vol 9 (No. 4), 1998, p 227–235
The Oxide Film Defect—The Bifilm / 369 23. Q.G. Wang, P.N. Crepeau, D. Gloria, and S. Valtierra, Advances in Aluminum Casting Technology, M. Tiryakioglu and J. Campbell, Ed., ASM Symposium, 2002, p 209–218 24. M.J. Couper, A.E. Neeson, and J.R. Griffiths, Int. J. Fatigue Fract. Eng. Mater. Struct., Vol 13, 1990, p 213–227
25. M. Cox, personal communication based on M. Cox, M. Wickins, J.P. Kuang, R.A. Harding, and J. Campbell, Mater. Sci. Technol., Vol 16, 2000, p 1445–1452 26. J.T. Berry, R. Luck, and R.P. Taylor, Shape Casting: The John Campbell Symposium (San Francisco, CA), M. Tiryakioglu and P.N. Creapeau, Ed., TMS, 2005, p 113–122
27. J.T. Staley, Jr., M. Tiryakiog˘lu, and J. Campbell, Shape Casting: Second International Symposium, P.N. Crepeau, M. Tiryakiog˘lu, and J. Campbell, Ed., TMS, 2007
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 370-374 DOI: 10.1361/asmhba0005222
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Shrinkage Porosity and Gas Porosity Qingyou Han, Purdue University
IN METAL CASTINGS, two types of porosity are identified: shrinkage porosity and gas porosity. Shrinkage porosity occurs when the shrinkage on solidification cannot be compensated (Ref 1). Gas porosity occurs when the gas concentration in the liquid is higher than its solubility (Ref 2). The causes of porosity formation for these two types of porosity are inadequate feeding and gas evolution. Under some conditions, one pore formation mechanism is dominant so that one type of porosity occurs, but in most situations, both of the causes play important roles in porosity formation. Inadequate feeding causes a local pressure drop so that the gas becomes more supersaturated, resulting in earlier formation of the gas porosity. During the solidification of local,
isolated liquid pools, shrinkage generates pores filled with gases. It is the interplay of these two mechanisms that gives rise to porosity in a casting. Figure 1 illustrates gas porosity and shrinkage porosity in an aluminum AA 5182 remelt secondary ingot. Gas porosity is usually spherical, and the shrinkage porosity is irregularly shaped. This distinction is based on the assumption that the irregular shape is due to the interdendritic fluid being drawn away from a partially formed microstructure to feed the volume contraction elsewhere.
Shrinkage Porosity Shrinkage. The formation of shrinkage porosity is strongly affected by the shrinkage behavior of the alloy and the feeding conditions for the casting. In general, liquid contracts on solidification because of the random arrangement of atoms in the liquid and the orderly arrangement of atoms (lattice structure) in the
solid state. Table 1 (Ref 3) lists the solidification shrinkage of some pure metals. The most dense solids are those that have cubic closepacked (face-centered cubic and hexagonal close-packed) symmetry. Contraction on solidification for these metals is in the range of 3.2 to 7.2%. The solidification shrinkage for the less closely packed body-centered cubic lattice is in the range of 2 to 3.2%. The exceptions to this general pattern are those materials that expand on freezing. These include cerium, silicon, and bismuth. No shrinkage porosity has been observed in these metals, provided the liquid contraction can be compensated by using risers. The shrinkage behaviors of pure metals and alloys are quite different. For pure metals, solidification occurs at the melting temperature (the liquidus temperature equals the solidus temperature). The freezing front is usually planar. During solidification, the freezing front progresses from the casting surface toward the center, forming centerline porosity. By using proper thermal control and risers, most of the solidification shrinkage of a pure metal can be collected in the risers as a shrinkage cavity.
Table 1 Solidification shrinkage for some pure metals Melting point Metal
Al Au Co Cu Ni Pb Fe Li Na K Rb Cs Tl Cd Mg Zn Ce In Sn Bi Sb Si
Fig. 1
(a) Gas porosity and (b) shrinkage porosity in an AA 5182 remelt secondary ingot
Crystal structure(a)
fcc fcc fcc fcc fcc fcc bcc bcc bcc bcc bcc bcc bcc hcp hcp hcp hcp fct Tetragonal Rhombohedral Rhombohedral Diamond
C
660 1,063 1,495 1,083 1,453 327 1,536 181 97 64 39 29 303 321 651 420 787 156 232 271 631 1,410
F
Liquid density, kg/m3
Solid density, kg/m3
1,220 1,945 2,723 1,981 2,647 621 2,797 358 207 147 102 84 577 610 1,204 788 1,449 313 450 520 1,168 2,570
2,368 17,380 7,750 7,938 7,790 10,665 7,035 528 927 827 1,437 1,854 11,200 7,998 1,590 6,577 6,668 7,017 6,986 10,034 6,493 2,525
2,550 18,280 8,180 8,382 8,210 10,020 7,265 ... ... ... ... ... ... ... 1,655 ... 6,646 ... 7,166 9,701 6,535 ...
Volume change, %
(a) fcc, face-centered cubic; bcc, body-centered cubic; hcp, hexagonal close-packed, fct, face-centered tetragonal. Source: Ref 3
7.14 5.47 5.26 5.30 5.11 3.22 3.16 2.74 2.6 2.54 2.3 2.6 2.2 4.0 4.1 4.08 0.33 1.98 2.51 3.32 0.64 2.9
Shrinkage Porosity and Gas Porosity / 371 For an alloy, solidification shrinkage occurs throughout its freezing temperature range. The shrinkage amount varies with the temperature, depending on the fraction of solids precipitated at given temperatures and the nature of the phases formed. Figure 2 illustrates the temperature versus fraction solid curves of three ironcarbon alloys: Fe-0.8wt%C, Fe-3wt%C, and Fe-4wt%C. The Fe-0.8wt%C alloy represents an alloy that solidifies into a single phase. The temperature versus solid fraction curve of this alloy is fairly straight, indicating the shrinkage is fairly constant throughout solidification. Most wrought alloys have temperature versus solid fraction curves similar in shape to that of the Fe-0.8wt%C alloy, except for a small amount of eutectic phase formation at the late stage of solidification. For the Fe-3wt%C alloy, 58% of the solid forms as primary austenite dendrites in the temperature range between 1150 and 1300 C
(2100 and 2370 F), and 42% of the solids form at 1150 C (2100 F) as the eutectic phases. The resultant solidification shrinkage of this alloy is highly nonuniform. During precipitation of the austenite dendrites, the shrinkage behavior of the alloy is identical to that of the Fe-0.8wt%C alloy. However, a large amount of shrinkage occurs at the eutectic temperature because of the precipitation of the eutectic phases. Most of the casting alloys, such as aluminum-copper, aluminum-silicon, aluminummagnesium, copper-base alloys, zinc alloys, and magnesium alloys, have behavior identical to Fe-3wt%C, except that the amount of eutectic phases and the eutectic temperature ranges may vary. In fact, the shrinkage behavior of the alloys containing eutectic phases is strongly affected by the types of eutectic phase formed. In the iron-carbon alloys, the formation of austenite and carbide (Fe3C) phases leads to a volume contraction, but the formation of
Fig. 2
Temperature versus fraction solid curves of three iron-carbon alloys: Fe-0.8wt%C, Fe-3wt%C, and Fe-4wt%C. These curves are obtained using a lever rule.
Fig. 3
Schematic representation of the feeding mechanisms in a solidifying casting. Source: Ref 3
graphite phase leads to a volume expansion. As a result, gray cast irons and ductile irons with composition close to or higher than the eutectic composition usually expand during solidification. This is because the volume fraction of the graphite phase is large enough that volume expansion due to graphite formation is larger than the volume contraction due to the formation of the other phases. A temperature versus fraction solid curve for Fe-4wt%C is also shown in Fig. 2. The eutectic fraction can be as high as 0.84, and the alloy expands during eutectic solidification. Most of the alloys shrink during solidification, so porosity usually forms in the castings of these alloys if the solidification shrinkage is not fed. When alloys expand during solidification, shrinkage porosity is unlikely to occur except for ductile irons. For ductile irons, a graphite nodule is usually surrounded by a shell of eutectic austenite during eutectic solidification. As a result, graphite expansion works against the austenite shell and pushes the austenite dendrite apart, leading to the formation of porosity among eutectic cells and dendrites (Ref 4). Feeding. Risers, or appendages, are added to a casting in order to feed the shrinkage in a casting. Solidification takes place simultaneously in both the casting and risers, and liquid flows from the risers to feed the casting. The flow is driven by solidification shrinkage, contraction of the liquid and the solid as they cool, and gravity. The possible feeding mechanisms are illustrated in Fig. 3 (Ref 3). Liquid feeding precedes other forms of feeding and is normally the only method of feeding for skinfreezing materials such as pure metals or alloys with a narrow freezing range. The liquid has low viscosity, and for most of the freezing process, the feeding pass is wide, so the pressure difference required for liquid feeding is negligible. Mass feeding involves the movement of a slurry of solid grains and residual liquid. It usually occurs in an alloy with a long freezing range when the fraction solid is smaller than the dendrite coherence point. Similar to that of liquid feeding, the pressure difference required for mass feeding is also small since the feeding pass is wide. Liquid feeding and mass feeding are the main feeding mechanisms for an alloy with a large fraction of eutectic phases at the end of solidification, such as the two high-carbon alloys shown in Fig. 2. During interdendritic feeding, liquid flows through tiny tortuous channels between dendrites to feed the shrinkage at the far end of the mushy zone. Inadequate feeding normally occurs in the interdendritic feeding region where the dendritic network exhibits a large resistance to fluid flow. Especially when the fraction liquid is so small that the remaining liquid in the mushy zone may be isolated by the solids, the solidification shrinkage of this liquid will inevitably produce pores, since feeding becomes impossible. Experimental
372 / Principles of Solidification measurements indicated that feeding stops before the end of solidification at a solid fraction of 0.9 to 0.99, equivalent to a liquid fraction of 0.01 to 0.1 (Ref 5). Taking a volumetric shrinkage of 6.7% for Al-8wt%Cu, for instance, the solidification of a fraction liquid of 0.1 will give 0.67% of porosity if all the pores are formed after feeding stops. Interdendritic feeding is important for shrinkage porosity formation in an alloy such as Fe0.8wt%C, shown in Fig. 2, that has a long freezing range. For such an alloy, the mushy zone is large and the fraction liquid at the point of final solidification can be extremely small. As a result, most long-freezing-range alloys exhibit poor pressure tightness due to porosity formation. Interdendritic feeding is also affected by the shrinkage behavior of the alloy. A large amount of shrinkage associated with eutectic formation at the end of solidification makes interdendritic feeding more difficult, especially when the fraction of the eutectic liquid is small. However, as the fraction of eutectic liquid increases, the feeding path becomes wider and feeding becomes easier.
Fig. 4
Hydrogen solubility in pure aluminum as a function of temperature at 101 kpa (1 atm) pressure. Source: Ref 6
Gas Porosity Gas porosity is generally caused by the evolution of gases during the casting and solidification processes. Gas porosity occurs when the gas concentration in the molten metal is higher than its solubility. The gases in molten metal may be the result of a reaction between the metal and the gases in atmosphere or in the solid materials that contain the molten metal. The type of gases that are responsible for porosity formation during casting varies from metal to metal.
Gases in Metal Aluminum. The only gas that has a measurable solubility in liquid aluminum alloys is hydrogen. Moisture from the refractory lining or atmosphere can react with molten aluminum: 3H2 O þ 2½Al ¼ Al2 O3 þ 3H2
(Eq 1)
The hydrogen that is released in this reaction can dissolve in liquid aluminum as atomic species according to: H2 ¼ 2½H
(Ref 6). The lines of best fit to the solubility in liquid, SL, and solid, SS, aluminum are: 2760 þ 2:796 and T 2080 log SS ¼ þ 0:788 T
log SL ¼
(Eq 4)
The unit for solubility in the previous equations is mL/100 g, and temperature is given in degrees Kelvin. Alloying elements affect the solubility of hydrogen (Ref 7). Generally, the common alloying elements decrease the solubility of hydrogen. Copper and Tin. The solubility of hydrogen in molten copper, tin, and copper-tin alloys is shown in Fig. 5 (Ref 8). In pure copper, hydrogen solubility is greater in the liquid than in the solid state (approximately 5.5 mL/100 g versus 2 mL/100 g), which is why gas porosity is a problem in copper castings. In addition to hydrogen, oxygen is also soluble. Reaction between hydrogen and oxygen produces water vapor according to: 2½H þ ½O ¼ H2 O
(Eq 5)
(Eq 2)
The gas pressure of water vapor is given by: The solubility of hydrogen in liquid aluminum is related to the partial pressure of hydrogen, PH2 , using Sievert’s law: ½H2 ¼ kPH2
(Eq 3)
where the constant k is the activity of hydrogen in the metal at that temperature and at 101 kpa (1 atm) pressure. Figure 4 shows the hydrogen solubility in pure aluminum as a function of temperature at 101 kpa (1 atm) pressure
½H2 ½O ¼ kH2 O PH2 O
(Eq 6)
High levels of dissolved oxygen lead to low levels of dissolved hydrogen, and vice versa. It is beneficial to melt copper in an oxidizing atmosphere and then deoxidize to minimize porosity formation in the casting. Other complex gases in copper-base alloys include CO in nickel-bearing copper alloys or SO2 if high-sulfur fuel is used for melting.
Fig. 5
Solubility of hydrogen at atmospheric pressure in pure copper, pure tin, and copper-tin alloys. Source: Ref 8
Iron and Steel. In irons and steels, gases that are responsible for porosity formation in casting include CO, H2, and N2. Hydrogen and nitrogen are soluble in molten steels. Oxygen is also soluble and reacts with carbon to form carbon monoxide, CO, by the reaction in steel as: ½C þ ½O ¼ CO
(Eq 7)
In steelmaking practice, the CO reaction is used to reduce the high carbon levels in pig iron produced by blast furnaces. The CO reaction is so vigorous and violent that it is called a carbon boil. After the carbon content is reduced into the specified range, the excess oxygen in molten metal is reduced by using a deoxidizing addition. The vigorous carbon boil will reduce both hydrogen and nitrogen in solution to negligible levels. In many steel foundries where steel is simply remelted, dissolved hydrogen and nitrogen will contribute to porosity formation in castings.
Gas Evolution during Solidification As illustrated in Fig. 4, hydrogen solubility in pure aluminum is greater in the liquid than in the solid state (approximately 0.65 mL/100 g versus 0.034 mL/100 g). The solubility falls by an order of magnitude upon solidification. Generally, there is a large solubility drop of the dissolved gases during phase transformation from liquid to solid. On solidification, the dissolved gases are rejected by the growing solid and are enriched in the remaining liquid in the mushy zone. Most of the solute elements in an alloy are also enriched in the remaining liquid, thus decreasing the solubility of the dissolved
Shrinkage Porosity and Gas Porosity / 373 gases in the mushy zone. When the gas contents are higher than their solubility, gas porosity forms. Based on the fact that the diffusion coefficient of the dissolved gas, such as hydrogen, is a few orders of magnitude higher than that of the other solute elements in the alloys, the lever rule can be used for calculating the local gas concentration in the liquid in the early stage of solidification, according to: CS fS þ CL fL ¼ C0
(Eq 8)
where CS and CL are the concentrations of a dissolved gas in the solid and in the liquid, respectively, C0 is the initial gas concentration in the molten alloy before solidification starts, and fS and fL are the local solid fraction and liquid fraction in the mushy zone, respectively. In Eq 8, CS is less than or equal to the solubility of dissolved gas in the solid. Thus, the concentration of the dissolved gas in the remaining liquid can be many times higher than the initial gas concentration in the molten metal at the end of solidification. In fact, the use of the lever rule for the prediction of dissolved gas concentration in the remaining liquid usually underestimates the dissolved gas concentration at the point of final solidification, such as the eutectic front. Because of their limited diffusion in the liquid, the dissolved gases are highly enriched at the eutectic front, and their concentrations can be a few orders of magnitude higher than that predicted using the lever rule (Ref 9). This accounts for the observation that gas porosity occurs in alloys with initial gas concentrations lower than the solubility of the gas in the liquid.
Porosity Formation Homogeneous Nucleation. Similar to the nucleation of any phase, the free energy, DG, for the formation of a spherical vapor bubble of radius r is (Ref 10): 4 G ¼ 4pr2 þ pr3 ðP PV Þ 3
(Eq 9)
where P is the hydrostatic pressure in the liquid, and PV is the vapor pressure, which is negligible in comparison with P since a very large negative pressure is required. The surface tension at the gas-liquid interface is represented by s. The maximum free energy, DG*, is calculated to be: G ¼
16ps3 3ðP PV Þ2
(Eq 10)
for pores with a critical radius, r*: r ¼
2s P PV
(Eq 11)
Hence, there is a critical radius for the bubble, below which it is not capable of surviving and above which it will tend to grow.
For the nucleation of a soluble gas in a liquid, the pressure required to nucleate a bubble is generally given by: Pg ¼ Pm þ
2s r
(Eq 12)
where Pg is the partial pressure of dissolved gas, and Pm is the local pressure. Assuming local equilibrium, the partial pressure of the dissolved gas, Pg, is related to the mass of the gas, n, in the pore by: Pg ¼ nRT
(Eq 13)
where R is the gas constant, and T is the temperature in the pore. To form a pore of radius r, n moles of gas are required. The pressure, Pg, needed to create a pore in metal is calculated to be on the order of 30,000 atm (3 109 Pa) (Ref 11). This means that homogeneous nucleation of pores is extremely difficult and is unlikely to occur in practice. Heterogeneous Nucleation. In castings, nucleation of pores can be expected to occur primarily at heterogeneous nucleation sites. The classical theory of homogeneous nucleation is extended to heterogeneous nucleation. The maximum free energy is: 16 3 2 ps P SðyÞ 3
(Eq 14)
SðyÞ ¼ ð2 þ cos yÞð1 cos yÞ2 =4
(Eq 15)
G ¼
where:
and y is the equilibrium contact angle of the bubble on the solid. The pressure required to nucleate a bubble is given by: Pg ¼ Pm þ
2s pffiffiffiffiffiffiffiffiffiffi SðyÞ r
(Eq 16)
When the contact angle is zero, the free energy for the formation of a pore will be zero so the pore can grow directly on the heterogeneous site. When the contact angle is 180 , Eq 14 and 16 reduce to Eq 10 and 12, which means the nucleation sites are ineffective. Heterogeneous sites are inherent in almost all normal castings, such as the solid-liquid interface and inclusions. Calculations indicate (Ref 12, 13) that the nucleation of a bubble on the planar interface is difficult. The cellular or dendritic interface provides a much more favorable nucleation site, partly because the interface is concave and partly because the soluble gases are concentrated in the interdendritic region. Inclusions are probably the most important sites for heterogeneous nucleation. It has been well demonstrated that the presence of inclusions greatly enhances the porosity formation in aluminum alloys. The easiest sites for bubbles to nucleate are the pre-existing gas bubbles entrapped in the cavities of the inclusions (Ref 3). No nucleation event is required, only growth. Such a site can act as a multiple
nucleation site if the bubbles detach, leaving a small bubble behind in the cavity. Equations 12 and 16 indicate that both Pm and Pg provide the driving forces for the nucleation of a pore. Porosity forms when: Pg > Pm þ 2s=r
(Eq 17)
It is the interplay of Pm and Pg that determines the formation of porosity in a casting. In the previous equation, Pg is a measure of the amount of dissolved gases in the liquid as described by Eq 3. If the concentration of a dissolved gas is higher than its solubility at Pm, which is the local pressure, a pore is likely to form. On the other hand, Pm is a measure of the ease of feeding the shrinkage locally under given external and hydraulic pressure. Difficulty in feeding occurs at locations where the local Pm is extremely low or negative. Under such a condition, the dissolved gases can escape from the liquid and form porosity to feed the solidification shrinkage. Neglecting the hydraulic pressure, the local pressure is given by the external pressure minus a pressure drop defined using Darcy’s law: V ¼
K @P mfL @x
(Eq 18)
where V is the velocity of the shrinkage-driven flow, @P/@x is the pressure gradient, m is the viscosity of the liquid, and K is the permeability of the mushy zone. For a mushy zone containing equiaxed grains, K follows the BlackKozenny equation (Ref 14): K¼
fL3 kC SV2
(Eq 19)
where kC is a constant, and SV is the solid-liquid interfacial area per unit volume. The local pressure, Pm, decreases with decreasing external pressure and increasing pressure drop or pressure gradient. The pressure drop increases with increasing velocity of the shrinkage-driven flow, solid-liquid interfacial area, and decreasing liquid fraction. Thus, it is extremely difficult to feed regions with a small liquid fraction that have a large solidification shrinkage.
Factors Affecting Porosity Formation Factors that affect porosity formation include the following: Dissolved Gases in the Melt. The higher the dissolved gas concentration in a molten alloy, the higher the Pg value. As a result, porosity will be more likely to occur. When porosity forms, the mass of the gas, f, in pores per 1 g of alloy is calculated using (Ref 15): 100f ¼ ðCL SL ÞfL þ ðCS SS ÞfS
(Eq 20)
374 / Principles of Solidification Thus, the amount of porosity increases with increasing concentrations of the dissolved gases, CL, in a molten alloy. Degassing is effective in minimizing porosity formation. Alloy Composition. Apart from affecting the initial concentrations of the dissolved gases in the melt, alloy composition also determines the solidification interval and the shrinkage characteristics of the alloy. Usually, porosity level increases with increasing solidification interval, reaches a peak value at the composition when the nonequilibrium eutectic phases start to occur, and decreases as the eutectic fraction increases. This phenomenon is associated with a decreased permeability due to an increased SV and a decreased fL as the solidification interval increases. When a small fraction of eutectic phase forms, a large demand for feeding occurs, resulting in a large shrinkage-driven flow and a large pressure drop. As the eutectic fraction is increased, the feeding channel becomes open, and feeding becomes easier. External pressure can be used to increase the local pressure, Pm, in the mushy zone and to minimize porosity formation. Squeeze casting is a casting process using external pressure during the solidification stage of a casting to achieve porosity-free components. Cooling Conditions. Fast cooling is usually beneficial in reducing porosity level if risers and chills are properly used for feeding the liquid shrinkage. Increasing the cooling rate can
increase the temperature gradient at the freezing front, making the feeding channel open and wide. Furthermore, the radius of a pore is usually proportional to the secondary dendrite arm spacing, which decreases with increasing cooling rates. A pore with smaller radius will have higher internal pressure (Eq 12) and contain more gas molecules for a given volume (Eq 13). As a result, the volume fraction of pores decreases with increasing cooling rate.
REFERENCES 1. T.S. Powonka and M.C. Flemings, Pore Formation in Solidification, Trans. Metall. Soc. AIME, Vol 236, 1966, p 1157–1165 2. E.J. Whittenberger and F.N. Rhines, Origin of Porosity in Castings of MagnesiumAluminum and Other Alloys, J. Met., 1952, p 409–420 3. J. Campbell, Castings, Butterworth-Heinemann, 1991 4. M. Burditt, Ductile Iron Handbook, American Foundrymen’s Society, 1992 5. V.L. Davies, Feeding Range Determined by Numerically Computed Heat Distribution, AFS Cast Met. Res. J., Vol 11, June 1975, p 33–44 6. C.E. Ransley and H. Neufeld, The Solubility of Hydrogen in Liquid and Solid Aluminum, J. Inst. Met., Vol 74, 1948, p 599–620
7. G.K. Sigworth and T.A. Engh, Chemical and Kinetic Factors Related to Hydrogen Removal from Aluminum, Metall. Trans. B, Vol 13, 1982, p 447–460 8. M.T. Roeley, Casting Copper-Based Alloys, American Foundrymen’s Society, 1984 9. Q. Han and S. Viswanathan, Hydrogen Evolution during Directional Solidification and Its Effect on Porosity Formation in Aluminum Alloys, Metall. Mater. Trans. A, Vol 33, 2002, p 2067–2072 10. J.C. Fisher, The Fracture of Liquids, Appl. Phys., Vol 19, 1948, p 1062–1066 11. R. Cole, Boiling Nucleation, Adv. Heat Trans., Vol 10, 1974, p 85–166 12. W.R. Wilcox and V.H.S. Kuo, Gas Bubble Nucleation during Crystallization, J. Cryst. Growth, Vol 19, 1973, p 221–228 13. P.M. Wilt, Nucleation Rates and Bubble Stability in Water-Carbondioxide Solutions, J. Colloid Interface Sci., Vol 112 (No. 2), 1986, p 530–538 14. A.J. Duncan, Q. Han, and S. Viswanathan, Measurement of Liquid Permeability in the Mushy Zones of Aluminum-Copper Alloys, Metall. Mater. Trans. B, Vol 30, 1999, p 745–750 15. K. Kubo and D.R. Pehlke, Mathematical Modeling of Porosity Formation in Solidification, Metall. Trans. B, Vol 16, 1986, p 359–366
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 375-378 DOI: 10.1361/asmhba0005223
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Castability—Fluidity and Hot Tearing Lars Arnberg, Norwegian University of Science and Technology Asbj rn Mo, SINTEF Materials and Chemistry
CASTABILITY of alloys is a measure of their ability to be cast to a given shape with a given process without the formation of casting defects. These defects can be cracks, pores, segregations, inclusions, or other internal or external faults that impair the properties and/or appearance of the casting. This article examines two phenomena closely related to castability: fluidity and hot tearing.
Fluidity Fluidity is usually defined as the maximum fluidity length, which is the distance a melt will flow in a standard mold with a constant cross section before the flow is terminated by solidification. It should be noted that this is different from a physicist’s definition of fluidity, which is the inverse of viscosity. Fluidity limits the geometry of a casting that can be successfully filled, and insufficient fluidity can result in misruns and other defects. Fluidity has been studied for more than a century, and theories about factors that limit fluidity as well as experimental methods for measuring fluidity have been developed. The mechanisms for flow termination have been studied by Flemings (Ref 1) and others and were found to be different for alloys with short freezing ranges than for long-freezing-range alloys with a mushy solidification mode. Shortfreezing-range alloys form columnar crystals on the mold wall, pinching off the flow as they grow, as shown in Fig. 1(a). For long-freezing-range alloys, equiaxed crystals or dendrite fragments are carried with the melt, as shown in Fig. 1(b), and choke the flow at the leading tip when the crystals become coherent and form a network. The solid fraction at which this phenomenon occurs depends on the dendrite morphology but will typically be 0.3 to 0.5 (Ref 3). The choking mechanism is much more efficient in terminating melt flow, and fluidity is therefore low in alloys with long freezing ranges.
Metal variables
a. Chemical composition b. Freezing range c. Viscosity d. Heat of fusion Mold and mold/metal variables a. Heat transfer b. Mold and metal thermal conductivity c. Mold and metal mass density d. Specific heat e. Surface tension Casting variables a. Metal head b. Channel diameter c. Casting temperature d. Inclusion content Many of these variables are interdependent: Freezing range, heat of fusion, and specific heat depend on alloy composition; mold variables depend on the casting process; and so on. A short freezing range will, as mentioned, improve fluidity. This is shown in Fig. 2, which shows how fluidity varies with alloy composition for lead-tin alloys. It can be seen that fluidity has maxima for the pure metals and at the eutectic composition where the freezing point is at a minimum. Similar composition versus fluidity curves have been constructed for many
alloys (Ref 2). The viscosity, and thus the fluidity, of liquid metals is strongly influenced by impurities; it has been shown that high-meltinclusion content significantly reduces fluidity (Ref 5). A higher heat of fusion will improve fluidity, and alloying can be used to improve fluidity. An example is the addition of silicon to aluminum alloys. Silicon has a heat of fusion that is approximately five times that of aluminum, and therefore, aluminum-silicon alloys have excellent fluidity. Molds with low thermal conductivity and low mass density can be used when fluidity is a problem. Similarly, insulating mold coatings can reduce heat transfer and improve fluidity (Ref 6). The geometry of the sprue and gating will influence fluidity. A larger metal head will give an increased pressure and a higher velocity of the melt stream, which will improve fluidity. A higher melt temperature will strongly improve the fluidity. For example, it has been found for aluminum alloys that an increase of
Factors Influencing Fluidity Fluidity is a complex technological property and depends on many factors (Ref 4). These can be categorized as follows:
Fig. 1
Stoppage of metal flow. (a) Pinching of metal flow in a skin-freezing alloy with a narrow freezing range. (b) Choking of flow at the leading tip in an alloy with a long freezing range. Source: Ref 2
Fig. 2
Influence of alloy composition on fluidity in lead-tin alloys. Source: Ref 2
376 / Principles of Solidification melt superheat of 1 K (2 F) typically increases the fluid length 1% (Ref 6). The effects of other melt variables, such as gas content, grain refinement, and various trace elements, have been investigated (Ref 4); however, their effects are unclear or are small compared to other effects such as melt superheat.
Fluidity Measurements Fluidity data are important in order to optimize alloys and casting processes, and there are many tests available to obtain such data (Ref 4). Unfortunately, data from the various tests cannot always be directly compared, although they normally give similar trends. Another difficulty with fluidity tests has been poor precision and reproducibility. Fluidity tests are based on pouring metal into one or several channels and then measuring the fluidity length. Three such tests are shown in Fig. 3. It has been demonstrated that the precision of the spiral test in Fig. 3(b) can be improved to a relative error of approximately 5% by using an automatic pouring system that gives a constant pouring temperature and metal head (Ref 6).
and inadequate interdendritic melt feeding of the solidification shrinkage. More specifically, at the later stages of solidification, the dendritic network grows to an extent whereby it has sufficient strength to transmit thermal contractions, and tensile stresses will build up in a restrained casting that is not free to contract. Because the grain boundaries are still covered by a thin liquid layer, the material is still quite weak, and melt feeding to compensate for the solidification shrinkage is simultaneously taking place. Various investigators have explained hot tearing as a too-severe thermally induced deformation and/or an inadequate melt feeding of the solidification shrinkage. In addition to being controlled by these two macromechanisms, a series of microscopic effects and phenomena play a significant role in hot tearing susceptibility, such as grain morphology, dendrite bridging (coalescence) (Ref 10), the local interplay between the coalescing dendrites and the localized liquid feeding (Ref 11), and the alloy composition, including the effect of various trace elements. Some conditions must be met for a hot tear to develop:
Modeling of Fluidity
The
Fluidity can be predicted with specialized software that simultaneously addresses the heat transfer in the casting and mold, and the solidification and fluid flow in the casting. Provided that accurate data are available for thermophysical properties, coherency, and surface and interface heat-transfer coefficients, the fluid length can be calculated, assuming that dendrite coherency terminates the flow at a given solid fraction (Ref 7).
Hot Tearing Hot tearing occurs in some alloys during the final stages of solidification. It is a serious defect and occurs in castings that are restrained from shrinking in the mold. The characteristics of a hot tear are summarized by Campbell (Ref 2): Ragged branching crack Main tear with offshoots along intergranular
paths
Oxidized fracture surface with dendritic
morphology Location at a hot spot Apparent randomness in appearance Highly alloy-specific The mechanism of hot tearing has been extensively investigated; see the overviews by Sigworth (Ref 8) and Eskin et al. (Ref 9). It is commonly believed that hot tears start to develop in the mushy zone at solid fractions close to 1, and the phenomenon is associated with excessive thermally induced deformations caused by the nonuniform cooling contraction of the casting
casting must be restrained from contracting. A coherent network must be established. It should be noted that dendrite coherency in this context; that is, when a network is withstanding tensile forces, is different from, and occurs later than, dendrite coherency in the context of fluidity when the shear strength of the network is important. A liquid film must still be present on the dendrite or grain boundaries. A significant temperature decrease must occur after coherency. The solid network must be sufficiently dense to prevent feeding of liquid into the tear.
Bridges between dendrites must be established in order to withstand the tensile stress buildup. In aluminum alloys, it has been found that bridging between dendrites is established at liquid fractions of approximately 0.1, where typically 90% of the grain boundaries are still covered by liquid films (Ref 12). At this stage, the permeability of the solid network is so low that feeding of liquid into the dilating network is prevented. On the other hand, if the solid dendrite network is bridged to a large degree and a low fraction of the grain boundary is covered by liquid, the network will deform plastically rather than tear. A hot tear will start as a pore in the bulk material or at the surface. Nucleation of a hot tear will occur under conditions similar to a pore, most likely an inclusion. Removal of inclusions or other defects will reduce the hot cracking tendency. Solidifying material in a hot spot is particularly vulnerable to hot tearing, because the surrounding material is colder, stronger, and contracting.
Fig. 3
Fluidity tests. (a) Vacuum fluidity test. (b) Fluidity spiral in sand mold. (c) Multichannel fluidity test in metal mold
Alloy Dependence Figure 4 shows how hot cracking susceptibility varies with alloy concentration for aluminumcopper alloys. It can be seen that the curve peaks at approximately 1wt% Cu. Similar l-shaped curves have been constructed for other alloys (Ref 2). The curves can be understood from the solidification paths of aluminum-copper alloys
Castability—Fluidity and Hot Tearing / 377 develop during the vulnerable stage when the grain boundaries are still largely covered by liquid. This alloy is therefore more susceptible to hot tearing. Finally, at even lower alloy concentrations, the fraction liquid at grain boundaries will be too low, and the material will again become more ductile and thus less prone to hot tearing.
Hot Tearing Tests
Fig. 4
Influence of alloy composition on hot tearing susceptibility in aluminum-copper alloys. Source: Ref 2
Various tests have been designed for measuring the hot tearing tendency in alloys (Ref 9). Most of these are based on casting in a mold that restrains the metal from contracting, typically, ring-shaped castings contracting onto a rigid core or straight “dog-bone” castings restrained at the ends. The tests are evaluated by inspection for hot tears. A variant of the straight-type casting has several straight sections of different lengths, where the longest sections are the most likely to tear. This test can be used to assign a hot tearing sensitivity index to alloys. High-temperature tensile testing has been used to assess the strength of alloys close to their solidus, and the results have been used to evaluate the hot tearing tendency. A problem with the tests, however, is that most of them are designed to be carried out during melting. Their relevance for evaluating the behavior of solidifying alloys is thus questionable. More recently, tests have been reported for measuring the development of tensile stresses in castings (Ref 13), and it is expected that these will be important for the further understanding and quantification of hot tearing phenomena.
Modeling of Hot Tearing and Hot Tearing Criteria Fig. 5
Temperature dependence of solidification of Al-1wt%Cu (dashed curve) and Al-20wt%Cu (solid curve)
shown in Fig. 5, which shows the temperature versus fraction solid curves for Al-20wt%Cu and Al-1wt%Cu. The alloy with the higher copper content has a liquidus temperature of 592 C (1095 F) and a freezing range of 44 C (80 F). The temperature decrease during solidification occurs early, and the eutectic temperature is reached at a fraction solid of 0.45. The solid phase at this point is not coherent and therefore will not transmit any tensile stresses. Furthermore, the structure is open enough to feed any opening between dendrites that may occur. After reaching the eutectic temperature, the alloy will solidify isothermally without any thermal contraction. Thus, this alloy will not hot tear. However, the Al-1wt%Cu alloy has a liquidus temperature of 657 C (1215 F) and will solidify almost isothermally until a fraction solid of approximately 0.9. At this stage, the solid network is coherent and will transmit stress. The temperature will then decrease more than 90 until the eutectic temperature is reached, and a significant stress will
Most models and criteria attempting to relate hot tearing susceptibility to controllable material and process parameters address one of the previously mentioned main mechanisms: excessive thermally induced deformations or inadequate feeding of the solidification shrinkage. Some investigators argue that hot tearing will occur if the liquid is no longer able to fill the intergranular openings caused by the solidification shrinkage. Others define a vulnerable region for high solid fractions in which the permeability is low and only thin continuous films of interdendritic liquid exist that are not able to sustain possible thermal stresses. Based on continuum mechanics models that quantify thermally induced deformations and associated stresses, various criteria are proposed that consider overcritical values of stress, strain, or strain rate as the triggering factor for hot tearing. Again, the overview by Eskin et al. (Ref 9) is recommended for details; only a few main points are summarized here. Rappaz et al. (Ref 14) were the first investigators to simultaneously model feeding difficulties and excessive deformations. They considered microporosity formation as the initiating
mechanism for hot tearing and included the additional need for interdendritic melt feeding caused by the tensile deformation of the solidified material in the direction transverse to the columnar dendritic growth. Further elaborations on this idea can be found in Ref 15 to 17. Volume-averaged two-phase conservation equations for mass and momentum in the coexisting liquid and solid phases in a coherent and isotropic mushy zone are formulated in Ref 18 to 20. Both the interdendritic feeding and the thermally induced deformations were simultaneously taken into account, and the possibility of dilatation (opening up) of the dendritic network was included. A hot tearing criterion was then introduced (Ref 20) by calculating a weighted sum of shear and dilatation deformation in a vulnerable solidification interval defined by a limit pressure for solidification shrinkage feeding and a limit solid fraction above which the material obtains sufficient strength. The main challenge with all modeling approaches so far is to formulate, by means of variables included in the model, hot tearing criteria that accurately differentiate the hot tearing susceptibility for different choices of material and process parameters in real casting operations. The formalistic approach in Ref 18 to 20 furthermore reveals the need for accurate constitutive relations for the rheology of the deforming dendrite network and for the interdendritic melt flow permeability. Although many models and criteria correctly reflect the trends in hot tearing susceptibility caused by changes in these parameters, reality shows that significant research and development work is still needed before modeling tools of real predictive power are available and widely implemented into industrial daily practice. (See also the article “Modeling of Stress, Distortion, and Hot Tearing” in this Volume.)
REFERENCES 1. M.C. Flemings, Solidification Processing, McGraw-Hill, 1974, p 219 2. J. Campbell, Castings, Butterworth-Heinemann, 1991 3. L. Arnberg, L. Backerud, and G. Chai, Solidification Characteristics of Aluminium Alloys, Vol 3, Dendrite Coherency, AFS, 1996 4. M. Di Sabatino, “Fluidity of Aluminium Foundry Alloys,” NTNU Doctoral thesis, 2005, p 185 5. M. Di Sabatino, L. Arnberg, S. R rvik, and A. Prestmo, Mater. Sci. Eng. A, Vol 413–414, 2005, p 272 6. M. Di Sabatino, L. Arnberg, S. Brusethaug, and D. Apelian, Int. J. Cast Met. Res., Vol 19 (No. 2), 2006, p 94 7. M. Di Sabatino, L. Arnberg, and F. Bonollo, Mater. Sci. Technol., Vol 23 (No. 1), 2005, p 3 8. G.K. Sigworth, AFS Trans., Vol 106, 1999, p 1053
378 / Principles of Solidification 9. D.G. Eskin, Suyitno, and L. Katgerman Prog. Mater. Sci., Vol 49, 2004, p 629 10. M. Rappaz, A. Jacot, and W.J. Boettinger, Metall. Mater. Trans. A, Vol 34, 2003, p 467 11. S. Vernede, P. Jarry, and M. Rappaz, Acta Mater., Vol 54, 2006, p 4023 12. Y. Ju and L. Arnberg, Int. J. Cast Met. Res. Vol 16 (No. 6), 2003, p 522 13. C. Davidson, D. Viano, L. Lu, and D. St. John, Int. J. Cast Met. Res., Vol 19 (No. 1), 2006, 59
14. M. Rappaz, J.-M. Drezet, and M. Gremaud, Metall. Mater. Trans. A, Vol 30, 1999, p 449 15. M. Braccini, C.L. Martin, M. Sue´ry, and Y. Bre´chet, Modelling of Casting, Welding and Advanced Solidification Processes— IIX, Shaker Verlag, 2000, p 18 16. J.F. Grandfield, C.J. Davidson, and J. A. Taylor, in Continuous Casting, W. Schneider and W. Ecke, Ed., DGM, Frankfurt, Nov 13, 2000, p 205
17. Suyitno, W.H. Kool, and L. Katgerman, Mater. Sci. Forum, Vol 396–402, 2002, p 179 18. I. Farup and A. Mo, Metall. Mater. Trans. A, Vol 31, 2000, p 1461 19. M. Mhamdi, A. Mo, and C.L. Martin, Metall. Mater. Trans. A, Vol 33, 2002, p 2081 20. M. M’Hamdi, A. Mo, and H.G. Fjær, Metall. Mater. Trans. A, Vol 37, 2006, p 3069–3083
ASM Handbook, Volume 15: Casting ASM Handbook Committee, p 379-381 DOI: 10.1361/asmhba0005224
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Semisolid Metal Processing Qingyue Pan and Diran Apelian, Worcester Polytechnic Institute
SEMISOLID METAL (SSM) PROCESSING, also known as SSM casting or semisolid forming, is a special die casting process where a partially solidified metal slurry (typically, 50% solid and 50% liquid instead of a fully liquid metal) is injected into a die cavity to form die cast components. The key to SSM processing is to generate an SSM slurry that contains a globular primary phase and exhibits thixotropic behavior. Semisolid metal processing not only produces complex dimensional details (e.g., thin-walled sections), as in high-pressure die casting, but also produces high-integrity castings currently attainable only with squeeze or lowpressure permanent mold casting. Semisolid metal processing offers several distinct advantages over conventional near-net shape forming technologies. These advantages include low manufacturing cycle time, increased die life,
reduced porosity and solidification shrinkage, improved mechanical properties, and so on.
Flow Behavior of SSM Viscosity Evolution During Continuous Cooling The origin of SSM can be traced to the experiments conducted by David Spencer at Massachusetts Institute of Technology in 1971 (Ref 1, 2). Spencer’s work originally was directed toward understanding hot tearing in alloy castings; however, in the course of his viscosity measurements of semisolid Sn-15wt %Pb slurries, he discovered that if a semisolid slurry is sheared above the liquidus temperature and slowly cooled into the two-phase region,
it can have an apparent viscosity that is orders of magnitude less than a slurry that is cooled to the same temperature (solid fraction) before shear, as illustrated in Fig. 1. Detailed microstructural analysis revealed that the difference in apparent viscosity values comes from the significant difference in morphology of the solid phase in the slurries (spherical versus dendritic). Moreover, at a given solid fraction, the apparent viscosity decreases with increasing shear rate or decreasing cooling rate, as illustrated in Fig. 2. This discovery has led to the development of semisolid processing.
Rheology of SSM An SSM slurry is a mixture of liquid metal and solid particles that have a predominantly
Fig. 1
Experimentally determined viscosity and shear stress versus fraction solid of Sn-15%Pb alloy cooled at 0.006 K/s with a shear rate of 200 s 1. Source: Ref 2
Fig. 2
Effect of shear rate on apparent viscosity of semisolid alloys. Source: Ref 3
380 / Principles of Solidification globular shape. The rheological behavior of SSM slurries is similar to that of viscoplastic materials, such as concentrated suspensions, pastes, foodstuffs, emulsions, and foams. This class of materials is characterized by the existence of a yield stress, which implies that the material behaves as a solid if the applied stress is below the yield stress; however, if the applied stress exceeds the yield stress, the material will flow and exhibit liquidlike behavior. The current understanding of the complex rheology of SSM slurries is based on the fact that under steadystate conditions, the semisolid slurry behaves as a shear-thinning (pseudoplastic) fluid, wherein the effective viscosity decreases with increasing shear rate. However, under rapid transient conditions, the semisolid slurry behaves as a shear-thickening fluid; that is, effective viscosity increases with increasing shear rate.
Yield Stress of Semisolid Slurries The yield stress of an SSM is the result of a continuous solid skeleton matrix formed by welded particles, physical crowding of particles, or other interparticle interactions during processing. The presence of a distinct yield stress in a semisolid material significantly modifies its rheological behavior. Both experiments and numerical simulations show that, depending on the flow velocity and yield stress of a semisolid slurry, there are five distinct die-filling patterns during semisolid processing: shell, disk, mound, bubble, and transition. The distinctive filling patterns are illustrated in Fig. 3. The measured yield stress of A356 and 357 aluminum alloys as a function of temperature in the two-phase region is shown in Fig. 4 and 5. It is found that:
The yield stress of commercial semisolid
aluminum alloys is a strong function of temperature (solid fraction). In the commercial forming temperature region, as temperature increases, yield stress decreases dramatically, varying between 103 and 10 1 kPa (150 and 0.015 psi). At a given solid fraction, the yield stress of semisolid slurries depends on microstructural indices (i.e., entrapped liquid content, shape factor of the alpha phase, and the alpha particle size). Low entrapped liquid content and small, round alpha particles tend to decrease the stress needed to initiate the flow, that is, the yield stress of the slurry.
Semisolid Microstructural Formation Solidification Microstructure Under Forced Convection Nearly all commercial alloys solidify dendritically with either a columnar or an equiaxed dendritic structure. During solidification, there are a number of processes taking place simultaneously in the two-phase region. These include ripening, interdendritic fluid flow, and solid movement. The dendritic structure is greatly affected by convection during the early stages of solidification, whereas under vigorous convection and slow cooling, the solidification structure tends toward a rosette or a spherical morphology. Figure 6 illustrates current understanding of microstructual formation under forced convection, specifically: Forced convection promotes crystal growth
Fig. 3
Flow patterns as a function of Bingham number, Bi, and Reynolds number, Re. The figure illustrates five distinct filling patterns (bubble, mound, shell, disk, and transition) in semisolid casting. Source: Ref 4
Fig. 4
Yield stress of GR A356 and MHD A356 alloys as a function of temperature. Source: Ref 4
Fig. 5
Yield stress of GR 357 and MHD 357 alloys as a function of temperature. Source: Ref 4
due to an enhanced mass transport in the two-phase region. Under forced convection, a laminar flow changes dendritic growth morphology to a rosette, whereas a turbulent flow changes growth morphology from a rosette to a sphere. Several mechanisms have been proposed to explain the observed nondendritic
Fig. 6
Schematic illustrating morphological transition under forced convection. Source: Ref 5
Semisolid Metal Processing / 381 REFERENCES 1. D.B. Spencer, Ph. D. thesis, MIT, Cambridge, MA, 1971 2. D.B. Spencer, R. Mehrabian, and M.C. Flemings, Metall. Trans., Vol 3, 1972, p 1929–1932 3. M.C. Flemings, Metall. Trans. B, Vol 22 (No. 3), 1991, p 269–293 4. Q.Y. Pan, D. Apelian, and A.N. Alexandrou, Yield Behavior of Commercial Aluminium Alloys in the Semi Solid State, Metall. Mater. Trans. B, Vol 35, Dec 2004, p 1187–1202 5. Z. Fan, Int. Mater. Rev., Vol 47 (No. 2), 2002, p 49–85
SELECTED REFERENCES A.N. Alexandrou and F. Bardinet, J. Mater.
Process. Technol., Vol 96, 1999, p 59–72
A.N. Alexandrou, E. Duc, and V. Entov,
Fig. 7
Schematic of semisolid metal (SSM) processing routes: thixocasting versus rheocasting. MHD, magnetohydrodynamic; GR, grain refinement
morphology. These include dendrite arm fragmentation, dendrite arm root remelting, and growth-controlled mechanisms. It seems that the observed rosette and spherical morphologies are more likely related to a growth phenomenon rather than mechanical fragmentation.
Globular Structure via Copious Nucleation Originally, to achieve an SSM slurry, it was thought necessary to cool the liquid into the two-phase region and to shear off and break the dendrites (i.e., melt agitation via mechanical or magnetohydrodynamic stirring). However, recent research and development show that it is not necessary to use shear forces to produce the globular semisolid structure. Instead, a slurry with an ideal semisolid structure can be produced directly from the melt if the temperature of the melt is such that many nuclei (copious nucleation) are generated, and the nuclei do not grow past a certain point (i.e., suppression of dendritic growth), nor melt back into the bulk liquid. This controlled nucleation and growth concept is the
genesis of various processes and methodologies to generate semisolid slurries from the liquid state.
Semisolid Processes: Thixocasting versus Rheocasting There are two primary semisolid processing routes: thixocasting and rheocasting. Thixocasting begins with a nondendritic solid precursor material that is specially prepared by the manufacturer using continuous casting methods. Upon reheating the material into the mushy (i. e., two-phase) zone, thixotropic slurry forms and becomes the feed for the casting operation, as illustrated in Fig. 7. Rheocasting (also known as slurry-on-demand) starts with material in the liquid state, and the thixotropic slurry is formed directly from the melt via special thermal treatment/management of the system; the slurry is subsequently fed into the die cavity. Of th
