TROUBLESHOOTING THE STEELCASTING PROCESS

TROUBLESHOOTING THE

STEELCASTING PROCESS

Edited by Dr. John M. Svoboda

Vice President-Technology and

Barbara Linskey Technical Assistant

Steel Founders’ Society of America, 1987 Cast Metals Federation Building

455 State Street, Des Plaines. Illinois 60016

Printed in the United States of America

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TABLE OF CONTENTS

Preface ............................................................... vii

Lecture I-Inclusions: Their Formation, Detection and Control. ........................................ 1

Dr. I.D. Sommervville and S. Dawson

Lecture Il-Gases in Steel Castings. ........................ 61 Dr. John M. Svoboda

Lecture Ill-Erosion & Expansion Type Steel Casting Defect.. ....................................... 103

Matt J. Granlund

Lecture IV-Linear Surface Discontinuities. ............. 123 Dr. W.J. Jackson

Lecture V-Adhering Sand Defects. ........................ 215 Dr. John M. Svoboda

Lecture VI-Shrinkage in Steel Castings. ................ 241 Ronald W. Ruddle

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PREFACE

The six lectures published in this book were presented at the 41st Technical and Operating Conference of the Steel Founders’ Society of America on November 15,1987 in Schaum- burg, Illinois. This lecture series was presented in university style format as a comprehensive short course covering steel casting defect analysis and prevention. The full day was devoted to the presentation and discussion of the papers.

In recent years, an extensive number of programs and publica- tions have been developed which cover many aspects of the quality assurance of castings, particularly the use of statistical process control. While these techniques have proven valuable, it remains the responsibility of the operating foundrymen to possess a thorough understanding of the mechanisms by which steel casting defects occur and the procedures and con- trols by which the defects may be prevented. While literally hun- dreds of papers have been published on these subjects, this is the first time that a basic “Handbook of Steel Casting Defects” has been organized and published. It was the intention of the Technical and Operating Committee that this text serve as a basic guide to all steel foundrymen.

The lectures have been prepared and presented by authorities on the subject of steel casting quality. Each lecturer was chosen because of his full understanding of the subject and his ability to impart his knowledge to those not necessarily too familiar with the technical terminology involved. The Society was indeed fortunate to be able to secure and present such a fine group of able speakers. A short biographical sketch of each lecturer is presented.

Technical and Operating Conferences are annual events under the direction of the Technical and Operating Committee for the benefit of the technical and operating personnel of the member companies of the Steel Founders’ Society of America. The “Lec-

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ture Series” was initiated at the 1958 T & O Conference and was designed to cover topics of prime interest in the operation of a steel foundry. Sessions have been held periodically since then.

The 1987 Conference was under the direction of Mr. John F. Moore. Dofasco Inc., Chairman of the Technical and Operating Committee. Other members of the Committee are as follows: R. F. Atkinson. APV Paramount Limited; C. T. Brandt, Missouri Pre- cision Castings; s. D. Hays, Fisher Cast Steel Products, Inc.; B. Johnson. Columbia Steel Casting Company, Inc.; J. A. Larson, Ingersoil-Rand Company; M. Rauguth, Maynard Steel Casting Company; W. Sanders, Advance Foundry Company; A K. Zaman. Rockwell International; T. Englat, Alloy Foundries Div.

November 1987

Dr. John M. Svoboda Vice President-Technology

Barbara Linskey Technical Assistant

Steel Founders’ Society of America

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ABOUT THE AUTHORS.. .

Dr. I. D. Sommerville NSERC Industrial Professor University of Toronto

On leaving school, Dr. Sommerville spent two years working in the steel industry before taking up full-time studies in metallurgy at the Royal College of Science and Technology, Glasgow, now the University of Strathclyde, from which he graduated in 1963 with the degrees of B.Sc. and A.R.C.S.T., both with first class honours, followed by a Ph.D. in 1966. After two years as a Research Fellow in the University of Sheffield, during which he investigated the thermodynamics of electro- slag remelting slags, he returned to the teaching staff at the University of Strathclyde, where he taught non-ferrous extraction, iron and steelmaking and advanced physical chemistry of extraction processes. During this period he carried out research in the area of slag thermodynamics, the kinetics of slag-metal reactions and water-modelling of continuous casting, being awarded the John Chipman Medal in 1981 for part of the slag-metal kinetic work. Also in 1981, he emigrated with his family to Canada to take the position of Senior Research Associate in the Department of Metallurgy and Materials Science at the University of Toronto, where he was appointed Associate Professor in 1983. In 1984, he was awarded an NSERC Industrial Professorship, the first Chair of its kind in metallurgy in Canada.

His current research interests are mainly focused on the development of improved slags and fluxes for the external processing of iron and‘steel, including the external desulfurization of hot metal and slags for ladle, tundish and mold dur-

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ing continuous casting. Other areas of interest are the application of plasma techniques in extractive metallurgy, the ration- alization of welding flukes, the electrical conductivity of sub- merged arc furnace slags, and the development of advanced refractories for steel casting systems. He is the author of around fifty papers. and a member of ISS-AIME and of Casteel Technology Associates, Inc.

Steve Dawson Research Associate

University of Toronto

Steve graduated with a bachelor of engineering in the metallurgical discipline from McGill University in December 1985. He then moved to the University of Toronto where he obtained a master’s of applied science in February 1987 for research dealing with the evaluation of liquid steel cleanliness. He is currently enrolled in the doctoral program as a member of the Ferrous Metallurgy Research Group at the University of Toronto.

During the final year of undergraduate studies at McGill, Steve won a Canadian essay competition sponsored by the Cana- dian Continuous Steel Casting Research Group for a review paper on horizontal continuous casting. He also received the Gordon Sproule Memorial Award which is given to one student each year in the graduating class of the metallurgical department. Since arriving at the University of Toronto in January 1986, Steve has received a University of Toronto Open Scholar- ship, an Ontario Graduate Scholarship, and a national scholarship from the Natural Sciences and Engineering Research Council of Canada. Steve has also received the 1987 Jerry Silver Memorial Award which is presented each year at the Iron and Steel Congress for the most outstanding presentation in the Continuing Research Section. Steve is a member of the ISS, the CIM, and the ASNT.

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Dr. W. J. Jackson Consultant

Dr. Jim Jackson retired in 1986 after 32 years with SCRATA. After wartime service as a radar mechanic in the Royal Air Force, he took his first degree at the University of Birmingham in Industrial Metallurgy. He then joined the British Oxygen Company working on the development of MIG and TIG welding, and later joined INCO as a technical assistant in the hot rolling department.

After gaining a Master’s Degree at London University, Dr. Jackson moved to SCRATA, Sheffield in 1954, and amongst the early papers published in his name were two which became classical reference papers on Subzero Impact Properties and Magnetic Properties of Steel Castings. He later became involved in high strength cast steels and his Ph.D. was awarded for work in this field. Other areas of involvement have been in steelmaking, in which he made many contributions to desulphurisation and deoxidation technology, including a text book “Steelmaking for Steelfounders” which received international acclaim. More recently, he has been involved with fracture mechanics, cracking mechanisms, engineering properties and non-destructive testing. He has always been actively involved with steel casting specifications, and has been a member of many committees, including those of SCRATA, British Standards, other national organisations and ISO.

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Dr. John M. Svoboda Vice President-Technology

Steel Founders’ Society of America

Prior to joining the Steel Founders’ Society in 1981, Dr. Svoboda’s background included several years as a technical consultant to the foundry industry with clients in the United States, Canada, Mexico and Brazil. As a consultant, he specialized in the area of quality control, including sand control, gating and risering, melting and metallurgy.

Before becoming a consultant, Dr. Svoboda was Director of Education for the American Foundrymen’s Society (AFS) Cast Metals Institute. In that position he developed a sixty-course educational program, prepared text, programmed learning courses and designed and developed laboratory facilities.

In 1985, Svoboda assumed the additional duties of President of Casteel Technology Associates, Inc., the technology development subsidiary of the Steel Founders’ Society of America.

Dr. Svoboda, who held the rank of Captain in the U.S. Army Corps of Engineers, held the positions of Project Metallurgist and Supervisor, Foundry Process Evaluation and Control for the Falk Corporation for ten years. He received B.S., M.SD. and Ph.D. degrees in metallurgical engineering from the University of Wisconsin. He is an active member of the American Found- rymen’s Society, Iron and Steel Society-AIME, American Soci- ety of Association Executives, American Society for Metals, American Society for Testing and Materials, the Materials Properties Council, and is a Fellow of the Institution of Diagnostic Engineers. In 1982, Svoboda presented the ex- change lecture to the Steel Castings Research and Trade

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Association in England and has been elected to membership in the British Iron and Steel Casting Institute. In 1987, he was elected to the Cope and Drag Club.

Dr. Svoboda co-authored, with Raymond Monroe, a paper on ladle desulfurization which received the prestigious Best Paper Award from both the Electric Furnace Conference and the AFS Steel Casting Division in 1984.

Ronald W. Ruddle Ronald W. Ruddle & Associates CTA Associate

Ronald W. Ruddle received his bachelor’s degree in Natural Sciences (Metallurgy) from the University of Cambridge in 1941 and his M.A. degree in 1945.

He was employed in Alcan’s aluminum foundry in Birming- ham, England (formerly Northern Aluminum Co.) from 1941 to 1942. He was engaged in the production of aluminum forgings from 1942 to 1946 at the DeHavilland Forge Co., South Wales, a subsidiary of the DeHavilland Aircraft Co. where he eventually rose to Chief Metallurgist.

He joined the staff of The British Non-Ferrous Metals Research Association in 1946 and was Head, Melting and Casting Section from 1947 to 1956.

Mr. Ruddle was appointed Technical Manager of Foseco Inc., Cleveland, Ohio in 1957, later becoming Vice-president, Tech- nology Foundry Division. He retired from the position of Direc- tor of Special Projects at Foseco in 1984, to establish his own consulting business.

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He is the author of more than 50 papers on various aspects of metal casting and refining, and of books on solidification, running and gating, risering and on the chemistry of copper and lead smelting.

Mr. Ruddle was Chairman of the A.F.S. Research Board and a member of Technical Council from 1965 to 1971. He was awarded the Society's Simpson Gold Medal in 1967 and was the Society's Hoyt Lecturer in 1971. He is a member of several other A.F.S. committees and has also been a member of a Na- tional Science Foundation Advisory Committee on solidification research.

Besides being a member of the A.F.S., Mr. Ruddle is a member of the A.I.M.E., ASM International, the Institute of Metals (British), and the Institute of British Foundrymen. He is a Fellow of the (British) Institution of Metallurgists.

Matt J. Granlund CTA Associate

Foundry Systems Control

Immediately upon completing his undergraduate studies at the University of Minnesota, Mr. Granlund entered the field of foundry sand technology and has remained there ever since. He has worked in the areas of sand system control and sand reclamation for Archer Daniels Midland-Ashland Chemical, Kordell Industries, Ablex Corporation Research Center and National Engineering Company over a span of thirty-two years.

Active in the American Foundrymen's Society, Mr. Granlund has served as a member andlor officer of many AFS sand related committees. In addition, he has published several sand related papers in the AFS Transactions.

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Mr. Granlund was recognized by his peers with the American Foundrymnen's Society Award of Scientific Merit in 1978 and the AFS Cast Metals Institute Director's Award in 1980.

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Lecture I

Inclusions: Their Formation, Detection and Control by I.D. Sommerville and S. Dawson

Ferrous Metallurgy Research Group Department of Metallurgy and Materials Science University of Toronto

INTRODUCTION

As normally produced, steel contains considerable quantities of second phase particles which are generally referred to as inclusions. While second phase particles are formed by intermetallic compounds, it is generally non-metallic particles which are intended when inclusions are mentioned. This latter group, which is far more prevalent and important than the former, can be subdivided into oxides, sulfides, carbides and nitrides.

While the effect of such inclusions on the properties of the metal is generally harmful, it should be realized that this is not invariably the case. Manganese sulfides, for example, can be particularly beneficial with regard to the free machining properties of steel while they are virtually intolerable in steels which must possess high ductility at low temperatures. An outline of the effects of selected second phase particles on some of the mechanical and physical properties of steel is provided in Table I, (1) which demonstrates that the ultimate effect of an inclusion can only be evaluated in consideration of the end use of the steel. Detailed studies addressing such factors as inclusion chemistry, size and origins have been published by Keissling (2) and Pickering. (3) Of all inclusion types, oxides are the most abundant and potentially comprise up to ninety- five percent of the total inclusion count. (4,5)

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THE FORMATION OF INCLUSIONS Inclusions can be divided into two groups: those of indigenous origin and those of exogenous origin. Indigenous inclusions are those formed by reaction of an element dissolved in the steel with an added element or with another phase. These inclusions consist primarily of oxides and sulfides and may be induced either by additions made to the liquid steel or by reductions in solute solubilities during the cooling and solidification of the steel. Conversely, exogenous inclusions are

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the result of slag entrainment or the mechanical incorporation of eroded or fragmented refractory materials such as protective shrouds, vessel linings, and mold materials. These primarily oxide based inclusions are readily distinguishable from their indigenous counterparts as they are typically larger in size, more complex in shape and structure, and often occur sporadically in preferred locations of an ingot or casting. It is, however, worth noting that not all inclusions can be strictly classified as either indigenous or exogenous: since complex inclusions can form as a result of exogenous particles acting as nucleation sites for the deposition of indigenous reaction products. (4.6)

Since oxide inclusions form such a high proportion of the inclusion population in steel, the major part of this review will be devoted to them. More specifically, the inclusions resulting from deoxidation and reoxidation processes will be distinguished and discussed in the following sections.

Inclusions Formed by Deoxidation

Since steelmaking is essentially a process of refining by selective oxidation, the oxygen content at the end is many times higher than the solubility limit for oxygen in the solidified metal. Consequently, in order to minimize porosity and inclusions in the cast product, this excess oxygen must be removed as far as possible while the steel is still liquid. The practice of lowering the oxygen level is known as deoxidation.

Deoxidation consists of the addition, dissolution, and diffusion of an alloying element (typically aluminum, silicon or manganese) which has a higher affinity for the dissolved oxygen than iron does at steelmaking temperatures. The diffusion and mixing are followed by reaction with the dissolved oxygen. nucleation and growth of the deoxidation products, and finally, their removal from the steel. Since their formation is considerably faster than their removal, they accumulate in the steel resulting in a large discrepancy between the total oxygen content of the steel as measured by chemical analysis and the dissolved oxygen content of the steel as measured by solid state electrolyte probes. The difference between the two values is accounted for primarily by the presence of oxide in-

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clusions, and decreases with time as these inclusions are removed.

The relative deoxidizing powers of various elements at 1600 °C are compared in Figure 1, due to Turkdogan. (7) Such a diagram can be used to predict the dissolved oxygen content of the steel in equilibrium with any content of deoxidizer, but gives no information on the total oxygen content of the steel, which is partly governed by physical factors. Also, since the precipitation deoxidation reactions are exothermic, they pro- ceed further with fall in temperature, so that this equilibrium dissolved oxygen value decreases as the steel cools and solidifies. Deoxidation products formed prior to the onset of solidification are called primary deoxidation products, and theoretically these can be removed completely, though this is seldom if ever achieved in practice. Deoxidation products formed once solidification has commenced are called secondary deoxidation products, and their removal cannot be complete, since some at least will be trapped by the advancing solidification front.

The difference between the various deoxidizing elements with respect to ultimate steel cleanliness has been investigated by several researchers. The most widely used deoxidizing elements. in order of deoxidizing strength, are manganese, silicon, and aluminum. Despite the extensive research conducted on deoxidation, there remains some controversy over the order of addition and the relative amounts of each element required to optimize the process. ltskovitch (8) found that coarse AI

2O

3 particles formed more readily when the initial

oxygen content was high and that these particles were more easily removed form the steel. Thus, ltskovitch recommended the addition of aluminum prior to the addition of any ferro- silicon. In contrast, Frage et al (9) have found that fewer oxide inclusions form when silicon is added first, followed by aluminum. The suggested explanation for this phenomenon is that silicon deoxidation results in larger inclusions which are subsequently reduced by aluminum to form even larger AI

2O

3

inclusions. Contradictory results of this nature suggest that the order of addition may depend on the secondary conditions of the process being conducted, such as temperature and

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chemistry. This premise is supported by the results of Yavoisky (10).

From Figure 1, it is clear that aluminum is a much more powerful deoxidizer than silicon, and it can therefore be used alone. Deoxidation to very low oxygen levels with silicon alone would require dissolved silicon contents in excess of the specified maximum for many grades of steel. It is fortunate, therefore, that although the deoxidizing power of manganese is considerably lower than that of silicon, it can be used to enhance the deoxidizing power of silicon. This is due to the fact that when the ratio of manganese to silicon in solution in the steel is four or greater, the product of deoxidation is not solid silica but rather a liquid composed of FeO, MnO and SiO

2 in which the

activity of SiO2 is considerably less than unity. The fact that a

critical manganese to silicon ratio exists is illustrated in Figure 2(a), while the way in which it changes with temperature is shown in Figure 2(b). The formation of liquid rather than solid inclusions also has repercussions on the rate and mechanism of their removal, as will be discussed in the last section of this review.

A similar effect can be obtained by the use of calcium together with aluminum as deoxidizers. As shown in Figure 3, (11) a lower dissolved oxygen content is obtained because the deoxidation product is liquid calcium aluminate rather than solid alumina, and consequently the activity of alumina is reduced. This is quantified further in Figure 4, (12) where the residual dissolved oxygen content decreases as the lime content of the aluminate increases. with the consequent decrease in the activity of alumina.

It is well established that primary deoxidation products are typically less than 10µm in size. Because inclusions much larger than this frequently occur in steel, it is necessary to account for the formation of these larger inclusions. The first such formation mechanism is the growth of primary deoxidation products. Approximately twenty years ago it was felt that the principal method of inclusion growth was by diffusion of atoms through the melt to the inclusion site and subsequent deposition. More recently however, it has been shown both by

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calculation and by experimental observation that within the first second after addition of the deoxidant 99.9% of the original dissolved oxygen is precipitated. (10,13-15) The inclusions therefore cannot grow by diffusion and precipitation of dissolved oxygen and thus a second mechanism must be considered to account for inclusion growth. The accepted mechanism is that of collision, followed by coalescence which results from the relatively high surface tension between the inclusions and the liquid steel. Silica, which has a melting point of 1728°C is present in the melt in a deformable state, and thus, when two silica inclusions collide they can develop into a single spherical inclusion. Inclusions resulting from combined silicon-manganese deoxidation also form spherical conglomerates. Conversely, alumina (MP ~2015°C) is a rigid solid at steelmaking temperatures and therefore sinters together to

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form large irregularly shaped inclusions which are known as clusters or galaxies.

Studies of inclusion sizes in killed steels have found that alumina galaxies typically range from 20 to 30 µm (3,16) while spherical silica conglomerates are 25 to 30µm in diameter. (13) Complex deoxidation products from combined deoxidation practices have also been found to be of a similar size. (16)

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Thus, residual deoxidation products are seldom large enough to cause serious defects or failures. Considerably larger and hence more dangerous inclusions are produced by reoxidation and by physical entrainment of slags and fragmented refractories. The inclusions formed by these mechanisms are discussed in the following sections.

Inclusions Formed by Reoxidation

Reoxidation is the increase in the total oxygen content of molten steel caused by the chemical interaction of iron or some element dissolved in the steel with oxygen supplied either by the atmosphere or by oxides such as refractories and fluxes. On the basis of this definition, reoxidation products should be regarded as indigenous inclusions. Thus, inclusions which form due to physical rather than chemical interaction of

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the steel. such as those due to mechanical erosion of refractories or to entrapment of slag, are not included in this category since these already existed as oxides and therefore should be classified as exogenous inclusions.

Reoxidation occurs by three separate mechanisms:

(a) direct oxidation of steel streams (as they fall through air).

(b) oxidation of the metal in the receiving pool by air entrained and of the liquid metal surface,

by the plunging stream. and

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(c) chemical interaction of the steel with solid or liquid oxide phases, such as refractories or slags.

The existence of the problem has been known for many years. As far back as 1945, Hultgren (17) drew attention to the presence of large inclusions in ladle samples of electric arc steel deoxidised in the furnace, which contained no large inclusions prior to tapping. In 1950 Brower et al (18) reported that reoxidation during tapping of open hearth heats was the main cause of alloy losses, while oxygen pickup during casting was reported by various authors between 1959 and 1970. Vingas and Caine (19) have ascribed macro-inclusions found in steel castings to reoxidation, while Van Vlack and Flinn (20-23) have discussed the reoxidation of steel deoxidised by aluminum during the manufacture of castings. However, it is really in the more recent literature, and particularly in that relevant to continuous casting, that reoxidation has become accepted as a major source of inclusions in the final steel. (24-30)

Reoxidation is particularly undesirable due to its potential to form large inclusions late in the steel casting process (teeming streams for example) where there is a high probability that the reaction products will be incorporated in the final product rather than floating out of the bath. With this in mind it becomes evident that reoxidation during the final transfer operation has the potential to nullify all of the steps taken previously to produce a clean steel. Atmospheric reoxidation between the ladle and tundish has been found to be responsible for increasing the number of large inclusions (100 to 200 µm) by a factor of 2.5 in continuous casting operations. (26) This size and number range has been confirmed by Tamamato et al (30) and Farrell et al, (31) both of whom found that reoxidation products were considerably larger than deoxidation products. and typically ranged from 100 to 200 µm.

The fact that reoxidation can occur in the mold immediately before solidification. when the possibilities of the products escaping are at a minimum, undoubtedly makes reoxidation an even more serious problem in the production of steel castings than in continuous or in ingot casting. In this case, reoxidation by reaction with air in the mold cavity and with

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moisture in the mold binder occurs in addition to that caused during the steel transfer from the ladle.

Inclusions due to reoxidation can be distinguished from those due to deoxidation by their size, as already indicated, and also by their composition. Farrell et al (31) have shown that reoxidation products tend to be richer in the weaker oxide forming elements such as silicon and manganese than the indigenous inclusions. More specifically, they concluded that in steels such as most fine grained qualities which contain sufficient aluminum for this to be in control of deoxidation, the presence of silicate inclusions is direct evidence of reoxidation. More recently Venkatadri (15) has concluded that any spinel-type inclusions are very probably the result of reoxidation leading to localised regions of high oxygen content. This is due, as illustrated schematically in Figure 5, to the fact that in the case of reoxidation there is a virtually infinite oxygen source so that once all the aluminum has been consumed, there is still plenty of oxygen remaining to react with silicon and manganese. Under these circumstances, the resulting inclusions are com- plexes containing silica, manganese oxide, and in extreme cases, iron oxide in addition to alumina. They can therefore include spinels such as galaxite and hercynite. The sequence shown in Figure 5 is the thermodynamic sequence. In fact, because iron is so much more available to the incoming oxygen, it seems highly likely that the initial product is iron oxide, which dissolves to give oxygen in solution which then reacts in accordance with the sequence shown.

In the case of steels deoxidized with silicon and manganese, the distinction between the products of reoxidation and deoxidation will not be so clear in terms of their composition, but there will still be a clear difference in their size.

Laboratory experiments using falling droplets have demonstrated that oxygen pickup by molten steel from air is extremely rapid, as shown in Figure 6. (32) Since the rate is controlled by mass transfer in the gas phase, oxygen is consumed as fast as it is supplied, and the rate is unaffected by the presence of either surface active solutes such as sulfur and oxygen or strong deoxidants such as aluminum and titanium.

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In the plant, reoxidation of the stream can be clearly observed by taking high-speed film, and Figure 7 (33) is taken from such a film. In-plant studies have also shown that carbon exercises a protective effect, mainly because it reacts with oxygen to produce a layer of carbon monoxide which essentially acts as a gaseous shroud inhibiting the ingress of further oxygen.

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In spite of the very rapid pickup of oxygen by steel streams falling through air, the vast majority of the reoxidation in most instances is caused by the entrainment of air into the receiving pool. Four different mechansims of air entrainment have been elucidated (34) depending on the turbulence level of the stream, as depicted in Figure 8, and it has been shown that the rate of air entrainment depends directly on the turbulence intensity of the stream, the height of fall and inversely on the nozzle d i ameter.

While interaction of the steel with gaseous oxygen probably accounts for most of the reoxidation observed in practice, its

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reaction with the oxide phases in slags and refractories can also make a significant contribution. In this case, the presence of a strong deoxidant such as aluminum is very important, since it has a much greater affinity for oxygen than iron, manganese or silicon. Thus the oxides of any of these elements will be reduced by aluminum dissolved in steel so that, for example, aluminum creates inclusions by the chemical erosion of siliceous refractories and converts siliceous inclusions in the steel to alumina or aluminates. Since contact

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with refractories in ladles arid the various components of the casting system clearly cannot be eliminated, the solution lies in the selection of refractories with compositions shown to be resistant to erosion. This will generally involve higher alumina and lower silica contents and greater use of basic refractories, the ultimate solution probably being dictated by economic rather than metallurgical aspects. Presumably, reoxidation by erosion occurs when aluminum killed steels are cast into sand molds, but since the composition of the molding sand is determined by other criteria, it is difficult to see what can be done in this case, other than minimising the residual aluminum content by careful deoxidation practice.

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Exogenous Inclusions

The most common sources of exogenous inclusions are slag entrainment and the degradation of refractories by physical interaction with the steel. Inclusions due to slag entrainment can be caused by gas bubbling in the ladle at an excessive flowrate, since pulsations at the slag-metal interface can lead to disintegration of the slag and entrainment into the steel bath. They can also be caused, particularly when the steel in the ladle has been reduced to low levels, by vortexing.

Physical erosion of refractories in the ladle, and more particularly in the runner and gating system as well as in the casting mold is also a source of large inclusions. Of these, it appears that the mold is the most critical area, and that appropriate alterations of mold material can have a significant effect on steel quality. Defects and inclusions due to variations in mold materials and practices have been reviewed in another paper at this conference.

The Size and Distribution of Inclusions

Kiessling (2) showed that a hypothetical steel containing 100 ppm of oxygen present as uniformly sized spherical alumina particles of specific gravity 4.0 should contain between 1012

and 1015 oxide inclusions per metric tonne. A similar number of sulfide inclusions may also be present. (3) The size, number, and spatial distribution of inclusions are intimately related and can vary greatly as is shown by the hypothetical range of values listed in Table 2.

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Further calculations based on Table 2 for a total oxygen content of 20 ppm show that at an average inclusion diameter of 10µm , there are 2 x 1010 inclusions per tonne of steel, with a mean spacing of 185 µm. In two dimensions, these values cor- respond to the presence of over four hundred inclusions (average diameter of 5 µm each) in each square millimeter of metallographic surface. Experimental verification of the calculations in Table 2 were provided by Bergh and published by Kiessling. (2) In this study, Bergh observed a mean inter- inclusion space of 30 µm and an average inclusion size of 3µm. On a volume basis this corresponds to approximately 1015 oxide inclusions per tonne. Similar experimental confirmations have been provided by Lindon and Billington (35) and by Lind- borg and Torsell. (13)

To fully complement the values listed in Table 2 it is necessary to establish typical oxide particle sizes which will determine a range for which the data are of practical value. The current Iit- erature (2,3,13-16,26,30,31,35) with some deviation list primary aluminum deoxidation products as being less than 5 µm while primary silicon deoxidation products are typically less than 10 µm in plain carbon steels. Reoxidation products of aluminum and silicon along with these elements can span from 50 to 300 µm depending on the concentration of other reactants present during the atmospheric exposure time. Finally, refractory erosion products typically range from 50 to 100 µm in aluminum killed steels while silicon killed steels contain erosion products in excess of 100 µm. The refractory erosion products in silicon killed steels are often larger than those in aluminum killed steels for two reasons. Firstly, the silicates which are eroded from the refractory by aluminum killed steels are prevented from coalescing because they are altered by reaction with aluminum in the steel. Secondly, it has been shown that aluminum in the steel decreases the amount of liquid in the erosion boundary layer and that this effect tends to decrease the size of erosion products which are formed by scouring the vessel lining.

Spanning beyond the realm of oxides, it has been established that the volume fraction of non-metallics in plain carbon steels is roughly 60-90% oxides, 5-35% sulfides, and 3-10% oxy-

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sulfides. The total mass fraction of inclusions present in steel can range from 0.012 wt % to 0.025 wt %. (36) The spatial distribution of sizes for all types of inclusions present in steel has been found by innumerable authors to approximate a log normal distribution. In conclusion it is important to note that the values presented in this discussion represent approximate estimates for plain carbon steels and variation from this range is possible.

THE DETECTION OF INCLUSIONS

The various methods available for determining metal cleanliness can only be evaluated in consideration of the inclusion content which is deemed to be acceptable. For example, in the production of drawn and iron (D & I) beverage cans from tinplate it is known that there must be fewer than five 50 µm inclusions per square metre of tinplate (equivalent weight approximately 2.5 Kg) to achieve canmaking goals of fewer than 1 flange crack in every 2000 cans. (37,38) An inclusion detection device to be used in the tinplate industry must therefore have the ability to differentiate between steels which border on this level of cleanliness. Furthermore, because a content of five 50 mm inclusions per square metre is an upper limit, it would be desirable if the evaluation technique could categorize cleaner (typical) steels where a single 50 mminclusion would occur as infrequently as once per square metre. Clearly, these circumstances require a sophisticated method of testing. In contrast, a component with a larger cross section which would not be as susceptible to inclusions may find that a simpler low cost and minimum sample preparation technique such as the macro-etch test is sufficient. Beyond the consideration of clean I i ness requirements, an ideal evaluation techri ique would also have rapid evaluation capabilities, the ability to analyse large samples with little or no sample preparation; high accuracy and reproducibility, and, good mechanical reliability.

It is the purpose of this section of the paper to introduce a number of evaluation techniques which range from commonly used laboratory techniques to novel methods for analysing the steel while it is still in the molten state. To facilitate the presentation of the various methods, they will be divided into

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two distinct categories: off-line and on-line analysis. The off- line techniques are characterized by the analysis of an extracted sample in a laboratory after the sample has been cast. The main disadvantage of this approach is the labour, effort, and time needed to discover possible liquid metal processing problems which can no longer be corrected. (39) Beyond this, off-line analyses are often not complete before energy and manpower has been wastefully spent on finishing a casting of insufficient quality. In contrast, on-line testing is that analysis which takes place in real-time, on the production floor. The obvious advantage of this approach is that results are obtained immediately and can therefore be used to prevent further processing of inferior quality steel, and potentially, to regulate certain aspects of the steelmaking/steel preparation process.

To date, off-line analysis has predominated in the steel industry primarily due to the hostile production environment, the chemically aggressive nature of liquid steel, and the limited time associated with in-situ evaluations. A review of off-line techniques used for the evaluation of steel cleanliness, is provided in the following discussion.

Off-line Techniques for Inclusion Detection

As a result of the harsh steelmaking environment it is not surprising that the earlier methods of evaluating metal quality were labaratory based. Although these techniques have evolved significantly, they are still plagued by time delay and demanding labour requirements. In this section, the various methods of off-line analysis will be reviewed and critiqued with emphasis being given to accuracy and sample represent ivi ty.

Metallographic Inspection

Several optical microscopy techniques based on comparison with standard charts have been proposed over the past fifty years, the most widely accepted of these being the Fox and J.K or Jernkontoret type counts. Other methods which are based on chart comparisons include Diergarten’s inclusion count, Walker’s inclusion count, and methods proposed by

19

Timken, General Motors, the SAE, and the ASTM. These techniques have been reviewed by other authors, (40-42) with the common conclusion being that chart comparison techniques lack the accuracy and reproducibility which is often required when evaluating high quality steels.

As a result of the insufficient correlations obtained using chart methods, a number of direct counting methods based on classical metallographic techniques have been developed. These include analyses based on linear fraction, areal fraction, point fraction, number in a given area, and number on a given length of traverse. Techniques based on these concepts have been reviewed in detail by both Blank, (40) and Allmand, (41) with an analysis of the statistical error associated with each method being given by Hilliard. (43)

In a review performed by Cottingham, (42) it was established that each of the direct counting methods had a relatively high degree of reproducibility. In fact, results obtained by different observers evaluating the same sample varied by less than five percent. It is evident from results of this nature that classical metallographic techniques are sufficiently accurate and reproducible to be used in the quality steel industry.

Having established that each of the direct counting techniques are suitable for inclusion counting, a preferred method must then be chosen based on relative efficiencies where the most efficient technique is that which requires the least effort in yielding the desired accuracy. With this definition of efficiency in mind, it is fitting to introduce and discuss the use of automatic image analysis as applied to the detection and sizing of inclusions.

Image analysis can be regarded as the process whereby certain quantitative information (eg. size, shape, and number) is extracted by means of computer processing from the digitized image of a certain specimen or sample. Image analysis then, relieves the operator of tedious counting, and at the same time reduces the potential error resultng from operator variation and judgement. The ability of these instruments to rapidly assess large sections of specimens also makes it possible to perform more measurements on each sample, and thus, to

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analyse a larger number of samples. It has been stated by Shehata and Boyd that, in view of the trend toward lower volume fractions, and more numerous and complex inclusion phases, automatic image analysis will soon be the only practical tool to obtain accurate quantitative (metallographic) measurements of inclusions in steel. (44)

A discussion of quantitative metallographic techniques cannot be considered complete until the effect of sampling and the associated statistical error is evaluated. A detailed analysis of these factors is provided by DeHoff in his paper en- titled, “The Statistical Background of Quantitative Metallog- raphy”, (45) while a brief review is provided in the following discussion. When a properly prepared micrograph is taken at such magnification as to show only one constitutent with no internal boundaries, it is not certain whether only one constituent is present. or if several constitutents are present. only one of which is occupying the entire field of the micrograph. Obviously, there can be very little confidence placed in a field of this nature, as the apparent areal, and volume fraction, is 100 percent. In contrast, one would feel much more confident in a volume fraction analysis resulting from a field displaying several hundred constituents or particles. Thus, it is intuitively evident that the reliability of an analysis is associated with the number of features observed or measured. In fact, it is statis- tically accepted that thirty similarly sized particles must be observed to account for the normally distributed probability of the plane of polish having sliced the particle at its top. middle, or bottom. Thus, if only one 75 µm inclusion is viewed metallo- graphically, it may be measured as a smaller particle depending on where it intersects the plane of polish. However, after thirty 75µm particles are observed, the probability is extremely high that at least one will have been intersected at its widest point.

The implication here is that an extremely large surface area must be observed if the infrequently occurring larger inclusions are to be found. If two inclusions of a given size are deemed to be intolerable, then the metallographic surface must be sufficiently large to contain sixty of these inclusions to account for the probability of not ‘cutting’ an inclusion at its centre. Clearly, the necessary sample size severely limits the

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applicability of metallography to finding the larger, and therefore more deleterious, inclusions.

A number of points can be made in summary of the metallographic techniques mentioned in this section. Firstly, it is apparent that chart counting techniques have a low degree of reproducibility and should therefore be reviewed prior to application to determine their suitability. Also, the manual counting techniques possess a relatively high degree of reproducibility, but are labour intensive and are therefore giving way to the more rapid image analyser. Of greatest importance, it has been established that while metallographic techniques are capable of determining the average size and standard devia- tions of the inclusion populations observed, there is an extremely low probability that the infrequently occurring macro- inclusions will be detected.

Preferential Dissolution Techniques

Preferential dissolution techniques are those methods which isolate the non-metallic constituents by dissolving away the host steel matrix. When the steel is fully dissolved, the resulting residue can be suspended in a suitable solution and analysed by some existing analytical method such as Coulter- counter, (46) or sedimentation analyser. The various methods available for isolating non-metallics from steel can be cate- gorised as either potentiostatic or chemical dissolution. A review of the many variations proposed to extract inclusions has been given by Bandi, (1) and by Koch and Sundermann. (47)

In potentiostatic dissolution, the steel sample to be dissolved is suspended in a cell compartment with an electrolyte, and is rendered anodic. The electrolyte in the anodic cell is separated from the cathode by a permeable membrane such as cellophane, or a sintered glass disc. The cellophane dia- phragm has been found to be the most effective due to its low electrical resistance, low cost, ease of handling, and noble nature in the presence of the electrolytes commonly used. (48) A particular electrolyte and specific voltage are selected to dissolve the steel while the second phases are retained as a residue within the confines of the membrane. If desired, vol-

22

tages can be adjusted to dissolve less stable phases and allow only the electrochemically stable phases to be isolated. Similar selectivity can be achieved by employing different electrolytes with over one hundred electrolytes having been proposed. (48,49) With most voltagele/ectrolyte combinations it is necessary to continually control the gas mixture, temperature, pH, and other cell parameters in order to ensure the selectivity of the dissolution. A typical potentiostatic cell design is illustrated in Figure 9.

The selectivity of potentiostatic dissolution techniques often results in the loss of such electrochemically unstable phases as FeS, MnO, Fe

3C, and other nitrides and carbides. (50) There

is however, a good isolation efficiency of the oxide based inclusions. Large samples can also be treated, with Kawasaki

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Steel Corporation dissolving four kilogram samples in ten litres of 10% ferrous chloride solution at a current of 5 to 10 A. Under these conditions, 20 to 25 days are required to dissolve the sample. (51) This mode of analysis is implemented as regular quality control practice for high grade steels, and although there is a considerable time delay, its ability to analyse large sample volumes (of any initial shape) makes it one of the most effective inclusion assessment techniques available today.

The chemically based dissolution techniques fall into two broad categories: (a) treatment in dilute acids which dissolves both the steel matrix and some of the chemically less stable phases; and (b) treatment with halogen based mildly alcoholic solutions. Because the dilute acid dissolution techniques are less selective, and not as widely employed, they will not be discussed further.

The halogen based separation is the oldest of the dissolution techniques dating to 1928. (52) Since that time, several chemical combinations and equipment designs have been evaluated. The most popular solvents include bromine in methanol, iodine in methanol. and iodine-methanol-methyl acetate. The iodine-methanol-methyl acetate is typically preferred due to its accuracy and reproducibility. (53) As was the case with potentiostatic dissolution, it is necessary to continually monitor all parameters such as solvent addition rate, temperature, stirring, moisture content (even trace amounts of water will result in reduced selectivity), and rate of gaseous product removal.

The slurried residue obtained from either potentiostatic or chemical dissolution methods is available for further analysis of particle size and chemical composition of the constituents. This can be accomplished by any one of a number of existing methods to yield quantitative information as obtained from large and therefore representative sample volumes.

Magnetic Particle Detection

Magnetic particle detection or magnetic leakage field inspection has become a widely used means of locating inclusions in

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sheet steel during the past five years. While, for castings, this technique is more suited to the detection of larger defects such as cracks and porosity, it still merits consideration for its ability to locate inclusions which lie near the surface.

Inspection with the aid of magnetic particles (typically iron oxide) relies on the fact that when a sample is magnetized, inhomogeneities (inclusions, cracks, and pores) will disrupt the continuity of the magnetic field and force a flux or leakage field to escape from the surface above them. The magnetic powder which is applied (either in solution or dry) to the surface of the test piece will accumulate at the leakage points and thus provide an indication of the size and location of the defect. A schematic diagram of the magnetic detection process is given in Figure 10. This figure shows the leakage field at an inclusion site caused by an externally applied magnetic fields (1); and, the fluorescent particles clinging to the inclusion site during ultraviolet inspection (2).

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The interacting parameters which ultimately determine the sensitivity of a magnetic leakage field analysis include the size of the magnetic particles, surface roughness, and sample thickness. The practical aspects of these factors have been discussed in detail by Wester et al (38) while more theoretical studies dealing with particle types and interpretation of defect indications in castings have been published by Hagemaier (54,55) and Forster. (56,57)

The primary advantage of magnetic particle detection is that it is capable of analysing large samples while requiring very Iit- tle sample preparation. The unfortunate disadvantage is that leakage fields can often not escape from the surface when dis- continuitites are deeply seated in thick sectioned castings. For this reason, the full potential of magnetic testing is only realized for certain casting geometries.

Radiographic Inspection

The basis of operation of radiographic testing is that the thickness, density, and atomic composition of a given testpiece will determine the rate at which penetrating energy, projected onto one surface of the piece, is attenuated on its way to the other surface. The intensity of the received radiation energy therefore indicates the presence of matter of an unex- pected density-either greater or lesser-within the body of the sample. The exiting radiation is projected onto a photo- graphic film to provide a permanent record of the internal quality of the sample. The testing procedure is diagram- matically outlined in Figure. 11 which was initially presented in a review performed by Strauss. (55)

When utilizing commercially available equipment and energy sources, radiographic sensitivity is limited to the detection of flaws which are greater than two percent of the thickness of the testpiece. Thus, for a casting with a 10 mm cross-section, the smallest flaw which can be detected is 200µm. Limitations of this sort should be fully realized before radiographic testing is applied to evaluate the internal quality of a casting. Ulti- mately, it may be that a ‘critically sized’ inclusion can only be detected by sectioning the casting-which would remove radiography from the realm of nondestructive testing.

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Ultrasonic Inspection

Like radiography, ultrasonic testing is a method which allows the investigator to nondestructively examine the internal quality of a casting. In addition, ultrasonics possesses the definite advantages of being portable, and being safe to use in both the laboratory and plant environment.

The most popular mode of ultrasonic testing is known as pulse-echo testing and is characterized by the use of a single transducer for both transmitting and receiving the energy. The set-up of a simple pulse echo test is shown in Figure 12. As can be seen from this figure, the transducer (typically piezoelectric) emits a sound pressure wave which is transferred into the sample with the aid of a coupling gel. The wave travels through the body, reflects off the back wall, and returns to the

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transducer where it is received. The magnitude of both the initial pulse and received pulse are displayed on an oscilloscope to provide an indication of the internal cleanliness of the component.

Several aspects of attenuation theory can be introduced with the aid of Figure 12. When sound is passed through a body which has no internal flaws (case a), the only mechanisms of attenuation are true absorption and beam divergence. Be- cause the amount of true absorption is very small (and consistent for components of uniform geometry), it is relatively easy to account for. A similar argument can be put forth to account for beam divergence.

When an obstructing object is placed in the path of the sound it will interact with the sound by redistributing (scattering) the wave energy. With commercially available equipment and polycrystalline test samples, the onset of scattering will occur when the inclusion is approximately one tenth of the wavelength (59) (although detection sensitivities as small as one percent of the wavelength have been obtained under ideal conditions). (60,61) When scattering of this type occurs (case b), the presence of an inclusion can be assumed due to the

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decrease in magnitude of the bottom echo relative to case a. As inclusion size continues to increase, the amount of scattering increases until, when the inclusion diameter is approximately one half of the wavelength, a sufficient amount of sound energy is deflected directly back to the transducer to create a discrete signal on the oscilloscope. This is known as diffraction, and is represented by case c. Response signals of this type often allow for the actual counting of the number of large flaws in the testpiece. While more confidence can be placed in a detection similar to that of case c, it is not always possible to obtain discrete signals. At a frequency of 10 MHz in steel, inclusions which are approximately 60 µm will contribute to a reduction in the bottom echo (case b) while an inclusion must be nearly 300 µm to produce a discrete signal. Unfortunately, it is not currently possible to obtain sufficiently high frequencies to produce discrete signals from inclusions which are on the order of 20 µm to 50 µm.

Further research has been done to evaluate the potential for determining particle size distributions in steel using ultrasonic attenuation measurements. (62,63) The results of these studies confirm that, at present, although ultrasonics is satisfactory for detecting, locating, and approximate sizing of defects, it is not yet capable of determining exact sizes or size distributions. In any case, ultrasonic inspection is a valuable technique when it is used within its potential, and when the results are interpreted properly.

Visual Inspection

Visual is the term used to define those methods of testing which are conducted with the naked eye, or at low levels of magnification. These tests are most frequently employed for determination of general quality with only secondary emphasis being placed on the detection of macro-inclusions. These tests are outlined in the ASTM Standards and include the macroetch test, the fracture test, and the step down method. (64)

The macroetch test consists of etching a large non-polished steel sample and subsequently examining the sample visually or at low (x 10) magnification. The sensitivity of this method

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corresponds to a minimum detectable inclusion size of approximately 500 µm. Inclusion monitoring is typically of secondary importance as macroetching is primarily used to monitor grain size and structure, segregation, and gross defects such as seams, laps, and shrinkage.

The fracture test, as the name implies, consists of preparing a sample with minimal surface irregularities (which might promote fracture) and fracturing the testpiece. The fracture surface is then examined at up to x 10 magnification for the presence of fracture inducing inclusions. The advantage of this technique is that a relatively large volume of steel can be manipulated to reveal the largest inclusion. The detection sensitivity of the fracture test is approximately 400 µm.

The step down method is used to determine the presence of macro-inclusions on a machined surface either visually with ‘good illumination’ or at low magnification. By preparing a surface, observing it, and then machining to a new level, the step down method is capable of determining areas of preferential defect accumulation, and the onset of shrinkage. As a result of the sample preparation and inspection techniques, this method is limited to a minimum detectable defect size of approximately 1500 µm.

The previous review of off-line methods of detecting inclusions in steel has shown that a number of techniques are available, each with its own set of advantages and disadvantages. It is the task of the foundryman to choose a technique which will best suit the needs of his operation, and more importantly, those of his product and customer.

In an attempt to overcome some of the inherent disadvantages of off-line testing, two novel techniques have been proposed to evaluate the steel while it is in the liquid state. The advantages of these techniques, along with a description of the operation of each will be reviewed in the following section.

On-line Techniques for Inclusion Detection

The catch phrase ‘on-line’ is the currently accepted industrial term referring to in-situ, real-time analyses. In this context, a

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truly on-line method of detecting or monitoring steel cleanliness must be implemented at some stage of the steel preparation process, and have the potential to yield results rapidly. A testing procedure which satisfies this definition would potentially provide information regarding possible liquid metal processing problems (feedback control) and also prevent the cost- inefficient casting and finishing operations which may be performed on metal of unsatisfactory quality (feedforward control).

Although on-line techniques were realized to be an integral part of the future of the steel industry in the late 1970’s, there are no commercially operable detectors to date. In contrast, on-line particle size anlaysis has been a reality in the mineral processing and chemical precipitation industries since the early 1970’s. The parallel requirements between these industries and the steel industry (that of detecting solid particle sus- pensions in a liquid medium) presents the steelmaker with the ability to select and adapt compatible existing methods of particle size analysis rather than postulate and develop new techniques. When the techniques available to the mineral proc- essor are evaluated in light of the hostility of liquid steel and the inherent differences between liquid steel and water, it becomes evident that only two principles have the potential for extension into the steel industry. The two satisfactory systems are the electrical resistance change (Coulter counter) technique and the ultrasonic attenuation technique.

Electrical Resistance Technique The adaptation of the Coulter counter approach to inclusion assessment in liquid metals has been widely reported by Guthrie et al since 1984. (39,65-59) This ‘electric sensing zone’ (E.S.Z.) technique (or LIMCA: Liquid Metal Cleanliness Ana- lyzer) was first applied to liquid aluminum after an initial modelling study on liquid gallium. The LIMCA apparatus has also been applied to zinc and lead, and is currently being adapted to cast iron, steel, and magnesium melts. Extensive developmental work has also been conducted on water models.

The operating principle of the resistive pulse/electric sensing zone technique is shown in Figure 13. The passage of a sus-

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pended particle through an electronically insulated orifice can be noted due to an increase in resistance beyond the normal value of the liquid metal. It has been established that the measured resistance change increases in direct proportion to the volume of the particle. This resistance change, in the presence of a constant applied current, will result in a voltage pulse-the duration of which is approximately equal to the transit time of the particle through the orifice. Thus, by counting the number of pulses, and measuring the magnitude of each pulse, the concentration and particle size distribution of a molten metal system can be obtained.

The measurement system for steel melts consists of a quartz or boron nitride probe, an electric circuit, a pre-amplifier, a peak detector, a pulse height analyser, and a recording system. Measurements are taken by drawing steel though the electric sensing orifice with the aid of a vacuum. The analysed metal is contained within the probe, and can be expelled by

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reversing the vacuum. Expelling the analysed steel allows the trial to continue without changing the probe.

Physically, the detection equipment which is to be immersed into the liquid steel consists of an electrode and a sampling probe. The electrode serves both to carry a heavy DC current into the melt, and to transmit the voltage pulse generated by the passage of inclusions through the sensing zone, into the preamplifier. Tungsten and molybdenum were initially tested as electrodes; however, graphite proved to be a sufficiently stable material, lasting indefinitely in transformer steel melts and for up to twenty minutes in carbon steels. With regard to the sampling tube, quartz was found to be appropriate at 1400°C (cast irons), while boron nitride was required for melts at temperatures of 1600 °C. (67)

The particle detection limit of the LIMCA is a function of both the orifice diameter and the applied DC current. Typically, the lower limit of detectability is considered to be one-twentieth of the orifice diameter. Thus to detect 50 µm inclusions, the orifice diameter could be as large as 1000 µm.However, this size orifice could lead to a reduced sensitivity as is suggested by Figure 14. In this figure, it can be seen from the slope of each curve increases as inclusion diameter increases. For example, at an orifice of 300 µm, the received voltage pulse does not change significantly over a range of inclusion diameters spanning from 15 µm to 25 µm.Mental extrapolation of the trends shown in this plot suggest that at an orifice diameter of 1000 µm,the sensitivity may not be sufficient to distinguish between inclusions in the 50 µm range.

The orifice diameter is also intimately related to sample size- as orifice diameter increases, sampling rate increases. At an orifice diameter of 250 µm,the sampling rate is approximately 7 g every ten seconds, (68) while at an orifice diameter of 1000 µm the sampling rate would be nearly 50 g every ten seconds. Thus, to analyse 2 Kg of molten steel with a 1000 µm orifice, over six and a half minutes of sampling would be required. Ulti- mately, it may be that the orifice diameter will have to be chosen as a compromise between sample size (large orifice) and detection sensitivity of resolution (smaller orifice).

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Ultrasonic Inspection

Although the ultrasonic evaluation of molten metal cleanliness was first reported in 1951 by Mountford, (70) the art is still in relative infancy. Little work has been done over the past thirty-five years with the exception of recently abandoned studies by Alcan, (39) and Mansfield of Reynolds Aluminum. (71-73) The progression beyond this research in aluminum and into liquid steel was considered in a technical proposal prepared jointly by Reynolds in conjunction with Adaptronics, (74) while preliminary results on passing sound through liquid steel have been reported by Mountford et al. (75) More recently, research carried out at the University of Toronto (76) has shown that it is possible to distinguish between steels with varying degrees of cleanliness, and further, that the presence of inclusions as small as 2µm have been noted in liquid steels.

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The common factor in each of the early studies on the ultrasonic evaluation of molten metals (70-74) was that the ultrasonic theory of solids was applied directly to the liquid state. This suggested that scattering was the only attenuating mechanism which was operative, and thus, that there was a minimum detectable particle size equivalent to approximately one tenth of the wavelength (see section headed Ultrasonic In- spection). As a result of this assumed limitation, the industrial attitude toward on-line ultrasonic evaluation was one of pessimism.

A review of the ultrasonic theory of liquids reveals that attenuation does not end with the cessation of scattering. Instead, there is a further attenuating mechanism which arises as a result of the friction and viscosity present in a liquid mechanism. Thus, attentuation in liquids is comprised of viscous losses, scattering effects, and true absorption, while only scattering and absorption exist in solids. The attenuation regimes which exist in liquids are best presented in graphical form as is done in Figure 15. This plot shows the characteristic attenuation of a single particle in a suspension of uniformly sized particles (monodisperse media) in liquid steel.

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It can be seen from Figure 15, that at 1 MHz in liquid steel (wavelength, X = 3000 µm), the onset of scattering occurs at approximately 300 µm(one tenth of the wavelength). Further- more, the diffraction loss regime (where particles deflect a sufficient amount of sound energy in a direction opposite to that of the incident wave to result in a discrete peak on the oscilloscope) begins at approximately one-half of the wavelength.

Beyond the scattering losses are the viscous losses which are unique to liquids. These frictional losses arise because the particles which are suspended in the sound field partake of the motion of the fluid to an extent determined by their mass and surface area, the frequency of sound, and the viscosity of the fluid. Because the motion of the particles lags behind that of the fluid, heat is generated at the surface of the particle which ultimately decreases the energy of the sound field. The characteristic ‘hump’ of the viscous loss curve is explained in the following manner. Extremely small particles (approximately 0.1 µm for the conditions of Figure 13) tend to move to and fro along with the liquid in the sound field resulting in little attenuation. As the particles become larger they lag more and more behind the movement of the fluid causing attenuation to increase. However, as the size increases, the total surface area of the particles decreases (for a constant volume concentration) and therefore the attenuation decreases. These opposing factors result in an absorption maximum.

In consideration of the viscous loss effect, it is apparent that the term ‘minimum detectable particle size’ has no significance in liquid metal analysis. In fact, theory predicts that a 1 µm particle will cause more attenuation than a 100 µm particle at 1 MHz in liquid steel. Based on concepts of this nature, further research was conducted at the University of Toronto to develop a probe for the evaluation of molten steel cleanliness.

The ultrasonic probe consists of a piezoelectric transducer, a one inch-diameter steel guide rod, and a reflector plate. Opera- tionally, the piezoelectric crystal generates and receives the sound which travels to and from the melt via the steel guide rod. Within the liquid steel, the sound reflects off the reflector plate, and finally, re-enters the steel rod. A schematic

36

representation of the hand-held probe is shown in Figure 16. One of the primary advantages of the ultrasonic probe technique is its ability to rapidly analyse large volumes of liquid steel. The volume of the sample chamber for initial tests (76) was set at 32 cm3 which is equivalent to 0.23 Kg of liquid steel. When the flow rate of steel through the sampling chamber (along with the ability of the probe to record 1000 measurements per second) is considered, it becomes apparent that sample sizes far in excess of other methods are obtained.

The pulse-echo probe was used in both the laboratory (25 Kg induction furnace) and steel plant (tundish) environments and produced reproduceable results in each case. Attenuation measurements are taken by immersing the head of the probe in the bath, allowing the end of the probe to wet, and recording the magnitude of the echo which is received from the reflector plate. The total time required for a test is approximately thirty seconds.

Typical results have shown that it is possible to distinguish between construction steels and merchant bar quality steels, and also, to determine the relative cleanliness of two con-

37

secutive heats of similar grade steel. The cleanliness rankings as determined by the probe were verified in each case by extensive metallographic analyses. In a final laboratory trial, it was clearly shown that alumina deoxidation products with an average size of 2.5 µm (as determined by metallographic inspection) could be detected with the probe. The current state of the research consists of attempting to relate a series of attenuation measurements to mechanical tests which are conducted on samples extracted from the melts. It is hoped that successful completion of this phase of the research will establish the potential of the probe for predicting the suitability of a steel while it is still in the liquid state.

The ultimate difficulty in relating the measurements recorded by the ultrasonic probe to the cast products may not lie with the probe itself, but rather with the concept of liquid metal analysis. While researchers can strive to develop a device which will accurately evaluate the liquid metal, they must realize that the metal quality will change between the point of analysis and solidfication. Factors such as inclusion coagula- tion and flotation, adherence to vessel walls, slag entrainment, reoxidation, and refractory wear will result inevitably in changes in metal quality. For this reason, on-line analysis is best thought of as an indication of metal quality, and not as a fully quantitative evaluation of the final cast product.

The previous review of off-line and on-line techniques for the evaluation of steel cleanliness has shown that several methods of anlaysis are available to the steelmaker. The chal- lenge, therefore, is not necessarily to develop new techniques, but to select the existing method which is best suited to the requirement of the finished product. In the final analysis, it must be realized that detecting inclusions is not a solution to the cleanliness problem; it is merely a recognition of the severity of the problem. The ultimate solution to improving metal quality is to control inclusion behaviour throughout the steelmaking and casting process. This aspect of steel casting is addressed in the final section of this paper.

THE CONTROL OF INCLUSIONS

The final inclusion content of any steel product, either wrought or cast, can be controlled by adopting

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(a) measures to minimize inclusion formation, and

(b) measures to maximize inclusion removal.

These will therefore be discussed in this concluding section of the paper.

Methods for Minimizing Inclusion Formation

With good deoxidation practice, the products of primary deoxidation should almost all be removed from the steel in the ladle, and generally those which are left, together with the products of secondary deoxidation, are usually so small that they do not cause defects in steel castings. Thus, measures to minimize inclusion formation are mainly concerned with reducing the extent of reoxidation experienced by the steel during the casting process. As discussed earlier, reoxidation products are formed closer in time to casting and solidification, so that the time available for their removal is greatly reduced, and they are generally much larger. For these reasons, they often result in serious defects in steel castings, and as illustrated in Figure 17, they accounted for over 80% of the macro-inclusions in carbon and low alloy steel castings, observed during a recent major project on the topic. (77)

Measures to minimize reoxidation can be classified as follows:

(a) stream compaction

(b) stream protection

(c) flow control during horizontal steel transfer.

Stream Compaction

Stream compaction refers to lowering the turbulence energy level in the stream, so that it is smoother, more laminar and less likely to flare and disintegrate. This reduces reoxidation in two ways. Firstly, since the surface area exposed to contact with air is reduced, less reoxidation occurs on the stream as it

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falls through the air. However, as explained earlier, this is a minor source of reoxidation, so that the second effect is much more important. This is to reduce the amount of air entrained by the falling stream, and hence the extent of reoxidation occurring in the impact zone of the receiving vessel.

Ladles with stopper rods invariably yield more compact, less turbulent streams than those fitted with slide gates, because

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the throttling of the latter, which is necessary during the major part of the casting process, introduces lateral velocity components into the stream, (78) so that when it is released from the constrictions imposed by the nozzle, it tends to flare and disintegrate. Another benefit of the stopper rod is that it helps to reduce the tendency to vortex formation and slag entrainment once the metal in the ladle has reached a low level.

Since stream turbulence increases with the height of fall, an obvious and simple way to improve stream quality is to reduce the height of fall, so that wherever possible bottom-pouring ladles should be used for the production of clean steel castings, rather than lip poured or teapot ladles. It may well prove worthwhile, also, to investigate ways in which the geometry of the casting arrangement could be adjusted so as to reduce the pouring distance.

Nozzle design also has a very important role in the attainment of compact, high quality streams. This is true both in the entry region and the collector region. The influence of swirl in the entry region has been demonstrated by the use of castellated nozzles (79,80) and also by Engh et al. (81) Where slide gates are used, the stream characteristics are determined by the throttling effect mentioned earlier rather than by condtions at the nozzle entry, and the use of extended collector nozzles has been found to be a very effective way of combatting this problem. (78,82) For optimum results, a collector nozzle length of six times the nozzle diameter was found to be necessary. (78) The shape of the collector nozzle is also important, as evidenced by the improvements obtained by the use of a trian- gular collector nozzle. (82) The cruciform nozzle, (83) which is somewhat similar to the castellated nozzle, but has a smooth entry section and a fluted lower section has also demonstrated impressive benefits. Thus, it would appear that the use of extended collector nozzles, with specialized cross-sectional designs, fitted to stopper rod ladles can confer significant improvements in pouring stream quality, helping to reduce further the incidence of defects due to reoxidation.

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Stream Protection

In the production of clean steel, stream protection is required in addition to stream compaction. This is achieved by the inter- polation of some shield between the molten steel and the atmosphere, where this shield can be either gaseous, as in inert gas shrouding, or solid, as when a ceramic shroud is employed. Neither approach is completely satisfactory on its own, and for the highest steel cleanliness ratings, a combination of both is required, in which inert gas is introduced into the ceramic shroud. Where shielding is physical as well as gaseous, compaction of the stream is less critical, so that the double protection is commmonly used in continuous casting in conjunction with slide-gate ladles.

Various arrangements for shrouding the ladle to tundish and tundish to mold streams are illustrated in Figure 18. (84) The simplest is the introduction of inert gas through a gas lens or ring located immediately below the ladle, so as to produce a curtain which is carried down by the stream into the tundish thus protecting the liquid surface of the tundish as well as the stream itself. While this method of protection has the advantage of not obscuring the view of the stream or interfering with the operation of the caster as in the case of a ceramic shroud, there is a disadvantage in that considerable draughts and winds can be present on the casting floor, so that the protective curtain around the stream can be partially removed. How- ever, provided the draughts are not too severe, reasonable protection can be obtained by this means, as illustrated in Figure 19. (78)

In general, gas shrouds can be used in many more applications than ceramic shrouds because they are less affected by geometric constrictions. One process which could perhaps be modified to be applicable to the production of clean steel castings is the IMPACT (Improved Metal Protective Atmo- sphere Casting Technique). (85) In this technique a collapsible bellows made from a coil spring and asbestos cloth is employed between the ladle and a covered tundish. The bellows or shroud completely surrounds the casting stream and argon is fed into this shroud under a slight positive pressure.

42

Argon is also blown into the volume of the tundish between the liquid surface and the cover. A similar arrangement is used to protect the tundish to mold stream during continuous casting.

43

Another possibility is the Pollard shroud, (86) where a hollow metal cylinder is placed around the stream, and argon or nitrogen is introduced into this shroud approximately half way down. Since the shroud does not touch either the bottom of the ladle or the top of the receiving vessel (i.e. tundish cover, ingot mold or casting mold), this gas is introduced at a sufficiently high flowrate (approximately 30 ft.3 min-1), so that not only is gas carried into the receiving vessel by boundary layer attachment, but also gas should flow out of the top of the shroud to avoid air being drawn in by the Venturi effect. The effectiveness of this latter feature is very questionable, since the Venturi effect is so strong, particularly at high casting speeds.

In the case of gaseous shrouding, the mechanisms of entrainment outlined previously indicate that perhaps more attention should be paid to the impact area on the pool and less to the point of emergence from the previous vessel, either ladle or tundish. Arrangements of this type also have the advantage of providing greater protection from oxidation of the turbulent upper surface of the receiving pool, which can cause a significant portion of the total oxidation. The obvious drawback to having the gas manifold immediately above the receiving pool is its susceptibility to damage by flaring streams and here again efficient stream compaction will play an important part.

The effectiveness of such an arrangement has also been demonstrated in practice. (80,87) In each of these independent studies, a gas shroud or manifold was placed on top of the mold and gas was injected to flood the mold cavity. This particular approach is advantageous in that it allows clear observation of, and easy access to, the stream, nozzle and the liquid pool surface.

It is also clear from a number of studies that entrainment of a protective gas can have a beneficial rinsing effect on the steel just prior to solidification. The rising swarm of gas bubbles produces a strong upward motion of the metal just around the jet impact point, which greatly increases the chance of inclusion separation. A similar effect can be produced by bubbling argon down the stopper rod, with the additional advantage of

44

reducing the frequency of nozzle blockage. This may be due to the fact that the presence of the gas bubbles results in a less streamlined flow with a more “blunt” velocity profile and a higher proportion of the higher velocity components near the wall. By virtue of their greater momentum, these have a more powerful sweeping effect, so minimizing adherence and build- up of solid particles on the wall and keeping the tube and nozzle clear.

The fact that the impact area is also of paramount importance in the entrainment of flux has been evidenced by the vast improvements resulting from the use of a teeming spout. (27) This again is consistent with the mechanisms of entrainment outlined earlier and since inward flow on the surface is the impact zone is mainly associated with streams of lower turbulence, this may provide an argument against concentrating all the effort on the minimization of stream turbulence.

Ceramic shrouds alone are not entirely satisfactory either, primarily because of the Venturi effect mentioned earlier. The rapid passage of the stream through the shroud produces a partial vacuum, perhaps as low as 0.1 atmos. depending on casting speed, which in turn causes aspiration of air through any imperfectly sealed joints, so that for the highest quality steel inert gas shrouding of all such joints is necessary to ensure that it is inert gas rather than air which is aspirated.

In general, wherever it is possible, both ceramic and gaseous shrouding should be employed. Where, because of geometric or other constrictions, gaseous shrouding is used alone, highly effective stream compaction is an essential pre- requisite, since it is impossible to protect a flaring and disintegrating stream effectively.

Fluid Flow Control During Metal Transfer

Where some form of runner system is involved between the receiving vessel and the mold in which soldification occurs, care must be taken to minimize exposure to air, and more particularly entrainment of air during this transfer step. In the case of continuous casting both the tundish and mold are covered by synthetic fluxes designed to protect the metal from

45

oxidation. In the production of steel castings, this approach may not be appropriate, so that some other form of protection, perhaps by introduction of inert gas above the surface, may be necessary. The geometry of the runner and gating system should also be designed to minimize turbulence and hence air entrainment by avoiding rapid changes of direction. A totally enclosed casting system in which the steel could be moved by the application of either increased or decreased pressure, as well as by gravity, would obviously be ideal but would also be expensive and probably only justified for the highest quality castings.

Calcium Treatment in the Ladle

Calcium treatment in the ladle is normally employed to modify both oxide and sulfide inclusions, accelerating coallescence and hence removal of the former and conferring shape control of the latter. As noted previously, this treatment also lowers the dissolved oxygen content in the steel. Also, because the calcium aluminate inclusions are liquid, casting of the steel without nozzle blockage is greatly facilitated. However, there is evidence that still another benefit is that it reduces the susceptibility of the steel to reoxidation during transfer between the ladle and the mold. Cramb and Byrne (88) have shown very clear differences in the tundish slag composition caused by calcium treatment in the ladle. In the absence of this treatment, reoxidation of the steel, caused mainly by rebounding of the impinging stream from the tundish bottom to produce an ‘open eye’ of steel exposed to the atmosphere, resulted in the tundish flux containing as high as 20% (FeO + MnO). When calcium treatment was employed in the ladle, this figure was reduced to <5%. The exact mechanism responsible for this marked reduction in the extent of reoxidation is not entirely clear, but it may be due to calcium vaporizing from the steel reacting with incoming oxygen and acting as a ‘getter’ to protect the steel from oxidation.

Mechanisms for Inclusion Removal

There are various mechanisms for inclusion removal which can be controlled and enhanced by the way in which the liquid

46

steel is handled between the ladle and the mold. Obviously these mechanisms are not entirely independent of one another, but for the sake of clarity they will be discussed in- dependently in the succeeding sections, while the way in which they interact will be indicated briefly.

FIotation

Since inclusions are appreciably lighter than steel, they will tend to separate under the action of gravity by the process of simple flotation. In a completely quiescient bath, the rate of this flotation is given by Stokes’ Law:

The only controllable variable is particle size, and since it appears squared in equation (1), it is of critical importance. This can be highlighted by calculation of the flotation rates of dif- ferently sized particles. For example, for a particle of 100 µmradius and density 4.0 g. cm-3, V

T = 1.0 cm s-1 and so the time

required to float one foot is 30 s. For a particle of 10 µ m radius and the same density, V

T = 0.01 cm. s-1 and so the time re-

quired to float one foot is 50 min. For lighter inclusions of density 3.0 g. cm-3, the times required to float one foot are reduced to 22 s for the 100 µm radius particle, and 37 min. for the 10µm radius particle. Clearly, then, particle size is much more important than particle density, and this is one of the main reasons that the formation of liquid inclusions can be so beneficial, since these can more readily coallesce to form larger globules. The characteristic particle sizes obtained during tapping and holding in the ladle when using various deoxidizing reagents are shown in Figure 20. The other point that emerges very

47

clearly even from such simple calculations is that the removal of very small inclusions by flotation requires unacceptably long holding times in the ladle, and so more effective means must be employed.

It should be stressed that Stokes’ Law only applies to a quiescent bath, and that in the presence of any appreciable level of turbulence in the metal bath, flotation under gravity is over- whelmed by currents in the steel, and the inclusions are then swept along by the moving metal. As a result of this, the inclusions only reach interfaces where they may be removed from the metal phase if the currents in the bath take them there.

Stirring

One of the simplest ways of enhancing inclusion removal is by stirring, (13,14,16) the energy for which can be supplied either inductively or by injection of an inert gas. Induction stirring in a ladle requires that the ladle itself be austenitic or at least have an austenitic panel, so that gaseous stirring is more common. The gas concerned is commonly argon, but may be nitrogen or carbon dioxide, and the use of a porous plug, either in

48

the base of the ladle or at the end of a top lance, provides smaller gas bubbles and a superior stirring pattern to that produced by a lance with a single orifice. Stirring appears to enhance inclusion removal partly by providing motion which carries the inclusions to an interface where they may escape from the steel and partly by causing agitation which en- courages collision and coallescence of inclusions to form larger particles, whose removal is favoured. In the case of gaseous stirring a third mechanism may be the attachment of inclusions to rising gas bubbles, so that they are essentially levitated out of the steel bath.

One method of introducing gas stirring which has been mentioned already is the passage of argon down the stopper rod. (78) This is practised by some plants in the tundish during continuous casting, and may be applicable to the production of castings by using it in a ladle where the casting rate is controlled by use of a stopper rod. Some of this gas may be carried down and facilitate inclusion removal in the runner and gating system, while there is also some evidence that it will help to delay the onset of vortexing once low metal levels have been reached in the ladle. Hence it may well help to reduce the extent of slag entrainment from the ladle, which can be responsible for major defects in castings.

It should clearly be understood, however, that there is an upper limit to the amount of stirring which can be applied in the ladle. The permissible injection rate varies with the method of introducing the gas and is higher when using porous plugs than when using a single orifice lance, since the former produces a swarm of relatively small bubbles, while the latter yields very large bubbles which ascend in close proximity to the lance. Increase of the gas flowrate beyond this optimum value will lead to excessively rapid metal movement at the slag-metal interface, resulting in entrainment of the top slag into the metal bath and the pushing back of the slag layer from the area in which the gas emerges to give an ‘open eye’ of steel exposed to the atmosphere. The single orifice lance is particularly prone to cause this problem, and since serious reoxidation can occur in this area, this is clearly counter- productive, and the stirring rate should be adjusted to avoid this situation.

49

Interfacial Effects

Nogi and Ogino (89) have clarified that the interfacial energy of an inclusion particle has an effect on the rate at which it floats out of liquid steel, and have expressed their results as a modification of Stokes’ Law:

where VT, g, r, P and η have the same signficance as stated previously, while β is a frictional coefficient. For non-wetting particles, β = 0, and for wetting particles, β = α The ratio of the rising velocities V

β=0/V

β=α = 1.5, indicating that a non-wettable

inclusion will rise 50% faster than a wettable inclusion of the same size and density.

A recent study by Wilshynsky, (90) using hollow glass microspheres in a water-model study to simulate the behaviourof inclusions in steel, has confirmed that interfacial effects have a very strong influence on the rate at which inclusions separate from steel. As received, these microspheres are hydrophilic or wettable; they are rendered hydrophobic or non-wettable by coating with an extremely thick layer of vinyl silane. In this latter condition, their separation is accelerated considerably.

However, it seems clear that the way in which interfacial tensions exert the major part of their influence is through their effect on the rate at which inclusions leave the steel at an interface. Figure 21 is taken from the classic work of Plockinger and Wahlster (91) on the rate of removal of various reoxidation products from steel during tapping, holding in the ladle and teeming. Comparison of the rates in Figure 21 with the particle sizes for these same deoxidation products shown in Figure 20, clearly indicates that particle size is not the major factor determining these rates of removal. This is partly due to the effect of the turbulence introduced during the tapping period, which, as discussed earlier, overrides flotation in accordance with Stokes’ Law. However, as Plockinger and Wahlster (91) have established, it is also due to the effect of the different inter-

50

facial tensions or energies between these various deoxidation products and the liquid steel.

Plockinger’s treatment applies primarily to the absorption of inclusions into a layer of slag or flux, but qualitatively the same effect appears to operate with respect to adherence of inclusions to solid refractory surfaces. When an inclusion approaches an interface, the main force favouring its removal from the steel is the dissipation of the interfacial energy between the inclusion and the liquid steel, so that the higher this value is the more rapidly the transfer across this interface takes place, and the more likely it is that the force favouring absorption into the second phase will be greater than that tending to retain the inclusion in the liquid steel. Inclusions with low interfacial tensions or small contact angles with steel are wetted by it, while those with high interfacial tensions or large contact angles are not, and the comparison between the behaviour of wettable and non-wettable particles at an interface is illustrated schematically in Figure 22. (92) This difference in behaviour also exerts a strong effect on the rate at which the particles coallesce to form clusters, as confirmed experimentally by Torsell and Olette. (93)

51

fluid Flow Control During Metal Transfer

Many water model and mathematical model studies of the continuous casting tundish (33,80,82,94-101) have confirmed the very strong influence of the flow pattern on the residence time in the tundish, and hence on the time available for inclusion separation. As indicated in the section on flotation, the flow pattern also has a strong effect on the opportunity for these inclusions to contact another phase, either liquid slag or solid refractory, and be removed by absorption or attachment. It has been demonstrated clearly that flow control devices such as dams and weirs or ported baffles can markedly increase inclusion removal by their effect on both residence time and the proportion of metal contacting an interface. Recently, the trend in tundish design has moved away from dams and weirs and towards ported baffles as the major flow control device. Three configurations of ported baffles in a tundish are shown in Figure 23, (101) while the results of these in terms of steel cleanliness are plotted in Figure 24. It is clearly advantageous, not only to direct the metal flow and hence inclusions to the upper surface but also to increase residence time further by directing the metal flow across the tundish. (102)

52

It would seem that a similar approach, perhaps using similar flow control devices incorporated into the runner system could be very beneficial in increasing the separation of inclusions during transfer of steel from the ladle to the casting mold. The simplest and least expensive means of establishing the optimum design for any given casting arrangement is almost certainly water modelling.

Absorption of lnclusions into a Flux Once flow control devices have imposed a flow pattern which carries inclusions to the upper surface of the metal, the next step is the provision of a flux which will absorb these inclusions once contact is made with them at the interface. There are two important aspects to the ability of a flux to absorb inclusions, and both are related to its composition. The first relates to the amount of the inclusion material which the flux can absorb before the solubility limit for this component is reached and the slag becomes saturated with it, while the second relates to the rate at which the actual transfer of the component across the slag-metal interface occurs.

With regard to the first aspect, the flux provided for inclusion absorption should ideally have a low content of the consti-

53

tuent to be absorbed, and if possible, should be able to absorb considerable quantities of it without a large increase in its viscosity or melting point. For example, if the inclusions concerned are alumina, a flux of composition 50% CaO, 50% SiO

2

would be able to absorb about three times as much as a flux of composition 50% CaO, 50% AI

2O

3. However, it should also be

remembered that the calcium silicate flux would not be in equilibrium with nearly such a low dissolved oxygen content in the metal as that in equilibrium with the calcium aluminate flux. Thus, provided that the proportion of inclusions to be absorbed is not too great, the calcium aluminate flux may well be preferable for the production of clean, high-quality castings.

With respect to the rate of absorption of inclusions, and again assuming that the inclusions concerned are alumina, a study

54

on continuous casting mold powders (103) has shown that this rate increases with increase in the lime to silica ratio from one to two, and decreases with increase of the alumina content of the flux. In general, the rate will be strongly influenced by the fluidity of the flux, so that in general dissolution of inclusions will be accelerated at the addition of fluidizers such as calcium fluoride which lower the melting point and viscosity of the flux. Unfortunately, such additions make the flux considerably more aggressive towards refractories, and may be unacceptable for this reason. In order to control this effect, the flux should contain at least some of the oxide of the refractory concerned, and preferably fairly close to the solubility limit for this component, if this can be combined with acceptable values of melting point and viscosity.

If it is desired to accomplish desulfurization as well as deoxidation, the flux should have a high sulfide capacity and here again fluxes based on calcium aluminate will be markedly preferable to those based on calcium silicate. If desulfurization is not desired, a calcium silicate-based flux will be more appropriate.

For the best results with regard to both oxygen and sulfur, the flux should be a synthetic mixture, of a composition delib- erately chosen to optimize removal of these elements from the metal. It should be free of reducible oxides such as iron or manganese oxides, so that if there is appreciable carry-over of slag containing these oxides, this slag will have to be reduced, possibly by additions of aluminum, before effective removal of oxygen and sulfur can be achieved. For the attainment of very low values of oxygen and sulfur the flux should also be free of silica, or contain only a very small amount.

Collection of Inclusions on Refractories

Experimental observations (104,105) indicated that inclusions in stirred melts can become attached to the surface of the con- tainer. The practical experience of nozzle blockage showed that the same effect could occur even from rapidly moving streams, and the work at Union Carbide (106,36) showed con- clusively that for this attachment to occur, the inclusions must

55

be at least partially solid. This effect can be turned to advantage and used as a means of inclusion removal, either by the use of filters or multi-ported baffles.

There is now a considerable literature on the use of filters as an additional means of producing cleaner steel, but essentially the problem of clogging or blockage of filters does not seem to have been overcome on a consistent basis for steels in the way that has been achieved for aluminum. This is largely related to the fact that aluminum has an appreciably lower inclusion content than is normally achieved in steels, so that perhaps multi-ported baffles may prove to be a more viable alternative for the latter, since the ports are many times larger than the orifices in filters. Such baffles have been referred to as inclusion collector boards, (107) and inclusions certainly do adhere to them as they do to walls, dams and weirs. However, if the ports are evenly distributed across the area of the baffle, with respect to both width and height, then the baffle no longer acts effectively as a flow control device. Various designs which would incorporate both functions are obviously possible.

Two points are worth mentioning in conclusion. Firstly, since the inclusions have to be at least partially solid, calcium treatment would not normally be employed if it is desired to achieve any appreciable inclusion removal by attachment to refractory surfaces. Secondly, the refractories used for such baffles, as for walls, dams and weirs must be of very high quality, so that erosion of the extra refractory surface exposed to the flowing steel does not introduce more inclusions than are removed by attachment or because of the superior flow pattern obtained.

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Hutchinson, Proc. NOH-BOS Conference, ISS-AIME, Vol. 59, 1976, p. 177. 81. T.A. Engh and K. Larsen, Scand. J. Met., Vol. 8, 1979, p. 161. 82. D.J. Harris and J.D. Young, Proc. Steelmaking Conference, ISS-AIME, Vol.

65, 1982, p. 3. 83. J.M. Svoboda, R.W. Monroe and G.J. Vingas, Proc. Electric Furnace Con-

ference, I.S.S.-A.I.M.E., Vol. 44, 1986, p. 251. 84. E. Elsner, H. Knapp, D. Ameling and G. Rudolph, Proc. International Con-

ference on Continuous Casting of Steel, The Metals Society, London, 1976, p. 243.

85. M.P. Kenney, Proc. Electric Furnace Conference, AIME, Vol. 25, 1967, p. 45.

86. N.L. Samways, B.R. Pollard and D.J. Fedenko, Proc. NOH-BOS Con- ference, AIME, Vol. 57, 1974, p. 71.

87. E.M. Calanog, R.S. Mulhauser and J.L. Wareham, Proc. NOH-BOS Con- ference, AIME, Vol. 58, 1975, p. 260.

88. A.W. Cramb and M. Byrne, Proc. Steelmaking Conference, ISS-AIME, Vol. 69, 1986, p. 719. (Also in lron and Steelmaker, Vol. 13, (5), p. 27).

89. K. Nogi and K. Ogino, Can. Met. Quart., Vol. 22, (l), 1983, p. 19. 90. D.O. Wilshynsky, M.A.Sc. Thesis, University of Toronto, 1984. 91. E. Plockinger and M. Wahlster, Stahl und Eisen, Vol. 80, 1960, p. 659. 92. 0. Repetylo, M. Olette and P. Kozakevitch, J. Metals, Vol. 19, (5), 1967, p.

93. K. Torsell and M. Olette, Rev. Metall. (Paris), Vol. 66, 1969, p. 813. 94. K.L. Kemeny, D.J. Harris, A. McLean, T.R. Meadowcroft and J.D. Young,

Proc. 2nd Process Technology Conference, ISS-AIME, 1981, p. 232. 95. G.L. Dressel, D.R. Shrader and D.A. Kukelow, Proc. Steelmaking Con-

ference, ISS-AIME, Vol. 66, 1983, p. 205. 96. S. Tanaka, M. Lye, M. Salcudean and R.I.L. Guthrie, Proc. International

Symposium on the Continuous Casting of Steel Billets, Met. Soc. C.I.M., 1985, p. 142.

97. R. Ahuja and Y. Sahai, Ibid, p. 73. 98. Y. Sahai and R. Ahuja, Proc. Steelmaking Conference, Vol. 69, 1986, p.

677.

Steelmaking Conference, Vol. 69, 1986, p. 699.

45.

59

99. Y. He and Y. Sahai, Ibid, p. 745. 100. J. Szekely and N. El-Kaddah, /bid, p. 761. 101. Y. Yoshii, Y. Habu, T. Nozaki, S. Itoyama, H. Nishikawa and T. Imai, Porc.

International Conference on Technology and Applications of High- Strength, Low-Alloy Steels, ASM, 1983, p. 377.

102. J. Tsubokura, I.D. Sommerville and A. McLean, Iron and Steelmaker, Vol. 12, (6), 1985, p. 43.

103. T. Emi, H. Nakato, Y. Ilda, K. Emoto, R. Tachibani, T. lmai and H. Bada, Proc. Steelmaking Conference, ISS-AIME, Vol. 61, 1978, p. 350.

104. N. Lindskog and H. Sandberg, Scand. J. Metall., Vol. 2, 1973, p. 71. 105. T.A. Engh and N. Lindskog, /bid, Vol. 4, 1975, p. 49. 106. J.W. Farrell and D.C. Hilty, Proc. Electric Furnace Conference, AIME, Vol.

29, 1971, p. 31. 107. L.A. Luyckx, Symposium on Deoxidation Practice - With and Without

Ladle Metallurgy, A.I.S.I., 1982, p. 25.

60

Lecture II

Gases in Steel Castings

by Dr. John M. Svoboda

INTRODUCTION The objective of this review paper is to consider some of the basic principles of physical chemistry and kinetics which relate to the understanding of the behavior of simple and complex gases in cast steel. A secondary objective is to present a system of classification of defects related to gases in order to provide a systematic method for the identification and correc- tion of these defects.

Generally, foundrymen think of gases in cast steel in terms of defects and problems with casting quality. Such nomenclature as pinholes, blows, porosity, blowholes adjacent to chills, surface pinholes, rock candy fracture, gas folds or stringers and metal penetration is used when discussing problems with gases. These defects account for a major proportion of defective castings produced and the resultant expenditure of time and effort in reducing the number of these defects.

Although the presence of gases in cast steel is usually thought of as an undesirable condition, we also use gases both as effective production tools, and to impart desirable metallurgical properties in many cases. Perhaps the most common example is the injection of gaseous oxygen to produce a carbon boil in processing of steel for casting applications. Other examples are argon flushing, the AOD process, injection of calcium alloys with argon to accomplish desulfurization, and the alloying of stainless steels with nitrogen.

61

An understanding of the basic thermo-chemical principles pertaining to the subject of gases in cast steel will allow the foundryman to use gases intelligently as production tools while reducing defects caused by these gases to a minimum. Some of these principles will be considered in the next section.

FUNDAMENTALS OF PHYSICAL CHEMISTRY

Classifications of Gases

For purposes of discussion, it is convenient to classify the types of gases of concern into three categories:

1. Monotomic gases (Ar, He) (inert gases)

2. Diatomic gases (O2, H

2, N2) (simple gases)

3. Complex gases (CO, CO2, H2O, NH3, SO2, H2S)

The monotomic gases, such as argon, have practically zero solubility in steel, and therefore behave in an inert manner toward these alloys. Because of this behavior, these gases can be used as inert carrier gases, and to provide protective cover atmospheres for certain applications.

The following discussion will be limited to those interactions between gases and metals which arise at low solute concentrations which can be conveniently described in units of parts per million (ppm). Emphasis will be placed on the conditions which can lead to gas defects in the castings. Figure 1 schematically illustrates the areas of interest, i.e., the transfer of gas across the liquid-solid, liquid-gas, and solid-gas interfaces; the mobility of gas solutes within the solid and liquid phases; and the formation and entrapment of gas bubbles.

Although each alloy system exhibits its own characteristics, discussion will be directed to general basic principles.

Solubility Curves (Simple Gases) (Ref. 1) The equilibrium dissolved gas content of a pure metal depends on the temperature and partial pressure of the react-

62

ing gas and, to a much smaller extent, on the total pressure of the system. The simple gases are defined as H2, N2, O2, etc. When a pure metal is brought to equilibrium with a gas phase of known partial pressure of H2, a condition exists for which the gas phase contains primarily H2 (with small amounts of H) and the metallic phase contains primarily dissolved mono- meric hydrogen, H, (with very small amounts of associated solutes, dimers such as H-H). The principal reaction describ- ing equilibrium between the two phases may be written as:

63

For which the equilibrium constant is

Equation (2) is a simple statement of “Sieverts’ Law”. The brackets are used to designate activity of dissolved hydrogen; however, since only low solute concentrations are considered

64

it is convenient to replace activity by concentration. Actually, there are a number of secondary reactions which might be considered within each phase such as the homogeneous dissociation of H

2 to gaseous monomers and the association of

dissolved H atoms to form clusters. But, major attention should be

_directed to reaction (1). By equation (2) the

equilibrium H content of a metal or alloy may be expressed in terms of the

_major variable.

The quantity ∆ G° is the standard Gibbs free energy for the solution of hydrogen by reaction (1). Noting that ∆ G ° = ∆ H ° - T∆ S° and rearranging,

Because ∆ S° and ∆ H °, the standard entropy and enthalpy of solution, are not dependent to any signficant extent upon pH2

or T, equation (4) is of the form

Figure 2 is a schematic representation of the effect of temperature and hydrogen and nitrogen partial pressures on the gas contents of a metal. It should be noted that alloying elements may raise or lower the solubility curve, (Figure 3). Another effect of alloying elements is to change the discontinuity at the melting point by creating a temperature region for equilibrium coexistence of liquid and solid. The curvature of the solubility relationship in the liquid-solid zone is determined by the liquidus-solidus boundaries of the phase diagram. It should be kept in mind that the previous discussion is applicable to the solution of any diatomic gas- hydrogen, nitrogen, oxygen, etc.

Nevertheless, the solubility of a gas in a metal can not be increased without limit by increase in the partial pressure of the gas. In each case a limiting solubility is reached at which the solute concentration results in precipitation of a condensed phase.

65

For example, in liquid iron, saturation with oxygen is reached at about 10-8 atm. pressure of O

2, above which liquid oxide

forms. (Figure 4) The dissolved oxygen content at the limit is about 2000 ppm. However, the addition of only 0.1 % aluminum can lower the solubility limit to about 2 ppm which, in this case, arises from precipitation of AI

2O

3. (Figure 5) Solubility

limited by precipitation of a non-metallic phase is an important topic to be treated in the discussion on compound gases. Figure 4 shows that, in general, an alloying element which tends to compound with the gas solute more strongly than the base metal will raise the solubility of the gas at a given partial pressure and lower the solubility limit determined by precipitation of a non-metallic phase. Such interactions occur with hydrogen, nitrogen, carbon, and sulfur as well as with oxygen.

Partitioning Between Liquid and Solid In addition to reaction (1) which describes equilibrium between the metallic phase and a gas phase, mutual equilibrium

66

between the liquid and solid phases in the absence of a gas phase may be represented by

Reaction (6) defines the partitioning behavior of a gas solute between the two metallic phases

The equilibrium constant, K6, depends upon temperature and

the particular alloy system; however, it is, for all important systems, less than one. Again, it should be pointed out that hydrogen is used for illustrative purposes only and that the other gases exhibit similar behavior. Because K

6 is less than

one, during the advance of solidification, the gas solute concentration in the liquid increases by the partitioning effect. Therefore, an initially low level of dissolved gas in the liquid can still result in overpotential for porosity formation in the last sections of a solidifying system. For example. delayed rimming action in steels.

67

Chemical Reactions between Metals & Compound Gases The solution of compound gases (such as CO, CO

2 H

2O, NH

3

SO2, H

2S) in metals is similar to the solution of simple gases in

that the dissolved gas is generally fully dissociated in the metallic phase. That is

The major difference is that the amount of each component dissolved depends upon the other.

68

For a given pX2Y, X increases as Y decreases. Figures 6 and 7 illustrate this effect for solution of H

2O and CO. At low levels

of dissolved O, water vapor is very effective as a source for dissolved H. Sim

_ilarly, at low levels of dissolved C, carbon mon-

oxide or_

carbon dioxide are very effective reag_ents for dis-

solved O. The solubility products of the compound gases increase

_with increasing temperature as in the case of the sim-

ple gases. And, the solubility products of the compound gases are greater in the liquid than in the solid. The effect of alloying additions on the solution of a gas such as CO is complicated

69

by the fact that one must consider the tendency for alloy - C clustering as well as alloy - O clustering. (Ref. 3) Because of

_

this it becomes necessary to_

work with activity coefficients calculated from interaction parameters. (Refs. 4-5)

Nucleation of Gas Bubbles In recent years the understanding of nucleation has increased very greatly. As a simple introduction to the subject, we might consider the free energy change resulting from the sudden clustering of n free solute atoms at constant temperature and pressure to form a spherical cluster of radius, r:

g is the free energy change per unit volume of a very large cluster and σ is the free energy change per unit of area of a cluster. Figure 8 shows a typical plot of ∆ G versus r for conditions favorable for formation of large clusters, g<O.

The maximum point of the curve is found from equating dG/dr to zero, which yields r* = - 2σ. __ All clusters of size greater

gthan r* can increase in size with dissipation of free energy to

70

relieve the overpotentials within the system. However, clusters of size less than r* are more probably to undergo dissociation with a dissipation of free energy. Therefore r* is viewed as a critical size cluster which must form through fluctuation processes, within the system in the same sense that a diffusing atom must overcome activation barriers in moving through certain systems. The previous treatment applies (with many limitations) equally to the case of gas bubbles and non- metallic inclusions. However, the presence of favorable foreign surfaces on which clustering can take place changes the nucleation from homogeneous to heterogeneous. Thereby, r* (or an equivalent radius, since the particle is not spherical under these circumstances) can be drastically reduced so that much smaller overpotentials and fluctuations are required to generate clusters of size greater than r*. During most metallurgical processes - particularly those involving porosity and inclusion formation - nucleation takes place heterogeneously. This is established by the fact that rarely are the concentrations of solutes much higher than those predicted by equilibrium considerations.

71

By mechanical stability considerations, the pressure difference across a spherical bubble can be shown to be 2σ/r where σ is the interface tension (generally equivalent to that previously defined). Therefore, just as the minimum value of g required for an overpotential necessary to have a significant probability for the generation of clusters of size, r* is

the overpressure within a system required for generation of gas bubbles of size r* is given by

In this sense, g and P are equivalent in that g relates to the undercooling or excess concentrations necessary for nucleation while P relates to the excess fugacity of gaseous solutes within the system.

Growth Although the interface reactions at a growing bubble or inclusion are important, in general, the growth of such phases in liquid metals depends primarily upon diffusion of solute elements to the growing interface. In the absence of stirring, an approximate solution of Zener (Ref. 6) to the three dimensional diffusion problem may be applied. (Figure 9) In this case local equilibrium at the reaction interface is assumed and therefore, the rate of growth of the particle is given simply by the solute gradient at the particle surface.

where ro is the radius of the particle, D is the solute diffusion

coefficient, and Co, C i are the solute concentrations in the precipitate and in the melt adjacent to the precipitate (interface equilibrium). By dimensional arguments alone, a solution to equation (11) of the form ro = (Dt) 1/2 is indicated, where D is a dimensionless parameter determined, by Co, C i, and C (the solute concentration at distances far removed from the parti-

72

cle). Or, more generally, is a measure of the excess potential or driving force for growth. Exact solutions to equation (11) for certain boundary conditions have been found by Zener; however, an approximate solution based on a finite depleted zone in the vicinity of the particle is found to give reasonable values. For example, if a depletion zone bounded by ro and r1 is defined by

73

If, on the other hand, strong agitation of the melt exists, then the assumption of a finite depleted zone as previously defined is unreasonable and it becomes necessary to express the solute gradient at the interface in the form (C - C

i)/E, where E

can be called the equivalent static boundary thickness (Fig. 10) (Ref. 7).

Then by integration of equation (11)

However, E has been assumed to be independent of ro. The dif-

ference between the two cases (equations 13 and 14) is best seen by noting that when E is of the same magnitude as r the growth of the particle with or without strong mixing is about the same. However, for larger particles surrounded by the turbulent liquid, E<<ro, and the growth rate is increased by mixing. Therefore, the growth of very small particles is not as greatly influenced by turbulence as the growth of large particles.

The autocatalytic effects of gas evolution from liquids are well recognized. The motion of the fluid induced by rising gas bubbles promotes growth of the held particles to a size necessary for breakaway from their held positions, and a boil can build up to catastrophic proportions. This effect is often noted when oxygen is injected into a steel under poor stirring conditions. With some delay after injection, a small boil initiates and gradually builds up as the overoxidized steel is mixed into the low-oxygen steel.

The boil can often be of such violence that large quantities of metal are ejected from the melt. With proper care - good mixing during injection or small intermittent injections - the boil can be controlled. The essential precaution is to avoid the establishment of very large potential differences in the relatively quiet melt.

Redistribution of Solutes During Freezing In the previous section, the partitioning of solutes between the liquid and solid phases was described. As Pfann (Ref. 8) has shown, the concentration of a solute in a liquid metal generally builds up (provided that the solute depresses the melting point) during the course of normal solidification. Only by extremely slow or fast solidification can the redistribution or zone-refining effect be reduced to insignificance.

Figure 11 indicates an idealized solute concentration profile at an advancing planar liquid-solid interface. Distribution of the original solute between the two phases must satisfy the conservation equation:

74 75

Where: CL and C

S are solute concentrations of liquid and solid

zones respectively; f is the volume fraction of the melt solidified; and C

L° is the initial solute concentration of the

melt. The integral on the left side of equation (15) must be taken over the entire solidified mass wherein C

S is a function

of position. Equation (15) is based on a generally uniform solute concentration throughout the liquid except for the gradient which exists over only a small volume element near the interface - that is C

L is a function of the extent of solidifica-

tion (or time) but not a function of position. By differentiating equation (15) and rearranging the following differential equation is obtained:

The quantity (CL - C

S) is dependent upon the equilibrium parti-

tioning of the solute, K, and upon the magnitude of the solute

76

concentration gradient in the liquid at the interface. The Burton-Prim-Slichter equation for the effective distribution ratio, C

S/C

L, is

where w is the velocity of the liquid-solid interface, ois the defined boundary layer thickness, and D is the diffusion coefficient of the solute in the liquid. Note that for large values of the dimensionless parameter, wσ/D, that the effective distribution ratio approaches unity (C

L - C

S = 0) and no segregation

occurs. On the other hand, for small values of wσ/D, the effective distribution ratio approaches the equilibrium partitioning ratio and C

L = C

S = (1 - K)C

L. For typical mixing conditions

and solute diffusion coefficients, the effective distribution ratio approaches unity for growth rates above 10-1 cm/sec. and approaches K for growth rates below 10-4 cm/sec. Therefore during the early stages of solidification, segregation is not as severe as during the later stages of a large solidifying section.

If the effective distribution coefficient is taken to be constant, then equation (16) can be integrated to yield

In addition to the limitations already stated, equation (18) only holds when the solute does not participate in other reactions - gas evolution or precipitation of an inclusion phase. Fur- thermore, it can be seen that C

L approaches infinity as f ap-

proaches 1. This is obviously impossible and, therefore, other reactions as previously stated must occur before solidification is complete. If the solute is non-gaseous, equation (18) is limited also by enrichment of the liquid to an invariant point, i.e. eutectic, or to the pure solute. The information available on zone melting and directional solidification may be drawn upon heavily in considering segregation effects during normal solidification of castings. Clearly, the formation of porosity and inclusions during the later stages of solidification is an important problem.

77

Entrapment of Gas Bubbles and Inclusions A rising particle - gaseous or condensed - reaches a ter- minal ascent relative velocity in a liquid metal which for small spheres is given by

where ∆ p is the difference between the density of the melt and the density of the particle, g is the gravitational constant, η is the liquid viscosity, and r is the particle radius. For typical metallurgical conditions the velocity in cm/sec. is about 105 times r2 (radius in cm). Therefore the entrapment of a free particle which has a radius greater than 10-3 cm is very unlikely. A free particle is one which is not attached to or blocked by a surface within the melt. In general the ascent velocity of an inclusion is about 1/2 to 1/4 that of a gas bubble. If the melt is severely agitated, it is possible for large particles to be held in swirls or to be caught in downward currents. However, the general motion of a fluid within which gas is being evolved is upward along the sides and downward at the center. There are two other factors which help to eliminate particles: coalescence of small particles to form a combined particle and the flotation of inclusion particles by gas bubbles.

If surface characteristics are favorable, an inclusion particle can attach itself to a rising gas bubble just as occurs during mineral benefication by the flotation process.

A growing gas bubble is held to a surface by surface tension forces. For a bubble growing on a horizontal surface, balanc- ing the surface tension retention force and the buoyancy force yields

where rc is the effective breakaway size of a spherical bubble

(approximated), a is the radius of the pore acting as retention zone, and δ, g and p are as previously defined. For typical metallurgical conditions r

c = a1/3 which means that for__

σ

78

for values of a ranging between .006 and 6 cm, rc varies from

0.1 to 1 cm. This accounts for the relatively uniform size of evolved gas bubbles. A bubble attached to a side wall tends to break away at slightly smaller sizes than given by equation (20). For side-wall bubble:

The mechanics of the disattachment are complicated. A large bubble tends to break away in a manner which produces a shower of very small free bubbles in addition to the major free bubble.

At a plane liquid-solid interface, individual bubbles may be entrapped or a string of bubbles (pinholes) may be entrapped depending upon the relative growth of the gas bubbles and velocity of the interface. If the velocity of the interface is very large and the growth rate is small the bubble tends to be frozen in place before it achieves a size necessary for breakaway. (Figure 12)

On the other hand, for relatively low interface velocity and high growth rate, the bubble tends to achieve breakaway size repeatedly and a channel of varying thickness results.

For typical growth conditions, a bubble can grow to breakaway size in about one second, therefore, the interface must be moving at a velocity somewhat above 1.0 cm/sec. for a bubble to be closed off individually.

The tendency for dendrite formation, greatly increases the probability of entrapment of gas bubbles. Furthermore, the size of bubbles which can be trapped is much greater.

Internal shrinkage zones - where liquid feeding is not adequate - are favorable regions for the formation and entrapment of large gas bubbles.

EXAMPLES The foregoing principles may best be illustrated by considering their application to the process of producing steel castings. (Ref. 3)

79

Activity and Solubility of Oxygen Oxygen dissolves in liquid iron, with its solubility limited by the formation of a liquid iron oxide phase. The oxygen pressure at which this phase appears has been calculated to be 0.8 x 10-8 atm at 1600°C (2912°F). This pressure is too low to measure directly, therefore it is often convenient to consider the solubility of iron oxide (rather than oxygen gas) in liquid iron.

The solubility of oxygen in high-purity iron depends on temperature. Considerable experimental difficulties are encountered in obtaining accurate solubility data because of the high reactivity of liquid iron oxide with any refractory con- tainer used in the experimental work. However, techniques have been developed by Chipman and co-workers (Refs. 9, 10) which can account for these problems, and the solubilityof oxygen in liquid iron in equilibrium with FeO is given by:

At times there is some confusion in using the equation since it has been the custom of some writers to assume that the dissolved oxygen is combined with the iron as FeO. There is little fault to be found with this assumption as long as we are dealing with the simple iron-oxygen system. However, in liquid steel there is ample possibility for many other oxides to be present and consequently it is unsafe to convert total oxygen to FeO. For those who prefer to express solubility of iron oxide in pure iron as a percentage of FeO, the relationship is:

The Carbon-Oxygen Equilibrium

Carbon and oxygen dissolved in liquid iron react to form the gases which produce the familiar carbon boil. The reaction of interest can be written as follows:

This equilibrium may be studied experimentally by considering the following two reactions:

Adding these we obtain the equation of interest (Equation 24). A considerable amount of experimental effort has been ex- pended in determining the equilibrium constants for Equations (25 & 26), which has provided the relationship for Equation (24):

It can be seen that under a given set of temperature and PCO2

conditions, a balance exists between [%O] and [%C]. This is usually expressed as the product [%O] x [%C] which equals 1.87 x 10-3 at 1800°K and 1 atm CO.

80 81

The above relationship, strictly speaking, is applicable to very dilute solutions of carbon and oxygen in iron. At increasing concentrations, appreciable departures from Henry’s law occur. Therefore we must replace [%C] with ac = fc [%0] and [%O] with a

o = f

o [%O]. Also, at very low carbon concentra-

tions the equilibrium gas may contain an appreciable fraction of CO2.

However, from most practica! discussions, the above correc- tions are unnecessary because the activity coefficients of carbon and oxygen change in opposite directions when their concentrations increase. The most important result of this is that the product [%C] x [%O] at constant CO pressure and constant temperature only exhibits a relatively small dependence on the C and O concentrations. For all concentrations which are of importance in steel melting, it is justifiable in practical discussions to assign this product a constant value at constant temperature (Refs. 11, 12). It may also be observed from Equation (27) that temperature dependence is also quite smal!. Rough calculations are often made considering the product to be independent of temperature as well. The following values may be considered approximately valid for any carbon content above about 0.02 percent.

The actual value of %C x %O, observed in steel-making practice is not the same as the equilibrium value. In Figure 13 are shown several curves for carbon versus oxygen derived from several studies of the open-hearth and electric furnace processes. In every case the product is substantially higher than that corresponding to equilibrium with carbon monoxide at 1 atm. The difference is, of course, associated with the mechanism of the overall process by which oxygen diffuses from slag through the metal to the point at which bubble formation occurs, this point being located at the furnace bottom. The excess oxygen content over and above the calculated value for equilibrium at 1 atm has been called O. This quantity proves to be a characteristic of the steel-making practice being

82

employed, and is in general lower in the electric arc than in the open-hearth furnace.

Hydrogen in Steel

Hydrogen is generally present to some extent but is always un- wanted in commercial steels. It can cause porosity and hairline cracks, particularly in large steel castings. Prolonged heat treatment over several weeks could be required to decrease the hydrogen by solid state diffusion to an acceptable level. This is obviously impractical and thus where steel is required with low hydrogen contents, every attempt should be made to minimize hydrogen absorption during processing of liquid steel.

The solubility limit of hydrogen in solid iron at its melting point is about 0.001 pct (i.e. 10 ppm). If the hydrogen content of the liquid exceeds this amount, it will be rejected during freezing and this leads to pinhole formation and ingot porosity.

The reaction between H2 and liquid iron can be written:

The effect of temperature on the equilibrium constant is given by:

For a partial pressure of one atmosphere equations (28) and (29) give the solubility of hydrogen in pure iron at 1600°C

83

(1873°K) at 0.0027 wt. pct. (27 ppm). If the pressure is reduced to 10-3 atmospheres, as in vacuum degassing, the equilibrium content would be about 1 ppm.

As in the case of nitrogen, interaction parameters are available which describe the effect of various alloying elements on fh so that the solubility of hydrogen in alloy steels can be readily calculated.

Absorption of hydrogen by molten steel can take place from moisture or hydrocarbons in the furnace atmosphere, from limestone, ore, millscale, scrap or alloy additions during refining, from the atmosphere during tapping and pouring, from damp refractories in the ladle, or from hydrocarbons in mold coatings or binders.

The dissolution of hydrogen in molten steel from water vapor may be expressed as:

Expanding equation (30)

This relationship between hydrogen and oxygen in liquid iron at 1600°C is shown in Figure 6 for a range of water vapor pressures from that typical for an arc furnace (0.015) to that for an open hearth (0.3). From this diagram it can be seen that the hydrogen content increases as the oxygen content decreases. At low oxygen levels typical of killed steels, or at the higher water vapor pressures, the equilibrium hydrogen content actually exceeds that which would be in equilibrium with pure hydrogen at one atmosphere pressure. Fortunately these equilibrium values are well in excess of those normally encountered in practice (i.e. about 2-6 ppm), due mainly to the continuous removal of hydrogen during refining by the rinsing

84

action of carbon monoxide bubbles. Further removal of hydrogen can be accomplished by argon rinsing in the ladle or AOD processing to give hydrogen levels below 2-3 ppm.

In summary, hydrogen absorption can be decreased by:

(i) Preheating scrap and other furnace additives.

(ii) Maintaining a good carbon boil.

(iii) Minimizing the time deoxidized metal is exposed to the atmosphere.

(iv) Ensuring the refractories are properly dried.

(v) Avoiding excessive use of hydrocarbon mold washes.

Nitrogen in Steel

Although the solubility of nitrogen in iron base alloys is in general small, the effects of nitrogen on the properties of steel may be quite profound. For most purposes nitrogen in finished steel is undesirable, particularly in the low-carbon grades, since on cooling to room temperature the solubility limit of nitrogen in the steel may be exceeded and this can lead to embrittlement and loss of ductility on aging. On the other hand, nitrogen can improve the workhardening properties and machinability of steels while in certain stainless grades nitrogen is important in order to stabilize the austenite phase. (Ref. 26)

At a partial pressure of one atmosphere, the solubility of nitrogen in pure iron is as shown in Figure 2. At the freezing point the solubility drops from about 0.045% to 0.013%. Any gas in excess of this level will be evolved and could cause porosity. Generally, however, the nitrogen content of steel is below the solid solubility limit and porosity due to nitrogen alone is not common.

85

The reaction between nitrogen and liquid iron can be written:

For pure iron, fN = 1, but in the presence of alloying elements

it may be more, or less, than unity.

The effect of temperature on nitrogen solubility in liquid iron is given by:

At 1600°C (T = 1873°K), the solubility of nitrogen in liquid iron under air, for which pN

2 = 0.79, is calculated from (32) and (33):

From (33), log K = -188.1 / 1873 - 1.246, i.e., K = 0.045, which from (33) corresponds to the nitrogen content of pure iron, when pN

2 = 1 atm.

Also from (33)

In hot metal or alloy steels, the effect of various elements on the behavior of nitrogen means that the activity may be actually greater or smaller than the actual concentration would imply. This means that f

N does not equal unity. The effect of

various elements on (wt %N) and on fN is shown in Figures 13

and 14 respectively. Both carbon and silicon decrease the solubility, and from a particular nitrogen activity, this means an increase in f

N, equation (32).

86

The effect of alloying elements may be represented approximately by equations of the following type:

where eCN

and eSiN

are called interaction parameters.

For example, the effect of chromium on the activity coefficient of nitrogen at 1600°C is given by eC

N = -0.045.

Thus for an alloy containing 22.2% Cr,

This means that nitrogen is only one tenth as active in this alloy as the same concentrations would be, in pure iron, for which f

N = 1. In other words, the nitrogen content of this alloy

would be 10 times greater than the nitrogen content of pure iron, at the same temperature and under the same nitrogen partial pressure. For pure iron under air at 1600°C, the equilibrium nitrogen concentration is 0.040% (Equation 35). For the 22.2%Cr alloy, this would be increased to approximately 0.4%.

The very low activity coefficient of nitrogen in high chromium alloys is one reason why it is particularly difficult to obtain low nitrogen levels in stainless steels. When such steels are produced in the electric arc furnace, nitrogen contents between 0.02 and 0.04% are not uncommon. This high value is also partly due to the dissociation of nitrogen molecules in the high temperature arc and to the nitrogen content of the alloy additions.

In carbon steelmaking, considerable amounts of nitrogen may enter the furnace with the scrap, but this is effectively flushed out of the metal during the carbon boil and the final nitrogens are often about 0.002%. In BOF steelmaking, the final nitrogen level can be increased if there are several turndowns prior to tap. This is due to refilling of the vessel with air, and the subsequent nitrogen absorption by the metal.

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Nitrogen pick-up from the atmosphere may also occur during tapping and pouring but the actual amount absorbed can be affected by the concentration of surface active elements, such as oxygen or sulfur, present in the steel. For example it has been reported that the nitrogen content of killed steels increased by 15-20 ppm during tapping and a further 15-20 ppm during pouring. On the other hand, no change was observed in the nitrogen content of rimming steel during tapping and pouring operations. This is probably attributable to the higher oxygen contents inhibiting nitrogen absorption. Laboratory studies have shown that sulfur has a similar inhibiting effect, as do selenium and tellurium.

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In addition to inhibiting nitrogen absorption, surface active elements also decrease the rate of nitrogen removal during vacuum degassing. For this reason, where steels with low nitrogen levels are required, this is best accomplished by careful selection of low-nitrogen alloy additives and minimum exposure of steel to the atmosphere during casting. This may involve the use of inert gas shrouds, or closed pouring systems.

Reoxidation Considerations (Ref. 27)

Refractory Interactions In the production of most steels, alloying additions are made to the ladle, and under these circumstances the ladle can no

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longer be regarded as simply a method of transporting molten steel from the furnace to the casting station. In actual practice, after preliminary refining in the furnace, steel of the re- fined composition is finally “made” in the ladle, possibly even with some additions to the mold. While it is true that slag/ metallrefractory interactions have always occurred in the ladle, such reactions become of critical significance when they involve elements which are present only in the parts per million range. Since it is these same elements which will determine whether or not the steel will possess the mechanical properties required in the final product, particular attention must be given to the reactions which occur from the time the steel leaves the furnace until it solidifies. For this reason the ladle should be considered as a steelmaking vessel, and treated accordingly.

Steelmaking ladles generally contain silica and alumino- silicate phases which are unstable when in contact with molten steel containing strong deoxidizing elements such as Ti, AI, Zr and rare earths. In addition, reaction between the slag, which is often basic in character to enhance sulfur removal, and the acidic refractories will result in the formation of a ladle glaze which may be rich in iron oxide, silica, and sulfur from the slag. Refractorylslag interaction with the formation of a glaze, and a subsequent glazelmelt interaction during the next heat in the ladle, will result in a loss of the deoxidizing elements for the steel with a consequent increase in oxygen content of the melt during the time it is held in the ladle.

In a similar way, when the sulfur dissolved in the steel has been lowered by the addition of strong desulfurizers, subsequent sulfur pick-up may occur from the ladle glaze formed during the previous heat. With sulfur present in the glaze at a higher activity than sulfur dissolved in the steel, there exists a driving force for sulfur transfer to the metal. This is a repetitive process from heat to heat. As the steel from one heat leaves the ladle, a new “sulfur-rich” glaze is formed by reaction between the ladle slag and refractory. When the next heat of steel is desulfurized, there is again the opportunity for sulfur pick-up from the glaze formed during the previous heat.

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It is clear that the ladle glaze in effect acts as a reservoir for both oxygen and sulfur. These elements, present in the slag, and subsequently stored in the glaze during one heat, are effectively pumped back into the steel during the next heat, after the addition of deoxidizingldesulfurizing elements. The increase in oxygen content of the steel can have a secondary, but nevertheless important, effect on the composition and therefore morphology, of the sulfides suspended in the melt. Since the strong sulfide formers are even stronger deoxidizers, and increase in the oxygen content of the melt implies that sulfides which were originally formed from strongly deoxidized steel, become unstable, and undergo transformation to oxy-sulfides andlor oxides, with sulfur reverting to the melt.

If reversion reactions of this type are excessive, iron-manganese sulfides may eventually form with consequent loss of shape control. In conventional steel-making operations, one would not expect to remove sulfur in an acid lined furnace, and for the reasons discussed here it is not at all surprising that low irregular recoveries, with erratic product performance, are observed when additions of desulfurizing elements are made to acid lined ladles. Improvements are to be anticipated when the working lining of a ladle consists of a stable, basic oxide such as magnesia rather than alumino-silicate. Ladle spraying with a magnesia or burned-dolomite base mix should also improve recoveries and performance, as will the use of insulating board linings.

Pouring Stream Effects In the casting of steels gas pick-up may occur at the following locations:

(1) by direct contact between the falling stream and the atmosphere;

(2) by reaction between the steel and bubbles of air carried down into the sprue with the falling stream;

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(3) by reaction at the mold-metal interface with the at-

Water model studies of casting streams and the effects of gas entrainment in continuous casting have been conducted by Maddever et al. (Refs. 14, 15) With smooth streams, gas dragged downwards by boundary layer attachment, is peeled off at the pool surface with very little entering the mold pool. Under these conditions gas pick-up will occur only by direct contact with the falling stream and at the pool surface. In the first cast the time of contact is only a fraction of a second, the surface arealvolume ratio of metal exposed to the atmosphere is relatively small and thus gas pick-up at this location should be of minor consideration. Similarly at the pool/gas interface, conditions are calm and relatively quiescent and here again gas pick-up should be relatively slight. Conditions are in marked contrast with a rough, broken or flaring stream. In this case copious quantities of gas are entrained by the falling fluid and carried into the bulk of the mold pool. Swarms of bubbles rising to the pool surface together with “wild” stream create extremely turbulent conditions at the surface. Thus with rough streams, conditions for gas pick-up are greatly enhanced. The ratios of surface area/volume of metal exposed to the atmosphere, during the fall, within the pool, and at the pool surface are all drastically increased and severe contamination of steel can occur.

By using polyethylene beads as tracers, it has been possible to observe and record photographically the mixing patterns in the mold associated with smooth and rough streams. In the former case, the particular matter remains on the surface except when it drifts directly under the incoming stream, in which case it is driven deep into the mold pool. In practice this could lead to the entrapment of mold slag within the internal structure of the casting. On the other hand, the surface of the metal next to the mold walls should be relatively clean and free from splash defects. In the case of the rough stream, the non-metallic material is forced outwards away from the incoming stream and towards the mold walls. Here it is carried downwards along the freezing interface and this in practice could lead to the formation of massive surface or sub-surface defects.

mosphere in the mold.

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Reoxidation Farrell et al (Ref. 16) have discussed in detail the formation of inclusions by reoxidation during castings. They have shown very clearly that in the case of aluminum-killed steel containing manganese and silicon, the first reoxidation products formed consist of large alumina galaxies. With continuing reoxidation, the inclusions formed contain decreasing amounts of aluminum and increasing amounts of manganese and silicon. The presence of silicate inclusions in aluminum- killed steel is strong evidence that reoxidation has occurred. The same conclusion can be made for steels deoxidized with aluminum followed by additions of calcium, zirconium or rare earth alloys.

In laboratory experiments pertaining to reoxidation, Green- berg (Ref. 17) has measured the rate of oxygen pick-up by iron droplets falling through air. Steel droplets were melted within a levitation coil, allowed to fall for various times through air in a reaction chamber, solidified in a copper mold located within a quench chamber and then analyzed for oxygen. It was found that the rate of reoxidation was unaffected by concentration changes in sulfur, titanium or aluminum. The measured rates were in good agreement with predicted rates for a reaction controlled by oxygen diffusion in the gas phase. This implies that the rate of reoxidation of steel exposed to the atmosphere is rapid and unhindered by the presence of sulfur or strong deoxidizers such as titanium or aluminum. The extent of reoxidation is minimized by decreasing the contact time between the steel and the atmosphere, i.e. the fall height, and by ad- justing conditions in the ladle andlor tundish to provide smooth pencil-like streams rather than coarse flaring streams with associated tearing, splashing, air entrainment and droplet formation. Improvements in steel cleanliness and heat to heat consistency should also be achieved with the use of inert gas shrouds and closed pouring systems.

Classification of Gas Related Defects One of the most important steps in the elimination of casting defects due to gases is the proper identification and classification of those defects. Many excellent papers have been written which can guide the foundryman in this task. (Refs. 18, 19) A practical and widely accepted method of defect classifi-

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cation has been developed by the International Committee of Foundry Technical Associations (CAITF) and is presented in the International Atlas of Casting Defects. (Ref. 20) Appropri- ate sections from this Atlas (using the CAITF classification numbers) are condensed in the following section.

Blowholes, Pinholes Smooth-walled cavities, essentially spherical, often not contacting the external casting surface are defined as blowholes or pinholes. The larges cavities are most often isolated, the smallest (pinholes) appear in groups of varying dimensions. The interior walls of blowholes and pinholes can be shiny or more or less oxidized. The defects can appear in all regions of the casting.

Possible Causes Blowholes and pinholes are produced because of gas entrapped in the metal during the course of solidification. Such gases, however, can also cause defects other than blowholes or pinholes.

1. Metallurgical Origin (endogenous gas holes) Excessive gas content in metal bath (charge materials, melting methods, atmosphere, etc.); dissolved gases are released during solidification. In the case of steel: formation of carbon monoxide by the reaction of carbon and oxygen, present as a gas or in oxide form. Blowholes from carbon monoxide may increase in size by diffusion of hydrogen or, less often, nitrogen.

2. Gas Arising from Mold or Core Materials (exogenous gas holes)

excessive moisture in molds or cores; core binders which liberate large amounts of gas; excessive amounts of additives containing hydrocarbons; coatings which tend to liberate too much gas.

3. Mechanical Entrapment of Gas (exogenous gas holes) insufficient evacuation of air and gas from the mold cavity (venting);

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insufficient mold and core permeability; entrainment of air due to turbulence in the sprue and runner system.

Remedies make adequate provision for evacuation of air and gas from the mold cavity (vents); increase permeability of mold and cores; use proper gating systems; assure adequate curing of no-bake molds; control moisture levels in green sand molding; reduce amounts of binders and additives used or change to other types; use coatings which provide a reducing atmosphere; keep the sprue filled and reduce pouring height; increase static pressure by enlarging sprue height; deoxide adequately; avoid reoxidation; reduce hydrogen and nitrogen contents by proper melting procedures; control the temperature and time of pouring.

Hints in Trouble-Shooting Diagnosis, while often difficult, is aided by the following:

Very large cavities (gross blowholes) are usually of external (exogenous) origin. Exogenous blowholes are often of variable dimensions. They may be isolated or in irregular groupings. Endogenous (internal origin) blowholes are usually small, uniform in size and distributed evenly through the casting or in one part of the casting. In steel, hydrogen blowholes have shiny walls, CO blowholes are bluish, and those from entrained air and gray and slightly oxidized.

Diagnosis is not always possible by a single examination of the defect. It generally requires investigation and studies, and then adjustment of metallurgical or molding factors.

Blowholes Adjacent to Inserts, Chills, Chaplets, Etc. Blowholes of varying dimensions within a casting, isolated or

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grouped, immediately adjacent to chaplets, chills, or other metallic inserts.

Possible Causes Moisture or oxides on the metallic parts. Absence of surface protection such as zinc or copper plating. Condensation of moisture on the metallic parts due to difference in temperature of parts (too cold), mold and cores. Use of parts which are dirty, oily, or coated with washes which tend to liberate gas.

Remedies Use only metallic inserts which are galvanized or copper plated. Carefully clean, dry, and preheat all metallic parts before placing them in a mold. Be sure the mold and cores are lower in temperature than the metallic inserts at time of use. Use chills having grooved surfaces.

Slag Blowholes Same appearance as blowholes but always accompanied by slag inclusions. Generally localized in the upper portion of the castings.

Possible Causes Oxidation reactions within the liquid metal which create liquid or solid oxides and thus gases (blowholes). Reaction of the liquid metal and its oxides with the furnace refractories, ladle linings or materials comprising the mold.

Surface or Subsurface Blowholes Cavities whose walls are generally smooth and rounded, often in the form of flattened bubbles with rounded or angular corners, located either singly or in groups at or near the surface of the casting.

The cavities are sometimes exposed to the surface, but most often are located beneath a thin layer of metal and cannot be seen until after blast cleaning or machining; they may sometimes appear as shiny spots on the casting at shakeout.

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Possible Causes Insufficient permeability of mold or cores. Pouring too slowly or too cold. Insufficient height of runners and risers. Air and gas bubbles which gravitate to the upper surface of the casting but are unable to escape, either because of lack of mold permeability or because of a presolidified skin on the casting.

Surface Pinholes Small cavities, often the size of the head of a pin, located in more or less extended colonies across the casting surface, and having one to two forms:

Generally spherical cavities which are removed by machining 1 or 2 mm (0.04 to 0.08 in.) from the casting surface (surface pinholes). Elongated (teardrop or rod-shaped) cavities which may be seen on a cut or fractured cross-section, located near the cast surface. If these cavities appear only after machining an apparently sound surface, after heat treatment, or if small holes seen originally on the as-cast surface become larger during machining, they are considered subsurface pinholes. The long axis of these smooth-walled, elongated cavities is always perpendicular to the casting surface and are situated between grain boundaries, as may be shown by macro etching. In general, the depth to which these pinholes extend beneath the surface does not exceed 4 mm (0.16 in), the remainder of the section being sound.

Occurance - Cast Steel especially in casting sections whose thickness is on the order of 15 to 30 mm (0.6 to 1.2 in.).

Green Sand Castings in the vicinity of surfaces formed by cores made with organic binders.

Possible Causes Metal I urgical

excessive oxygen or hydrogen content in the metal bath (due to charge or the melting process); excessive tem perat u re.

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From Mold and Core Materials too much moisture in mold andlor cores; in the case of organic binders, too high a urea content.

Formation Surface Pinholes

Reactions between carbon and slags rich in iron oxide (un- saturated in silica) result in the formation of carbon monoxide, which is evolved as surface pinholes. The latter may be further enlarged by the diffusion and liberation of hydrogen.

Sub-Surface Pinholes Surface oxidation of the metal stream entering the mold cavity, by reaction with the mold atmosphere or with constituents in the molding materials. Reaction of the metal oxides thus formed with the carbon in the liquid metal, resulting in the release of carbon monoxide. These reactions, in the zones adjacent to the surface, can cause cavities which may become even further enlarged by the release of hydrogen. When the gas pressure attains sufficient value, bubbles form as thin channels, displacing the intercrystalline liquid between the cells which are forming in the vicinity of the surface; they emanate from near the surface, through which they may emerge in some cases.

Remedies Molding

Reducing moisture content of mold and core sand mixtures; Employ mold and core binders containing as low a urea content as possible.

MetalIurgy Maintain the hydrogen content of the bath as low as possible, which may be accomplished by strong agitation of the bath during refining. Deoxidize the bath carefully with manganese and silicon, and, during pouring, with aluminum. Avoid excessively long runners. Make the sprue and runner system of refractory hollowware. Pour rapidly but without turbulence.

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Conchoidal or Rock-Candy Fracture The fractured surface of the casting shows smooth, slightly curved facets like that of rock candy.

Possible Cause This defect occurs in steel castings when the aluminum and nitrogen contents are too high, causing aluminum nitride to precipitate at grain boundaries (particularly in heavy sections and low alloy steels).

Remedies Limit the addition of aluminum to the value sufficient to deoxidize the steel. Also, fix the nitrogen by titanium or zirconium additions (with due regard for deleterious effects on mechanicaI properties).

Surface Folds, Gas Runs Irregular fold marks distributed across a surface of the casting. For cast steel, the defect most often occurs on thin, horizontal surfaces.

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Possible Causes High viscosity of the liquid metal. Low pouring temperatures and slow pouring. Gases formed by reaction between the metal and the mold wall. Formation of (oxide) skins during pouring.

Metal Penetration A projection of particular geometic shape, comprised of an intimate mixture of sand and metal having a spongy appearance and strongly adhering to the casting. Generally occurs at locations where the sand is the hottest (cores, concave sections) and is of lowest density.

Although many factors contribute to the occurrence of metal penetration, as well as burn-in and burn-on, it has been shown that excessive oxygen content in ferrous metals is a significant contributing factor (Refs. 21, 22). (Figure 15)

Summary Basic principles of physical chemistry relating to gases in cast metals have been presented as well as a suggested method of defect classification. It is hoped that these comments will stimulate further interest and research in this important area of metal casting technology.

References 1. G.R. St. Pierre, "Physical Chemistry Fundamentals Governing the

Behavior of Gaseous Elements in Metals," AFS-T&RI Third Advanced Seminar, Asheville, N.C., October 1966.

2. J.D. Fast, INTERACTION OF METALS AND GASES, Academic Press, New York, 1965.

3. J.M. Svoboda and A. McLean, PHYSICAL CHEMISTRY OF FERROUS MELTING, Cast Metals Institute, Des Plaines, Illinois, 1975.

4. J. Chipman, Jnl. Iron & Steel Inst., Vol. 180, 1955, p. 97. 5. M. Weinstein and J.F. Elliott, Trans AIME, Vol. 227, 1963, p. 382. 6. C. Zener, Jnl. of Applied Physics, Vol. 20, 1949, p. 950. 7. C. Wagner, PHYSICAL CHEMISTRY OF STEEL MAKING, (J.F. Elliott, Ed.),

8. W.G. Pfann, LIQUID METALS AND SOLIDIFICATION, (R. Daddin, Ed.), ASM,

9. J. Chipman and K.L. Fetters, "The Solubility of Iron Oxide in Liquid Iron,"

John Wiley and Sons, New York, 1958, p. 237.

1958, p. 218.

ASM TRANSACTIONS, 1941, Vol. 24, p. 953.

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10. C.R. Taylor and J. Chipman, "Equilibria of Liquid Iron and Simple Basic and Acid Slags in a Rotating Induction Furnace," AIME TRANSACTIONS, 1943, Vol. 154, p. 228.

11. C.E. Sims, Ed., "ELECTRIC FURNACE STEEL MAKING," Vol. II, Inter- science Publishers, New York, 1963.

12. C. Bodsworth and H.B. Bell, PHYSICAL CHEMISTRY OF IRON AND STEEL MANUFACTURE, 2nd Ed., Longman, London, 1972.

13. A. McLean and D.A.R. Kay, International Symposium on Micro-Alloying, 1975.

14. W. Maddever, et al, Can. Met. Quarterly, Vol. 12. no. 1, pp. 79-88, 1973. 15. W. Maddever, M.A. Sc. Thesis, University of Toronto. 1974. 16. J.W. Farrell, P.J. Bilek, and D.C. Hilty, ELEC. FCE. PROC.,TMS-AIME, 1970,

17. L.A. Greenberg and A. McLean, Trans. ISI Japan, 1974, Vol. 14, pp. 395-403. 18. J.V. Dawson, J.A. Kilshaw, and A.D. Morgan, "The Nature and Origin of

Gas Holes in Iron Castings," MODERN CASTINGS, June 1965, p. 144. 19. R.C. Mazumdar, "Causes of Blowholes and Pinholes in Cast and Ductile

Iron Castings," MOLTEN METAL, Jan./Feb. 1975, p. 8. 20. INTERNATIONAL ATLAS OF CASTING DEFECTS, English Edition,

American Foundrymen's Society, Des Plaines, Illinois. 1974. 21. J.M. Svoboda, "Effect of Oxygen Content on Metal Penetraton in Steel Cast-

ings,'' ELECTRIC FURNACE PROCEEDINGS, AIME, Vol. 25, 1967, p. 28. 22. J.M. Svoboda, and G.H. Geiger, "Mechanisms of Metal Penetration in

Foundry Molds," AFS TRANSACTIONS, Vol. 77, 1969, p. 281. 23. W.J. Jackson, STEELMAKING FOR STEEL FOUNDERS, Monograph 4,

Gases in steel, SCRATA, Sheffield, England, 1973. 24. G.J. Davies, SOLlDFlCATlON AND CASTING, John Wiley & Sons, New

York, 1973, p. 16. 25. R.A. Flinn, FUNDAMENTALS OF METAL CASTING, Addison-Wesley

Publishing Co., Inc., Reading, MA, 1963. 26. S.J. Pawel, "The Role of Nitrogen in the Localized Corrosion Resistance of

Cast Duplex Stainless Steels," RESEARCH REPORT NO. A88, SFSA, 1987. 27. J.A. Griffin and C.E. Bates, "Development of Cating Technology to Allow

Direct Use of Steel Castings in High Speed Machining Lines," RESEARCH REPORT NO. 100, SFSA, 1987.

28. R.W. Monroe, "Mold-Metal Interactions: Gas Holes in Steel Castings," STEEL FOUNDERS' RESEARCH JOURNAL, No. 3, 3rd Quarter 1983, p. 5.

29. R.W. Monroe and J.M. Svoboda, "Making Quality Steel Castings: A Review of Twenty Years of SFSA Literature," SPECIAL REPORT NO. 23, SFSA, 1984.

30. J.M. Svoboda, "Melting and Deoxidation of Cast Steels," STEEL CASTING METALLURGY, SFSA, 1984, p. 103.

Vol. 28, pp. 64-86.

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Lecture III

Erosion and Expansion Type Steel Casting Defects by Matt J. Granlund

INTRODUCTION In recent years, the reduction of erosion defects has received increased emphasis because machinability and service requirements have become more important. Fortunately, a larger number of investigators who have studied erosion defects agree on the mechanism of formation than do investigators studying any other major defect in steel castings.

These investigators agree that the expansion of the silica grain is the main factor in all mechanisms proposed for the formation of rat-tails, buckles and scabs. While there is good agreement on the erosion defect mechanism, there are many mechanisms proposed for expansion defects. Again the ever increasing emphasis on quality has made the control of these defects imperative.

Since the composition of the steel is a minor factor in the production of these defects, this lecture will cover the chemical and mechanical aspects of the molding procedure. This lecture does not present any new data on either of the two defects categories, but does summarize the most important points presented through the years. Because of their effect on profitability, the control of these types of defects is an on- going concern to all steel foundries.

EROSION DEFECTS

Definition and Other Considerations

Erosion defects are surface defects that consist of a rough, lump-like protrusion of metal above the intended surface of

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the casting. The lump replicates the depression that was caused by the erosion in the mold. Loose sand may or may not become embedded in the surface of the casting. Other names used for these defects are erosion scab, cut, wash and sand inclusions.

Erosion type defects originate from two sources, erosion of the gating system or erosion of the mold surface. The erosion of the gating system produces loose sand that washes into the mold cavity. This sand is usually carried to the cope surface or is trapped along the vertical walls. When sand inclusions are encountered in a casting it is helpful to get several gating systems, shot blast them and examine them for possible defects. All metal enters the mold through the gating system and the sand in this area usually is subjected to the most severe conditions. Often this source of inclusions is overlooked.

The erosion of the mold surface produces a depression on the drag surface of the mold. The sand that has been “plowed out by the metal” is replaced by the metal. The replacement causes a lump-like defect, usually on the drag side of the casting. The lump requires some additional cleaning effort to get the casting back to the original shape. Even more important than the scab that must be removed, is the location of the loose sand that was moved out by the metal stream. If the sand goes off as small agglomerates of sand they will be embedded in the surface of the casting or the sand could under certain conditions pass off into the risers. This embedded sand may or may not scrap the casting depending on the severity, but it will have a detrimental effect on any machining that would be required. If the sand comes off in one large piece and is embedded in the casting it can produce a defective cast i ng .

Mechanism of Mold Erosion Defects

The mechanism of erosion defects is the dislodging of sand grains or sand masses from the mold surface by metal impingement, metal turbulence or gas agitation. The molten metal occupies the cavity left by the removed sand, causing an erosion defect. Even with the knowledge of the cause of the

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defect and the remedies to prevent it, erosion defects are still widespread because of the severe conditions molten steel applies to the sand mold. Figure 1 outlines the mechanism of erosion defects.

Some investigators have suggested that erosion and metal penetration are in some way related. Both erosion and metal penetration appear to be proportional to the pore size and that many of the remedies for erosion also act to reduce metal penetration. The mechanism that produces erosion defects is not the same as the mechanism that produces penetration defects. Most erosion defects take place before the mold cavity is filled, while most metal penetration defects form after the completion of the mold pouring.

Many conditions aggrevate erosion defects: excessive metal flow in one area of the casting, direct metal impingement on the sand in the mold cavity, excess metal velocity, high pouring temperature, poorly compacted sand and insufficient bond strength in the sand.

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Tests show that erosion, in the absence of scabbing, is not caused by expansion. Metal penetration into the pores between the sand grains and lifting poorly anchored sand grain: out of the mold surface causes erosion defects. Reducing the size of the voids between sand grains by harder ramming, use of finer sands or use of inert filler materials will improve casting results.

Causes of Erosion Defects

The variables that can cause erosion defects are:

Casting and Pattern Design and Flask Equipment Gating System Pouring Molding Sand Core Sand Mold Density and Mold Coatings

Casting and Pattern Design and Flask Equipment A design that requires gating through a thin section, excessive metal flow over a given area of mold surface, and gates ar- ranged so they impinge the metal directly against the mold or core increase the chance of erosion. Wherever possible the casting should be fed through a heavier section and the metai should be introduced into the mold in an equalized flow to avoid excessive localized heating.

Insufficient fillets on gates, abrupt section changes tending to make the molten metal squirt into the mold and flasks too small to permit proper gating increase the potential for erosion defects. Bars that are improperly placed and patterns with vertical walls too close to flask prevent uniform ramming, which give poor rammed properties to the mold surface. Soft spots are prone to metal gouging of the mold surface. Non- uniform density will also have an effect on the expansion characteristics of the sand.

Gating System

Increasing the number of gates or the cross-section of the gates decreases the metal velocity and decreases the tur-

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bulence in the metal stream. A high loss of velocity in the ingate may reduce erosion in the cavity itself, but it may cause more erosion in the gating system where most of the energy is absorbed. Bottom gates produce less erosion than do top gates.

Pouring

The erosion that takes place when metal hits the pouring cup or the sprue is increased by an increase in pouring height. The higher the lip of the ladle, the greater the impact on the sand. Impact on the bottom of sprue and impact on walls of sprue can dislodge sand grains. Use of a pouring basin to take the shock from the poured stream, ceramic tile down sprue and runners and a reservoir at the bottom of the sprue to take the

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shock of the falling stream, all aid in controlling the erosive ef fect of metal stream. The erosion on the sprue wall often is greater than at the bottom of the sprue. The sprue should be full at all times to minimize the erosion on the sprue wall.

Both the sprue height and the sprue diameter have an effect on erosion defects. Savage and Middleton (1) showed that anincrease in sprue height increased erosion and an increase in sprue diameter increased erosion, due to higher velocities and flow rates. Figure 2 shows their finding on effect of sprue diameter and sprue height on erosion.

An increase in pouring temperature will increase erosion. The greater the pouring rate, the less the erosion, provided the stream velocity in the gating system is not increased. The

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greater the weight of the metal poured, the greater the erosion potential. Washing type of erosion is produced where a great amount of metal runs over the sand surface.

Savage and Middleton, in the same paper, showed the effect of metal poured weight on erosion. Figure 3 shows the effect of weight poured and also the bulk density of the sand used in the test. The casting was a plate 4” x 4” and 5/16” thick with a catch basin for the metal at the end of the plate.

One investigator concluded that irrespective of the type of sand that was used, resistance to impact erosion was lower than the resistance to washing erosion.

Molding Sand

The composition of the base molding sand does affect the amount of defects produced. The higher density sands, such as Zircon, Olivine and Chromite reduce the defect. It is probably due to their lower expansion rate and finer grain size rather than their weight. Finer sands do reduce erosion due to smaller pore size.

Because new mixes of sand tend to reduce erosion, the practice of using new sand facings in problem steel castings has been used for years. The deterioration of the bentonite with repeated heating and cooling is listed as a major cause of erosion in system sands.

Increased clay content, to a point, will decrease erosion tendency, due to a stronger bond between sand grains. Western bentonite is more resistant to erosion than is Southern bentonite.

Increased moisture in the sand increases the erosion tendency. The increased moisture gives more steam production which gives more turbulence to the metal flow. This is very important if the sand has become finer for any reason. If the fines have built up in the sand system the permeability will be lowered dramatically and any increase in steam production can

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become a problem with a very tight sand. Moisture, above the optimum amount, gives a sand that is easier to ram and higher densities are possible. High density tends to lower the erosionresistance, but this increase can in some cases cause expansion problems.

Completeness of mulling is a very important factor in developing the maximum strength out of the bond available and plays a very important part in resistance to erosion. Poor binder cov- erage in poorly mulled sand gives poor strength development. Without uniformly distributed bond there are weak spots in the mold. Yearly (2) pointed out the necessity of good clay distribution. Figure 4 shows the same base sand with the same amount of clay and water, but with different degrees of mulling. In the well mulled sand density seems to be a minor factor in minimizing the erosion. The success of green sand relies on maintaining a permeability high enough to allow the free escape of gases formed at the metal-mold interface. There is an indication that coarse sands at low densities will be more prone to erosion even though an adequate amount of clay is present.

Pore size affects erosion resistance. A number of investigators consider pore size to be the main variable in erosion control. Finer sand produces smaller pores in the rammed sand and a lower permeability. With smaller pores and a smoother surface the metal does not have as much of a chance to wash away the mold face. Table 1 shows data on the use of various fillers to combat erosion effects. The increase of siliceous fireclay to 30% did not help the resistance and actually slightly increased the erosion.

110 111

Dextrine and starch additives help in erosion resistance because they delay the beginning of expansion scabbing and therefore the mold can be rammed tighter.

Core Sand

Excessive moisture in or on the core can cause a core blow which can cause an erosion scab. Core wash not properly dried or wash that did not have sufficient penetration into the core can cause erosion to take place. Undercured cores, with less than maximum strength developed, will produce wash defects. The core sand needs proper amount of binder plus catalyst and correct mixing cycle to give maximum erosionresistance. Therefore good maintenance of mixing equipment impacts on casting quality.

Mold Density and Mold Coatings

The higher the bond strength, the higher the resistance to erosion. Increased strength can come from more clay and water. more complete mulling or from better ramming. The molding machine must be in good shape to be able to apply the same energy time after time to the sand. Uneven or soft ramming en- courages erosion. Resistance to erosion increases as the mold density is increased to a point, than expansion scabbing will increase because of excessive density. The critical point for scabbing is at a lower bulk density with calcium bentonite than with sodium bentonite, according to Savage and Middle- ton (1). Greater mold density will increase erosion resistance due to smaller voids and stronger bond, both at room and high temperature.

Mold washes have reduced erosion defects but are difficult to apply in the most important area: the sprue and gating system. Much of the erosion takes place in the gating system and the sprue is often impossible to coat with wash. The use of ceramic tile is encouraged where metal impingement will affect the gating system.

Time until pour is important in both green sand and dry sand molding. When the surface of a mold dries to a depth of a few

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sand grains only, it is more friable than that of either green or fully dried molds. If cores are set in this type of sand, grains can be dislodged from the surface and these grains will be washed away into the mold.

EXPANSION DEFECTS

Definitions and Other Considerations

The expansion of the silica grain is the cause of the expansion type of defects. Figure 5 shows the expansion curves of several refractories. The large and sudden expansion of the quartz grain causes the mold to expand. Radiant heat converts the moisture into steam and the steam moves away from the cope surface. Further back the steam reaches a cooler part of the mold, condenses and forms a wet layer of lower strength.

The sand grains on the surface expand due to the heat. One of three things will happen: 1) If the sand grains in the dry layer

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can move closer together as they expand, no expansion defects occur. Soft ramming and the use of buffers that burn out, help to reduce the expansion. 2) If the expanding dry layer can compress the surrounding sand there will be no defect. 3) If the dry layer cannot expand freely it will increase in length by bulging downward from the cope face.

A rat-tail is a long, narrow furrow or a small step occurring on the casting surface as a result of sand expansion and very minor buckling of the mold surface. Rat-tails are the first step in expansion defects. A buckle, a depression on the surface of a casting, caused by expansion of the sand mold and pronounced buckling of the mold surface is the second step in these types of defects. A scab, a raised rough area that usually consists of a crust of metal covering a layer of sand or a raised rough area of essentially solid metal, is the third step. Cope spall or complete failure of the cope surface is the final stage in expansion defects.

In B. C. Yearly’s 1964 article (3), he mentioned that it wasn’t until 1933 that the literature contained any references to expansion as a possible cause of scabs and rat-tails. That year Nilsson (4) suggested that scabs were caused by the expansion of the silica grain and that it could be overcome by the use of wood flour or other combustible fillers.

Scabs, buckles and rat-tails occur when the mold surface is unable to accommodate the expansion of the sand on heating by the metal poured. Rat-tails and buckles are the same type of defects in an incipient form which did not have time to develop into scabs before the mold was filled and the casting solidified. Scabs and buckles usually occur on the cope surface or on flat vertical surfaces. Rat-tails occur primarily on the drag surface.

Rat-tails are usually narrow depressions or grooves on the casting surface and sometimes are so small that they show up as faint irregular lines. Rat-tails occur through heat conduc- tion, whereas buckles and scabs develop primarily from the effects of heat radiation. Expansion defects never occur after the mold is filled, they always occur during the filling of the mold.

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Two types of expansion scabs can form on the surface of a casting. The first scab consists of a crust of metal that forms when molten metal flows into the cavity created when the surface layers of the mold have buckled, partly separated from the mold and then cracked. The buckled surface stays intact and the metal can be pried off of the surface to show the original buckle. The other type of scab is essentially solid metal that has taken the place of sand that has spalled off. Ex- pansion scabs occur primarily on flat cope surfaces or high up on vertical walls.

Mechanisms of Expansion Defects

A number of mechanisms have been proposed. The mechanisms are presented in the Steel Founders’ Society of America Special Report No. 2-Titled “Scabs, Buckle and Rat-tail Defects on Steel Castings”. Rather than present the total discussion on mechanisms, the reader is referred to this long report. The major mechanisms presented are:

1. The Sand Expansion Mechanism of Buckles and

2. The Wet-Layer Mechanism of Buckle and Scab

3. The Wet-Layer Mechanism of Rat-Tail Formation 4. The Water Evaporation Mechanism of Rat-Tail,

Scabs Formation

Formation

Buckle and Scab Formation

The wet-layer mechanism, as presented by Van Eegham (5), is shown in Figure 6. He pointed out the influence of the hot-wet tensile strength in the condensation zone. Patterson, Boenish and Gable (6) showed that the scab area of a test casting was directly proportional to the compressive strength and the distance of the condensation zone from the mold interface and inversely proportional to hot-wet tensile or the tensile strength of the condensation zone.

Figure 7 shows the schematic representation of the scab defect in both restrained and unrestrained conditions. In the unrestrained condition, if the expansion is not great enough to totally loosen the cope sand a vein-like defect will form around the edge of the cope sand. Investigators have concluded that

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the strength of the hot surface layer and the strength of the mold is roughly 10 times that of the condensation layer.

Causes of Expansion Defects

The variables that can cause expansion defects are:

Casting and Pattern Design and Flask Equipment Pouring

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Gating System Molding Sand Core Sand Mold Density and Mold Coatings

Casting and Pattern Design

Large interrupted flat surfaces, inadequate radii in the fillets and large smooth uninterrupted surfaces increase the potential for expansion defects. Many consider casting design to be a major variable in cause of expansion defects.

Large castings or castings that fill slowly give prolonged periods of exposure of the cope surface to radiant heat. In smaller castings the problem is usually caused by intense local heating for short times. Inadequate fillets, bars or flask too close to the pattern surface increase the potential of expansion defects. Uneven ramming, due to bars or nearness of pattern to the flask wall increases expansion problems.

Gating System

Gates that prevent adequate pouring speed and/or cause interrupted metal flow will increase expansion defects, as will a gating system that causes nonuniform heating of the sand surface. A greater number of ingates or the same number of enlarged ingates will allow the mold to fill faster.

Pouring

A high pouring temperature and long pouring time (slow filling of the mold) will increase scab and buckle formation. The energy transferred by radiation varies by the fourth power of the temperature of the radiating body, therefore pouring temperature is a very important item. With increasing pouring time the defect increases in severity and in size. It is generally agreed that changes in steel composition have little effect on expansion defects.

Moulding Sand

Silica sand will be more prone to expansion defects than chromite or olivine due to its high expansion rate. Clay bonds

contract at elevated temperature and help to reduce buckling and scabbing. An increase in the amount of clay increases the contraction and reduces the tendency to these defects. West- ern (sodium) bentonite gives a greater resistance to scabs and buckles than will Southern (calcium) bentonite. Western bentonite that is chemically altered or treated is more prone to expansion defects, while activated calcium bentonite seems to approach sodium bentonite in expansion defect resistance. The dry and hot strengths are increased with increased clay con tent.

Wet tensile strength testing is a good measure of expansion defect resistance. Van Eeghem (5) found that the main difference between scabbing sands and non-scabbing sands was

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the higher wet tensile of the latter. Patterson and Boenisch (6) concluded that a desirable bentonite is one that provides the greatest possible increase in wet tensile strength and the least increase in compressive stresses under heat.

The bentonite added to the sand not only adds adhesive properties, but it also influences the density of the rammed mold and thereby the expansion characteristics. Western bentonite by itself is too water sensitive to be used as the only bond in steel sands. Figure 8 shows that small additions of water move the sand from the scab-free area (lower left hand corner) to one in which scabs formed (upper right hand corner). Figure 9 shows that cereal additions to Western bentonite bonded sand increases the range of water before scabbing is en-

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countered. The higher the bentonite content the greater the effect of the cereal addition. These materials add cushioning material to the sand mix. High moisture promotes the defects Dried sand molds have a much higher resistance to the defects than do green sand molds.

Water addition above temper point is considered to be a univer. sal cause of expansion defects. A high moisture content makes the sand easier to ram and the mold hardness and the rammed density gets to a point where the hot strength and the expansion are increased enough to cause expansion defects. High water in the clay causes high clay shrinkage, which should aid in reducing expansion, but high water clay sands have low shear strengths which allows high compacted densities.

New sand additions must be maintained to control the amount of impurities in the sand. Reclaimed sand can be used in place of new sand to lower the cost of new sand.

A high percentage of the individual grain must be properly coated with the bonding clay to minimize buckles and scabs. Good mulling is necessary to achieve this state of clay disper- sion. Only the clay that is between the sand grains at their points of contact will produce strength and aid in expansion control.

Core Sand

Non-uniformity in density of the core and under-cured cores will increase expansion potential. Improper core wash application, poor penetration of the wash into the core, improper composition of the wash and improper drying of core wash all will aggrevate core wash scab defects. A poor grain distribution that gives very high core densities will cause expansion defects.

Mold Density and Mold Coatings

Excessive ramming to high mold densities promotes scab and buckle formation as does non-uniform ramming. Gaggers, bars or flasks too close to the mold surface will cause expansion defects.

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Many investigators have shown that the severity of expansion defects can be reduced by reducing the mold hardness. High pressure molding is used to make a mold that is dimensionally accurate, a mold that is rammed to a high density and to give a sound casting. Any defects encountered in the molding method must be overcome by additions to the sand.

Many molds in steel foundries are washed with a mold coating. An optimum wash density for any given sand and rammed density is necessary to get the best possible casting results. Insuf- ficient penetration of the mold wash, excessive mold wash or improper drying of the mold wash will cause problems.

CLOSING COMMENTS

Both of these types of defects are aggravated by the increased temperature encountered in steel castings. The steps suggested to rid steel castings of these defects will have to be applied in a different manner in each foundry. One point that is important in all investigations is that the sand must be controlled and consistent and it must be molded in a consistent and controlled manner. Control in sand, as in the metallurgy of the casting, is important. To do sand testing is a waste of time if the data is not used to control the system and the method and to predict what the casting results will be.

The two sets of defects show the need for compromise in all aspects of foundry sand. The mold must be made tight enough to resist the erosive forces of a steel metal stream, and it must not be rammed to such a high density that expansion defects will appear.

The bibliography at the end of the paper is a summary of some of the more important papers presented on the subject. The list is not complete by any means. The two SFSA special reports list 116 articles on the subject, and that is not a complete list of everything presented on the subject. Only articles in English and readily available are listed.

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BIBLlOGRAPHY

1. Savage, J. and Middleton, J.M., "Final Report on Mold Erosion", Journal British Steel Casting Res. Assoc., No. 68, 17-26, August, 1962.

2. Yearly, B.C., "The Importance of Mold Density and Mulling Efficiency to the Success of the Green Sand Process", SFSA, Steel Foundry Facts, No. 290, pp 2-11, January, 1969.

3. Yearly, B.C., "Influence of Clay and Water on Expansion Defects Part I",Foundry, Vol. , No. 11, pp 73-77, November, 1964.

4. Nilsson, S.C.V., "Temperature Aftects Molding Sand", The Foundry, Vol. 61, pp 10-13, March, 1933.

5. Van Eegham, J.R., "Wet Tensile Strength in the Condensation Zone-Its Relation to Scabbing Tendency of Synthetic and Natural Bonded Sands", Modern Castings, Vol. 54, No. 12, pp 490-496, December, 1968.

6. Patterson, W., and Boenisch, D., "The Scab Diagram for Green Sands". Giesserei, October, 1964.

7. Vingas, G.J. and Zrimsek, A.H., "Steel Foundry Green Facing Sands", AFS Transactions, Vol. 71, pp 50-74, 1963.

8. Boenisch, D. and Patterson, W., "Discussion of Scabbing Tendencies of Green Sand", Modern Casting, Vol. 50, No. 10, pp 94-108, October, 1966.

9. "Eroded-Sand Defects in Steel Castings", SFSA, Special Report #5, September, 1970.

10. Graham, A.L., Dietert, H.W. and Rowell, V.M., "Mold Surface Failure: Where Defects Begin", Modern Castings, Vol. 30, No. 5, pp 36-46, November, 1956.

11. Hofmann, F., "Causes of Sand Expansion Defects and Mold Cavity Enlargement with Clay Bonded Molding Sands", AFS, Cast Metals Research Journal, Vol. 2, No. 4, pp 153-165, December, 1966.

12. Kennedy, V., "The Wet Strength of Green Silica Sand in Relation to Scab- bing and Sand Erosion Defects", Jour. Steel Research, SFSA, No. 39, pp 2-10, June, 1967.

13. Levelink, H., and van den Berg, H., "Green Sand Scabbing Tendency Testing by Shock Heating", AFS Transactions, 1962.

14. Locke, C., and Berger, M.J., "The Flow of Molten Steel in Sand Molds", SFSA Research Report No. 26, January, 1951.

15. Nicholaides, P. and Rose, R., "Study of Mold-Metal Reactions", SFSA, Journal Steel Casting Research, No. 32, pp 3-8, January, 1964.

16. "Scab, Buckle and Rat-Tail Defects in Steel Castings", SFSA, Special Report #2, April, 1970.

17. Yearly, B.C., "How to Stop Mold Wall Erosion", Foundry, Vol. , No. 8, pp

18. Szreniawski, J., "Effect of a Metal Stream on the Inner Surface of a Sand Mold", Foundry Trade Journal, Vol. 106, pp 323-330, March 19, 1959.

19. Middleton, J.N., "Factors Affecting the Surface Quality of Steel Castings", British Foundryman, Vol. 57, pp 1-19, January, 1964.

20. Parkes, W.B., "Sand Control with Particular Reference to the Prevention of Scabbing", AFS Transactions, Vol. 60, pp 23-37, 1952.

21. Wallace, B.T., "Mold-Metal Interface Studies", Steel Casting Institute of Canada, Research Report No. 11, 1953.

77-81, August, 1966.

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Lecture IV

Linear Surface Discontinuities by Dr. W. J. Jackson

It is a well-known fact that all engineering materials are imper- fect, the reason being that either perfection is technically impossible to achieve, or it is economically not feasible to achieve. During manufacture of an engineering component therefore, a degree of control is exercised in order that a compromise can be made between the degree of perfection required (fitness for purpose) and the cost of manufacture. Steel castings are no exception, and it is acknowledged that imperfections and discontinuities which occur during manufacture can be controlled to a greater or lesser extent, which will depend upon the requirements of the casting when placed in service. The determination of the latter requirements is largely the domain of the engineering designer, while the control of discontinuities to the required level is the responsibility of the steel founder.

Discontinuities may be embedded within the steel and therefore cannot be seen; they may also be surface breaking and can be seen, if not unaided by eye, then with the assistance of an appropriate non-destructive testing method. Surface discontinuities, as well as differing in size, can differ in shape; they can be round or elongated, the latter usually being referred to as linear. One definition of a linear discontinuity is one having a length at least three times greater than its width. It is this type of linear surface discontinuity that is the subject of this chapter.

Linear surface discontinuities do not arise from a single cause; there are several, and those encountered during investigations over the last 25 years at the Steel Castings Research

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and Trade Association are shown in Fig. 1. Their origins and means of controlling them are discussed under appropriate headings, concluding with a summary of their effect upon the performance of a casting in service.

SOLIDIFICATION AND UNDER-RISER EFFECTS

Under-riser discontinuities are usually found at the riser pad by NDT methods after burning off the feeder heads. Typically, the discontinuities may extend into the casting surface where attempts to remove them by grinding or arc washing may resuIt in further propagation.

Many factors contribute to the incidence and severity of under- riser discontinuities, including composition (especially carbon, sulphur and phosphorus), temperature of the casting during riser removal (preheat), riser diameter, cutting process heat input, shrinkage or segregation from the riser extending

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into the pad, contamination from riser feeding aids, and the heat treated casting condition prior to head removal.

These causative factors may be generalised as compositional and stressing effects. The fact that under-riser discontinuities occur after burning and rarely, if ever, when saw cut, suggests that thermal stresses introduced via the cutting operation contribute to their development. These tensile stresses and strains developed during cooling after burning are accommodated within the riser pad heat affected zone (HAZ). Accord- ing to the hardenability response of the pad, the transformed HAZ matrix may be of insufficient strength and ductility to withstand the imposed thermal stresses and strains. Discon- tinuities may extend for considerable distances beneath the hardened martensitic HAZ formed, the propagation resulting from tensile stresses developed as the pad cools.

Under-riser discontinuities are liable to form, being aided by any additional embrittling or weakening effects arising from,

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for instance, segregates and intergranular precipitates. In fact, riser segregation may be considered a major source of discontinuities during burning operations in carbon and lowalloy steels which would not normally exhibit them (Fig. 2).

Segregation

Several mechanical, chemical and thermal factors influence the degree, direction and extent of segregation in steel castings including the solidification rate and temperature in terval, convectional effects, pouring temperature and section thickness. (1) Typically, during the initial stages of solidification, low carbon dendrites form whilst the remaining liquid metal enriches in carbon and high carbon liquid is drawn into the casting by the vacuum caused by solidification contraction. (2) The degree of this carbon segregation varies according to the carbon level, with low carbon steels being less prone to carbon segregation due to their shorter solidification inter- vals. In large low carbon steel castings, heavy segregation of carbon, sulphur and phosphorus near the riser pad centre is a major factor increasing under riser cracking susceptibility. For instance, carbon segregation in a 0.2%C steel as high as 0.6 to 0.85% at the riser pad centre can occur. (3,4) Underhead segregation of aluminium, silicon and boron will also occur. (4)

The segregating tendency of alloys, residuals and impurities increases proportionally to the carbon segregation, increasing markedly for the elements classified as impurities or residuals (e.g. sulphur, nitrogen, tin, arsenic, phosphorus) compared with the normal alloys (e.g. molybdenum, manganese, silicon, nickel, chromium). (5)

Increased tendency to hot tearing adjacent to the riser-casting contact area has been noted in medium carbon-manganese steels containing 1.25% or more copper. Lead and tellurium, as well as copper may give rise to the related phenomena Of hot shortness and liquid metal embrittlement. (6)

Shrinkage

Gross shrinkage should be avoided by the provision of adequate feeder heads. (7) Microshrinkage, however, can arise

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unexpectedly in regions that are otherwise adequately fed. It can occur in the interfacial regions of dendrite grains without any obvious centre of nucleation (Fig. 3), or if adjacent to a mould wall it can be nucleated by gas or more commonly, by microscopic particles of sand or slag which act as “mini hot- spots”. A plane of interdendritic shrinkage breaking out to the surface can give NDT indications (MPI, LPI) having the appearance, and form of cracks (Fig. 4). An area exhibiting dendritic shrinkage in an 18Cr-8Ni steel casting is shown breaking out to the surface in Fig. 5.

Areas of macrosegregation, being the last to freeze and hence cut off from an liquid feed metal, will invariably contain microshrinkage (Fig. 6). In under-riser regions that show cracks after burning off and grinding the pad, it is always possible, therefore, that some of the cracks are microshrinks, or were originally present as microshrinks and were propagated by thermal stressing.

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To guard against under-riser microshrinkage is difficult, as the solidification phenomenon is complex and to the practical foundryman, unpredictable. One precautionary measure, however, is to avoid making the riser too tall in relation to its cross-section, otherwise secondary shrinkage will occur, penetrating into the body of the casting.

Surface microshrinkage and sub-surface microshrinkage (revealed by grinding or scaling off during heat treatment) in areas remote from feeder heads can often be positively identified by breaking open the alleged crack and observing the tips of the dendrite growth (Fig. 7). It should be mentioned that such an interdendritic shrinkage cavity is frequently the source of a hot tear, which can remain as such or be “healed” by an inflow of liquid metal (as mentioned in the section on hot tearing). Surface initiated microshrinks are also difficult to guard against, but in general the surface adjoining a riser should itself slope upwards to the riser, thereby facilitating the passage of any potentially nucleating gas or non-metallic particles into the riser.

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Very often the distinction between microshrinkage and hot tearing is difficult to make. When it is, the remedy is the same-improve feeding and uni-directional solidification towards the riser.

Risering Effects The segregating tendency of carbon, sulphur and phosphorus increases with riser diameter, with the amount of vertical

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metal movement, and contamination by feeding aids, but is reduced by adding molten steel of the initial or lower carbon content after pouring. Since some segregation is invariably present near the cut, it is recommended that casting stresses be reduced by heat treatment prior to head removal. (8)

The cost savings incurred by use of necked-down risers compared to full section risers results mainly from reduced riser cutting and pad grinding operations. Riser size selection from the viewpont of effective feeding and riser efficiency is not related to the use of necked-down risers. (9) In other words, a riser that is designed for maximum feeding efficiency cannot be made more efficient by adopting the necked-down principle.

Detailed analytical results for necked-down and full section risers are given elsewhere. (10)

Many variations in riser feeding aids like hot topping compounds, insulating and exothermic sleeves can be encountered. For instance, aluminium is usually incorporated as a fuel in exothermic materials, but its source and purity may vary. (11) Also, filler materials and insulating additives employed in exothermic and insulating sleeves offer additional variables. Binders for insulating materials may be any variety of organic or inorganic materials. Fillers may consist of slag wool, alumina, or various types of clay. Iron oxide, barium nitrate and manganese oxide are common oxidisers and variations in these may occur, depending on formulation. As such, riser feeding aid contaminants may comprise carbon, barium, aluminium, manganese, silicon and boron.

Control of Under-Riser Effects

Too shallow a feeding head may result in the primary segregation, which forms under the shrinkage cavity in the head, being sucked into the casting during the final stages of solidification. (12) Practical methods of reducing segregation are to keep pouring temperatures as low as possible and to reduce the amount of harmful impurities (e.g. phosphorus, sulphur). (1 3)

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In general, riser removal methods comprise mechanical (sawing, drilling), electrical (arc, plasma) and thermal (flame, powder, lancing) methods. (14) With regard to UK steelfounders, a detailed survey indicated that oxyacetylene or oxy- propane methods were used for removing risers up to 24in diameter in most foundry grades of steel. (3) With larger risers (especially in excess of 36in diameter), oxygen lancing, drilling and burning methods were employed or, in extreme cases, sawing. (Risers on stainless grades were removed by powder cutting.) It was found that:

(1) surface discontinuities did not usually occur in carbon steel castings with less than 0.35%C and riser sizes less than 24in-these risers could be burnt off while they were cold and in the as-cast condition.

(2) in heavy castings in low carbon steel with riser sizes of 30 to 72in, and in alloy steels and carbon steels with more than 0.4%C, discontinuities did not occur unless precautions were taken, e.g.

(a) preheat before burning, or burn while still hot from knock-out in the mould.

(b) anneal or normalise and burn off while still hot

(c) drill larger (e.g. 72in) risers, followed by anneal and

(d) after burning; transfer casting immediately to furnace

(150-300°C) or cold.

burn off while still hot (>250°C).

for annealing.

Another approach by a UK steelfoundry was to relate the under-riser cracking susceptibility in terms of the head diameter and carbon equivalent (CE). (15) In essence, the CE merely expresses the hardenability response of the flame cut riser HAZ, being dependent on compositional and thermal (heat input, riser site) factors. The resulting nomogram shown in Fig. 8 provides selection of preheat temperature and prior heat treatment to be adopted prior riser removal.

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This procedure has proved successful in practice, although it is noted that defects like underhead shrinkage may act as stress raisers and so alter the control conditions. In accordance with the CE concept, cracking increases with increasing carbon, manganese, nickel, chromium, molybdenum, copper and with riser size, since this promotes higher cooling rates and increased segregation in the riser pad. The CE values should be the actual CE based on the probable analysis of the metal beneath the riser. This can, of course, significantly differ from the ladle analysis as a result of segregation during solidification and riser feeding aid contamination, as mentioned in the previous section.

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HOT TEARING

Hot tears probably still represent the most widespread and troublesome source of linear surface discontinuities in the steel foundry, The elimination or, at least, control of hot tears to within acceptable limits in carbon, low alloy and high alloy steel castings of both simple and complex design is essential in providing cost effective production castings. The literature on hot tearing is extensive, but, in many cases it is contradictory.

Hot tears are interdendritic ruptures that occur either during the final stages of solidification (above solidus) whilst there is still some interdendritic liquid, or immediately after solidification (below solidus) when low melting point segregates are present. The tears, invariably of jagged appearance, may be skin deep or extend through the entire section thickness, Fig.

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9. In some cases, they may be discontinuous and consist of a number of short, unconnected ruptures, or may be fully or partially healed by the residual melt (Fig. 10). As a result of the high temperatures and oxidising conditions that exist at the time of tear formation, the tear surface is invariably decarburised and oxidised (Fig. 11). Carbon segregating from the decarburised region appears as high carbon areas (Fig. 12).

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It is generally accepted (16-19) that hot tearing requires the simultaneous presence of:

(i) restraint, provided by hindrance of the casting thermal contraction due to mouldlcore or casting design.

(ii) hot spots, provided by section changes, junctions, gates or risers where stress concentration may occur.

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If the casting contracts freely or the casting is of uniform section and without hot spots, hot tearing is unlikely to occur. Some stress is necessary and can be mechanical, such as self-restraint by the action of the mould, or thermal as a result of temperature differents in the steel causing unequal contrac- tions. Some stress always remains in the casting, but thank- fully does not reach sufficient levels to cause cracking, or even distortion. But at the last stages of solidification, when there are only very small liquid channels separating the solid grains, only a minute stress is necessary to cause separation of the grains. When there is no liquid remaining, and the grains are at extremely high temperatures, and the steel is consequently of very low strength, small stresses can cause fracture; the crack is through solid steel and can be called a cold crack, even though the steel was very hot when the stress was relaxed by de-cohesion.

Sometimes, the rupture has been produced merely by shrinkage separating the grains, and stress has played no part. If surface breaking, the MPI or LPI indication on NDT surface examination will be elongated and have the appearance of a crack and will be logged as a crack or linear surface discontinuity. If the shrink is not surface breaking and is picked up by radiography, it will be logged as a hot tear, implying that stress played a part in separating the grains. The distinction is too fine in this case, to be able to answer the question “is it a tear (stress involved) or a shrinkage cavity (no stress involved)?” Whether stress is involved or not, improved feeding in that region is likely to prevent the defect re-occurring.

Two basic mechanisms purport to explain hot tear formation, one concentrating on stress-strain effects whilst the other em- phasises interdendritic liquid film surface tension effects. The stress-strain mechanism (21-23) envisages that the hot zone approaches a liquid film stage at temperatures near the solidus in which, under conditions of restraint, tensile stresses or strains develop. Separation then commences through these liquid films at low tensile stresses and when the strain in the adjacent solid regions exceed a critical value, metal separation occurs by tearing of the solid steel, and a hot tear develops. Variants of the mechanism claim that the tears

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either initiate at the extreme surface of the casting (20,21) or at a small distance beneath the surface. (22,23)

The surface tension mechanism (24) proposes that under conditions of restraint, the liquid films separating the dendrites in hot spots are subjected to tensile loads. If the critical tensile stress is exceeded, the meniscus of the liquid film at the mould face, maintained by ferrostatic pressure and capillary action, recedes from the casting surface and a hot tear develops.

Compositional Effects

Carbon and Low Alloy Steels

Phosphorus and Sulphur-Steelfounders are unanimous that phosphorus and sulphur in the steel melt have a deleterious effect on hot tear tendency. Both phosphorus and sulphur tend to segregate and promote low melting point films (17), thereby providing-a longer temperature range in which a contracting casting can build up the stresses to cause hot tearing. Analysis of a healed hot tear in a 0.05%S, 0.05%P steel revealed 0.3%S and 0.4%P, indicating that there was substantial phosphorus and sulphur enrichment of the residual melt. (22) There is some evidence to suggest that the effects of phosophorus and sulphur are additive, and that increasing the phosphorus level intensifies the effect of sulphur. German work indicates that phosphorus and sulphur levels must be held as low as possible, preferably less than 0.03% each. Moreover, combined phosphorus and sulphur levels of more than 0.07% with normal deoxidation and melting procedures are very likely to cause hot tearing. (25)

Carbon-Steelfounders are not unanimous in their opinions on the effect of carbon on hot tearing susceptibility. Typically, within the 0.02 to 1.0%C range, investigators report that increasing carbon levels either increases (23,26) or decreases (21,27) hot tear tendency where, in some cases, specific carbon ranges prove more or less tear susceptible. (23,27,28) It may be expected that high carbon steels with their wider solidification ranges should be more tear susceptible. (1) Fur-

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thermore, steels containing 0.16 to 0.23%C have a narrower solidification range, and at least 20% liquid solidifies at constant temperature, i.e. at the peritectic transformation. Thus, during the final solidification stages, such steels contain approximately 80% solid in contact with 20% liquid and, if tearing occurs, there will be a greater opportunity for the residual melt to heal the tear. It is evident, however, that these considerations are not borne out by the results of technical literature, where the effect of variation in carbon content is highly contradictory. It seems safe to conclude that variations in carbon content have only a minor effect on hot tearing propensity.

Manganese-Tearing susceptibility appears to be related to the Mn/S ratio rather than Mn alone, with higher Mn/S ratios reducing the hot tear tendency. (26) With insufficient manganese, sulphur can precipitate in the form of low melting point FeS or Fe-FeS eutectic films, thereby causing a widening of the solidification range and a greater tendency to hot tearing. (17) An optimum manganese content with various sulphur levels exists for optimum tear resistance. (16) With very low sulphur steels (i.e. ~0.005%), the manganese content has Iit- tle effect but, as sulphur increases, increasing amounts of manganese are required to give the best tear resistance, e.g. at least 28:1. (17,28)

Silicon-There is general agreement (17,26,29-31) that variations in silicon content between 0.1 to 1.1% have little or no effect on hot tear tendencies.

Tin and Copper-Both tin and copper may be present in carbon and low alloy steels as residuals from the scrap used in melting. It appears (17,27,29,30,32) that copper contents from 0.5 to 3% have little effect on hot tear propensity.

Other Elements-The limited data available on the effects of alloy additions on tearing in low alloy steels may be summarised as follows:

(i) nickel up to 3% has little effect on hot tearing (29,30);

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(ii) the following alloying elements appear to improve hot tear resistance (29): up to 3%Cr, up to 1%Mo, up to 0.25%Ti; up to 1.16%V;

(iii) the beneficial effects of titanium and vanadium are thought to be due to grain refinement, which minimises formation of continuous fiIms during solidification;

(iv) the effect of alloy combinations is unknown but it is generally considered (29) that such combinations have no significant effect other than that of the individual elements.

High Alloy Steels

In general, the high alloy steels are considered by steelfounders to be more free from hot tearing than carbon or low alloy steels. In fact, 18Cr-10Ni steels do not tear as severely as do carbon or 12%Cr steels. (33) Nevertheless, sulphur and phosphorus should be kept to low levels to obtain maximum tear resistance. (17) One recent investigation involving a 20Cr- 10Ni steel has highlighted the importance of a low pouring temperature and a balance of composition such as to ensure the absence of delta-ferrite. (34)

Deoxidation and Melting Practice

Deoxidisers that are added singly or in combination include aluminium, silicon, titanium, calcium, zirconium and rare earth metals, which combine with the oxygen to form oxides and, in some cases, also combine with sulphides. The type, composition and distribution of oxides and sulphides in cast steel varies with the amount and type of deoxidants, as well as the sulphur and oxygen contents of the steel. When the oxygen content is low and the sulphur level is high, the sulphur is retained in the last metal to solidfy and is finally precipitated as low melting point grain boundary films, with a consequent deleterious effect on hot tearing properties. (17)

In aluminium deoxidised steels, when the aluminium level is such that type II sulphides (Fig. 13) are formed, the suscep-

140

tibility to tearing is high. (21,29,35) At low sulphur levels (i.e. ~0.010%), the amount of type II sulphides will be small and discontinuous, and may not be expected to have any marked effect on tearing. With higher sulphur levels, the presence of type II sulphides may be deleterious as larger proportions of continuous low melting point films form during solidification.

It has also been recommended (25) that deoxidation be aimed at producing solid deoxidation products (using aluminium, zirconium, magnesium, or small additions of titanium; and that the use of calcium, barium and large additions of titanium be avoided, because liquid deoxidation products are formed.

It is doubtful if evolved gases have any real effect on hot tearing. (17) Possibly evolved gases, which react with constituents in the steel melt to form reaction products that are precipitated at selected sites (e.g. grain boundary precipitation of AIN, MnO in austenitic Mn steels), may affect hot tearing properties.

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Pouring Conditions and Gating/Risering Practice

It is generally accepted (23,29) that lower pouring temperatures and fast pouring rates reduce hot tear tendency by minimising temperature gradients and hot spot formation. Also, low pouring temperatures tend to promote finer equiaxed grains, which are considered to be less prone to tearing than steels with large columnar grains. (16,23) Low pouring temperatures as close to the liquidus as possible, prevent or considerably reduce tearing in production castings. (17,36)

Gating practice may also markedly affect temperature gradients in the casting, especially in the early stages after pouring when hot tearing occurs. (17) lngates are preferred hot tear sites since the adjacent mould absorbs much heat and thus delays solidification at this point. This condition may be severe when a single ingate and slow pouring are employed, since this combination produces an intense hot spot.

It is desirable to employ multiple ingates to permit the metal to enter the mould at as many points as possible (17) This practice minimises hot spots and promotes flatter, less severe temperature gradients in the casting. It should be stressed, however, that this may not be conducive to efficient feeding, where steep temperature gradients towards the feeder head are required. A compromise is thus required, and when possible, multiple ingates should be encouraged.

Adequate feeding is essential to prevent internal shrinkage and, at the same time, to heal any tearing. The use of excessively large feeder heads should be avoided since this increases hindrance to contraction, accentuates hot spots and aids sulphur and phosphorus macrosegregation. (37)

Mould and Core Material Effects

Since mouldlcore collapsibility characteristics within the hot tearing temperature range may determine the restraint offered to contraction of the casting, variations in mouldlcore density and binder type may contribute significantly to the incidence and severity of tearing. (38)

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Increased mould density, which often results in higher hot strength, is a major factor in increasing the restraint of the casting. (16) There is a good correlation between hot tear tendency (bore cracks) and core density with clay and oil bonded sands, (39) and mould density of CO2 process sands has a marked effect on tear severity. (40) Reduction of mould density, therefore, is one practical method of minimising hot tear tendency, and the rammed density should be no higher than is necessary to meet other requirements (e.g. metal penetration, sand erosion, swelling of the casting). The various types of clay bonded or naturally bonded sands have little effect on tearing propensity. (38) On the other hand, oil and some resin bonded sands generally produce much more tearing than clay bonded sands when the casting size is small. The reduced tearing tendency of larger castings with organic binders may be attributed to the thicker sections increasing the amount of heat available to burn away the organic bond, thereby permitting the sand to collapse.

Core sand mixes in particular may be very potent causes of hot tears, especially in thin section castings. Since organic binders maintain high rigidity shortly after pouring (e.9. when hot tearing occurs), it is important to employ a minimum amount of binder (oil or resin) consistent with adequate handling properties and resistance to erosion.

The resistance to contraction and tearing susceptibility increases through the following sand mixes: (37)

(i) green sand (ii) dry sand - clay bonded (iii) sodium silicate bonded (CO2 process and ester hardened) (iv) resin bonded shell sand (v) alkyd resin/oil (perborate or isocyanate types) bonded

sand (vi) oil bonded sand (vii) phenol formaldehyde resin/isocyanate (pep-set) sand (viii) furan resin bonded sand.

Thus, clay bonded sands are to be preferred over oil bonded sands for reducing mouldlcore restraint, with green sand

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moulds being capable of deforming more easily than dry sand moulds. At times, green sand cores or part green sand cores are employed to minimise tearing in difficult castings. Furan resins are not recommended for binding cores with difficult light section castings. Sodium silicate bonded sand cores and cores with a green sand top half, provide lower tearing propensity with thin section castings.

Mouldlmetal friction effects (e.g. flash, fins, burrs, metal penetration (should be controlled since these may have a considerable localised, restraining effect on the contraction of the casting. (1,25) Similarly, mouldlcore reinforcements (e.g. irons, metal rods, etc.) should be kept to a minimum since these can seriously reduce collapsibiIity. (17)

Chills and Exothermic Pads

Two methods are available for hot spot modification in order to prevent hot tearing. (17)

(i) application of chills, internal or external, so that the cooling of the potential hot spot is accelerated and the temperature in that region is equalised with that of adjacent sections.

(ii) enlarging the hot spot, by tapering or padding the section

In general, external chills are preferred to internal chills for preventing tears since internal chill performance is affected by metal temperature and can give inconsistent results.

To prevent internal shrinkage, the external chill thickness should be between half and full thickness of the section to be chilled, depending on whether the casting is of a bar or plate design. (17) For tear prevention, however, the chill may be thinner than this since, in many positions such as fillets, the function of the chill is merely to form a solid skin that is strong enough to withstand the stresses and strains imparted by the hindrance to contraction. Typically, the metal chill thickness need only be 1/16th of the thickness of the section, and even milder refractory chills are employed for this purpose.

to avoid heat concentration into one small section.

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Experience with bore cracking in valve bodies (41) has shown that external chills tapered to a feather edge (e.g. into the bore of the casting) prevented tearing, and that it is best to stagger the chills so that they do not terminate in the same circle. As a general rule, the chill length should be three to four times the section thickness to be chilled and the distance apart equivalent to the length of the chili. (1,37)

The use of brackets as mild chills (cooling fins) apparently often only displace the tearing from one location to another. (17,37) Although thin brackets are more effective, care is required in their application since oversize brackets can accentuate tearing by restricting casting contraction. Moreover, it is considered (37) that the use of brackets is not as effective as correct chiIIing procedures.

Casting Design

Design is of paramount importance and as a result of design modifications, the control of stress levels, stress concentration and hot spots together with adequate feeding practice can markedly improve hot tear resistance. A number of design rules have been formulated for both simple joining sections (42) and for stress active sections, (43) which are too extensive to be summarised here. The engineering designer should be made aware of the fact that some designs are impossible to reproduce satisfactorily.

Control of Hot Tearing

Perhaps the most detailed recommendations for the control of hot tear defects are those adopted by the French, (44) being based on steelmaking, foundry and design variables in the form of a decision diagram for carbon and low alloy steels. The diagram in Fig. 14 may be summarised as follows:

(a) Control of restraint. (i) Casting design-observe design rules for simple and

stress active sections. (ii) Mould/core material section-ensure good collap-

sibility characteristics at temperatures near to the Ii- quidus and keep mouldlcore reinforcements to a minimum.

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(iii) Mould/metal friction-avoid flash, fins, burrs, metal penetration.

(b) Control of hot spots.

(i) Gating/risering practice-use feeding systems that promote directional solidification. Ensure good feeding at potential hot spots. Employ multiple ingates to promote equalisation of metal temperatures.

(ii) Pouring conditions-pour as quickly as possible and as cool as possible.

(iii) ChiIls/exothermic pads-ensure proper use of chiIIs and pads to modify hot spots.

(c) Control of melting and deoxidation.

(i) Hold phosphorus and sulphur as low as possible, each should be less than 0.03%.

(ii) Combined phosphorus and sulphur levels should not exceed 0.05%.

(iii) Deoxidise the melt to obtain type Ill sulphides (generally 2lb Al/ton).

(iv) Selection should be made to obtain solid deoxidation products (i.e. aluminium, zirconium, cerium, small titanium additions).

(v) Deoxidisers such as calcium, barium and large titanium additions are not recommended owing to the formation of liquid deoxidation products.

(vi) Certain deoxidisers may inhibit excess phosphorus and sulphur levels.

(a) If P=0.03%, some benefit may be obtained by deoxidation with small amounts of titanium (e.g.

(b) If S=0.03%, some benefit may be obtained by deoxidation with large amounts of silicon (e.g.

0.07%).

0.25%).

(d) Control of cooling times.

(i) In the solidification range the skin at a hot spot should cool rapidly, the cooling rate of the skin should be above 3.5°C/sec.

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(ii) If the cooling requirement is not met, some benefit may be obtained by use of deoxidation procedures that give solid deoxidation products (e.g. zirconium, cerium).

(iii) The ratio of cooling times of the heavy sections to the adjacent thin sections must be as small as possible (a critical ratio is about 3.5)-if troublesome, modify with chills or exothermic pads.

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SURFACE BREAKING INCLUSIONS AND FILMS

Surface Breaking Inclusions

Non-metallic inclusions, whatever their origin, because of their lower density, tend to float to the top of the mould immediately after the casting has been poured. ideally, such inclusions should be swept into the feeder head but when the shape of the casting, and the location of the feeder when present, are unfavourable, entrapment will occur. Such inclusions, when present before the onset of solidification, will often be aligned with the subsequent dendritic pattern, but will not be seen in the as-cast surface. On machining, however, they will reveal themselves and on the application of a surface inspection technique will lead the operator to think that small surface linear discontinuities are present (Fig. 15). The ultimate effect on the service performance of the casting will be as stress raisers and could lead to failure of the casting.

148

Such inclusions could originate from the furnace, the ladle refractory, loose sand in the mould, etc. (45) but they are more likely to be of an indigenous source, arising from re-oxidation of the steel. In these instances, the products of oxidation are high in aluminium but frequently contain some silicon and sometimes manganese. Aluminium re-oxidation products usually take the form of a “skin” or “film” and this contributes to their linear appearance (Figs. 16 and 17).

Re-oxidation occurs at all times when steel is in the liquid state. With respect to deoxidants added to the ladle, it occurs during holding in the ladle, during pouring from ladle to mould, when the steel is in the runner system (often with entrained air) and when the steel is in contact with the mould cavity. Often the inclusions are oxides/silicates of aluminium or zirconium, depending on the additions made, and can be in any steel. For example, the alumina shown in Fig. 17 was in a plain carbon

149

150

steel having an analysis of 0.115%AI. The galaxies of alumina and needles shown in Fig. 18 were shown by EPMA to be mainly AI

2O

3 and AIN and occurred in CN-7M steel. The galax-

ies in Fig. 19 were shown by EPMA to be high in zirconium, the CN-7M steel having an analysis of 0.015%AI and 0.019%Zr. In Fig. 20, energy dispersive analysis showed the material in the discontinuity and interdendritic inclusions (the array of black dots) was essentially AI

2O

3 with some chromium oxide and

silicates.

In heavy castings made in furan sand moulds, sulphur can be picked up at the surface during solidification and takes the form of grain boundary (type II) sulphides (Fig, 21). Such intense outlining of grain boundaries with eutectic sulphides almost invariably leads to surface cracking, as does stringer- like phases of sulphide (Fig. 22).

151

152

Control of Surface-Breaking Inclusions

Re-oxidation of liquid steel is occurring at all times and the greater the exposure of surface to air, the greater will be the amount of re-oxidation products. Steps should be taken at all times to limit re-oxidation:

do not hold for excessive times in the ladle, with or without slag cover;

maintain a short pouring stream with a single and smooth flow, avoiding splashing or a ragged stream;

keep running systems as short as possible;

design and position the ingates for non-turbulent entry of steel into the mould cavity, to avoid swirling about during filling;

153

avoid excessive additions of strong deoxidants; if porosity is the problem, look at melting practice to minimise gas content of the melt;

in “difficult” steels, avoid strong deoxidants altogether and rely on calcium and silicon.

Surface Film Associated Discontinuities

In high alloy steel castings, a variety of surface film associated defects, sometimes occurring as laps or folds (Fig. 23) and sometimes as non-metallic films or inclusions (Fig. 24), are occasionally observed. The severity varies from smooth grooves to crack-like indications, as illustrated in Fig. 25. Judging the appearance by visual inspection can lead to erro- neous conclusions, for sub-surface oxide films can penetrate through the wall thickness (Fig. 26).

Beneath a lap a small area of microshrinkage is sometimes present (Fig. 27). Furthermore, a lap can sometimes be the result of cracking in a mould (46); the mould crack is reproduced as a vein on the surface of the casting, which, when ground off reveals by magnetic particle inspection a linear surface discontinuity.

154 155

156

In an investigation into the nature of surface films in high alloy steel castings carried out by the Steel Founders’ Society of America (47), it was found that surface films in both defect- containing and defect-free castings consisted of FeO, Cr

2O

3and Fe/Cr oxide mixtures, the latter also invariably containing fayalite (Fe

2SiO

4). The surface films within the defect itself,

however, were quite different. These comprised high melting point CrMnSi refractory oxides with FeO, Fe

3O4 and Fe2SiO

4 at

the defect mouth, suggesting a metal-surface oxide reaction during mould filling. At this stage, mould moisture level, pouring temperature, Mn/Si level and oxygen content were considered likely candidates and experiments were made in order to study these variables.

157

Results indicated that both burn-on and fluidity were governed by pouring temperature, with high pouring temperatures promoting long runs and considerable burn-on. The incidence of surface defects, however, was governed by Mn/Si levels; pouring temperature, mould moisture and gas content (oxygen and hydrogen) had no discernible effect.

More heats were made to study further the Mn/Si ratio and it was found that high Si-low Mn heats promoted castings with few, if any, folds and sub-surface oxide films. On the other hand, low Si-high Mn heats contained folds and sub-surface oxide films. Increasing the manganese level of low silicon heats produced an increased amount of sub-surface oxide film, even though it tended to improve the external casting appearance. High Si-high Mn heats, however, produced poor external surfaces, but less sub-surface oxide films than low silicon heats, regardless of the Mn level. High silicon heats produced castings in which folds tended to heal.

The heat composition in terms of the Mn/Si level, therefore, is the predominant factor governing surface defect formation, with high Si-low Mn heats promoting defect-free castings. Although pouring temperature (1510°, 1620°C), mould moisture (2.5-4.5%) and gas content (hydrogen and oxygen) have no discernible effect on defect formation, some mould moisture is required to produce the iron oxides. Under highly oxidising conditions, it is believed that silicon contributes appreciable oxidation resistance such that with high silicon levels, protective SiO

2 or Cr2O3

films form, which prevent iron oxide formation. On the other hand, high manganese levels promote a non- protective oxide film, where continued oxidation permits the formation of detrimental iron oxides which lead to the formation of surface film associated defects.

Control of Surface Film Associated Discontinuities

Manganese is not really needed in high alloy steel castings since it is not an effective deoxidiser and is not needed to tie up sulphur, whilst low manganese levels are not considered to have a detrimental effect on mechanical properties. (47) Moreover, the recommendations indicate that if silicon levels are held between 1.0 and 1.5% and manganese levels held to a

158

maximum of 0.1%, high quality surfaces and freedom from sub-surface films result. Reduction in manganese, however, depends upon both the percentages and manganese levels of the stainless scrap and returns used for melting.

CHICKEN-WIRE CRACKING

This type of surface discontinuity is usually associated with nickel-bearing low-alloy steels as a variable network of fine, shallow, crack-like indications on the casting surface, which in some cases, can only be detected by magnetic particle inspection. The discontinuities are commonly called chicken-

159

wire cracks, but they are also known as network cracks, or craze cracking. They can occur all over the surface of a casting (Fig. 28) or in certain regions only (Fig. 29). The network can be so coarse that the defect can sometimes appear to be a series of linear surface discontinuities (similar to Fig. 30). The incidence and severity of chicken-wire cracking appears to bear little relation to steelmaking, foundry and design variables.

160 161

Mould-Metal Reactions

Perhaps the most important cause of poor surface quality, burn-on and metal penetration in steel castings is fayalite (FeSiO

3) which occurs as a result of mould-metal reactions

under prevailing oxidising conditions as follows (48)

The molten FeSiO, formed by wetting and dissolving the silica sand mould face, enlarges the openings between the sand grain interstices thereby promoting metal penetration, whilst any residual FeO/FeSiO

3 remaining attached to the solidifying sur-

face produces the burn-on appearance. Under certain conditions, attack of the solidifying/cooling steel surface by mould- metal reaction products may occur giving rise to crazing.

Olivine, chromite and zircon sands undergo similar mould- metal reactions, although the reaction products may be somewhat different and the extent of the reaction may vary. (49) In fact, mainly FeSiO, for silica and zircon sands, FeO for olivine and chromite sands. (50)

The fayalite reaction during pouring of pure iron into sand moulds at 1650°C in oxidising environments may be summarised as follows (48):

(i) the solidified iron shell at the mould-metal interface forms a liquid oxide comprising essentially FeO, taking place at the solidification temperature of iron at 1380°C.

(ii) this liquid FeO may either (a) oxidise further to Fe

3O

4-this occurs only if oxygen

can reach the surface faster than it can diffuse through the FeO into the metallic iron.

(b) react with the excess SiO present-this occurs at temperatures below 1210°C to form fayalite (FeSiO

3).

Elements dissolved in steel may modify these mould-metal reactions. In particular, the more oxidisable elements like carbon, silicon and manganese will reduce the oxygen available

162

for iron oxidation, being depleted from the solidified skin to an extent dependent upon time, temperature and oxygen partial pressure. An increase in carbon level retards fayalite formation since carbon in the skin layer is oxidised before iron, with silicon and manganese having a similar but lesser effect, (48) i.e. in oxidising mould environments, elements with a high affinity for oxygen (e.g. carbon, manganese, silicon), (a) pass out of the metal to form an oxide and (b) react with the sand. In particular, low melting point silicates form with manganese bearing steels in silica sands, whilst the chemically less reactive sands (e.g. chromite, olivine) react weakly with the metal oxides. (49)

With regard to sand additives, bentonite (which consists of complex silica-alumina compounds and other oxides) tends to promote fayalite formation. (48) At steel pouring temperatures, the bentonite is liquid and provides a medium for solution of liquid iron oxide. The sand grains at the mould-metal interface become coated with liquid bentonite which aids the iron oxide to flow around the individual sand grains. The FeO-SiO, reaction is initiated around these sand grains, the bentonite promoting fayalite formation by dissolving the iron oxide quickly and increasing its fluidity, the melt forming precipitating fayalite on cooling.

On the other hand, organic binders which burn out at lower temperatures will have little effect on the basic reaction between silica, steel and oxygen, except those binders that liberate reducing gases. This retards fayalite formation since reducing gases inhibit iron oxidation at the skin surface. In fact, heating Armco iron and 0.8%C steel pellets embedded in silica sand for 1h at 1400°C in either reducing (H) or inert (N) atmospheres will result in no mould-metal reaction; with oxidised Armco iron and 0.8%C steel pellets, fayalite will form when treated in an inert atmosphere.

The effects of mould variables were studied, using clay and silicate bonded sands containing 8% manganese ore, 50% manganese ore and sand bonded with shale tar. (51) Appar- ently, the reducing atmospheres resulting from the pyrolysis of shale tar suppressed crazing, the weakly oxidising atmo-

163

sphere with 8% manganese ore made crazing a risk, whilst 50% manganese ore made it a certainty on the majority of castings. Interestingly, the pouring of Ni-Cr-Mo heats into green and dry sand moulds revealed no evidence that mould moisture aggravated crazing tendencies.

Mouldwash-Metal Reactions

Reactions of liquid steel with three types of mouldwash have been investigated in the USA (52, 53):

(i) neutral (inert) mould washes-these provide an inert film barrier, thereby preventing the FeO-SiO

2 reaction.

(ii) reducing mould washes-these provide a reducing atmosphere, thereby preventing FeO formation by reacting with the available oxygen.

(iii) oxidising (reducible) mould washes-these provide an oxidising atmosphere which oxidises FeO to higher, less reactive iron oxides and/or provides a residual metallic film at the interface. All three mould washes provided some improvement in the as-cast surface finish compared to uncoated moulds, better surfaces being generally obtained with the reducing or neutral mould washes. Par- ticularly poor results, however, were obtained with some of the oxidising (reducible) mould washes. (52) Hexa- chlorobenzene coatings provided the optimum surface quality. This was presumed to result from the provision of an effective reducing gaseous blanket of C-CI compounds.

Metallographic examination of the as-cast surfaces revealed surface attack on all castings irrespective of whether grain boundary oxidation also occurred, the nature and thickness of the oxide coating varying according to mould wash type. The overall surface attack observed with the reducing types of mould wash comprised a relatively thin oxide coating (~0.025in) with FeO-SiO

2 as the primary constituent again be-

ing FeO-SiO2. In some cases, this penetration completely sur-

rounded the grains, being apparently characteristic of the more aggressive oxidising washes.

164

To summarise the American work, (52,53) both surface and grain boundary oxides were identified as primarily FeO-SiO

2,

with small scattered amounts of mould wash oxides, MnS and reduced metal from the reducible oxide washes. In addition, all castings including the reference casting with no mould wash exhibited an oxide 'penetration' in the form of rounded non-metallic inclusions within the decarburised layer, both at and near the oxide surface. These were identified as deoxidation products comprising AI

2O

3, SiO

2 and MnO. These oxide

globules formed extensively with certain mould washes and, in some cases, formed a line of inclusions ahead of the oxide layer.

Swedish investigations (54) have shown that in Ni-Cr-Cu-V steels, network cracks appear to a depth ranging from 0.2 to 5mm, when poured into olivine sand moulds with a zircon wash. The crack surfaces comprised the dominant phase of FeOMnO with Fe

2O

4, FeOSiO

2 and FeS also present. In accord-

ance with previous Swedish work on network cracks in low alloy NiV steels, the mechanism of attack was postulated to be grain boundary oxidation. Proposed preventive measures included the use of protective atmosphere and reducing mould washes.

Stage at which Chicken-wire Cracking Occurs

Soviet work has confirmed the observation of others, in that the appearance of chicken-wire cracking can be sporadic in nature. (51, 55-57) Castings poured from the same heat displayed different crazing tendencies, but crazing usually occurred on the surfaces of thick sections free from burn-on, but coated with a continuous layer of scale. The crazing initiated during the cooling/solidification stages by preferential development of scaling grooves at the grain boundaries, which acted as stress raisers. Craze cracking then commenced at these grain boundary locations when the stress exceeded at certain critical value, being aided by any intergranular precipitation of non-metallic inclusions. This stress may arise from quenching (55) or transformation (50) effects or from the development of tensile stresses associated with decarburisation. (51) Crazing occurred only when the casting was cooling/

165

solidifying in the mould, and any subsequent heat treatment merely opened up previous cracks which followed the grain boundaries, the surfaces becoming decarburised. Aluminium nitride was suspected of being a cause of crazing, but deoxidation with 0.15%Ti or Zr had no discernible effect on cracking tendencies. (51) It was elsewhere considered that suI- phides, oxysulphides, alumina, silicates, titanium nitride and aluminium nitride could aid crazing by weakening the grain boundaries. (57)

Oxidation studies carried out in England 30 years ago showed that heating of polished metal surfaces in air to give a liquid oxide layer resulted in oxide penetration to the solid metal. (58) The mechanism involved a surface oxide layer forming with diffusion of oxygen into the metal, producing layers of oxide particles. This indicates that chicken-wire cracking is most likely to initiate when, at an early stage of cooling, liquid oxide (or fayalite) is present at the mould/metal interface.

Electron-probe Micro-analysis

In recent years, a large number of samples of Ni-Cr-Mo steel containing chicken-wire cracks have been submitted to SCRATA for examination. In some instances the cracks have been very deep (30mm) although more frequently, the average depth is less (5mm) (see Fig. 30). Microscopical examination has shown that there is a marked similarity in crack morphology in all the castings examined:

(i) the crack opens to the surface;

(ii) there is an infilling of slaggy material;

(iii) surface decarburisation has taken place;

(iv) an envelope around the crack contains a globular precipitated phase, heavier near the surface and becoming less as the crack continues away from the surface.

A considerable amount of probe analysis has taken place to identify these phases, with a view to determining the mechan-

166

ism of formation of the crack. Typical large and small cracks are shown in Fig. 31. The small crack on the right has branched and the grey phase with which it has become filled continues on the surface on both sides, the funnel shape indicating that metal has been progressively removed. The dot-like precipitates can be seen in the decarburised zones at the surface and upper parts of the crack.

A wide range of constitutents was found, but to illustrate the results obtained on one particular crack, the example shown in Fig. 32 is given. For area 1 (Fig. 32(b), electron micrograph):

(a) the fine network of dark phase (Fig. 33(a)) detailed in Fig. 33(b) consisted of mixed Fe-Si-Cr-Mo-O compounds (as confirmed by EDAX analysis and elemental X-ray maps, not reproduced here).

(b) the grey phase within the defect proper but at the matrix interface consists of individual dark and light grey phases, comprising iron oxides with trace amounts of silicon, chromium and manganese (see Fig. 34).

167

For area 2 in Fig. 32(b) (see Fig. 35):

(a) the grey phase within the defect proper, but near to the matrix interface, consists of iron oxides with traces of silicon, chromium and manganese (see Fig. 36).

(b) the fine network of dark phase detailed in Fig. 37 consists of mixed Fe-Si-Cr-Mn-O compounds.

For area 3 in Fig. 32(b) (see Fig. 38), the white phase in the grey phase consisted of iron, manganese, silicon and chromium (Fig. 39).

168 169

170 171

172 173

Examination of the dot-like phases at the crack tip in another Ni-Cr-Mo steel (see Fig. 40) gave the following results:

Many other analyses have been made of cracks, their infilling and environs, and whilst the precise nature and extent varied with location, iron and oxygen were the main constituents, and silicon, chromium and manganese were always present. Furthermore, chromium and phosphorus were concentrated in some phases well above the average or bulk level. No observation could be made with respect to nickel, which is surprising because all examples of this characteristic type of craze cracking have been in Ni-Cr-Mo steels. Nevertheless, the observations are consistent with the theory of penetration by oxygen, probably in the form of a silicate (fayalite).

Control of Chicken-wire Cracking

At present the precise mechanism of craze cracking is not understood. Whilst an oxidation phenomenon is almost certainly involved, what triggers off craze cracking in some parts of a casting and not in others is not at all clear, although variable surface stresses may be involved. However, given that oxygen (via a low melting point phase) is the primary element responsible, then all efforts should be made to obtain reducing conditions at the mould/metal interface (no mould paint or organic bonded sands); restricting oxygen available (in the metal; entrained air; argon shield); (59) chilling to reduce the effect of low melting point phases.

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LINEAR SURFACE DISCONTINUITIES ASSOCIATED WITH MICROSEGREGA ION AND PRECIPITATION

Microsegregation

Microsegregation in cast steel manifests itself by changes in solute concentration between dendrites, i.e. across the interdendritic region. It can occur without precipitation in two forms, viz. equilibrium and non-equilibrium. Equilibrium segregation occurs as a result of inhomogeneities in the matrix giving rise to sites for which solute atoms have a lower free energy. These sites occur at interfaces as well as defect sites, dislocations and stacking faults. At any given temperature there is a unique solute concentration that is asymptotically approached as time goes to infinity and at a rate governed by diffusion.

Non-equilibrium segregation depends upon rate processes and, in general, disappears as time approaches infinity, provided that diffusion is allowed to reach full equilibrium. This form of segregation includes the growth of precipitates as covered by the previous section on precipitation.

This section deals with equilibrium segregation, in particular with phosphorus segregation and temper brittleness. Both of these phenomena can give rise to embrittlement which results in fracture by cracking in an intergranular mode.

Temper brittleness is a form of intergranular failure in low- allow steels, occurring when the steel is slowly cooled from impurity elements phosphorus, tin, antimony or arsenic to the existing grain boundaries. The effect of these elements depends upon the alloy content of the steel. Antimony, tin and arsenic only exert an effect in steels containing nickel and chromium, by phosphorus. (6) Copper and tin have also been reported to be embrittling elements, (6) 0.37% copper having an even greater effect than 0.06% phosphorus. (6) In practice, however, temper brittleness is mainly the result of phosphorus segregation at prior-austenite grain boundaries. (61) After the casting has received its first anneal, the prior austenite grains

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are finer than the as-cast austenite grains on which AIN precipitates, and whilst an intergranular fracture mode predominates, the coarse angular facets associated with Rock Candy are not normally seen.

It has been hypothesised that segregation of phosphorus at temperatures around 500-550 °C depends upon the presence of nickel and chromium, (62) or upon the presence of molybdenum through an Mo-P interaction (Mo

3P formation leaves a

low level of free P). (63) Molybdenum, therefore, de-embrittles the phosphorus-rich grain boundaries (63) and so can carbon, (64) through site competition. The action of added chromium is to precipitate carbon so that it no longer segregates and hence phosphorus segregation returns. (64) A similar occurrence can take place when, with increased tempering time, molybdenum forms carbides (Mo

2C) and phosphorus segrega-

tion returns tot he same level as when no molybdenum was present in the first place. (65) It is apparent that very complex movements of atoms, depending on temperature and time, can take place at grain boundaries and that such movements can account for the difference between non-reversible and reversible temper brittleness. In the latter form, temper brittleness can be removed by tempering at a higher temperature, e.g. 650 °C, which is outside the critical range, and rapidly cooling through the critical range. The effect on impact transition is to cause a displacement of the embrittled steel to higher temperatures, i.e. a shift to the right.

From consideration of the “relative strength of grains and boundaries” theory, it can readily be seen that if sulphur is reduced to very low levels (say 0.005% and below) the grains themselves become stronger; if phosphorus becomes simulta- neously higher at the grain boundaries as a result of segregation, then intergranular fracture can occur by phosphorus embrittlement. Thus, by consideration of the foregoing theories, very low-sulphur low-carbon steels, especially those containing chromium, can become intergranularly embrittled by phosphorus segregation, even though the bulk analysis for phosphorus is low. More and more evidence is becoming available that very low phosphorus contents are beneficial for achieving

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high toughness and minimising embrittlement in wrought steels. (66)

Other deleterious side-line effects have been reported in wrought very low-sulphur steels. Some very low-sulphur steels have been found to exhibit hardening characteristics considerably in excess of those expected. (67,68) In welding, this higher hardening capacity gives rise to an increased susceptibility to hydrogen induced cracking. (67) Furthermore, the effect of sulphides and other non-metallic inclusions in providing ‘sinks’ for hydrogen and lowering its activity has been recognised for a long time. (70)

Aluminium nitride and temper brittleness and instigators of intergranular fracture. Both are the result of complex mechanisms and contemporary research has aided the understanding of these mechanisms. A phenomenon that is less well understood is now appearing as new steelmaking and ladle refining treatments are coming into use which allow sulphur contents to be reduced to ultra-low levels (e.g. 0.005%S or even lower). In these instances there is a very real possibility that phosphorus, although at low levels by present understanding (e.g. 0.010 to 0.015%P) can cause embrittlement and intergranular fracture, leading to the development of linear surface discontinuities.

Solid State Precipitation

During the cooling of steel, phase transformations take place, with the formation of new phases in accordance with the phase or equilibrium diagram, or new compounds can precipitate, usually with undesirable morphologies.

Precipitation takes place preferentially on sites of high interfacial energy, such as boundaries between similar grains, boundaries between different phases already present from solidification, metallic or non-metallic, and on certain crystal- lographic planes. Precipitates of a compound of a discrete nature weaken rather than strengthen the interface and the weakness can become apparent as small stresses cause decohesion and separation on a macro-scale, i.e. a disconti-

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nuity, usually of a linear nature. Such precipitates can be typically nitrides, ferrite or carbides, and will be discussed under the three following headings.

Rock Candy

Aluminium nitride precipitates in the solid state on the austenite grain boundaries formed from cooling after solidification. Since the grain size is large in the as-cast condition, fractures following these grain boundaries will show very coarse angular facets (Fig. 41), which accounts for the popu- larly used term “rock candy fracture”. The consequence of fracture in an intergranular manner such as this is to lower the upper shelf energy in the impact energy transition curve, as illustrated in Fig. 42. (70) At low temperatures, fracture will normally take place by cleavage in steel free from aluminium nit ride precipitates.

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The presence of these precipitates can not only lead to the development of linear surface discontinuities, but can also result in the spontaneous separation of a casting (usually of thick wall or complex shape) into two halves when residual stress has been sufficiently high to cause this. (72)

In commercial casting compositions, the aluminium nitride precipitates mainly in the temperature range 1250-950°C, therefore, heating above 1250°C should cause it to be taken back into solid solution. Upon the first austenitising heat treatment (usually in the region of 950°C), some of the AIN will be taken into solution, but not all, and on slow cooling, AIN embrittlement will reappear, possibly showing smaller facets. If the austenitising temperature has been high enough to remove all of the AIN precipitate, slow cooling will cause it to reappear on the boundaries of the now smaller grains, and if AIN embrittlement occurs it will show less coarse facets. Aluminium nitride induced fracture shows a considerable relief or depth in the facets, which show markings of AIN dendrites (Fig. 43(a)) and often appear dull, unlike the case of temper embrittled samples (Fig. 43(b)). Aluminium nitride precipitates are very thin and are not readily seen in the optical microscope. Their apparent absence, therefore, cannot be taken to imply freedom from AIN embrittlement. When visible, they appear as short rod-like precipitates (Fig. 44).

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Research into the mechanism of aluminium nitride embrittlement allowed the explanation of a previously unexplained phenomenon, i.e. steels made by the Bessemer or Converter process (with high nitrogen contents) did not exhibit AIN embrittlement. This is because the high sulphur levels (0.06-0.08%) resulted in grain weakening, allowing a (premature) ductile fracture to take place. It is also possible to cause a transition from ductile to intergranular fracture by lowering the tempering temperature (of a hardened steel)-the grains become harder and stronger and it is easier for fracture to take place at grain boundaries. (70)

The amount of AIN available for grain boundary precipitation is a function of both the aluminium and nitrogen contents of the steel. The time available for precipitation is also a significant variable and rapid cooling through the precipitating temperature range will prevent AIN embrittlement. In thick section castings, whatever the imposed external conditions, it may not be possible to attain a sufficiently fast cooling rate and it becomes increasingly important to maintain both aluminium and nitrogen at very low levels. (71-74)

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The phenomenon of AIN embrittlement, or Rock Candy fracture is discussed more fully elsewhere. (75-77) The following precautions should be taken to reduce the risk of AIN embrittlement:

(i) in steelmaking, keep the nitrogen below 0.01 %.

(ii) when deoxidising with aluminium, avoid "over-kiIling". Residual aluminium contents in excess of 0.05% are unnecessary; in large castings, a maximum of 0.03% should be maintained and if necessary, deoxidation supplemented with titanium or zirconium after the aluminium addition.

Ferrite Films

In steels where the carbon and manganese contents are just below the eutectoid composition there will only be a small precipitation of ferrite on cooling through the short Ac

3-Ac

1

range. This will precipitate at the austenite grain boundary or at other interfacial sites, such as non-metallic inclusions (see

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Fig. 45). In cooling from solidification, the austenite grains will be large and the ferrite precipitates at their boundaries will be of considerable length; after the first heat treatment, i.e. full re-crystallisation on cooling from a temperature above the Ac

3,

the austenite grains will be smaller, but the ferrite films will be thinner and just as likely to cause embrittlement. Whilst ferrite is a ductile phase, it must be remembered that in thin films, surrounded and held in position by strong pearlite, the ferrite is under triaxial stress conditions and unable to deform, i.e. it is brittle.

Ferrite films have been known to cause spontaneous through- thickness cracking, as stresses develop whilst cooling on the foundry floor. The only way to overcome this problem, assuming design is inflexible, is to change the chemical composition of the steel so that at the natural cooling rate, ferrite is not precipitated at grain boundaries (nor cementite, which, of course, is equally embrittling). This phenomenon is common in 1 1/2%Mn type steels and an addition of molybdenum usually affords a way out of the problem.

Carbide Precipitation

In low alloy steels, where dendritic segregation occurs during solidification, relatively high (up to two times) concentrations of strong carbide forming elements can occur (e.g. chromium, molybdenum and vanadium). Carbon will be held in these regions as carbides (see Fig. 46) and will be difficult to remove. Practical heat treatments are not long enough, or at sufficiently high temperatures, to cause the homogenisation of chromium, vanadium and molybdenum. Double heat treatment can slightly improve mechanical test results, but the cracking propensity is not removed.

In 13%Cr steels carbide precipitation at phase boundaries can be present after heat treatment (Fig. 47). In the example shown, the carbides do not form long chains, but they are nevertheless embrittling in a 13%Cr steel. High temperature homogenisation is necessary to take these carbides into solution and an accelerated cooling rate to prevent them precipitating in this string-of-pearls morphology.

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In 13%Mn steels, carbide precipitation is very heavy in the as- cast condition (Fig. 48(a)). Heating to 1050°C and rapid cooling is necessary to take and retain carbides in solution. Failure to do this correctly can lead to cracking (Fig. 48(b)). It is widely believed that such cracking occurs upon water quenching, but this is not normally the case; at the start of the quench the steel is fully austenitic and very ductile. Cracking, if it occurs at all, usually occurs on heating the as-cast steel, and the term

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“clinking” is used to describe this. “Clinking” is a heating phenomenon, whereas “quench cracking’’ is a cooling phenomenon; the latter is usually a phase transformation effect, whereas clinking is not necessarily so. Austenitic manganese has a lower thermal conductivity than ferritic steels, consequently, on placing in a hot furnace, or on a hot hearth, or heating from cold at too high a rate, thermal stresses readily build up which cannot be accommodated by the carbid- embrittled steel. The remedy, of course, is to avoid rapid heating, particularly of complex shapes or where there are big changes of section.

Chromium-nickel austenitic steels also have low thermal conductivity, but are rarely embrittled by such coarse precipitates and do not suffer from clinking. They can, however, contain chromium carbide chains (Fig. 49) and sigma-phase in the as- cast condition; in large amounts (Fig. 50) sigma-phase will result in considerable embrittlement and propensity for cracking. Carbid eutectics (Fig. 51) lead to a similar problem. The remedy lies usually in correcting the balance of composition. (78,79)

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The solution is not always that simple, however. In some high alloy Cr-Ni steels, the deleterious precipitates can be nitrides, or carbo-nitrides, or intermetallic phases of a more obscure nature than sigma-phase. One example of cracking in a CN-7M steel casting is shown in Fig. 52. Metallographic examination revealed that the defects were associated with stringers of globular precipitates (Fig. 53). The existence of unidentified precipitates has been reported elsewhere, although straight- forward MC

6 precipitates, indirectly related to a high silicon

content in the steel, have also been found. (80) The reason for CN-7M steel being so ultra-sensitive to surface discontinuities has not been satisfactorialy resolved, but it is likely that there is more than one mechanism operating.

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EMBRITTLEMENT BY HYDROGEN

In the bulk steel industry it is well known that during solidification and cooling of ferritic steel ingots, particularly in Ni-Cr, Ni-Cr-Mo and Ni-Cr-Mo-V types, (81,82) hydrogen can cause discontinuities known as hairline, shatter or flakes. (83-86) The prevalence of these discontinuities in ingots, as distinct from steel castings appears to arise from massive ingot section thicknesses (rather than gross tonnage) and rapid cooling (e.g. poured in metal moulds) promoting hydrogen retention or cracking and embrittlement. Nevertheless, linear surface discontinuities and fracture on the foundry floor does occur in steel castings from time to time.

During solidification and subsequent cooling, the solubility of hydrogen in steel decreases markedly, thereby giving rise to supersaturated solutions. In practice, the hydrogen content of liquid steel is considerably less than the solid solubility limit. As a result, gross porosity in steel castings due to hydrogen expulsion during solidification seldom occurs in practice. Dur- ing cooling in the mould, however, a point may be reached when the steel becomes saturated with hydrogen and, on further cooling, is expelled from solution. That is, on transformation from gamma-phase to alpha, hydrogen becomes instan- taneously less soluble but more easily diffusible. This is a major cause of the suspectibility of ferritic steels to hydrogen cracking and embrittlement, the matrix being easily supersaturated with highly mobile hydrogen.

Microstructural features like grain boundaries, inclusion interfaces, pores, voids, etc., can act as effective traps (sinks) for hydrogen. In general, trapping effects become appreciable at temperatures below about 150 °C, the atomic hydrogen being 'desorbed' from the matrix to form molecular hydrogen. In fact, trapping effects are believed to be responsible for castings being less prone to hydrogen cracking and embrittlement than wrought products owing to the higher volume fraction of microstructural traps; and high sulphur steels being less susceptible to hydrogen cracking and embrittlement by providing abundant inclusion interfaces for hydrogen recombination.

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Cracking and Embrittlement Effects

The simultaneous presence of hydrogen and stress within a susceptible microstructure at temperatures between - 100 and 200 °C provide the basic conditions conducive to hydrogen cracking and embrittlement. In general, maximum susceptibility occurs at slow strain rates around room temperature, especially in hardened restrained transformation products of Iimited ductility. (87,88)

Normally, for cracking to occur, the hydrogen must be concentrated at the embrittlement locations (tips of cracks) by diffusion. It is this requirement for time for diffusion (e.g. incuba- tion period) which allows hydrogen to build up to the appropriate critical levels-hence the terms delayed or cold cracking. This form of embrittlement is reversible, in that if no cracks have formed and the hydrogen is removed, full ductility may be restored.

As in the case of brittle fracture, hydrogen embrittlement is sensitive to temperature and strain rate: maximum embrittlement occurs at slow strain rates around room temperature (Fig. 54); as temperatue increases, ductility improves until at -200°C the effect of hydrogen is lost; as temperature decreases, the effect of hydrogen diminishes until, below the ductile-brittle transition temperature for steel, any embrittlement due to hydrogen is masked by the inherent brittleness; high strain rates reduce the hydrogen effect-in fact the effect of hydrogen is lost completely in an impact test.

Hydrogen cracking can occur spontaneously after austenite has transformed to a hardened, restrained microstructure. In general, martensites tend to be of limited ductility compared to ferrite-carbide aggregates (e.g. pearlite, bainite), and since hydrogen reduces the amount of plastic deformation that the matrix can withstand under stress, such transformation products are particularly susceptible (Fig. 55). Stressing conditions inevitably arise from restraint in the mould and from geometric design.

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Control of Hydrogen

Hydrogen is best removed at source during steelmaking. Pro- cedures vary widely with arc furnace refining practice, and induction furnace melting practice. (89,90) Hydrogen can also be removed by treatment in the ladle. (89-92) However, the control of hydrogen in liquid steel is so vast a subject, it is considered to be outside the scope of this chapter.

After solidification, excess hydrogen can only be removed by diffusion, i.e. immediate* heat treatment at subcritical temperature of 600-650°C. (93) This is the optimum temperature range for hydrogen diffusion in ferrite**, and also aids decomposition of any retained austenite and stress relief by plastic relaxation. The soaking time depends to a large degree on the initial and final desired hydrogen content, as well as the section thickness being treated. Although there is no universal 'safe' hydrogen level for carbon and low alloy steels a consistent hydrogen level of around 3ppm for all but high strength steels (e.g. less than 750 MPa) should be satisfactory. (84) * i.e. before the casting has cooled below -200°C if any risk

of hydrogen cracking or embrittlement is to be avoided. ** the diffusivity of hydrogen at 600 °C in α is the same as that

of γ at 1000 °C.

LINEAR SURFACE DISCONTINUITIES ARISING FROM QUENCHING

A commonly used heat treatment practice, particularly for low and medium alloy steel castings where advantage is taken of their hardenability, is quenching and tempering. Quench cracks can arise in hardened, restrained sections during or after quenching when the residual stress levels, perhaps aided by local stress concentrators, exceed the critical stress for crack initiation. Typically, these cracks run from the casting surface into the centre, as illustrated in Figs. 56(a) and (b). They are essentially straight cracks, free from decarburisation and oxidation products, running through hardened transformation products unrelated to interdendritic regions (Fig. 57). The flame hardening of alloy steels requires special precautions (98) and incorrect flame hardening will also result in surface cracking (Fig. 58).

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Compositional and Cooling Rate Effects

The aim of quenching is to achieve the desired mechanical properties by cooling at a sufficient rate to suppress the higher temperature diffusional decomposition of austenite to pearlite or bainite aggregates, thereby promoting full martensite hardenability. All alloying elements, except cobalt, have the beneficial effect of retarding the diffusional decomposition of austenite such that martensite can be formed in alloyed steels under less severe quench conditions than for plain carbon steels. In other words, for a given quench rate, thicker sections of alloy steels can be fully transformed to martensite.

Selection of quench medium is important, and should only be severe enough to obtain the desired through section hardness to give the required mechanical properties after temperature (Fig. 59). Both the steel composition and section size dictate the quench speed to give a fully hardened through thickness section, with complex sections being particularly prone to cracking and distortion. Of course, agitation of the casting (if possible) and circulation of the quenchant also affect the quench speed, whilst environmental considerations (smoke, fume, fire) may also affect quench selection.

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Thermal Sources

The stresses due to thermal volume changes (e.g. thermal contraction on cooling) depends upon both the degree of strain mismatch and the metal stresslstrain relationship during quenching as follows:

(i) for a given stresslstrain relationship, high temperature dif- ferentials (e.g. hot spots) promote high strain mis- matches-for instance, temperature gradients are promoted by low thermal diffusivity, large section thickness and high quench severity (H).

(ii) for a given strain mismatch, high modulus of elasticity (E) promotes higher residual stresses. Since residual stress levels cannot exceed the yield strength without plastic relaxation, the higher the yield strength the higher the possible residual stress. If the yield decreases rapidly with increasing temperature (e.g. becomes plastic), then the strain mismatch will be smaller at higher temperatures owing to plastic relaxation.

Transformation Sources

The volume expansion associated with martensite transformation proceeds progressively during cooling from the Ms to Mf temperatures. Austenite transformation to bainite or pearlite also produces a volume expansion, but of lesser magnitude.

Stress Concentration and Relaxation Effects

The yield behaviour with temperature and its relation to the Ms temperature is important with regard to the nature and level of residual stress. If the steel has some plasticity at the Ms, the strain mismatch can be relieved by plastic relaxation. On the other hand, i f the Ms is low and the yield is high, then substantial stresses can develop, especially in the presence of stress raisers. These stress concentration effects can be relieved by minimising changes in casting section thickness and sharp re- entrant angles by means of gradual tapers and generous fillets respectively.

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As a rule, the higher the Ms temperature, the less the expansion associated with martensite formation (e.g. the specific volume change is reduced) and, as a consequence, there is a reduced tendency to form quench cracks. For a given quench rate (assuming full hardenability), high carbon steels in particular promote quench cracking tendency by drastically lowering the Ms temperature as well as forming harder martensitic transformation products of more limited ductility. This lowering of the Ms temperature makes residual stress relief by plastic relaxation more difficult. In fact, an extensive evaluation of the quench cracking susceptibility of low alloy Ni-Cr- Mo steel confirms the detrimental effect of Ms depressors (especially carbon) and embrittlers (e.g. phosphorus) as follows (94):

(i) carbon strongly promotes cracking susceptibility;

(ii) manganese, chromium and phosphorus appear to be detrimental, but to a lesser extent than carbon;

(iii) sulphur, silicon, nickel, molybdenum and aluminium in the ranges studied showed little, if any, effect;

(iv) normal boron additions (<0.003%) had little effect on cracking tendency.

Even if no cracking occurs on quenching, the casting can still crack at room temperature due, for instance, to isothermal transformaton of retained austenite to martensite (e.g. additional volumetric strain). If the casting is heated to only 100°C, some local plastic relaxation can occur to partially relieve excessive internal stresses. In practice, this implies that after quenching, the casting should be transferred immediately to the tempering furnace. Alternatively, if the Mf is high enough the casting may be removed from the quench whilst still warm.

Means of Quench Cracking Control by Selection of Quench Practice

Given a certain steel type, casting geometry, shape and size, the risk of quench cracking on heat treatment can be effectively minimised by careful selection of quench practice (95-97).

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Water quenching is usually restricted to plain carbon and low alloy steels with less than 0.25%C where, if possible, castings should be of simple shape with no sharp stress raisers. The major drawbacks of water quenching are

(a) persistence of the vapour cooling stage and localised vapour blankets in the boiling stage

(b) irregular vapour blankets which break up at sharp corners, resulting in temperature gradients across the casting

(c) the liquid cooling stage commences at much lower temperatures than with oil, making the casting more prone to cracking or having higher residual stresses.

The most significant feature of oil quenching is the reduced cooling rate during the liquid cooling stage. Mineral oils are used and can be broadly classified as ambient (slow) oils, warm (normal) oils and hot (fast) oils.

Synthetic polymer quenchants usually comprise polyalkalene glycols (PAG) which dissolve in water at room temperature. As the solution is treated, the polymer becomes insoluble in water at temperatures above about 77°C. When the solution is cooled again, the polymer goes back into solution again and is fully miscible. It is this property of inverse solubility which im- parts the unique cooling mechanism to polymer quenchants.

The cooling rate of the polymer quenchant can be readily varied to suit specific requirements by changing the concentration of the solution which affects the deposited polymer film thickness. The quench speed is also influenced by bath temperature, with operative temperature ranges from 15-65°C. Successful application depends upon surface condition, geometry and section thickness.

Control of Quench Cracking

Quench cracks in steel castings arise from excessive stressing action in transformation products of limited ductility (e.g. martensite), being associated with thermal and transformation

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events during cooling. For a given steel type, means of control may be to:

(i) ensure quench medium is not too severe-it is possible that the through section hardness may be obtained with a less severe quenching action. In particular, equalisation of section centre and surface temperature gradients is critical, and requires careful consideration of quench type, concentration, temperatures, etc.

(ii) eliminate notch-like stress raisers-minimise changes in casting section thickness and sharp re-entrant angles by means of gradual tapers and generous fillets respectively. The presence of embrittling (e.g. phosphorus) or weakening (e.g. sulphur) elements as a result of, say, segregation may aid quench crack initiation and propagation.

(iii) avoid prolonged periods between quenching and tempering-this helps minimise quench stresses building up to critical levels for cracking. Preferably, transfer to tempering furnace while still warm according to section thickness and Mf considerations.

LINEAR SURFACE DISCONTINUlTlES ARISING FROM WELDING

Linear surface discontinuities in steel castings are usually rectified by fusion welding. Fortunately, cast steels have better weldability than equivalent wrought steels, and lower preheat temperatures can be used for cast than for wrought steels. (99) Nevertheless, unless welding is properly carried out, welds themselves can exhibit cracks, which show as linear discontinuities on the surface of the casting when the weld is ground off. Such cracks may be primarily in the weld deposit or in the heat affected zone (HAZ) of the parent metal and, of course, this applies to welded joints as well as repair welds. The subject is immense, and cannot be adequately covered in this chapter but since the mechanisms of cracking in welds have a close relationship to some of the mechanisms discussed earlier, brief discussion of the main features of weld cracking will not be out of place.

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Cracks will occur in weld metal and parent metal when stresses exceed the tensile strength of the metal. Cracking is generally associated with stress concentration at or near discontinuities, or near to “notches” resulting from design.

Linear surface discontinuities arising from welding can be classified as either hot to cold cracks. Hot cracks develop at elevated temperatures. They commonly form during solidification of the metal at temperatures near the melting point. Cold cracks develop after solidification of a fusion weld as a result of stresses. Such cracks in steel are sometimes called delayed cracks when they are associated with hydrogen embrittlement. Hot cracks propagate between the grains while cold cracks propagate both between grains and through grains.

Cracks may be longitudinal or transverse depending on their orientation with respect to the weld axis.

Transverse weld cracks are perpendicular to the axis of the weld and, in some cases, extend beyond the weld into the base metal. This type of crack is more common in joints that have a high degree of restraint.

Longitudinal weld cracks are found mostly within the weld metal, and are usually confined to the centre of the weld. Such cracks may occur as the extension of cracks formed at the end of the weld. They may also be the extension, through suc- cessive layers, of a crack that started in the first layer. Crater cracks are formed by improper termination of a welding arc; they are usually shallow and are sometimes in the form of a star-like cluster. Underbead cracks are generally cold cracks that form in the HAZ of the parent metal. More specifically, cracks that occur in either weld metal or parent metal will be discussed under these headings.

Cracking in the Weld Metal

Solidification or centre-line cracking takes place by a mechanism similar to hot tearing. The cracks are often associated with high transverse strains in restrained welds. To overcome this it may be necessary to (a) improve the fit-up (i.e. avoid

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large gaps in fillet welded joints), (b) use preheat to avoid thermal stresses, (c) reduce welding speed or alter manipulation to improve the contour of the deposit, (d) sequence welds to balance shrinkage stresses, (e) fill craters at the end of a weld run to a slightly convex shape prior to breaking the arc, and (f) avoid rapid cooling.

Parent metal composition can also affect solidification cracking, principally carbon, sulphur and phosphorus. The higher the content of these elements in the parent metal, the greater will be the risk of cracking in the weld metal. Contaminants remaining on the parent metal (e.g. cutting oil), high dilution from the parent metal and hydrogen from the electrode can also contribute to weld metal cracking. To overcome dilution it may be necessary to butter the joint faces prior to welding and use the electrode negative. Low hydrogen electrodes should be used and be properly baked and kept warm to avoid moisture in the coating.

Cracking in the Parent Metal

A ductile parent metal, by localised yielding, can withstand stress concentrations that might cause a very hard or brittle metal to fail. Parent metal cracking is usually longitudinal and takes place in the heat affected zone. If high HAZ hardness is due to high hardenability of the parent, pre-heat is usually capable of preventing cracking. Solid solution hardening elements such as phosphorus can often be countered by increasing the pre-heat, but separate phases such as sulphides will always contribute to embrittlement and cracking in the HAZ. If the sulphides are predominantly Type Il, then increasing the pre-heat is futile and crack-free HAZs are almost impossible to avoid; the only possible recourse to try is using low heat input to deposit very thin layers.

A high hydrogen parent metal, particularly if welded with rutile or unbaked electrodes, will severely embrittle a heat affected zone. To counteract hydrogen HAZ cracking, attention must be given to (a) welding consumables and their use under controlled conditions, (b) pre-heat that is uniform over the entire section to be welded, (c) precautions (a) and (b) to be taken

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even for assembly tack welds, (d) a defect must be fully removed by chipping, gouging, grinding or machining and the weld preparation must be completely cleaned from any materials used to non-destructively test for soundness.

Not all parent metal cracking is due to hydrogen, of course. Other factors such as Type II sulphides, brittle phases, residual stresses, high hardenability, inherent brittleness and low ductility (e.g. 28%Cr steel) play major parts in causing cracking. Some remedies are too obvious to mention but consideration should always be given to (a) welding when the parent is in the optimum condition of heat treatment, i.e. when hardness is lowest, ductility is highest and grain size smallest; in practice, welding in the as-cast condition is usually avoided; (b) avoiding Type II inclusions by giving attention to deoxidation practice; (c) welding on to a base free from porosity, inclusions or other flaws, (d) welding a “stress-free” casting; stress relieve before welding if necessary, (e) avoiding build-up of stresses during welding; in effect, this involves avoidance of heat concentration in any one area by making short runs in dia- metrically opposite areas rather than filling up a complete area at a time.

It should be remembered that, when repair welding under-riser discontinuities, the surface exposed by riser removal will rarely be of the same composition as that of the body of the casting; under-riser segregation will be present to a certain extent and on pads left by large risers, carbon and sulphur may segregate by several orders of magnitude. In fact, the discontinuities may have arisen by thermal dressing in the first place. Very high preheat temperatures will be necessary for both air- carbon arc cutting and fusion welding of highly segregated areas.

Linear surface discontinuities may be found with magnetic particle inspection, which have no readily apparent cause. Investigation may associate areas containing these discontinuities with high hardness and in such instances, it may be concluded that the cracking occurred as a consequence of striking an arc on the casting during a welding operation, either accidentally, or intentionally to clean the electrode end.

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These problems arise mainly on hardenable alloy steels and when discovered the casting should preferably be re-heat treated.

There is another form of parent metal cracking not mentioned so far, and this is re-heat cracking, sometimes called stress- relief cracking. Whilst it occurs during service at elevated temperatures, it can also occur during stress relief after welding. This form of cracking takes place only in steels which contain elements (such as chromium, molybdenum and vanadium) which cause secondary hardening in the HAZ. Grains within the HAZ are strengthened by secondary hardening during stress relief and the grain boundaries become relatively weaker. Impurities in the steel which occur at the grain boundaries accentuate the weakening so that de-cohesion can occur at certain stress levels. The presence of stress concentrators such as pre-existing cracks, lack of root fusion defects and partial penetration welds, will accentuate re-heat cracking.

A full heat treatment after welding will avoid re-heat cracking, but this is frequently not possible. Restricting the heating rate to stress-relieving temperature can be helpful, as can inter- pass relieving. Grinding the toes of welds before they cool can also be beneficial. Any notch-like defects should be prevented or removed before post-weld heat treatment.

A partially fused root run will frequently allow a crack to start and progress through the entire thickness of the weld. Condi- tions for full fusion welds should always be carefully followed. Lack of side-wall fusion and slag entrapment should similarly be avoided and in repair welding it is important that defect removal permits a weld preparation to be left that allows adequate access and contour necessary for the production of defect-free welds.

The control of cracking during welding forms part of the technology of welding and is too large a subject to be dealt with in this chapter. Useful guides to the welding consumables and conditions of pre- and post-heat treatment are provided by the Steel Founders’ Society of America. (100-103)

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EFFECT OF LINEAR SURFACE DISCONTINUITIES ON PERFORMANCE OF THE CASTING

At the surface, linear discontinuities are potential sites for initiation of failure by fatigue andlor brittle fracture. Not all linear surface discontinuities are immediately dangerous, however. Many engineering structures operate with known surface defects; they can be quite innocuous until they grow to a critical size which will lead to fracture. Therefore, provided they are regularly monitored and are of subcritical size, surface defects can be tolerated.

During the last two decades, a considerable amount of research has been carried out into the effects of flaws and imperfections on service performance. Extensive testing programmes on actual castings were carried out in the USA (including programmes on surface discontinuities (104-109)) from which it became evident that castings could withstand considerable dynamic loading even though they contained discontinuities; and that quench cracks and hot tears were not necessarily the most serious kind of discontinuity (104,105,107,108) (Fig. 60). Obviously, the length, depth and position with respect to applied stress had an important bearing on the assessment, as did also the strength and metallurgical condition of the casting. In tensile loading linear surface discontinuities were found to be the most severe of all imperfections as illustrated by the results given in Table I. In impact loading by tension and bending modes, hot tears were shown to raise the transition temperature considerably and lower the upper- shelf impact energy (Fig. 61).

The development of fracture mechanics has made it possible to quantify the acceptable amount of linear surface (and also embedded) discontinuities. The property of a material known as fracture toughness can be measured and when this is known, together with the applied stress in the section containing the discontinuity, a calculation can be made to determine the size of discontinuity that can be tolerated. When a critical size is encountered, fast fracture occurs in a brittle manner. The application of fracture mechanics to steel castings has been studied in the UK (111-117) and in the USA (118-120) and

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206 207

designs based upon fracture mechanics calculations have been used in engineering constructions. (112,121) Despite this scientific approach to engineering principles, factors of safety still have to be used in the final calculations because of practical limitations, e.g. surface discontinuities are rarely of the “ideal” shape used in fracture mechanics calculations; interactions of various defects are not easily accounted for; there are limitations in non-destructive testing techniques in accurate measurement of size and shape of a discontinuity; the exact applied stress is rarely known (although it can be estimated); and the stability of a discontinuity cannot be guaranteed (it could grow to a critical size).

It may be concluded that linear surface discontinuities are undesirable in any casting and their criticality is dependent upon so many material, design and service factors, many of which are not accurately determinable, that they can rarely be pronounced as innocuous. For this reason, many purchasers of steel castings and inspecting authorities specify that linear surface discontinuities are unacceptable.

CONCLUDING REMARKS

The major causes of linear surface discontinuities arising in the steelfoundry have been discussed, together with available means of controlling them. Unfortunately, the time has not yet arrived when linear surface discontinuities can be “designed out”. Nevertheless, a clearer understanding of the mechanisms by which they occur will assist in resolving problems when they arise.

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Lecture V

Adhering Sand Defects by Dr. John M. Svoboda

INTRODUCTION

The problems of mold-metal interface reactions and metal penetration phenomena have been studied by many investigators over a period of more than fifty years. Many fine contributions have been made toward the understanding of the mechanisms involved, and the identification of the significant process parameters. However, a study of the literature will soon reveal that a substantial lack of agreement exists among the investigators. The author personally feels that this is due, to a large degree, to an attempt by most investigators to identify one, more or less universal, mechanism which covers all situations. For example, many papers do not make a distinction between the effect of differences in solidification made between steel and cast iron when discussing adhering sand defects. Also, only a few recent papers discuss adhering sand defects which occur with chemically bonded sand mixes. Because the penetration phenomenon is quite complex, such a simplified, single explanation usually proves inadequate. However, if one critically evaluates all of the mechanisms proposed, we find that all seem to have their place in the overall scheme of things.

Before discussing the details of how these defects occur, it is probably best to define the nature of the various defects which make up the family of “Adhering Sand Defects”. While the various investigators seldom use the exact same definitions, most tend to follow the system of nomenclature proposed by Committee 4F-Mold Metal Interface Reactions of the American Foundrymen’s Society (1). These are as follows:

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BURN ON (Figure 1) A casting defect that consists of a thin layer of sand tightly adhering to the surface of the casting.

Burn on cannot be removed by normal abrasive cleaning methods. It is usually removed by surface grinding. When casting surface finish is important or the degree of the defect is very severe, burn on may render the casting unusable.

Burn on most commonly occurs at the hottest sand/metal locations or where heavier metal sections surround thin sand-mold or core areas.

Burn on is caused by mold/metal reaction. This reaction forms iron silicate (fayalite) that penetrates between the surrounding sand grains, causing a thin layer of sand to be tightly bonded to the casting surface.

BURN IN (Figure 2) A casting defect characterized by a layer of sand tightly adhering to the casting surface. Burn in is more severe than burn on. The adhering sand layer may appear vitreous.

Burn in cannot be removed by normal foundry abrasive cleaning techniques. The defect can usually be removed by extensive grinding, however in some cases, it leads to scrapping of the casting.

Burn in generally occurs at “hot spot” areas of the casting or where especially thick metal sections surround thin sand mold or core areas.

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Burn in is the result of a mold/metal reaction. This reaction forms iron silicate (fayalite) between the surrounding sand grains, causing it to fuse into a sandlmetal mass that is tightly bonded to the casting surface.

PENETRATION (Figure 3) This casting defect consists of an intimate interlocking mass of sand grains (silicates) and metal tightly bonded to the casting surface.

Penetration is characterized by the presence of identifiable base metal, intermingled with sand grains and fused silicates, forming a continuous interlocked structure of sand and metal on the surface of the casting.

The defect is tightly bonded to the casting surface and can only be removed by extensive chipping and grinding. Penetration is often so severe that castings are beyond the point of economical rework and must be scrapped. Machin- ing of castings that exhibit this defect is often difficult and may result in excessive tool wear.

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Penetration is more prevalent at a castings “hot-spots” and in heavy section areas or “transition zones” from thin to heavy metal sections.

Penetration is the result of severe mold/metal reaction between the casting and the surrounding sand media. An intimate mixture of sand and metal becomes fused to the casting as a tightly adhering mass.

SHELLING This defect consists of a “shell” of silicates, sand and metal that is distinct from the casting surface, but con- nected to it by metallic veins. The defect is unique in that it is separated from the casting surface by a refractory coating or mold wash.

Shelling is rigid and hard, usually magnetic, with the out- ward appearance of a solid mass of metal penetration. The “shell” may have different metallurgical characteristics than the base metal contained in the casting.

Shelling can often be removed from the casting surface by prying or chipping, depending on the severity of the defect. It is more prevalent on inside corners and concave surfaces of heavy section ferrous castings.

Shelling is the result of a mold/metal reaction, external to casting. This creates iron silicate (fayalite) within the interstices of the surrounding sand that forms a “shell” con- nected to the casting surface by veins through the remnants of a refractory coating or mold wash.

While the author does not agree with these definitions in their entirety, they form a good basis for discussion. In addition, the first three generally conform to the definitions presented in the “International Atlas of Casting Defects.” (2)

Perhaps a more convenient method of classifying adhering sand defects is to consider the mass transport mechanisms by which the defect is generated. This, then, allows a more

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direct approach to developing methods of prevention. The following discussion will pertain only to adhering sand defects which are observed in carbon, low alloy and high alloy steel castings.

REVIEW OF PENETRATION MECHANISMS

Although many theories and mechanisms have been proposed the principal, most thoroughly investigated, mechanisms are as follows: A. Liquid State Penetration B. Chemical Reaction Penetration C. Vapor State Penetrations

Liquid state penetration is a mass transport process in which liquid metal enters the pore spaces of the compacted mold due to pressure or capillary forces. This mechanism has been extensively studied by Hoar and Atterton (3-4) Pettersson (5), Lyashchenko (6) Estratov (7) and others.

Chemical reaction penetration may be defined as a process where the metal is oxidized by the atmosphere present at the mold-metal interface and the oxide formed subsequently reacts with the molding aggregate. This process has been extensively studied and the phase relationships documented. In- vestigators include Goodale (8), Asanti (9), Colligan, VanVlack and Flinn (10), Gertsman and Murton (11), Kolorz and Orths (12), Markhasev and Lugovskaya (13) Savage and Taylor (14) Chechulin (15) and others.

Vapor state of pentration is a three step mechanism con- sisting of the vaporization of a volatile metallic species at the interface, the diffusion of the vapor into the pores of the molding aggregate, and the condensation or decomposition of the vapor to deposit metal or oxide. Principal investigators of this mechanism include Emmons and Bach (16), Sanders (17), and Svoboda and Geiger (18-19).

General, overview discussions are given in references 20-26 which describe the scope of the problem in the steel casting

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industry. While all three of the above mechanisms of formation are of academic interest, it is very important to put into perspective the relative frequency of occurrance observed in commercial casting practice. As shown in Figure 4, the most frequently observed type of defect is caused by liquid state transport (75%) with chemical reaction transport (20%) and vapor state transport (5%) accounting for the balance. Therefore, the major portion of this paper will concern itself with the first two mechanisms.

LIQUID STATE PENETRATION

The most commonly observed type of adhering sand defect- liquid state penetration-may be said to occur according to the following mechanism:

Liquid metal enters the voids between the said grains due to pressure and/or capillary forces.

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A typical example of a polished specimen of this type of defect is shown in Figure 5. The chemical analysis of the metal, hence the microstructure, on either side of the mold-metal interface is essentially the same as shown in Figure 6, indicating that the metal cross the interface without chemical altera- tion.

In order for a metal to penetrate the void in the mold aggregate, the metal must be of sufficient temperature to possess adequate fluidity. If the metal does not wet the molding aggregate, the pressure required to cause penetration is given by the following:

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This relationship is illustrated schematically in Figure 7. Since the greater the value of penetration pressure (P), the greater the resistance to penetration, it is desirable to maximize the surface tension term (6) and minimize the pore radius (r). The effect of the contact angle term (e) will be considered later.

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If the contact angle between the molding aggregate and the metal remains relatively constant, the penetration pressure is proportional to the surface tension of the metal. This is iI- lustrated in Figure 8 for metals exhibiting a wide range of surface tension values.

The amount of superheat and the freezing range of the alloy also influence penetration behavior. Excessive superheat and wide freezing ranges both increase the amount of penetration for a given molding aggregate. This effect is illustrated for iron-carbon alloys in Figure 9. While the freezing range for a given alloy grade is fixed, the operator has some degree of control of the pouring temperature, i.e. amount of superheat. Shop trials have shown that minimizing pouring temperature is a very effective way of minimizing adhering sand defects, as shown in Figure 10.

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While the surface tension for a given alloy at a given temperature is essentially a constant, the addition of trace elements can greatly lower the surface tension. Elements which behave in a surface active manner are especially significant as shown in Figures 11-12. Indirect measurements of commercial grades of steel indicate similar behavior for oxygen (27).

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From these observations it becomes obvious that the use of any of the clean steel technologies to minimize the oxygen and sulfur contents would be beneficial. In addition, care must be exercised that localized contamination of the metal surface does not occur from a high sulfur catalyst used with a chemically bonded sand (28).

It can also be seen from Equation 1 that the penetration pressure (P) can be maximized by reducing the size of the pores or voids between the sand grains. This is normally accomplished in one of four ways (22):

1. Harder ramming 2. Using a finer sand 3. Addition of a fine “filler” 4. Application of a mold coating

The effect of ramming to a high density is illustrated in Figure 13. With green sand, this higher density can be achieved with machine molding and “drier” sand mixes.

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It is equally important that chemically bonded sands also be processed to achieve density. The sands must be used before they exceed their bench lives. The proper control of amount and type of binder and catalyst, as well as sand temperature, is necessary to achieve the desired densities. An example of this is given in Figure 14 (ref. 29). Also, some mechanical compaction or vibration is necessary even though chemically bonded sands are more flowable than green sand mixes.

The use of a finer sand, or finer distribution, is beneficial up to a point. For small or medium work an AFS 50-60 sand gives adequate results. For more severe conditions, however, a finer sand is desirable. Fine grained silica sands, unfortunately, are not very satisfactory, and an alternative is to use a more ther- mally stable sand such as zircon or chromite. (22)

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In years past, it was customary to add a “filler” such as silica flour (minus 200 mesh) to the mix to create a smaller effective void size. Health considerations now restrict the use of silica flour for this application. Some use is made of zircon flour, however, the use of any fine “filler” increases the binder requirement greatly, and the usual alternative is to use one of the specialty sands such as zircon, chromite or olivine.

The fourth method of increasing effective density is the use of a refractory mold coating. This subject will be discussed in the next section, however, it must be recognized that a coating is only as good as the base sand to which it is applied. Therefore, it cannot compensate for a mold or core of unsatisfactorily low density.

From the above discussion it can be seen that one of the most effective methods available to eliminate adhering sand defects is the use of the specialty sands such as zircon,

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chromite, olivine, or other similar materials. These sands have been covered extensively in the literature (refs. 30-34) however a few comments are in order.

These sands are considerably more expensive than silica and therefore should be used judiciously as facing materials rather than as the normal base aggregate. Also since they have a higher density, and zircon a naturally finer grain size, the binder requirements are higher, further emphasizing careful usage.

One topic of special note is that the contamination of chromite by silica, either naturally or during processing and reclamation, severly reduces the effectiveness of chromite in preventing adhering sand defects. This sometimes is called “glazing” and has been well documented in the literature (refs. 32-34). Poyet and Chevriot (28) have documented the con- taminating effect of SiO

2 (Figure 15) and recommend a max-

imum allowable limit of 2%.

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CHEMICAL REACTION PENETRATION

The Chemical Reaction Mechanism may be described as follows:

Metal surface oxidized by gas atmosphere. Oxide wets and dissolves silica sand. Oxides and silicates are drawn into mold by surface energy effect.

Iron oxide (FeO) readily wets and dissolves silica sand (e = 21°). The iron silicates formed tightly bind the aggregate grains to each other and to the casting surface. This process leaves a thin layer of aggregate grains tightly bonded to the casting surface but seldom reaches any great depth, probably due to the rapid drop in temperature in the mold as the distance from the interface increases.

The liquid silicates also plug the pore spaces. There is usually only a small amount of metal in the penetrated layer, the bind- i ng agents being oxide-aggregate reaction products.

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The chemical reaction penetration mechanism was first reported by Goodale (8) who states that the molten metal surface oxidizes readily during and after pouring and a liquid oxide layer forms on the casting surface. This oxide layer will wet the silica aggregate and since it is very fluid at steel pouring temperatures it is easily drawn into the porous aggregate compact. The oxide further reacts by dissolving silica to form iron silicate. (Figure 16)

The extent that an alloy will oxidize and the degree to which the oxides will react with the molding material are dependent in part on the alloy composition. If large amounts of reactive solutes are present, chemical reactions with the mold will be more severe. The phase equilibria and reaction kinetics are discussed extensively for the various metal-mold material systems in the literature cited (refs. 35-39). In general, high concentrations of reactive components, such as manganese in steel, lead to more severe penetration problems.

It is clear from the above description that the two most practical methods of reducing the amount of chemical reaction at the mold-metal interface are:

Provide a non-oxidizing atmosphere. Use a refractory less reactive than silica.

While it is impossible, in a practical sense, to exclude all oxygen from the mold-metal interface, there are some measures that can be taken. With green sand molding, the moisture contained in the sand mix provides a strongly oxidizing atmosphere. Consequently, operating at the lowest practical moisture content for the given mix will prove beneficial. This also allows a greater degree of compaction which helps prevent liquid state penetration. Use of organic chemically bonded systems is also beneficial in reducing the amount of oxidation.

Perhaps the most practical method of prevention of the “burn- in/burn-on” type of defect is the use of a refractory that is less reactive with the iron-oxide formed at the interface than is silica. The commoh refractories available are those mentioned previously-zircon, chromite and olivine. These can be used

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effectively as facing sands described in the literature. Olivine is particularly effective with high manganese steels as the Ii- quid MnO formed at the interface is especially destructive to silica.

The alternative, which is the most common approach, is to use a mold/core coating. These generally are in the form of a liquid slurry with a refractory filler such as zircon flour in a carrier which can be either water or organic solvent based. Use of a coating allows a layer of highly resisitant refractory to be deposited on the surface of the mold or core. In recent years, the sciences of minerology and rheology of mold coatings have become very sophisticated and are covered well in the literature (refs. 40-41). The effect of a mold coating is shown in Figure 17.

VAPOR STATE PENETRATION

The vapor state penetration mechanism is one of the most in- teresting and controversial mechanisms of mass transport involved in the penetration phenomenon. The process may be basically described as a vaporization-diffusion process where a volatile metallic species vaporizes at the mold-metal interface, diffuses into the porous compacted mold, and condenses or decomposes on the sand grains. The vaporization and diffusion processes are thoroughly discussed in Ref. 16-19 and 42.

While this mechanism is of significant interest, it must be reemphasized that it is only applicable to very large castings with solidification times measured in hours rather than minutes. It has been shown (18-19) that the theromodynamic

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conditions under which vapor transport occurs are very specific as to temperature and oxygen concentration in the mold gas atmosphere, again the conditions observed with large castings. Additionally, the mass of metal transported as a vapor is relatively small, however, the condensed film establishes conditions which allow the second "'wave" of Ii- quid state penetration which normally occurs.

Given these qualifying remarks, it must also be stated that when this type of penetration does occur, the resulting adhering sand defect is very severe and almost impossible to remove. This mechanism will be summarized below.

As defined above, vapor state penetration is a process con- sisting of the formation of a volatile metallic species at the mold-metal interface, the coupled diffusion of this species and heat into the aggregate, and subsequent condensation of the species in the aggregate. The most probable metallic species are either the elemental metal vapors, or volatile metal-oxygen molecules such as (FeO)

2. Manganese vapor and

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the volatile oxide (MnO)2 play a particularly important role due

to their very high vapor pressures. Although large masses of metal are not transported by this mechanism, it is important because the metallic material coats the aggretate grains as a very thin film which alters the wetting characteristics of the metal-aggregate system, and the heat transfer characteristics of the aggregate. These films are shown in Figures 18-19. These factors tend to promote the other two penetration mechanisms. As discussed previously, liquid state penetration is a process where metal enters the void spaces in the aggregate due to pressure andlor capiliary effects and proceeds until solidification occurs. The penetration pressure is a function of the surface tension of the metal, the pore size in the aggregate and the contact angle between the metal and the aggregate.

The deposition of metallic material on the aggregate grains by vapor transport changes the wetting characteristics from a case of nonwetting (e > 90°) for the clean-metal-aggregate system to a case of wetting (e < 90°) for the metal-metallic vapor deposited material system. This allows a second wave of liquid pentration to occur quite easily. This effect is illustrated in Figure 20.

The deposition of the metallic film on the aggregate grains also alters the heat transfer properties of the mold, primarily because at higher temperatures, the largest part of heat transfer in a mold is due to radiation. This allows a given point in the mold to reach a higher temperature than it normally would, and this allows the liquid state transport to reach a greater depth.

Although the liquid state penetration mechanism ceases to be operative when solidification occurs, vapor transport continues to be operative and significant amount of mass transport occurs at temperatures as low as 1000F. Any metal deposited in the compact by liquid state transport still offers a surface for vaporization. While vapor state transport is not a serious problem in small castings which cool quickly, it becomes a very serious problem in large castings which take

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hours, or days, to cool. Cored areas which are surrounded by heavy sections of metal cool very slowly, and vapor transport plays a very important role in the penetration of these areas.

SUMMARY OF MECHANISMS

Based on the previous discussion, the author feels that the subject of adhering sand defects can be best approached by analyzing the mechanisms of mass transport as follows:

Liquid State Chemical Reactions Vapor State

Practical aspects of the problem follow.

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METHODS OF PREVENTION

While the previous discussion will interest the research- oriented metallurgist, the average foundry metallurgist is more interested in the practical aspects of the problem, and more concerned with methods of preventing penetration. The methods must, of course, be simple and economically attrac- tive.

A discussion of methods may be best approached by considering the variables which the foundryman can manipulate. A substantial amount of penetration may be eliminated by the use of dense sand mixes with a high hot strength. The loss of ease of shakeout must be sacrificed in cases of expected severe penetration. Care must be exercised to ensure hard ramming, especially in fillets and protrusions. The use of more refractory and less reactive aggregates, such a zircon sand and chromite is beneficial because the amount of chemical reaction penetration is decreased. The high cost of such aggregates can be offset by selective use as a facing sand only in areas where severe penetration is anticipated.

With reference to all aggregates discussed, the void fraction cannot be reduced to a point which is effective in inhibiting vapor transport because the mixes become impossible to mold into accurate molds. Also, other foundry problems related to permeability, such as blow holes and scabbing, are aggravated.

Mold and core coatings are effective in minimizing penetration. In less severe applications a silica flour is sufficient, although it may present a health hazard, while severe applications require the more refractory zircon flour coating. Chromite flour coatings are not very effective. In all cases, it is very important that the coating application be uniform and be dried carefully. It is very easy to lose all benefits of coating through careless application.

In severe applications, multiple coats are required. The increased effectiveness of multiple applications drops off rapidly however, and anything greater than three coats is un-

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necessary. Optimum benefit is usually gained with two coats. When vapor state penetration is important, that is, with large castings and enclosed cores, the addition of a strong deoxidizer to the sand is beneficial. A good drying practice for molds and cores is essential in order to minimize the oxidizing effect of excessive amount of water vapor.

Since the penetration problem is aggrevated by a high oxygen content in the metal, all precautions taken to minimize the dissolved oxygen content will be beneficial. These include a good deoxidation practice, low pouring temperature, minimum excess superheat in the furnace, and precautions to minimize reoxidation during metal transfer and pouring.

The use of the lowest pouring temperature compatible with metal fluidity is desirable because vapor state penetration is a strong function of temperature; that is, the vapor pressure of the metallic vapors is an exponential function of temperature and the diffusivity is a function of temperature to the 3/2 power. Furthermore, high pouring temperatures lead to long solidification times thus allowing more time for liquid state penetration to occur. The high pouring temperatures also tend to cause thermal failure of the aggregate, and also cause an increase in the reaction rates involved in chemical reaction penetration.

The use of aggregates with a higher heat capacity than silica, such as zircon sand and chromite ore, are beneficial since they serve as a mild chill and help offset the effect of temperature. The use of Fe, Cu or graphite inserts (chills) in selected areas such as fillets, roots of gear teeth and other areas where the temperature is likely to remain high for a long time are beneficial since they promote an early skin formation and also help offset the other effects of high temperatures.

In conclusion, the reader must realize that the solutions to the penetration problem are very complex. There is no magic cure- all for the problem, and only careful attention to all of the many variables will produce good results.

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References 1. Meeting Minutes, AFS Committe 4F, 1987 2. M.T. Rowley, Ed., INTERNATIONAL ATLAS OF CASTINGS DEFECTS,

CIATF, American Foundrymen's Society, 1974. 3. T.P. Hoar and D.V. Atterton, "Penetration of Molten Metal into Compacted

Sand", JISI, v. 166, p. 1, (Sept. 1950). 4. T.P. Hoar, D.V. Atterton and D.H. Housman, "Influence of Ramming and of

Sintering on the Penetration of Molten Metals into Compacted Silica-Sand Mixtures", JISI, v. 175, p. 19, (Sept. 1953).

5. H. Pettersson, "An Investigation of Penetration of Steel into Molding Sands", TRANSACTIONS AFS, v. 59, p. 35, (1951)

6. N.N. Lyashchenko, "Effect of the Gaseous Medium on the Micro Con- figuration of the Surfaces of Castings", GASES IN CAST METALS, B.B. Gulyaev, ed., Consultants Bureau, New York, p. 225, (1965).

7. Y.A. Estratov, "Evaluation and Filtration of Gases in Casting Molds", GASES IN CAST METALS, B.B. Gulyaev, ed., Consultants Bureau, New York, p. 192, (1965).

8. P.L. Goodale, "Notes on Behavior of Sand Molds in Steel Foundries", TRANSACTIONS AFS, v. 38, p. 471, (1930).

9. P. Asanti, "On the Interface Reaction of Chromite, Olivine and Quartz Sands with Molten Steel", AFS CAST METALS RESEARCH JOURNAL, v. 4, p. 9, (Mar. 1968)

10. G.A. Colligan, R.A. Flinn, and L.H. Van Vlack, "Effect of Temperature and Atmosphere on Iron-Silica Interface Reaction", TRANSACTIONS AFS, p. 452, (1958).

11. S.L. Gertsman and A.E. Murton, "An Investigation of Metal Penetration in Steel Sand Cores", TRANSACTIONS AFS, v. 58, p. 595, (1950).

12. A. Kolorz and K. Orths, "On the Influence of Penetration of Molten Steel Due to Mould Material, Mould, and Metal", GIESSEREI, No. 53, p. 733, (Oct. 27, 1966).

13. E.I. Markhasev and E.S. Lugovskaya "Influence of Metal-Mould Reactions on the Formation of Burn-On", RUSSIAN CASTINGS PRODUCTION, p. 25, (Jan. 1962).

14. R.E. Savage and H.F. Taylor, "Fayalite Reaction in Sand Molds for Making Steel Castings", TRANSACTIONS AFS, v. 58, p. 564, (1950).

15. V.A. Chechulin, "Chemical Reactions of Gases Evolved From the Mold With Steel and Cast Iron", GASES IN CAST METALS, B.B. Gulyaev, ed., Consultants Bureau, N.Y., p. 214, (1965).

16. R.C. Emmons and J. Bach, "Steel Penetration", FOUNDRY, (Apr. 1955). 17. C.A. Sanders, "Are Carbonyls Causing Penetration and Burn-On in High

Density Molds?", MODERN CASTING, p. 149, (May, 1965). 18. J.M. Svoboda and G.H. Geiger, "Mechanisms of Metal Penetration in

Foundry Molds", TRANSACTIONS AFS, p. 281, (1969). 19. J.M. Svoboda and G.H. Geiger, "Diffusion of Metal Vapor Species in

Porous Aggregates", TRANSACTIONS AIME, Vol. 245, p. 2363, (Nov. 1969). 20. J. Frawley, W. Moore and A. Kiesler, "Simulating Mold-Metal Reactions in

a Small Laboratory Test", TRANSACTIONS AFS, 1974, p. 561. 21. J. Rous, "Burning-In Defects in Steel Castings", SLEVARENSTVI, No. 10,

1981, p. 413.

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22. METAL PENETRATION AND SURFACE ROUGHNESS, SCRATA Technical Bulletin No. 4, SCRATA, Sheffield, England.

23. C.G. Wagner, “Observations on the Penetration of Steel Into High Density Clay Bonded Molds With Controlled Atmospheres”, TRANSACTIONS AFS, 1979, p. 573.

24. METAL PENETRATION, BClRA Broadsheet No. 102. 25. SAND BURN-ON, BClRA Broadsheet No. 109. 26. J.M. Svoboda, “Metallurgical Aspects of Mold-Metal Interface Reactions”,

STEEL FOUNDRY FACTS, No. 354, p. 40, February 1983. 27. J.M. Svoboda, “Effect of Oxygen Content on Metal Penetration in Steel

Castings”, ELECTRIC FURNACE PROCEEDINGS, AIME, p. 28, 1967. 28. P. Poyet and R. Chevriot, “lnfluence of the Presence of Chromite in

Reclaimed Silica Sand Used in the Steel Foundry”, FONDERIE, March 1980, p. 93.

29. R.L. Naro and R.D. Tenaglia, “Influence of No Bake Binder Processing Variables on Metal Penetration in Steel Castings”, TRANSACTIONS AFS, 1979, p. 145.

30. SPECIAL REFRACTORY SANDS, SCRATA Technical Bulletin No. 2, SCRATA, Sheffield, England.

31. K. Suzuki, “Resistance of Metal Penetration of Chromite for Steel Castings”, TRANSACTIONS AFS, 1976, p. 183.

32. E.L. Kotzin, THE APPLICATION OF NON-SILICA AGGREGATES IN THE PRODUCTION OF STEEL AND HIGH ALLOY CASTINGS, Paper No. 26, 47th International Foundry Congress, Jerusalem, Israel, 1980.

33. J. Biel, K. Smalinskas, A. Petro and R.A. Flinn, “Variable Affecting Chromite Sand Performance in Molds”, TRANSACTIONS AFS, 1980, P. 683.

34. A. Petro and R.A. Flinn, “Mold-Metal Interactions Between Chromite Sand and Cast Steel”, TRANSACTIONS AFS, 1978, P. 357.

35. G.P. Kim, “Chemical Burn-On on Iron Castings”, RUSSIAN CASTINGS PRODUCTION, July 1976, p. 292.

36. A.A. Bagrov, A.M. Lyass and I.V. Valisovskii, “Influence of Fe-C Alloy Com- position on Burn-On”, RUSSIAN CASTINGS PRODUCTION, June 1976, p. 251.

37. A. Lyass “Chemical Burn-On on Iron Castings”, RUSSIAN CASTINGS PRODUCTION, May 1975, p. 202.

38. A. Draper and J. Gaindhar, “The Role of Mold Atmosphere in the Penetra- tion of Steel in Sand Molds”, TRANSACTIONS AFS, 1975, p. 593.

39. M. Nakayama, T. Kinoshita and K. Matsuda, “Metal-Mold Interface Reac- tion Rate in Steel Casting”, TRANSACTIONS OF THE JAPAN FOUN- DRYMEN’S SOCIETY, Vol. 1, May 1982, p. 66.

40. MOULD AND CORE PAINTS FOR STEEL FOUNDRY USE, SCRATA Technical Bulletin No. 3, SCRATA, Sheffield, England.

41. G.J. Vingas, “Mold and Core Washes for Heavy Section Steel Casting”, PROCEEDINGS, 38th SFSA Technical & Operating Conference, 1983, p. 324.

42. B. Ozturk and R.J. Fruehan, “Formation of SiO(g) and SiS(g) from Coke”, IRON & STEELMAKER, July 1987, p. 43.

43. A. Muan and E.F. Osborn, PHASE EQUILIBRIA AMONG OXIDES IN STEELMAKING, Addison-Wesley Publ. Co., 1965, p. 62.

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Lecture VI

Shrinkage in Steel Castings by Ronald W. Ruddle

INTRODUCTION Unsoundness arising from metal shrinkage has been a severe problem with steel castings, almost as long as steel castings have been made. This unsoundness has been the source of both expense and aggravation to steel founders and, because of its severity, has placed the steel foundry industry at a competitive disadvantage relative to most other sections of the foundry industry, for example, the iron founders. Therefore, shrinkage, and methods for combating its effects, are matters of prime importance to steel founders. In fact, shrinkage problems have become more acute, as the years have gone by, because the users of steel castings have steadily raised their soundness requirements. The increasing level of technology generally, has forced the pace here, especially in the “high- tech” industries, such as aerospace, military hardware, auto- mobile racing, etc.

The intent of this paper is to review and summarize information on the causes and types of shrinkage unsoundness, the conditions under which shrinkage unsoundness forms, the effects of shrinkage unsoundness on mechanical properties, and methods available for the elimination of shrinkage problems.

THE COMPONENTS OF SHRINKAGE Shrinkage unsoundness arises as the result of the concurrent action of several different volume changes which take place during the freezing of a steel casting. These are:

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1. The volume contraction which takes place when steel freezes. (Freezing shrinkage)

2. Volume contraction of already solid steel taking place while still-liquid metal is freezing.

3. Change in the volume of the casting as the result of mold wall movement.

The volume contraction of the liquid steel in cooling from the pouring temperature to the liquidus temperature also requires consideration. Although this volume change does not normally contribute to the formation of shrinkage voids in the casting, it does add to the volume of feed metal which must be supplied by the risers and therefore, must be allowed for in establishing riseri ng practice.

The magnitude and relative contribution of each of these factors will now be discussed. Unfortunately, reliable data in regard to these volume changes is lacking, as will be seen.

Note that, in this discussion, all volume changes are expressed as percentages of specific volume at the liquidus or solidus. This procedure is preferred to the normal procedure of expressing these changes relative to the volume at room temperature, since it is the volume change at high temperature which is important in discussing shrinkage.

Freezing Shrinkage Data on the freezing shrinkage of pure iron is available, if not markedly concordant. There seems to be no good data for the freezing shrinkage of steels. Pehlke, Jeyarajan & Wada (1) have recently reviewed all the available information for both iron and steels. They quote the following equations for the densities of solid and liquid iron near the melting point:

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In these equations, T is the absolute temperature (temp. in deg. C or +273). The density units are kg/cu.m.

If these equations are evaluated at the melting point (1538 deg. C or 1809 deg. K), the volume change on freezing is found to be about 3.34%. This value is in reasonable agreement with most of the values quoted by earlier researchers.

No reliable data seem to be available for steels.

Volume Contraction of Liquid Steel From equation (2), it may be shown that the volume contraction of liquid iron just above the melting point is about 1.19% per 100 deg. C or 0.66% per deg. F.

Pehlke et al recommend the following equations for liquid 0.5 and 1.0% carbon steels respectively, for temperatures near the liquidus temperature.

These two equations lead to the value of 1.05% per 100 deg. C (0.58% per deg. F) for the liquid shrinkage of both steels.

Volume Contraction of Solid Steel Equation (1) leads to a value for the volumetric shrinkage of solid iron at temperatures close to the melting point, of 0.87% per 100 deg. C (0.48% per 100 deg. F).

No good data exist for steels.

The contribution of this factor to the total solidification shrinkage is quite small. For example, for a steel whose freezing range is 60 deg. F, the volume contraction of the solid steel is about 0.5 x 0.48 x 60/100 = 0.14%, an amount which is barely significant compared with the contribution of the other factors to the total freezing shrinkage.

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Volume Changes Caused by Mold Wall Movement Bertolino and Wallace (2) studied the dimensional changes occurring in 0.3% carbon steel castings of various sizes, cast in molds made from a range of different materials. Increases in the volume of steel castings ranged up to about 5% for castings made in green sand, and up to about 1.3% for castings made in dry sand molds. Volume changes in silicate- bonded molds were less than 1%. The results obtained with silicate molds may be assumed to be typical of rigid molds generally, such as resin-bonded no-bake molds.

Bertolino & Wallace give much information on the influence of factors such as riser height, pouring temperature, and casting shape and size, on the magnitude of mold wall movement effects, and their paper should be consulted for further details.

Summary of Shrinkage Data From the above, the following volume change data seem the best currently for use in discussion of the shrinkage of low carbon steels.

Liquid Shrinkage: 1.2% per 100 deg. C (0.67% per 100 deg. F)

Freezing Shrinkage: 3.3%

Solid Shrinkage: 0.87% per 100 deg. C (0.48% per 100 deg. F)

Shrinkage caused by Mold Wall Movement: 0 - 5%

Thus the total shrinkage involved in the formation of porosity in low carbon steels is at least 3.5% and may range as high as 8.5%, depending on the freezing range and the amount of mold wall movement occurring. In calculating the size of risers, these numbers should be increased by 0.5 to 1.5% to allow for liquid shrinkage, depending on the pouring temperature, for a range of 4.0 to 10.0%.

The above values may provisionally be employed for high carbon steels and alloy steels in the absence of better data.

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TYPES OF SHRINKAGE POROSITY IN STEEL CASTINGS Shrinkage manifests itself in various ways in steel castings. In an incompletely fed casting in a low carbon steel. shrinkage may take several forms of which the following are the most common. One of these is the formation of massive cavities of fairly regular shape. located at the heat centers of the casting: Fig. 1 is an example of this type of shrinkage. These cavities form at the locations of the last liquid metal to freeze. They are often found at the junctions of sections, such as "L", "T", "X" and "Y" junctions. Depending on the solidification conditions and the alloy content of the steel. the walls of these cavities may be smooth or rough showing the outlines of dendrites.

An equally common type of shrinkage is "centerline" shrinkage, found on the central axes of long bar or rod castings or in the central planes of long plate castings. Again. these are locations where the last liquid metal freezes. Fig. 2 is an example of centerline shrinkage in a bar casting.

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In plate-like sections, it would be more accurate to refer to this type of shrinkage as “central-plane’’ shrinkage, since the shrinkage is no longer disposed along a line but along the central plane of the casting section. The appearance of “central plane” shrinkage in radiographs and cut surfaces of plate-like castings, is dependent on the plane of the cut or radiograph. If this is close to perpendicular to the shrinkage plane, the shrinkage appears as in Fig. 2. However, if the plane of the cut or radiograph is essentially the same as the shrinkage plane, shrinkage takes on a filamentary appearance as in Fig. 3.

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Risers, ingot castings and similar shapes, normally exhibit “piping”, which as the name implies, is a pipe-shaped cavity with one end exposed to the atmosphere. The shape of pipes in pure metals and low alloys, such as low carbon steel, approximate to the solid of revolution formed by rotating a parabola about the “Y” axis (Fig. 4). Pipes in higher alloys are less regular in shape.

An allied type of shrinkage is sometimes found on the surfaces of steel castings, notably at heat centers such as occur at sharp internal corners and edges. This kind of shrinkage resembles a pipe in that it is a tapered cavity with the wide end open to the atmosphere. However, the shape is generally quite irregular; an example is shown in Fig. 5. These punctures of the frozen skin may be accompanied by local dishing of the surface of the casting. Defects of this kind are sometimes called “external shrinks”.

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The particular characteristics of shrinkage in low carbon steels are that shrinkage cavities and pipes usually have relatively smooth walls and that centerline shrinkage is closely confined to the central axis or plane of the casting section. These characteristics are evident in Figs. 1 and 2.

Increase in the alloy content of the steel modifies the shape of these cavities. High carbon steels and alloy steels exhibit shrinkage cavities and pipes whose shape is markedly less regular than those in low carbon steels, these cavities often showing filamentary offshoots. Centerline shrinkage in the higher alloys is less strongly concentated about the central line or plane and is more diffuse in appearance. At the same time, there is a tendency for a new type of shrinkage - micro-

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shrinkage - to appear. Microshrinkage consists of tiny “pinhole-sized” cavities, typically about 1 mm. in diameter, located between the arms of the dendrites. An example of microshrinkage is shown in Fig. 6. Higher magnification reveals the interdendritic nature of microshrinkage (Fig. 7). Although microshrinkage tends to appear in the interior of the casting, away from the mold walls, it is not necessarily confined to the neighborhood of heat centers but, instead, may affect large regions of the interior of the casting. Microshrinkage is especially prevalent in regions of equiaxial crystallization which commonly occur in the central regions of castings. The occurrence of microshrinkage and the size of the cavities both increase with the size of the casting. All these changes are related to the freezing range of the steel and to the rate of freezing as discussed in the next section.

FACTORS AFFECTING THE FORMATION OF SHRINKAGE CAVITIES The formation of shrinkage cavities is not as simple a matter as might be expected. Several studies (3-8) have shown that the driving force required to form (homogeneously nucleate) a shrinkage cavity is immense. Consider the nucleation of a spherical cavity whose radius is r. The pressure in the cavity is represented by p. Surface tension forces, T, act on the walls of the cavity. It can be shown that these quantities are related by the following equation:

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This equation indicates that the pressure required to form a cavity is inversely proportional to the radius r. Consequently the pressure is huge when the cavity is very small. This analysis is an oversimplification of the situation but it does, rather clearly, show the difficulty of forming cavities in solidifying metals by homogeneous nucleation. It has been calculated that the pressures involved are of the order of thousands of atmospheres (6).

Therefore, it is now believed that cavities cannot form by homogeneous nucleation in freezing metals, but must form by a process of heterogeneous nucleation (5). In other words, the presence of a particle capable of acting as a nucleus for cavity formation is necessary.

It has also been concluded that, in the absence of suitable nuclei, the casting would be sound and shrinkage compen- sated for by the collapse of the casting walls (or drawing down of metal in the riser). This effect is thought to be, at least in part, responsible for the very high soundness of the aluminum castings produced by the Cosworth process (9, 10).

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Normally, there are plenty of nuclei so that shrinkage cavities have no difficulty in forming. For obvious reasons, they form in the last pools of liquid metals to freeze. This accounts for the tendency, in low carbon or other steels of short freezing range, for single large cavities to form at heat centers or along center- lines of sections.

In steels of longer freezing range, such as high carbon steels and many alloy steels, the dendritic nature of freezing becomes more marked and the zone of partial freezing becomes wide. Small pools of liquid metal become isolated between the arms of the dendrites and when these pools finally freeze, tiny cavities are formed there. This is the way in which microshrinkage is produced. Thus, microshrinkage is favored by increase in the freezing range. The occurrence of microshrinkage in AISI 4330 and 4340 steels was the subject of a lengthy investigation (11-14) at MIT some years ago. Among the find- ings was that the occurrence of microshrinkage is sensitive to the temperature gradient during freezing. This matter is discussed in more detail below.

EFFECT OF SHRINKAGE ON MECHANICAL PROPERTIES In the past and still today, many steel castings are made with shrinkage cavities in some heat centers (generally centers of heavy sections) and with centerline shrinkage in many parallel-walled sections. The question arises “How damaging is shrinkage and how important is it that shrinkage be eliminated?” In considering this question, it must be borne in mind that steel castings vary from castings which are lightly stressd and failure of which would not be of great moment, to highly stressed castings, failure of which would have disastrous and even Iife-threatening consequences.

Discrete cavities in heat centers may not be particularly harmful in some castings if these regions are not highly stressed and are not machined. Nevertheless, it must be recognized that these cavities do substantially reduce casting strength and may initiate cracking in service.

The effects of centerline shrinkage are more debatable. As Briggs (15) pointed out many years ago, it is probable that

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some centerline exists in many parallel-walled casting sections of castings. Since, under normal conditions of loading, one surface of these sections is generally in tension while the other surface is in compression, the central line or plane is virtually unstressed. As the result, the presence of centerline shrinkage has little weakening effect on the section. Sims (16) has called attention to the excellent service record of railroad side-frame castings despite the presence of centerline shrinkage; the same is probably true of railroad bolster castings. Brinson and Duma (38) have pointed out the elimination of centerline shrinkage by means of padding, results in only a 5% increase in the yield point.

There is much evidence to support the view that moderate amounts of centerline shrinkage have little effect on mechanical properties. Fig. 8 shows that test bars varying form radio- graphically sound to Class 5-6 shrinkage, are little different in either tensile or yield strength (17). Bending strength is like- wise not greatly affected by centerline shrinkage. Even if the region of maximum stress contains Class 2 shrinkage porosity, the loss of strength is only about 20% (Fig. 9).

However, impact properties are significantly affected by the presence of centerline shrinkage, as Fig. 10 shows (18). Fatigue strengths are also considerably affected. Torsion fatigue strengths may be reduced by up to 32% and bending fatigue strength reductions of about 15% have been noted. (18)

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Clearly, it is not possible to give a simple answer to the question posed at the beginning of this section in regard to the permissible amount of shrinkage porosity. All that can be said is that a foundry should weigh very carefully the extent to which shrinkage is acceptable in a given case, and balance this against the ease with which shrinkage can be prevented from occurring. If prevention is easy and inexpensive, it should always be done. If elimination of all shrinkage would necessi- tate the use of difficult and expensive techniques, the foundry must attempt to strike a balance between, on the one hand, the service conditions of the casting and the consequences of failure (both engineering and liability), and, on the other hand, the expense involved. In recent years, soundness requirements have steadily increased; in view of this and of the need to compete with high quality castings from abroad, it may, in future, be advisable to err on the side of shrinkage minimization. Furthermore, if steel foundries are content to make castings containing substantial amounts of shrinkage, they reduce these castings to the level of gray iron castings and the reasons disappear for preferring steel in the application being considered.

Relatively inexpensive means are now available for the elimination of centerline shrinkage (see next section) and therefore, it is suggested that there no longer seems to be reason for continuing to make castings containing more than small and insignificant amounts of centerline shrinkage. Even if centerline shrinkage present in a casting is relatively harmless, its mere presence is certainly not a commercial asset and could be a tremendous liability in the event of legal action following a failure!

The effects of microshrinkage on the mechanical properties of steel castings was investigated as part of the MIT study (11-14). Some of their results are reproduced in Fig. 11. Note that the measure of microporosity employed in this work was an arbitrary one; the actual percentage porosity was probably about 1/20th of that indicated. In this figure, the strength properties are plotted aaainst distance from the chilled ends of

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plate castings. Microshrinkage levels are low or zero near the chilled end but increase rapidly with distance from the chill. As may be seen, microshrinkage, at least to the degree experienced in this work, has little effect on tensile strength or yield point, but does significantly reduce elongation and reduction of area. Thus, the importance of microshrinkage in a given case, depends mainly on the importance of ductility to the service conditions of the casting.

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PREVENTION OF SHRINKAGE The basic principles of feeding practice for the elimination of macroshrinkage in steel castings are to supply feed metal to the heat centers present during the solidification of the casting and to ensure that freezing of the casting is “directional” towards these heat centers. Normally, low carbon steel castings freeze by “skin formation”; a thin shell of solid metal forms on the mold walls and thereafter freezing progresses by the gradual thickening of this shell or skin. If we consider the freezing of a plate section, this means that freezing takes place in a direction towards the center of the section (Fig. 12A). It is this mechanism which gives rise to centerline shrinkage. Directional solidification means that there also exists a component of freezing direction towards the heat center and attached riser, as shown in Fig. 128. Here, the solidification fronts are no longer parallel to the mold walls, but are in- clined to them. Thus, the two fronts advancing from opposing walls, form a “V”, the wide end of which is towards the heat center (Fig. 13). In this way, feed metal is readily able to flow from the riser and heat center down towards the apex of the “V”, and complete feeding and elimination of centerline shrinkage takes place.

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The principal tools available to the foundryman for the proper feeding of steel castings are:

Risers of the correct size and volume Proper location of risers Metal or graphite chills Padding Local use of “chilling” mold materials

Riser location Finding the optimum locations for risers on a complicated casting remains an art although the foundryman now has much technology to guide him. Risers should normally be attached to heat centers and, if necessary, chills or chilling mold materials should be disposed so as to ensure directional solidification towards the risers. In general, top risers are preferred because they tend to feed more efficiently than side risers and because they are more economic. However, it frequently happens that heat centers are located in the lower regions of castings; in these instances side risers should be employed. Sometimes, castings have heat centers which, by virtue of their location, are unfeedable. In such instances, the only recourse is to eliminate the effect of the heat center by chilling.

Riser Size Calculation It is not the intent of this paper to describe in detail the way in which the correct size of riser to employ, in a given instance, can be determined. This subject has been addressed in the

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literature on many occasions and the reader is referred to the references cited below for further information.

There are now many methods for the calculation of riser sizes, most of which work fairly well and some of which readily lead to optimized risers. Some of the more important techniques are briefly reviewed below. All require use by a knowledgeable person; none are foolproof.

The various methods break down into the following groups:

(a) Empirical techniques in which the correct size of riser is found from a plot of experimental data. The best known of these methods is that of Pellini and his co-workers (19) who suggest determining riser sizes from the data plotted in Fig. 14 and a “shape factor” defined as:

Casting length/(casting width + casting thickness)

The method works reasonably well with simple shapes, but suffers from the inability of the above simple shape factor to deal with more complex shapes. The method cannot handle feeding aids.

(b) Semi-empirical methods which combine some theory with practical experience. Caine’s method (20) is an example. Caine’s method is expressed in the following equation:

In this equation, F is the square root of the relative freezing times of riser and casting and Z is the volume ratio of riser and casting. A, B and C are constants determined by matching Caine’s equation to experimental results, as in Fig. 15. Caine’s method has been a good guide to the estimation of the size of conventional sand risers, but. like Pellini’s method, it cannot readily handle feeding aids.

(c) Methods based on Chvorinov’s Rule. Almost 50 years ago Chvorinov (21) showed, from heat flow considerations, that

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the freezing time of a casting is proportional to the square of its volume to area ratio:

In this equation, t is the freezing time of the casting and V and A are respectively, its volume and area. The ratio V/A is now often called the modulus of the casting. At first sight, it would seem that, if the modulus of the riser exceeds that of the casting section to which it is attached by a suitable amount - say 15 to 30% - the riser should provide a satisfactory feed. Unfortunately, this is not always so. The reason is that, not only must the riser take longer to freeze than the casting, but it must also contain enough liquid metal to feed the casting. Drainage of feed metal from the riser, in effect, reduces its freezing time and its effective modulus.

To allow for this effect a number of variants of the basic modulus method have been introduced. Wlodawer (22), who tremendously expanded usage of the modulus method, placed limitations on the volumes of his risers. Adams and Taylor (23), Ruddle (24), Marchant (25) and Prabhaker and Nehru (26) have all used a complex cubic equation to allow for the volume effect. This overcomes the objection voiced above but the equation is difficult to solve (this difficulty has been eliminated with the advent of the microcomputer).

The second term is the volume component of riser size and the third term is the volume component.

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Another, somewhat similar technique has been advanced by Creese (27).

(d) Computerized riser size calculation. In the last few years, several computer programs (28) have been introduced to calculate riser sizes. These programs all use the computer to do the tedious calculations necessary to determine section volumes, areas and riser sizes, using the methods described above. The “Crusader” program, developed by SCRATA uses Caine’s equation (20). The “Feedercalc” program (28) is based loosely on the modulus method and yet another program uses the cubic equation referred to above (29).

(e) Computer modelling. Computer simulation of the freezing of steel castings is rapidly coming to the fore. This method of predicting the course of solidification in castings should eventually permit the correct size of riser to use in a given case, to be readily and accurately determined. Sev- eral such programs now exist (30, 31) and others are in course of development (32, 33). However, it should be pointed out that none of them, at this time, enable the optimum riser to be calculated; this must be found by trial and error. These simulation programs also need extended prov- ing against experiment before they can be fully accepted.

Most of the methods currently in use for the calculation of riser sizes, although extremely valuable aids for the foundry engineer, are not exact and in many instances, ignore the influence of some important factors. For example, most methods of calculating riser sizes fail to take into account the position of the gate or gates attached to the casting or section. Most castings are bottom gated, to minimize turbulence and damage to the mold resulting from the free fall of metal introduced via top gates. Unfortunately, bottom gating combined with top risering, results in a highly adverse initial temperature gradient in the solidifying metal. If directional freezing is to take place, this temperature gradient must be reversed as soon as possible. This is normally done by making the top risers somewhat oversize so as to slow down freezing

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of the top of the casting while permitting the lower regions to cool normally. Eventually, this results in a reversal of the unfavorable temperature gradients.

Pouring rate is another important variable. Fast pouring mini- mizes temperature gradients, slow pouring enhances them. Therefore, slow pouring may be desirable if the initial temperature gradients favor directional solidfication, otherwise fast pouring is to be preferred.

Other factors not taken into account in most risering methods, include radial heat flow from the cylindrical surfaces of risers and the divergent heat flow from the edges and corners of both casting and riser - both of these effects expedite freezing.

Chills Chills are useful in the rigging of steel castings for several purposes. In the first place, they are useful to extend feeding ranges. Second, they may be used to eliminate heat centers such as those produced by bosses, ribs, etc. A third use for chills is to steepen the temperature gradients occurring during freezing.

The functioning of chills was studied in detail by Pellini et al (34-35). Chills cause rapid solidification of the adjacent steel; Pellini et al estimate that a steel chill causes rapid freezing of a thickness of steel about equal to the thickness of the chill. They warn that chills may cause hot tearing at the boundaries of the chilled region as depicted in Fig. 16.

Myskowski and Bishop developed data in regard to the use of chills to neutralize the hot spots generated by ribs. Fig. 17 shows how a 1/4 in. chill eliminates the effect of a 1 in. thick rib on a 2 in. thick plate; note that the temperature distribution produced by the 1/4 in. chill (Fig. 17C) is almost identical with that in a plate made without a rib (Fig. 17A). Comparison of Figs. 17B and 170 reveals that a 1 in. thick chill causes premature freezing of the rib area. This highlights the fact that, in this kind of application, chill thickness must be quite precisely correct. The original papers of Pellini et al give much data which can be used to find the correct size of chill to employ to neutralize ribs, bosses and appendage sections.

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The other applications of chills are discussed under the headings of “Feeding Range” and “Microshrinkage”.

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Chilling Mold Materials Sometimes, especially in the manufacture of Iight-sectioned castings, the intense chilling effect produced by a metal or graphite chill is unnecessary and may even be undesirable. In these instances, local use of a mold material which is more chilling than ordinary molding sand, may have the desired effect in eliminating hot spots andlor promoting directional freezing. Chromite sand and zircon sand are both useful in these applications.

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Feeding Range and Padding It is well-known that a riser attached to a parallel-walled section, such as a bar or plate, is limited in its ability to feed the section. This topic was investigated in detail by Pellini and his co-workers at the NRL (36). They found that there is a sound zone adjacent to the riser and another one at the end of the casting section (assuming the presence of an end). In the first zone freezing is directional because of the proximity of the riser; in the second, the high cooling rate produced by the edges and corners combined with that of the end face, again produces directional freezing for a limited distance (Fig. 13). If the section is long, there is no component of freezing towards the riser between these two regions and centerline shrinkage results.

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The NRL study covered the standard cases of bar and plates, with and without ends: the use of chills to extend freezing ranges was also examined. These four cases are depicted in Fig. 18 for plate castings. Pellini et al generalized their results in the form of equations. Steel Founders Society of America later generated some slightly different data. The SFSA work included a study of "semi-plates" of different width to thickness ratios.

Later Johnson and Loper (37) made an experimental study of bars and plates whose thicknesses ranged from 1/2 in. to 1 1/2 in., and combined their data with that of the NRL study, Johnson and Loper developed equations in terms of section modulus rather than thickness. Johnson and Loper point out that none of the data differ greatly and suggest that a single curve might cover both bars and plates, as indicated in Fig. 19.

Flemings et al (13) point out that feeding range is strongly dependent on the degree of soundness required. Citing an unchilled 1/2 in. thick plate studied in their work, they state that feeding distance is 6 in. if ordinary commercial soundness is considered; it is 4 in. for plates examined by ordinary radiographic techniques and may be as low as 2 in. or even 1 in. if

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highly sensitive radiographic techniques are used. Both Pellini et al and Johnson and Loper used ordinary radiographic methods capable of revealing cavities about 0.02 in diameter. The MIT study used a much more sensitive microradiographic technique.

Feeding range data are a most valuable tool for the rigging engineer; indeed, these data are almost indispensible when complex castings are being rigged. Many times, the rigging engineer will find that his risers will not feed to the end of a section. In this event, he may use padding to extend feeding further down the section.

Unfortunately, rules for padding are not as well established as feeding range rules. The early data of Brinson and Duma (38) is often considered to lead to excessive thickness of padding in thin sections and insufficient padding of heavy sections. The writer has obtained good results from the following equation for padding taper (39).

Metal padding is a reasonable technique to use to extend feeding distances providing the padding need not be removed after the casting has been made. However, if subsequent removal of the padding is essential, the cost becomes excessive. In attempts to avoid this difficulty, moldable exothermic materials and, more recently, insulating materials have been employed to line at least one mold wall of the section and thus to retard solidification locally (40). These materials, especially the insulating type, have been successful in greatly extending feeding down plate sections (41).

Risering Economics The need to use risers of substantial size in the production of steel castings results in feeding practice which is costly, relative to that for castings in other competitive metals, such as iron. For this reason, it behooves the steel founder to pay close attention to the cost of the technique he adopts and to use all possible means to minimize this cost. Cost-cutting

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measures which should be considered include the use of insulating riser sleeves and hot topping, breaker cores to reduce riser removal costs, and the use of chills to eliminate some risers, as described below.

Use of insulating or insulating/exothermic riser sleeve usually permits drastic reductions to be made in the size of the required risers; yields rise correspondingly from the 25-50% range to the 65 to 90% range with substantial cost benefit (23, 42).

Breaker cores may significantly lessen riser removal cost for risers up to about 5 in. dia. If breaker cores are to be effective, they must be properly designed and made in a mold material which will not distort during freezing. The design of breaker cores has been investigated by SFSA (43). The work of Sciama (44) indicates that for steel castings, the diameter of the hole in the core must be at least 45-50% of that of the riser.

It is sometimes possible to replace some of the risers attached to a casting by a chill. This is a particularly effective procedure when a long bar section, for example, the rim of a wheel casting, is to be fed by a series of risers. As the NRL research showed (45), it is usually possible to replace every other riser by a substantial chill as shown in Fig. 20.

The cost of feeding steel castings has been considered in detail by Ruddle (23, 42, 46) and others.

Microshrinkage Microshrinkage was discussed to some extent above. How- ever, little was said about the conditions required to eliminate it in steels subject to this defect.

The MIT study (11-14), mentioned earlier, showed that directional solidification, although necessary if microshrinkage is to be prevented, is not alone sufficient to eliminate the defect in steels of the 4330 and 4340 types. It is additionally necessary that, during solidification, the temperature gradient towards the riser be extremely steep, in excess of 200 deg. F/in. (13). In a sand casting, this implies that the casting must freeze very quickly, at a rate which can only be produced in a

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small casting by heavy chilling. Fig. 21 shows microporosity (again in arbitrary units) plotted against distance from the chilled ends of 1/2 in. thick plate castings. This figure also contains data for unchilled plates. As Fig. 21 shows, the heavy chill suppresses microshrinkage for a distance of about 1 in. down the 5 in. plate. It was not found possible to completely suppress microshrinkage in 3 in. long plates.

Hence, it appears that fine microshrinkage can only be prevented with the greatest difficulty in small castings in steels of relatively long freezing range. Prevention of microshrinkage in larger sand castings in these steels would appear to be a vir- tual impossibility, if highly sensitive methods of detection are employed.

Casting Design Much of the difficulty encountered in making sound steel castings is the result of poor casting design which, itself, often arises because part designers are lacking in knowledge of the casting process and because of insufficient communi- cation between the designer and the foundryman in the early

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stages of the design process. It cannot be too strongly empha- sized that it is most desirable that discussions take place with the designer before the design is set. If this is done, undesirable features of casting designs, such as unfeedable heat centers, inadequate radii, reverse tapers, intersections of several ribs at one location, etc. can almost always be avoided.

SUMMARY AND CONCLUSIONS This review of the current state of knowledge on shrinkage in steel castings is summarized below:

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1. Reliable data are lacking for most of the volume change components involved in the freezing of steel and which are responsible for shrinkage defects. However, reasonable estimates can be made from data for pure iron and other information.

2. Shrinkage porosity can take many forms in steel castings. Discrete cavities and centerline shrinkage are common in low carbon steels. In alloy and high carbon steels of relatively long freezing range, gross shrinkage tends to be more dispersed and microshrinkage is also usually present.

3. Shrinkage porosity forms during freezing, by heterogeneous nucleation, i.e. cavities are nucleated by particles of foreign material.

4. Centerline shrinkage, at least in moderate amounts, does not greatly affect tensile and bending strength. Impact and fatigue strengths are affected to a greater extent. Micro- shrinkage mainly affects ductility.

5. Methods for the elimination of shrinkage have been reviewed. Many techniques exist and most are capable of giving good results in the hands of an experienced and knowledgeable individual. It is desirable to pay particular attention to the cost of implementing these measures.

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TROUBLESHOOTING THE STEELCASTING PROCESS