a hypothesis for cast iron micro structures

A Hypothesis for Cast Iron Microstructures

JOHN CAMPBELL

The various microstructures of cast irons are reviewed, including carbidic and graphite forms(flake, compacted, spheroidal, and undercooled, etc.), exploring whether the presence ofexternally introduced defects in the form of oxide double films (bifilms) in suspension in meltsseem to provide, for the first time, a uniform explanation for all the structures and theirproperties. Silica-rich oxide bifilms provide the substrates on which oxysulfide particles form,nucleating graphite. The presence of the film provides the favored substrate over which graphitegrows, which leads to the development of flake graphite. The addition of limited Mg to formcompacted graphite destroys all but a remnant of the silica-rich bifilms. The oxide film remnantis stabilized by the presence of the graphite nucleus, which causes the graphite to grow unidi-rectionally in a filamentary form. The addition of excess Mg destroys all traces of the oxidebifilms, leaving only the original nuclei, around which graphite is now free to entirely enclose,initiating the spherical growth mode. Undercooled graphite is the true coupled growth form,nucleated at even lower temperatures in the absence of favorable film substrates in suspension;the graphite adopts a continuous growth mode in a matrix of austenite. Carbides in mottled andwhite irons form on the oxide bifilms that often lie along grain and interdendritic boundaries,which explains the apparent brittleness of these strong, hard phases. In most cases of non-spheroidal growth modes (flake and misshaped spheroids), it is proposed that the impairment ofthe mechanical properties of irons is not strongly determined by graphite morphology but by thepresence of oxide bifilms. Spheroidal graphite iron has the potential for high properties becauseof the absence of bifilms.

DOI: 10.1007/s11663-009-9289-0� The Minerals, Metals & Materials Society and ASM International 2009

I. INTRODUCTION

THE forms of graphite in cast irons have been thesubject of intense interest andhuge research effortsmainlysince the 1950s, but a full understanding has been elusive.Readers are referred to the review by Loper[1] for a wide-ranging synopsis covering many details not included inthis study. Here, a different review is made of theliterature, exploring the possibility of a unifying approachbased on the hypothesis that oxide films (as doubled-over‘‘bifilms’’) are present in liquid irons.

Recently, a comprehensive understanding of themicrostructure of Al-Si alloys has been proposed interms of bifilms, explaining both the mechanism ofmodification and the structures of hypoeutectic andhypereutectic alloys.[2] Bifilms (double films), usuallyoxides, are introduced into molten metals at every pouror stirring event. These surface films, which are doubledup during the process of entrainment into the bulk, seemto be of major significance for the development ofsolidification structure.[3,4] Their outer surfaces seem to

be favored substrates for the precipitation of many, ifnot all, second phases during solidification, whereas theinner unbonded interface acts as a crack. The films aregenerally so thin they are effectively invisible to casualobservation, which leads to a liquid that is invisiblycracked and, on solidification, to a solid whose cracksmay or may not be visible.Nakae and Shin,[5] among many others, have drawn

attention to the analogous features of Al-Si and Fe-Calloys. This article is an extension of the bifilm hypoth-esis, which is apparently valuable to an understanding ofthe Al-Si system as well as to a possible understandingof the various morphologies of carbon in the form ofgraphite and carbides in the Fe-C alloy system. Natu-rally, the presence of invisible defects in opaque liquidsis not easily confirmed directly, although it will not beimpossible in the longer term. In the meantime, thisarticle surveys the experimental evidence from theliterature to ascertain whether indirect evidence issupportive of this hypothesis.

II. THE EVIDENCE FOR OXIDE BIFILMSIN CAST IRON

De Sy[6] has shown that liquid cast iron generallycontains significant quantities of oxygen in solution inexcess of its solubility. He concluded, on the basis ofcareful and rigorous experiments, that the undissolvedfraction of oxygen was present as SiO2 particles.Interestingly, by heating to 1550 �C, he confirmed the

JOHN CAMPBELL, Emeritus Professor, is with the Department ofMaterials and Metallurgy, University of Birmingham, Edgbaston,Birmingham B15 2TT, UK. Contact e-mail: [email protected].

This article is based on a presentation given in the ‘‘3rd ShapeCasting Symposium,’’ which occurred during the TMS Spring Meetingin San Francisco, CA, February 15–19, 2009, under the auspices ofTMS, the TMS Light Metals Division, the TMS SolidificationCommittee, and the TMS Aluminum Processing Committee.

Article published online September 1, 2009.

786—VOLUME 40B, DECEMBER 2009 METALLURGICAL AND MATERIALS TRANSACTIONS B

expectation that the SiO2 solids dissolved because theybecame less stable than CO but reappeared on coolingonce again. Hartman and Stets[7] report not only thepresence of SiO2 in suspension but also olivine,2(Mg,Fe)OÆSiO2 is found in those irons that containMg. Hoffman and Wolf[8] find a variety of oxides,including SiO2, FeO, and MnO among others, butsuperheating and holding at high temperature elimi-nated many of these. In an elegant study of thethermodynamics, Mampaey and Beghyn[9] show howmainly SiO2, together with some FeO, forms in a typicalmelt when cooling from 1480 �C to 1350 �C.

It seems reasonable to speculate that these oxidesalmost certainly would not be compact spheres, cubes,rods, etc. but would most likely be in the form of films.Only films would have a sufficiently low Stokes velocity(one or two orders of magnitude lower than particles asa result of their greatly increased drag) to remain insuspension for long periods of time associated with theseexperiments, as well as the long periods during whichirons are held molten in holding furnaces.

In any case, of course, the film morphology is to beexpected. During melting in the cupola as droplets ofiron rained down, the natural enfolding of the surfacefilm of SiO2 of each droplet would ensure a naturalpopulation of SiO2-rich bifilms. Additional treatmentsor handling, such as pouring actions, stirring in induc-tion furnaces, and the oxide introduced from the surfaceof the charge (whether steel, pig, or foundry returns),would increase this already large, natural population.Even so, it is well known that iron from electric furnacesis more liable to chill formation problems in thinsections than cupola iron, and this effect has been widelyaccepted as the loss of nuclei (in agreement withproposals made in this article), especially during anextended time in holding furnaces or pouring systems.

Silica-rich oxides in irons are not stable in someconditions. For instance, they will be expected to go intosolution if the melt is held above approximately 1450 �Cfor any length of time.[9] Also, they are rapidly reducedto Si metal, and their oxygen is removed as MgO whenMg is added to the melt. A different population of oxidebifilms definitely exists in ductile irons that have beencast turbulently. In this case, the presence of Mgstabilizes the magnesia, MgO, in the surface film,although Si might also contribute, thus forming amagnesium silicate MgOÆSiO2, which is equivalent toMgSiO3. Both magnesium oxide and magnesium silicateare extremely stable and represent permanent damagefolded into the liquid metal and subsequently trans-ferred to the casting. Although the bifilms are known tohave an initially compact morphology as a result of theturbulence during their formation and are relativelyharmless as cracks, the subsequent straightening of thesebifilms by various natural processes, such as the growthof dendrites, which creates extensive planar cracks, iscommon; this process results in extraordinary structures(Figures 1 and 2)[10] and the phenomenon of brittlenessin so-called ductile iron in the form of plate fracture, asdiscussed earlier by the author.[3]

Thus, at least two different populations of oxidebifilms seem to exist in suspension in liquid iron, which

include the following: (1) the hypothesized silica-richbifilms as a natural population in equilibrium with themelt, the amount of silica-rich phase being predictableby thermodynamics, and (2) the known magnesia-richbifilms in ductile irons are the result of mechanicalaccidents that involve the turbulent entrainment of the

Fig. 1—A so-called ductile iron casting illustrating brittle fracture asa result of dendrite-straightened bifilms.[10]

Fig. 2—Optical micrograph through the fracture surface of ‘‘plate-fractured’’ ductile iron shows misshapen graphite nodules growingon bifilms straightened by dendrite growth.[10]

METALLURGICAL AND MATERIALS TRANSACTIONS B VOLUME 40B, DECEMBER 2009—787

surface film into the bulk liquid, which results innonreversible damage.

In passing, it is worth drawing attention to thepossibility that other nonoxide bifilms are to be expectedin cast irons. For instance, nitride bifilms probably formin cast irons giving rise to the ‘‘nitrogen fissure’’ defectsassociated in the past with high nitrogen binders.[3] Also,of course, carbon bifilms exist in irons in those cases inwhich hydrocarbon gases from the mold disassociate onthe hot surface of the melt, so that if surface turbulenceoccurs, the resulting carbon film can be folded in. Suchlustrous carbon defects are not necessarily permanentbecause the carbon can slowly dissolve, which explainswhy such features are observed in only relatively thinsection castings but not in thicker sections.[3] In thisreview of the microstructures of cast irons, only oxidebifilms seem to have significant roles.

III. GRAPHITE NUCLEI

Mizoguchi et al.[11] have demonstrated that austeniteis ineffective in nucleating graphite. In fact, they findthat undercoolings below the liquidus of between 200 �Cand 400 �C are required to trigger nucleation byaustenite. A more unfavorable nucleus would be difficultto imagine. The question therefore arises, What doesnucleate graphite? This question is all the more intrigu-ing following the work by Mampaey and Xu,[12] inwhich they found that a single population of nucleicould explain both gray and ductile irons.

There is a growing consensus that both flake andspheroidal graphite nucleate on similar, if not identical,nuclei (for instance, Warrick[13]) composed of particlesof complex oxides and sulfides. This was the conclusionreached in the first study after the development in theU.K. of the microprobe analyzer in 1974.[14] Manyconfirmations of this finding have since been made (forinstance Skaland[15]), which suggests that the oxysulfidemix of the various elements will have a spectrum oflattice spacings ensuring that at least part of thecompound will match graphite, and therefore possiblyconstitute a favored substrate. As an example of anexcellent recent study, while working on preconditioningtreatments for gray irons (treatments involving smalladditions of elements such as Al prior to inoculation -possibly to enhance the population of naturally occur-ring nuclei in uninoculated irons), Riposan[16,17] definesa three-stage model for the nucleation of graphite:

(a) Small oxides (<2 lm diameter) are formed in themelt (from the preconditioner, usually containing Aland/or Zr, leading to the oxides consisting mainly ofalumina or zirconia).

(b) Complex sulfides (<5 lm diameter) nucleate on theoxides (mainly based on MnS with low Mn/S ratiobut complicated by additions, particularly Ca, Sr,and Ba). The complex sulfides form a shell aroundthe central oxide.

(c) Graphite nucleates on parts of the sulfide shell.

Riposan[16,17] finds that MnS shells with Mn/S highratios are numerous in the matrix but do not seem to

nucleate graphite. Only those shells with low Mn/S ratioseem to act as nuclei. The mechanisms proposed byRiposan[16,17] are highly convincing and are in agree-ment with the general consensus that few nuclei existprior to inoculation but are enhanced in effectivenessand number by preconditioning and inoculation, andtheir final composition consists of mixed oxides andsulfides. Many authors have reported the beneficialeffects of S, for instance Chisamera et al.[18]

In this article, the mechanisms proposed to explainthe various morphologies of graphite are based on thepossibility of nucleating particles (probably based onoxysulfides as described by Riposan[16,17]) effective forall types of graphite, but the subsequent growth of thegraphite is affected by the presence or absence of oxidebifilms of different types. This approach is the first tosuggest separate functions of the graphite nucleatingparticles and graphite growth substrates.

IV. FLAKE GRAPHITE IRON (FGI)AND INOCULATION

Uninoculated iron is characterized by poor control ofthe graphite morphology. Flakes occur but are relativelyfew in number and uncontrolled in size. The relativelyfew opportunities for the carbon to precipitate lead torelatively large regions of the iron elsewhere beingsupersaturated with carbon. Thus, carbide precipitationis likely in places. The mechanical properties of the ironare generally poor. In general, as will become clearduring the progress of this account, it seems that somenuclei exist prior to inoculation, but their number andeffectiveness cannot be relied on. The gradual introduc-tion of the inoculation process occurred from about1920 onward, and its development continues to thepresent day (for instance Skaland[15] and Hartunget al.[19]). The inoculation process was found to increasegreatly the number of nuclei available, giving a copiouscrop of graphite flakes of good uniformity of size, with areduced tendency to carbide formation and a conse-quent benefit to the mechanical properties and machin-ability of the iron.The inoculation process is often carried out by

addition to the melt of granules of a graphite promotersuch as a ferrosilicon, which contains perhaps 50 to75 wt pct Si. Interestingly, the melting point of such aninoculant is close to 1210 �C, which is below liquid ironcasting temperatures, usually in the range 1350 �C to1400 �C. Thus, the added particles melt. However, theparticles take time to melt and time to disperse, formingtemporary supersaturated regions of liquid rich insilicon surrounding the melting and dissolving inoculantparticles. Hurum[20,21] was the first to draw attention tothis phenomenon, but it has been studied by severalothers since (for instance Fredriksson[22]). Hardinget al.[23] show how this region is effectively undercooledby several hundred degrees Celsius as a result of itsconstitution of nearly undiluted graphite promoter, thusproviding a region in which a high driving force existsfor the nucleation of graphite. Inoculation with graphiteor SiC, and so on, acts similarly to promote regions of


high-constitutional undercooling, which encouragesgraphite nucleation and growth on any suitable avail-able substrates that happen to be in this region. Theoverwhelming driving force explains the wide variationof successful nuclei, which, in other circumstances,would be expected to be of only mediocre, if any,effectiveness.

Because of the additional trace alloying elements inthe inoculant (particularly group IIA elements of theperiodic table, including Mg, Ca, Sr, and Ba), the MnSshells around the oxide centers are modified in theseregions in which the additions are concentrated (prior totheir dissipation and dilution in the melt). All this istaking place in the region highly constitutionally under-cooled with respect to graphite, thus naturally providingthe driving force for nucleation of graphite in preciselythe location needed for maximum effectiveness.

Bearing in mind that many (if not all) second phasesprecipitate on bifilms as preferred substrates,[4] it seemsreasonable to assume that these new graphite nucleiwould also preferentially nucleate on bifilm substrates.Thus, graphite would, in turn, nucleate from thosenuclei already sitting on the oxide bifilms. Figure 3schematically shows a graphite nucleus formed on anoxide bifilm approaching an undercooled region pro-vided by a dissolving inoculant particle. On entering thisregion, the nucleus experiences a massive driving forceas a result of hundreds of degrees of effective underco-oling, which forces graphite to form around the nucleus.

Harding et al.[23] point out that once nucleated in theregions of high driving force for initiation, the graphiteparticles attached to their nuclei now will emerge fromthese regions into the general melt where they willbecome unstable and start to redissolve. The observa-tions by Loper and Heine[24] confirm that graphite cannucleate in both hypoeutectic and hypereutectic irons at1400 �C, well into the liquid range, high above theexpected liquidus temperatures. Mampaey[25] confirmsthat graphite forms in the melt prior to the appearanceof austenite. (These observations are contrary to expec-tations based on the equilibrium diagram based, ofcourse, not only on equilibrium behavior but also on thebehavior of alloys of perfectly uniform composition,neither of which applies during the mechanism ofinoculation of cast irons.)

Given sufficient time, all the inoculant particles willhave melted and dispersed, leaving no pockets ofundercooling, and all the embryonic particles ofgraphite nucleated in the undercooled regions will havere-dissolved. This is almost certainly the phenomenonknown to all foundry personnel as ‘‘fade’’ of theinoculation effect. However, graphite embryos on theirnucleating particles will take time to go into solution,so that if the melt is cast without undue delay, manywill survive to reach the regime close to the freezingtemperature where they will now start to regrow, asobserved experimentally by Loper and Heine.[24] Feestet al.[26] find that although the Si-rich inoculant parti-cles disperse relatively rapidly, the graphite formedrapidly in these regions is slow to redissolve. This isreasonable because the graphite initially grows quicklyin the highly constitutionally undercooled region, but itwill subsequently dissolve in the open melt ratherslowly because the graphite will find itself only mod-estly above its equilibrium eutectic temperature, so thatthe matrix would be effectively nearly saturated incarbon.The newly forming graphite cannot grow completely

around the nucleating particle because the particle itselfhas itself grown on the planar bifilm substrate so that atleast one of its faces is inaccessible (Figure 4(c)) (workby Riposan[16,17] indicates that only parts of the nucleantparticles are active nucleation sites, so that initial growthis rather irregular). The silica-rich bifilm will form a‘‘next best’’ substrate for graphite, so although insuffi-ciently favored to cause nucleation, it is sufficientlyfavored to support the subsequent growth of thegraphite. Thus, in gray irons, the graphite extendsacross the bifilm, leading to the fairly flat morphology offlakes in gray iron. The flakes grow in regions ahead ofthe solidification front (i.e., above the general eutecticfreezing temperatures) because of the energeticallyfavored growth of graphite on the oxide substrates insuspension (Figure 5).[27]

The growth morphology of graphite, extending inthe directions in its basal plane, would favor thestraightening of the bifilm (Figure 3(d)). The bifilmwould be expected to be extremely thin, possiblymeasure in nanometers, its minimal rigidity exertingnegligible constraint of the advancing graphite crystal.The freedom from restraint would explain the devel-opment of relatively perfect crystals of graphite asobserved growing ahead of the coupled eutectic graph-ite (Figure 5).

Fig. 3—The mechanism of inoculation: (a) a bifilm in suspension inthe melt together with nuclei already attached from trace contami-nants or preconditioners; and (b) the bifilm floating into a region ofhigh constitutional undercooling surrounding a dissolving inoculantparticle, initiating a graphite flake.

Fig. 4—(a) A bifilm with precipitated nucleus from natural contami-nants or preconditioners; (b) additional nuclei provided by inoculant;(c) graphite nucleation on nuclei in regions of high effective carbonsupersaturation as in Figure 3(b); (d) growth of graphite flakes alongthe length of the bifilm, straightening the bifilm, with a consequentcentral planar crack in the graphite flake.


The mechanism proposed above explains the growthof flake graphite from nucleating particles introduced byinoculation. Particles that seem to be nuclei for theinitiation of flakes have often been observed, and theseparticles seem likely to be a universal phenomenon inboth flake and nodular irons.[16,17,28]

Eventually, the advancing solidification front willovertake those flakes growing on bifilms floating freelyin suspension in the liquid. Thus eventually, these freelyfloating flakes will become incorporated into the solid(Figure 5). Thus, it would be expected to be common tosee gray irons with two separate populations of flakes:(1) those formed by free growth in the liquid (primaryflakes) and (2) those formed by coupled growth at lowertemperatures (eutectic flakes). The coupled growthmode is discussed below. A bimodal distribution ofgraphite flakes is therefore to be expected in most grayiron microstructures, and it is clearly observed inFigure 6.[29]

Less obvious but important bimodal distributions arealmost certainly common, as may be inferred from the

work of Enright et al.[30] On the first occasion that theseauthors used the highly sensitive fractal analysis tech-nique to assess microstructures of cast irons, theyobserved this result. Thus, unless this is an extraordi-nary coincidence, the populations of primary andeutectic flakes should be expected to be mixed in most,if not all, flake graphite irons, even though, of course,such mixtures may not always be obvious to casualobservation.In agreement with the mechanism described in this

article, Goodrich[31] attributes the Type C iron (ASTMA247), which is characterized by large, straight flakes,with some branching, to the result of the growth of theflakes in the liquid, unencumbered by the presence ofaustenite. Primary flakes seems to be a good name forthese forms. He calls these proeutectic flakes. Theyoriginate in suspension in the melt and therefore canfloat to the upper regions of a casting. The morecommon Type A graphite flakes are similar, displayingonly minimal irregularity, suggesting a similar originand behavior in the melt. Loper and Fang[32] use deepetching to reveal what they call pre-eutectic flakes, whichexhibit elegant hexagonal symmetry and are apparentlylargely free from defects. For many other irons, thepresence of a dense mesh of austenite dendrites con-strains the size and shapes of flakes and prevents anysignificant buoyancy effects.[1]

More usually in castings, the graphite flakes are seento branch relatively frequently. In terms of the bifilmsubstrate, this is straightforwardly understood from theirregular structure of the bifilms. During their entrain-ment from the liquid surface into the bulk melt, theytend not to entrain as nicely parallel double films, butinstead they entrain as randomly folded, messy struc-tures. Thus, folds leading to parts of the double film atirregular angles to the main bifilm fold are to beexpected and would account for the branching ofgrowing flakes.Experience of variable performance is also to be

expected. For instance, on the Monday morning afterthe melt has been held for the weekend, operatorscommonly find the iron has poor graphite structure,particularly if the iron has been held in an acid-linedinduction furnace. Riposan[16,17] shows that this is atleast partly the result of the gradual loss of the graphitenuclei. Part of the effect would also be expected to beloss of bifilm substrates. Similarly, iron heated to hightemperature suffers a similar degradation of graphitestructure, almost certainly as a result of the dissolutionof the bifilms because of the instability of SiO2 aboveapproximately 1450 �C in the presence of carbon. Itwould be interesting to know whether the melt, afterlosing its silica-rich bifilms at high temperature, wouldregain its good solidified structure when cooled onceagain, because although de Sy[6] reports that the silicareappears in the melt on cooling, without some kind ofsurface turbulence, the form of the silica may not be abifilm, nor even a film, but it may be a compact particle.As such, it is not likely to be a good substrate for thedevelopment of a good flake structure. Even so, the finalpouring of the melt into the mold may provide sufficientturbulence to address this problem, which makes the

Fig. 5—Extrinsic initiation of straight graphite flakes in the liquid,ahead of coupled growth of eutectic graphite and austenite.[27]

Fig. 6—Two populations of flakes: extrinsically formed on bifilmsand intrinsically grown as a coupled (undercooled) eutectic.[29]


problem essentially invisible to those attempting tostudy the effect.

In agreement with the prediction that graphite growson oxide bifilms, the effect of oxygen addition to themelt during the pouring of iron into the mold isdemonstrated by several authors. For instance,Basdogan et al.[33] and Chisamera et al.[18] foundoxygen to be highly effective in converting carbidicirons into beautifully ‘‘inoculated’’ flake graphite irons.Liu and Loper[34] found that oxygen was necessary tonucleate kish graphite on the surface of gray iron melts.Larger quantities of kish were formed at temperaturesbelow 1400 �C, below which SiO2 is stable, and kish wasnot observed in Si-free melts. Moreover, Johnson andSmart[35] describe a critical experiment in which they usesophisticated Auger analysis to prove that two or threeatomic layers of oxygen (and interestingly, sulfur) arepresent on fracture surfaces of gray irons that wereadjacent to a graphite flake (fractured and observed inhigh vacuum). In contrast, the hollows in the fracturesurfaces of spheroidal graphite matrix that had con-tained spheroids exhibit no oxygen. This behavior isexactly predicted by a bifilm hypothesis: If flake graphiteformed on one side of the oxide bifilms, then the fracturesurface would necessarily reveal the oxide, but oxideswould be absent in the case of spheroidal graphite iron(as discussed below). For most intermetallics precipitat-ing on one side of a bifilm, the halves of the bifilmappear on both fracture surfaces. However, gray ironhas a curious behavior in that nothing seems to attach tothe graphite, so that both halves of the bifilm seem tohave attached to the matrix. It is right to question,therefore, what role has the bifilm played in the fracturebecause its central unbonded interface has not seemed toact as the decohering surface on this occasion. In fact, itseems likely that the bifilm will have nucleated thedecoherence, but the continuing decohering actionseems to follow the graphite interface. This nonattach-ment (perhaps we should say ‘‘active detaching’’)phenomenon between graphite and other phases isobserved elsewhere in bubble trails in gray irons as wellas in lustrous carbon films detached from the matrix andadhered to the sand mold.[36]

Although all the above discussion relates to silica-based bifilms, evidence indicates that alumina-based, orpossibly Al-containing Si-based bifilms (for instance,based on mullite or other stable alumino-silicate com-pound) exist. Carlberg and Fredriksson[37] find that castirons based on Fe-C-Si exhibit fine graphite structures,whereas those based mainly on Fe-C-Al display coarsegraphite flakes. Chisamera et al.[18] confirms that con-ventional gray irons that contain Al develop coarsegraphite flakes. It would be reasonable to expect thatsilica-rich and alumina-rich bifilms would have differ-ent mechanical and chemical characteristics and, there-fore, would develop different qualities of precipitatedgraphite.

The relatively poor mechanical properties of gray ironseems likely to be more to do with the presence of bifilmcracks down the centers of graphite flakes (or the sidesof graphite flakes if graphite grows on only one side ofthe bifilm— the impression now is that the flake has

decohered from the matrix) rather than any intrinsicweakness of the graphite itself. It seems possible thatgraphite is strong in tension perpendicular to its basalplane, because although the interplane bonding is notstrong, which leads to the familiar easy shear in its freecondition, there are approximately 1015 bonds per molein the graphite surface. The huge number of bonds willensure high strength in tension in a situation in whichthe graphite is encapsulated, preventing easy failure bypeeling or shear. A crack down the center of a flake isobserved in Figure 7.[38]

In their studies of crack initiation and propagation inirons, Voigt and Holmgren[39] report many centerlinecracks in graphite flakes plus some decoherence from thematrix.The above discussion relates to those graphite flakes

growing freely in the melt giving rise to large, randomlyoriented flakes which tend to float or settle irregularly,creating what has been called in the past an ‘‘anoma-lous’’ eutectic.

V. COUPLED EUTECTIC GROWTHOF GRAPHITE AND AUSTENITE

In this section, we move the focus from ‘‘anomalous’’to the truly regular, ‘‘classic’’ eutectic form, in whichaustenite and graphite grow in a coupled mode.In the absence of suitable nuclei that have formed on

oxide substrates in suspension in the melt, the carbon insolution will be unable to precipitate. Thus, the melt willcontinue to undercool until the undercooling finallybecomes sufficient to provoke precipitation on someother (less favorable) substrate. Only relatively few suchnuclei will operate, activating in those parts of the meltthat are especially cool, such as those regions close to

Fig. 7—Graphite flake exhibiting a central crack (the solid state pre-cipitation of surrounding temper graphite is also fractured off).[38]


the mold walls. The subsequent evolution of heat willinhibit other nuclei from becoming active. All thecoupled eutectic forms described below are thereforenot especially influenced by the (rather rare) nucleationevents. They are mainly influenced by the continuousgrowth process.

At modest undercoolings, the coupled growth takesthe form of rosettes, which are often called ‘‘cells’’(Figure 8).[40,41]

Thus, it seems likely that a single initiation event on anucleus, which is often sited on the mold wall, expandsthe coupled growth front as a hemisphere to form therosettes, or cells (Figure 8). The cells are beautifullyregular structures with interflake distances now strictlycontrolled by diffusion in the boundary layer immedi-ately ahead of the advancing front. Whether the rosetteform is a strictly coupled growth, not requiring thepresence of bifilms, is not clear. For instance, it may besome kind of aggregate of bifilms that, the graphitehaving nucleated, subsequently develops along radiallyoriented bifilms to generate the appearance of a singlegrowth phenomenon. Those bifilms present betweenflakes, at less than a diffusion distance from neighbors,will lose any graphite to their neighbors by a naturalcoarsening process, and so these bifilms become invis-ible. Thus, a clearly formed rosette structure may becapable of generating from a random morass of bifilms.A careful examination of graphite flakes to determinewhether they contain bifilms, and an examination ofinterflake regions to check for bifilms, will help to settlesuch questions.

At lower undercoolings, finer, more highly under-cooled eutectic, which is sometimes designated Type Dor E according to ASTM specification A247, seems ingeneral to have been avoided for general engineeringcastings. This is possibly because the interflake diffusiondistance is now so small that only ferrite can be formed,limiting the strength of such irons.

During coupled growth, flakes have to realign theirgrowth direction continually because of the intrusion oftheir neighbors into their growth space. Because the

growth direction of graphite is mainly parallel to thebasal (0001) plane, this means that the crystal has todevelop faults to allow it to change direction. Probablyall coupled eutectic graphites experience this effect.However, in particular, this explains the ‘‘coral’’ type ofgraphite morphology (to be discussed later), which ishighly faulted.[42]

We would expect, therefore, that fine graphite mor-phologies, types D and E graphite, would be highlyfaulted, containing high defect densities, whereas rosette(or cell) graphite would represent an intermediate caseas a result of its larger spacing. Primary flake graphite,as mentioned above, would be expected to contain theleast faults.

VI. SPHEROIDAL GRAPHITE IRON (SGI)(DUCTILE IRON)

When sufficient magnesium is added to the melt, theoxide bifilms are completely eliminated. In the case ofsilica-rich bifilms, the silica will be reduced by magne-sium to (1) silicon metal, which will dissipate intosolution in the matrix, and (2) solid magnesium oxidethat will precipitate probably on the pre-existing nucleithat originally sat on the films, augmenting theseoriginal particles. The reaction is simply

SiO2 þ 2Mg ¼ Siþ 2MgO

The total loss of bifilms means that only solids remain-ing in suspension in the melt are the original particulatenuclei, possibly augmented by additional MgO. If sulfuris also present in the melt, then the MgO is likely tocontain a component of MgS. These compact nuclei arenow the only nucleation sites available for the precip-itation of graphite. The precipitating graphite growsover the compact nucleus, wrapping completely aroundit so as to form a compact initiating morphology.The disappearance of the bifilms and the initiation of

spheroids are shown schematically in Figure 9. The‘‘wrapping around’’ process (Figure 9(d)) may consist ofrenucleation of many separate microscopic grains ofgraphite on favorable fragments of the oxysulfide sur-face. The growth mode is probably some kind of additionof carbon atoms to spiral growth steps generated by0001h i oriented screw dislocations (Figure 10).[43] In thisway, the radial structure of graphite nodules developsfrom the graphite grains growing radially out from thecompact nucleus to form the familiar approximatelyspherical nodule (Figure 11).[44]

The careful study by Riposan[16,17] on flake graphiteirons shows that the particulate nuclei do not seem, ingeneral, to be completely wrapped around by graphite,but graphite forms preferentially on isolated regions ofthe particle. On the one hand, if such a behavior alsoexists for ductile iron, as seems possible, then thisunpromising start to nodularity will emphasize thepossible importance of plastic constraint of the sur-rounding austenite matrix as will be discussed below. Onthe other hand, the addition of Mg, in addition toeliminating the silica-rich bifilms, might also somehow

Fig. 8—A scanning electronic microscope (SEM) image of a deeplyetched rosette of flake graphite, expanding to form a cell (courtesyof Fras et al.[40,41]).


affect the ability of the nuclei to work more efficiently, sothat a complete wrapping around effect might bepromoted. Subsequent research will clarify this point.

Later, as the graphite grows, Stefanescu[45] concludesin his review that all the evidence points to nodulesinitially growing freely in the liquid, subsequentlydeveloping a shell of austenite, and finally contactingand becoming incorporated into an austenite dendrite.A minor modification of this development may beenvisaged, in which the graphite nodule does not grow ashell of austenite until it contacts an austenite dendrite.At that moment, a shell of austenite would be expectedto wrap itself rapidly around the nodule. Painstakingmetallography would be required to clarify this detail.Anyway, whatever the finer details of the encapsulationprocess, the shell of austenite seems a key featureassociated with the growth of spheroids.

It seems possible that the spherical morphology of thegraphite nodules may be encouraged by the mechanicalconstraint provided by the nodule having to force itsgrowth against the resistance provided by its surround-ing shell of austenite.[46] Many studies have clearlyrevealed the deformation of austenite dendrites by thegrowth of internal nodules (Figure 12).[47]

This lumpy morphology has been attributed tovarious mechanisms, all of which are likely to contributeto some degree, as follows:

(a) Ruxanda et al.[48] and Stefanescu[45] assume theprotrusions to be the natural growth shapes arisingfrom cooperative growth of austenite and graphiteby diffusion from the liquid.

(b) Buhrig-Polackzed and Santos[49] indicate in a sche-matic illustration that the contact between nodules

with their austenite shells and the austenite dendritesresults in their mutual assimilation, to create alocal bump on the dendrite. (Some subsequent sur-face smoothing driven by surface energy would beexpected to occur rapidly.)

(c) Deformation of the dendrite by plastic flow, locallyexpanding the surrounding solid to accommodatethe increasing volume occupied by the graphite hasto be important. This effect seems to have beengenerally overlooked, but it is important and worthyof examination, as discussed below.

The pressure developed in a thick spherical shell(Figure 13) expanding plastically because of internalpressure is quantitatively expressed by[50]

P ¼ 2Y ln b=a ½1�where P is the internal pressure, Y is the yield stress, andb and a are the external and internal radii of the shell,

Fig. 10—The probable structure of a graphite nodule.[43]

Fig. 11—Graphite nodules in an austempered iron indicating nucle-ation on a small central inclusion.[44]

Fig. 9—(a) The melt with bifilms and sundry contaminant or preconditioned nuclei; (b) the elimination of the silica-rich bifilms by addition ofMg; (c) the survival of existing nuclei and additional nuclei from inoculation; (d) the nucleation of graphite, wrapping completely around exist-ing nuclei (particularly if they happen to pass through constitutionally supercooled regions); (e) growth of spheroids.


respectively. The logic is as follows: If a perturbation tothe spherical graphite shape were to occur, having anecessarily smaller radius r, the pressure to extendplastically the growth at this location would (accordingto Eq. [1]) be increased (Figure 12(b)). Thus, growth ofthe extension of smaller radius would be discouragedbecause additional pressure would be required tostabilize the perturbation. The easier spherical growthmode, simply expanding the uniform radius a, wouldtherefore be encouraged. It seems, therefore. that thereis some qualitative justification for believing thatmechanical forces stabilize the spherical growth modeof the nodule.

However, it is useful to ascertain whether there isquantitative justification for this mechanism. If we takeY to be approximately 6 MPa[51] for austenite at themelting point of iron, and a = 2 nm and b = 20 nm,then we find P = 30 MPa approximately. Even atvalues of a = 20 lm and b = 200 lm, P is of courseunchanged, which indicates that there is a substantialrestraining pressure on the growth of the nodule duringmost of its life.

With regard to the possible asymmetric effect of aperturbation of radius r, taking r = a/2 to a/10 locallyincreases P to approximately 35 to 60 MPa, respectively.Thus, a rounding effect caused by mechanical smoothing

of the forces to expand the austenite shell seems to beimportant. Although a creep model rather than theabove plastic model might give a somewhat moreaccurate result, the above result can be relied on to giveus an order-of-magnitude estimate of the effect. Clearly,more work is required to confirm this preliminaryindication.Johnson and Smart[35] use the sophisticated and

respected perturbation analysis by Mullins and Sekerkato suggest that interfacial energies are of importance inspherodizing graphite nodules up to a diameter ofperhaps 50 nm, after which the spherical form can nolonger be stabilized. Thus, much speculation by earlierauthors that interfacial energies may be important indefining the shape of spheroidal graphite seems irrele-vant. This leaves mechanical stabilization as a potentialcritical mechanism. This conclusion is reinforced by therecent evidence[16,17] that the early phases of growth ofthe graphite around the nucleus are anything butregular. Clearly, a strong spherodizing effect is neededthat neither the nucleation mechanism nor surfaceenergy can provide.In general agreement with this conclusion, Jiyang

et al.[52] used color etching to reveal the austenite shellsaround graphite. They found that if the shell formedquickly and completely, then the nodule developed as asphere, whereas slow-developing or nonenvelopingshells led to misshapen nodules.In passing, it may be significant that on the addition

of Mg causing dissolution of the silica-rich oxide bifilms,any residual air trapped between the films is expected tobe released. The air trapped in a bifilm will be expectedto lose its oxygen by continued oxidation of the matrix,followed by nitrogen that would react to form nitride,[53]

leaving mainly argon. In this way, it seems possible thatclouds of fine argon bubbles will be released into themelt. It seems likely that some Mg vapor will also diffuseinto the bubbles. It is not easy to define the sizes of suchbubbles with any accuracy. For instance, a bifilm of100 lm square and an average gas gap of 1 lm wouldyield a pore of approximately 20 lm diameter. A similararea bifilm of an average 10-nm gas gap would form apore approximately 5 lm in diameter. A small bifilm of

Fig. 12—The distortion of dendrites as a result of the internal growth of nodules (after Hillert[47]).

Fig. 13—A thick shell expanding plastically because of internal pres-sure. A perturbation radius r is not favored because a higher localpressure is required. Thus, sphericity is encouraged.


10 lm square and 10-nm average spacing would create a1-lm-diameter pore. Thus, it seems that a fog of bubblesin the range of approximately 1 to 20 lm is to beexpected.

It is intriguing, but not perhaps relevant, that a theoryproposes that nodules nucleate from Mg bubbles insuspension in the melt.[54,55] However, attractive thishypothesis might be to explain graphite coatings insidepores in solidified castings, as a result of the reduction instrain energy involved, any strain energy relief in theliquid state is zero, and the reduction of surface energy toencourage such precipitation in the liquid state seemsnegligible. Furthermore, the theory requires the incor-poration of solid particulate oxide nuclei into the bubble,but a successful penetration of a particle into the bubblewould depend critically on a reduction in interfaceenergies. This is unlikely for particles formed by precip-itation in situ in the liquid, which will be in perfect atomiccontact with the melt (i.e., will be well ‘‘wetted,’’ which isa necessary condition for the particle to be a nucleatingagent). Such particles will be energetically rejected bybubbles. Thus, although a mechanism for the presence ofextremely fine bubbles may be provided by the currentanalysis, it seems to be irrelevant to the formation ofgraphite nodules. It does not offer support to the gasbubble nucleation hypothesis.

Another interesting aside can be noted. The transfor-mation of the planar cracks sandwiched inside thebifilms into clouds of fine bubbles that may float andescape from the alloy is the essence of the process bywhich apparently brittle gray iron becomes ductile.(Ductile iron only becomes embrittled once again, asnoted below, if oxide bifilms are reintroduced byturbulence by handling of the melt or poor filling systemdesign of the casting.)

Finally, curious observations, such as that reportedby Yamamoto et al.,[56] in which flake is converted intonodular iron by simply purging the melt with finebubbles of nitrogen, argon or carbon dioxide becomeexplicable. From experience in the light metals indus-tries, it is known that purging with gases can eliminate

bifilms from melts. Thus spherodization seems to beachievable via a purely mechanical route, replicatingthe condition achieved chemically by the addition ofMg.

VII. COMPACTED GRAPHITE IRON (CGI)

If the addition of magnesium is more carefullycontrolled to some level intermediate between spheroi-dal and flake iron, compacted (‘‘vermicular’’ based onthe Italian for ‘‘worm-like’’) graphite is the result(Figure 14).[57]

In our bifilm model, it is clear that most of the oxidebifilms will be dissolved by the addition of Mg.However, small patches may remain if the Mg additionis not too high; the tiny patches on which the originalnuclei sat will be resistant to dissolution because theywill be stabilized by their attachment to the nuclei.(Naturally, it will have been energetically favorable forthe nuclei to attach to the bifilm, so that the combina-tion of nucleus and film will enjoy a reduction in overallenergy, stabilizing the combination.) Only half of thebifilm will be retained in this way, its distant ‘‘twin’’ halfnot enjoying the protective influence of the nucleus willdissolve and disappear. Only the small part of 1/2 of thebifilm together with its unbonded interface, the rem-nants of the layer of air, will remain (Figure 15(b)).The subsequent nucleation of graphite on the nucleus

will result in rapid spreading of growth around thenucleus. On arrival at the nonwetted interface of theresidual patch of bifilm, this spreading will be arrested(Figure 15(c)). The subsequent growth of graphite isforced to occur not radially but in general unidirection-ally away from the bifilm residue (Figure 15(d)).Clearly, the growth cannot now be a flake because no

bifilm is present. Originally, it was thought that it couldnot be spherical because it could not encapsulate thecomplete nucleating particle, but the observations byRiposan[16,17] and others indicates that the nucleatingparticle is often not entirely surrounded in the case of

Fig. 14—CGI viewed by (a) SEM deep etching and (b) optical metallography.[57]


flake iron, and this may also be true of nodular iron.Naturally, the subsequent growth form is not possible topredict and has to depend on completely differentfactors. With regard to the microstructure of growth,Cole[58] had observed a fine, unidirectional spiral struc-ture similar to the worm-like growth mode clearlyobserved in Figure 14. Liu[59] finds that the growthdirection is along the C axis (0001 direction perpendic-ular to the basal planes) and seems to develop by a spiraldislocation mechanism as witnessed by the coarse andirregular spirals that they observe.

The great sensitivity of the compacted graphitemorphology to magnesium concentration is corrobo-rated by the proposed bifilm mechanism. If the Mg levelis too low, then residual bifilms will encourage flakegraphite, whereas if the Mg level is too high, then thelimited stability enjoyed by the residual bifilm patcheswill be overcome, and the last remaining patches ofbifilm will be dissolved, encouraging the growth oftotally spherical grains.

Clearly, the final form of CGI is not stronglyinfluenced by the inoculation and nucleation events. Itdevelops its continuous worm-like growth because of acontinuous growth process. As an important feature ofthe continued growth of CGI, several workers find thatthe graphite seems to stay in contact with the liquid (forinstance, References 60 and 61), thus transferring theexpansion of the graphite to the liquid, reducing feedingrequirements,[62] in contrast to ductile iron in which thegraphite transfers its expansion to its surroundingsolidified shell, expanding the casting and increasingfeeding requirements.

VIII. CORAL GRAPHITE

Coral graphite is, as its name suggests, a fibrousmorphology of graphite, formed in rather pure, nonin-oculated irons. Because of the purity of the alloys, and thefact that it can be grown in alloys containing Ce, no silica-rich oxide bifilms can be present. The purity and absenceof inoculation also means that no oxysulfide nuclei arepresent. Thus, neither flakes nor spheroids can form.

Thus, coral graphite occurs at high undercoolings(clearly at least one nucleation event is required, but the

phenomenon is dominated by growth). After initiationin undercooled regions, the growth of the eutectic islikely to be so fast that it will cover large undercooledregions at the mold wall, thus advancing subsequentlyon a substantially planar growth front away from thewall. This seems to be a coupled growth mode.For interested readers, even though nowhere in their

report do Nakae and Shin[5] mention coral graphite,they present micrographs that clearly illustrate coralgrowth when the iron is sufficiently pure, or Ce is added,and when the growth rate is sufficiently high. They showthe close similarity between the coral eutectic structuresof Fe-C and Al-Si alloys.CGI and coral forms of graphite (Figure 16) are both

fine, fibrous filamentary morphologies. It is tempting toconsider that the resemblance reflects an underlyingsimilarity.

IX. MISSHAPEN SPHEROIDS

The presence of ill-formed spheroids, particularly ifpresent in large numbers, is widely known to beassociated with the reduction in mechanical properties

Fig. 15—The formation of CGI by addition of just enough Mg to (a) melts with existing bifilms with attached nuclei, to (b) eliminate most ofthe silica-rich bifilms, but not the remnant of film attached to the nuclei. (c) Inoculation promotes graphite growth on the nuclei; (d) growth con-tinues unidirectionally.

Fig. 16—An SEM image of deeply etched CGI closely resemblingcoupled eutectic coral structure, illustrating the overall similarity ofthese structures (Courtesy of Thomas Prucha, AFS, 2009).


of nodular irons. Hughes[44] describes how a goodductile iron can achieve at least 90 pct nodularity butless good irons can fall to as low as 50 pct or less andsuffer reduced properties. The 50 pct or so componentof flakes or other non-nodular shapes does not partic-ularly affect properties such as proof strength butgreatly reduces those properties sensitive to failure, suchas tensile strength and ductility. Hughes,[44] in commonwith widely held belief, attributes this loss of fractureresistance to the sharp notches at the root of flakes.However, this belief presupposes that the flakes act ascracks. This would not be true if graphite had hightensile strength perpendicular to its basal plane but isprobably only true if the flakes are formed on bifilms;the bifilm provides the unbonded interface, as a crack. Itis the presence of the crack provided by the bifilm thathas to be viewed as the principal cause of failure.

In terms of the bifilm hypothesis, graphite would beexpected to grow on oxide bifilms. During the Mgtreatment, the silica-rich bifilm content of the meltwould be effectively eliminated. Thus, in a liquid nowcleaned from transient silica-rich bifilms, spheroidswould be created in suspension. However, many Mgaddition techniques are extremely turbulent, so thatlarge quantities of Mg oxide and/or Mg silicates areexpected to be created by the turbulent jumping andsplashing of the liquid surface, thereby folding in theoxide surface of the liquid to create entrained bifilms.These new bifilms will be permanent defects formedfrom highly stable magnesia or magnesium silicate.Prior to pouring, some of these will fortunately floatout, adding to the Mg-rich slag. Thus, given a reason-able time between Mg treatment and the pouring of thecastings for separation of Mg-rich bifilms (an interestingand clearly important process variable that seems notwell researched), not all of these defects will find theirway into the castings to impair the structure andproperties.

Unfortunately, on pouring into the mold, additionallarge quantities of Mg-rich oxide bifilms are likely to bereintroduced, particularly if the mold filling system is arather poor design.

On contact with an oxide bifilm, the subsequentgrowth of the existing suspended spheroids will beredirected along the plane of the film. The symmetricalspherical constraint previously provided by the sur-rounding austenite is also destroyed, aiding the non-spherical development. Thus, the spheroid will grow tobecome significantly misshapen.

This effect can be observed in Figure 2. In this figure,several oxide bifilms have been straightened by thegrowth of dendrites so as to lie along 100 planes. Theseplanes, now containing a planar bifilm crack, lead todisastrous brittleness known as plate fracture (Figure 1)described elsewhere by the author.[3] Meanwhile, thenodules attached to the bifilms are clearly poorlyshaped. Additional misshapen nodules are evidentelsewhere in the structure. These are expected to belying on random areas of bifilm not straightened bydendrite growth.

The appearance of poorly shaped spheroids is thereforepredicted to be associated with the growth of nodules on

bifilms as a result of either (1) insufficient dwell time forthe damage introduced during Mg addition to float out,or (2) poor casting practice, in which an otherwise nicelyinoculated and spherodised melt is recontaminated withoxides. It is interesting to predict that perhaps more timeafter the spherodizing treatment to allow themelt to clear,together with a properly designed filling system, orcounter-gravity filling system, should completely elimi-nate poor nodularity.A step in the right direction is presented in the work of

Takita et al.,[63] who observe that nodules are convertedfrom misshapes to spherical by the use of a filter to takeout the ‘‘inclusions produced by inoculation.’’ Thispositive step contrasts with that taken by Liu et al.[64]

who, after the Mg addition, added ‘‘post inoculants.’’These were highly successful to increase the nodule countbut led to a disastrous fall in nodularity. This was almostcertainly a result of adding the inoculants through themelt surface, which would have also entrained thesurface oxide to create Mg-rich oxide bifilms that, oncontact, would generate nonspheroids.Furthermore, of course, the significant reduction of

properties associate with malformed spheroids cannotbe the direct result of the shape of the spheroids, becausethey occupy such a small volume fraction of the alloy.The loss of properties is predicted to be the result of thepresence of the bifilms in the melt, occupying a vastlygreater cross sectional area than the spheroids. Theseextensive bifilms simply act as cracks in the casting,significantly reducing properties.

X. EXPLODED NODULAR GRAPHITE

‘‘Exploded’’ spheroids (Figure 17)[65] are commonlyobserved in irons subject to graphite flotation andespecially if the composition of the iron is sufficientlyhypereutectic.[66,67]

This undesirable morphology is not easily explainedat this stage as a result of relatively little experimentalwork to clarify the problem. Cole[58] suggested they hadsuffered remelting as a result of being carried byconvection in and out of hot zones of the liquid. Thisseems most unlikely, however, because if the nodule hadgrown uniformly in a compact morphology, then theuniform graphite would be expected to have a substan-tially uniform rate of dissolution; solidification andremelting would be expected to be reversible.Because exploded nodules appear exclusively in the

flotation region of hypereutectic irons two far morelikely factors are as follows:

(a) Nodules growing in a sufficiently hypereutecticmelt[66] will experience an enhanced driving force forgrowth because of the carbon supersaturation thatdevelops as the melt cools. This will encouragegrowth instabilities leading to ‘‘dendritic’’ ratherthan ‘‘planar’’ growth, leading to exploded ratherthan smooth spheroid surfaces.

(b) Austenite will be less likely to form in hypereutecticirons, especially in conditions of carbon enhance-ment by segregation during cooling in the liquid


phase. The nodules may have nucleated early in theliquid phase and grown without the benefit of themechanical constraint of the austenite. When notpressurized to remain spherical, the nodule willbe free to grow more like a dendrite, developinginstabilities that grow into projections to its growthfront, finally developing the characteristic explodedforms. Evidence for mechanical restraint as apowerful effect is presented in the section on nodulargraphite above.

XI. CHUNKY GRAPHITE

Chunky graphite is often observed concentrated in thecenters of heavy sections of nodular iron castings.‘‘Chunky’’ is not a particularly helpful descriptive adjec-tive for this variety of graphite. Its ‘‘chunkiness’’ is onlyapparent under the microscope at high magnification;otherwise, it simply appears to be fine, irregular,branched, and interconnected fragments (Figures 17(c)and 18).[68]

Once again, the properties of nodular iron arereduced. However, it seems the loss of properties ispossibly more associated with the short diffusion dis-tances between branches of the graphite filaments,promoting the development of ferrite.[65]

Liu et al.[59,65] find evidence that chunky graphitegrows along the C-axis direction, as does nodulargraphite. Furthermore, they report observations onspheroids that exhibit gradual degeneration, whichslowly take on the growth forms of chunky graphite.

Thus, they conclude that chunky graphite is a degen-erate form of spheroidal graphite, and their workimplies that chunky graphite grows out from spheroids.Itofugi and Uchikawa[69] confirm the identical growthorientations of spheroidal, compacted, and chunkygraphites.All these workers observe the characteristic form of

chunky graphite, as an apparently ‘‘stop/start’’ growthin the C-direction consisting of nearly separate pyrami-dal ‘‘chunks’’ linked by a narrow neck, like beads on abranching string. The individual chunk sections arecomposed of layers parallel to the basal plane, but only

Fig. 17—(a) Spheroid and (b) malformed spheroid; (c) chunky graphite (after Liu et al.[65]); (d) SEM iron image of an exploded spheroid;(e) electron image.[58]

Fig. 18—Graphite nodules and areas of fine, chunky graphite in thethermal center of a 200-mm cube.[68]


nanometers thick. This characteristically lumpy growthmay be the result of a pulsating or irregular advance ofthe growth front, with the austenite advancing to grownearly over the top of the graphite, forming the nearlypinched-off neck of the graphite, only to be overtakenagain because the carbon in solution will now build upin the liquid ahead of the front, accelerating the nextphase of growth of the graphite until the local carbonconcentration is depleted once again, and so on.

Observations by Kallbom et al.[68] seem consistentwith an origin associated with bifilms. They observe thechunky graphite to be concentrated in the center ofheavy sections, which is explained by the growth of thefreezing front pushing bifilms ahead by their observa-tions of ‘‘stringers’’ of graphite nodules. These featuresare almost certainly sheets of oxides decorated withgraphite nodules that have been nucleated on the oxide(analogously to those observed in Figure 2). Theyconclude that a collaboration among Mg, S, and O isimportant for chunky graphite formation.

Although the above authors provide evidence thatsuggests the presence of bifilms, it seems possible that atthe same time there is likely to be an absence of nuclei.This is because so much time is available for particles insuspension to float out from the centers of heavysections where chunky graphite is commonly found. Inthe absence of nuclei, spheroids cannot form. Thus, if agraphite nodule can be initiated somewhere in or nearthis region, perhaps happening to float into this region,it will continue to grow. However, it will not enjoy thebenefit of the presence of an austenite shell in a region ofhigher temperature and enhanced segregation, so that itsgrowth mechanism seems likely to be asymmetrical.Because the growth will constitute an extension of thespheroid, the spheroidal mechanism of growth along theC-axis is likely to be continued.[59] In the absence ofsignificant numbers of nuclei, the whole region will fillwith a variety of continuous, branching growth, extend-ing along the C-axis direction. The extended size ofchunky graphite regions, which are much larger thancells of other types of graphite,[69] corroborate theabsence of nuclei in these regions.

If the above mechanism is correct, then explodedspheroids can be viewed to be a similar growth mode tochunky graphite, both originating from spheroids.However, why continued growth should not occur inall directions as in the exploded variety, but only in afew select directions (possibly one direction) in chunkygraphite, is not clear at this time.

It is hoped that in the near future, the correctexplanation for the origin of chunky graphite might beelucidated by subsequent careful experiments. The keyword here is ‘‘careful.’’ For instance, the experiment byAsenjo et al.[70] that involves the placement of inocu-lants in different branches of a runner system todifferent sized cavities in a mold to compare the effectsof mold inoculation in different heavy castings was aclever concept but regrettably flawed in execution. Thiswas because, in common with most iron casting, therunners were not designed to be pressurized and fill ona single pass. Thus, a reverse flow is likely to havecontaminated the mold cavities, and all the cast

material would have suffered from turbulence and airentrainment, therefore containing unknown quantitiesof oxide bifilms. Clearly, in the future, much greatersophistication of melting and casting will be requiredfor experiments designed to clarify the solidificationmechanisms for cast irons.

XII. IRON CARBIDE (CEMENTITE)

Work by Rashid and Campbell[71] has demonstratedthe nucleation and growth of carbides on oxide bifilmsin vacuum-cast Ni-base super-alloys. It would beexpected, therefore, that an analogous reaction wouldoccur in Fe-C alloys, because the austenite formingduring solidification also possesses a closely similar face-centered-cubic structure, and other conditions, such asthe temperature and solidification rates, are all similar.Carbides in irons seem to form preferentially at grain

boundaries and often seem to be associated both withresidual graphite (sometimes as nodules, malformednodules, or flakes aligned with the boundary) and poresall forming on the same boundaries. This is a clue totheir bifilm associations.Faubert et al.[72] have studied carbides in heavy-

section austenitic ductile iron (ADI). Toward the top oftheir castings, they find degradation of properties moreserious than they would have expected from the carbidesthemselves. They suspected that the real impairment wascaused by the presence of films that had floated into thisregion. Almost certainly, the ‘‘films’’ would have been‘‘bifilms’’ (it seems impossible to devise a mechanism bywhich a single thickness of film can be introduced into amatrix). Bifilms would segregate to grain boundariesand possibly actually constitute the boundary. Thepresence of the bifilm is not only inferred from (1) thecracked carbides but also (2) from the linear rows ofnodules viewed in micrographs from this work, (3) fromthe pores as the residues of air bubbles trapped betweenthe films, and (4) from the graphite flakes sitting in theboundary (called by the authors, unflatteringly, ‘‘degen-erate’’ graphite). Both graphite and carbides areexpected to form on the wetted, outer surfaces of thebifilms. The presence of the central unbonded region(including the pores) between the films, constituting thecrack through the interiors of the carbides, explains theapparent brittleness of the carbides. These intermetalliccompounds would otherwise be expected to be strongand resistant to failure by cracking at the modeststresses that can be induced by solidification andcooling.Stefanescu[43] quotes the work of Hillert and

Steinhauser,[73] in which the growth of iron carbideeutectic (ledeburite) occurs by the spreading of carbide(cementite) across a plane, followed by the developmentof a rod type of eutectic at right angles. It is tempting toconsider that the original planar expansion would havebeen facilitated by growth across the surface of a bifilm.The bifilm would originally have been randomly crum-pled but would have been straightened by the progressof the carbide across its face, thus creating an essentiallyplanar crack that would constitute a serious defect in the


carbide. Associated branching cracks would have arisenfrom irregular folds in the bifilm.

XIII. GENERAL

Overall, both nucleation and growth mechanismsinfluence graphite morphology, although these mecha-nisms dominate to different extents in different circum-stances. For instance, nucleation dominates theformation of spheroids and deformed spheroids, whereasfor flake graphite, both nucleation and growth areinfluential. For compacted, undercooled coral andchunky graphites, nucleation occurs once to initiate eachcell, after which continuous growth leads to developingcontinuous branching morphologies. Indeed, the finer,fairly continuous forms of coral, chunky, and compactedgraphites are so similar that they often seem to beconfused in the casting literature. They do seem similar inthe sense that they all appear to be more-or-less coupledgrowth forms, advancing together with the austenite.

If, as proposed here, the various forms of graphite aresignificantly influenced by the presence or absence ofbifilms, then it would explain the historical resistance ofthe phenomenon to explanation so far. This seemstypical of bifilm phenomena.

Furthermore, it seems likely that the principle causeof reduced mechanical properties in all cases of non-spheroidal forms of cast irons is the presence of variouskinds of oxide bifilms that act as cracks. Ductile iron isductile because of the absence of oxide bifilms, notbecause of its spherical graphite morphology; but whenoxide bifilms (mainly magnesia-rich) are entrained bypoor casting technology, ductile iron can be seriouslyreduced in ductility, with castings at times even failingdisastrously by the brittle ‘‘plate fracture’’ mechanism.[3]

Naturally, this short account presents only an outlineof a new approach to the structures of cast irons. Muchmore research is needed to prove, enhance, or disprovethese proposals. It is hoped that improved control andimproved castings will result from these efforts.

XIV. CONCLUSIONS

A hypothesis is proposed as follows:

1. Cast iron melts normally contain double films(bifilms) in suspension.

2. Inoculation produces oxysulfide particles thatnucleate on silica-rich oxide bifilms.

3. Graphite nucleates on the oxysulfide particles andgrows, spreading over the bifilms, straightening thebifilms, and forming flakes of crystallographicallynear-perfect graphite. The presence of the bifilms,as cracks, trapped inside or alongside graphiteflakes accounts for the poor tensile properties offlake irons.

4. Compacted graphite forms on oxysulfide nuclei thatoccupy bifilm residues.

5. Spheroids form in the absence of silica-rich bifilms,nucleating on oxysulfide particles. The spherical

growth morphology may be natural but is addi-tionally encouraged by the mechanical constraint ofthe austenite matrix; the absence of bifilms explainsthe high mechanical properties.

6. Coral morphology nucleates on unknown nuclei atlow temperatures, expanding to form cells of cou-pled growth with austenite, consisting of highlyfaulted continuous branching filaments of graphitein the austenite matrix; bifilms play no part in thisgrowth mode.

7. Misshaped spheroids seem to be spheroids that haveencountered a Mg-rich bifilm, subsequently grow-ing along the bifilm and losing sphericity (Figure 2).In addition, the presence of the bifilm destroys thesymmetrical mechanical constraint of the austenitethat favors sphericity.

8. Chunky graphite occurs in heavy-section ductileiron regions. The formation mechanism is notclear at this time. The central regions compriseeither (1) regions containing Mg-rich bifilms as aresult of poor casting techniques or (2) regions fromwhich nuclei have floated out, leaving regions nearlydevoid of nuclei and creating graphite akin to coralmorphology. As such, it may be a coupled eutecticform, and its ‘‘beads on a branching string’’ mor-phology may result from an unstably advancinggrowth front.

9. Exploded spheroids may be the result of growth inthe liquid, without the benefit of the mechanicalconstraint of an austenite shell.

10. Carbides form at low temperatures on oxide bifilms.The presence of bifilms in the carbides explains thebrittle behavior of these strong intermetallics andtheir common association with both pores andresidual graphite fragments.

ACKNOWLEDGMENTS

The author is grateful to those who have assistedwith the micrographs; to Carl Loper and Riposan andcolleagues for inspiring research; and, last but notleast, to the painstaking, assiduous, and doubtfulreviewer of this article from whom I learned much.

REFERENCES1. C.R. Loper: AFS Trans., 1999, vol. 107, pp. 523–28.2. J. Campbell: Mater. Sci. Technol., 2009, vol. 25, pp. 125–26.3. J. Campbell: Castings, Elsevier, Oxford, UK, 2003, pp. 178–81;

161–62; 158–60.4. J. Campbell: Mater. Sci. Technol., 2006, vol. 22 (2), pp. 127–45;

(8), 999–1008.5. H. Nakae and H. Shin: Int. J. Cast Met. Res., 1999, vol. 11 (5),

pp. 345–49.6. A. de Sy: AFS Trans., 1967, vol. 75, pp. 161–72.7. D. Hartmann and W. Stets: AFS Trans., 2006, vol. 114, pp. 1055–

58.8. E. Hoffman and G. Wolf: Giessereiforschung, 2001, vol. 53 (4),

pp. 131–51.9. F. Mampaey and K. Beghyn: AFS Trans., 2006, vol. 114, pp. 637–

56.10. R. Barton: Foundry Trade J., 1985, pp. 117–26.


11. T. Mizoguchi, J.H. Perepezko, and C.R. Loper: AFS Trans., 1997,vol. 105, pp. 89–94.

12. F. Mampaey and Z.A. Xu: AFS Trans., 1997, vol. 105, pp. 95–103.13. R.J. Warrick: AFS Trans., 1966, vol. 74, pp. 722–33.14. M.H. Jacobs, T.J. Law, D.A. Melford, and M.J. Stowell: Met.

Technol., 1974, vol. 1 (11), pp. 490–500.15. T. Skaland: AFS Trans., 2001, vol. 105, pp. 77–88.16. I. Riposan, M. Chisamera, S. Stan, P. Toboc, E. Ecob, and

D. White: Mater. Sci. Technol., 2008, vol. 24 (5), pp. 579–84.17. I. Riposan, M. Chisamera, S. Stan, C. Hartung, and D. White:

Mater. Sci. Technol., in press.18. M. Chisamera, I. Riposan, and M. Barstow: AFS Trans., 1996,

vol. 104, pp. 581–88.19. C. Hartung, C. Ecob, and D. Wilkinson: Casting Plant Technol.,

2008, vol. 2, pp. 18–21.20. F. Hurum: AFS Trans., 1952, vol. 60, pp. 834–48.21. F. Hurum: AFS Trans., 1965, vol. 73, pp. 53–64.22. H. Fredriksson: Mater. Sci. Eng., 1984, vol. 65, pp. 137–44.23. R.A. Harding, J. Campbell, and N.J. Saunders: Solidification

Processing 97, Sheffield, UK, 1997, pp. 489–93.24. C.R. Loper and R.W. Heine: AFS Trans., 1961, vol. 69, pp. 583–

600.25. F. Mampaey: AFS Trans., 1999, vol. 107, pp. 425–32.26. E.A. Feest, G. McHugh, D.OMorton, L.S. Welch, and I.A. Cook:

Proc. Warwick University Conf., 1980, pp. 232–39.27. Y.X. Li, B.C. Liu, and C.R. Loper: AFS Trans., 1990, vol. 98,

pp. 483–88.28. L. De Rong and Y.J. Xiang: AFS Trans., 1991, vol. 99,

pp. 707–12.29. M. Hillert and S. Rao: Proc. Brighton Conf., 1967, pp. 204–12.30. N. Enright, S.Z. Lu, A. Hellawell, and J. Pilling: AFS Trans., 2000,

vol. 108, pp. 157–62.31. G.M. Goodrich: Metals Handbook, vol. 15, Oxford, OH, 2008,

pp. 785–811.32. C.R. Loper and K. Fang: AFS Trans., 2008, paper 08-065(05),

p. 8.33. M.F. Basdogan, G.H.J. Bennett, and V. Kondic: Proc. Warwick

University Conf., 1980, pp. 240–47.34. S. Liu and C.R. Loper: AFS Trans., 1990, vol. 98, pp. 385–94.35. W.C. Johnson and H.B. Smart: Proc. Sheffield Conf., 1977,

pp. 125–30.36. J. Campbell: Int. J. Metalcast., 2007, vol. 1 (1), pp. 7–20.37. T. Carlberg and H. Fredriksson: Proc. Solidification and Casting of

Metals Conf., 1977, pp. 115–24.38. S.I. Karsay: Ductile Iron II, Engineering Design Properties

Applications, Quebec Iron & Titanium Corporation, Canada,1971.

39. R.C. Voigt and S.D. Holmgren: AFS Trans., 1990, vol. 98,pp. 213–25.

40. E. Fras, M. Gorny, and H.F. Lopez: AFS Trans., 2007, vol. 115,pp. 1–17.

41. E. Fras, M. Gorny, and H.F. Lopez: Metall. Mater. Trans. A,2007, vol. 38A, pp. 385–95.

42. P. Zhu, R. Sha, and Y. Li: Proc. Mater. Res. Soc., 1985, vol. 34,p. 3.

43. D.M. Stefanescu: Metals Handbook, ASM, Oxford, OH, 1988,vol. 15, pp. 168–81.

44. I.C.H. Hughes: Metals Handbook, ASM, Oxford, OH, 1988,vol. 15, pp. 647–66.

45. D.M. Stefanescu: Metall. Mater. Trans. A, 2007, vol. 38A (7),pp. 1433–47.

46. A.N. Roviglione and J.D. Hermida:Metall. Mater. Trans. B, 2004,vol. 35B, pp. 313–30.

47. M. Hillert: in Recent Research on Cast Iron, H.D. Merchant, ed.,Gordon and Breach, New York, NY, 1968, pp. 101–27.

48. R. Ruxanda, L. Beltran-Sanchez, J. Massone, and D.M.Stefanescu: AFS Trans., 2001, vol. 109, pp. 37–48.

49. A. Buhrig-Polackzek and A. de Santos: Metals Handbook, ASM,Oxford, OH, 2008, vol. 15, pp. 317–29.

50. P. Chadwick: Int. J. Mech. Sci., 1963, vol. 5, pp. 165–82.51. J. Campbell: Trans. TMS-AIME, 1967, vol. 239, pp. 138–42.52. Z. Jiyang, W. Schmidt, and S. Engler: AFS Trans., 1990, vol. 98,

pp. 783–86.53. R. Raiszadeh and W.D. Griffiths: Metall. Mater. Trans. B, 2008,

vol. 39B, pp. 298–303.54. A.A. Gorshov: Russ. Cast. Prod., 1964, pp. 338–40.55. H. Itofugi: AFS Trans., 1996, vol. 104, pp. 79–87.56. S. Yamamoto, Y. Kawano, Y. Murakami, B. Chang, and

R. Ozaki: AFS Trans., 1975, vol. 83, pp. 217–26.57. D.M. Stefanescu, R. Hummer, and E. Nechtelberger: Metals

Handbook, ASM, Oxford, OH, 1988, vol. 15, pp. 667–77.58. G.S. Cole: AFS Trans., 1972, vol. 80, pp. 335–48.59. P.C. Liu, C.R. Loper, T. Kimura, and H.K. Park: AFS Trans.,

1980, vol. 88, pp. 97–118.60. J.Y. Su, C.T. Chow, and J.F. Wallace: AFS Trans., 1982, vol. 90,

pp. 565–74.61. F. Mampaey: AFS Trans., 2000, vol. 108, pp. 11–17.62. J.D. Altstetter and R.M. Nowicki: AFS Trans., 1982, vol. 90,

pp. 959–70.63. M. Takita, Y. Yang, and H. Nomura: Int. J. Cast Met. Res., 1999,

vol. 11 (5), pp. 357–62.64. S.L. Liu, C.R. Loper, and T.H. Witter: AFS Trans., 1992, vol. 100,

pp. 899–906.65. P.C. Liu, C.L. Li, D.H. Wu, and C.R. Loper: AFS Trans., 1983,

vol. 91, pp. 119–26.66. G.X. Sun and C.R. Loper: AFS Trans., 1983, vol. 91, pp. 841–54.67. A.P. Druschitz and W.W. Chaput: AFS Trans., 1993, vol. 101,

pp. 447–58.68. R. Kallbom, K. Hamberg, and L.-E. Bjorkegren: Proc. World

Foundry Congr., Harrogate, UK, 2006, paper 184/1-10.69. H. Itofugi andH. Uchikawa:AFSTrans., 1990, vol. 98, pp. 429–48.70. I. Asenjo, P. Larranaga, J. Sertucha, R. Suarez, J.-M. Gomez,

I. Ferrer, and J. Lacaze: Int. J. Cast Met. Res., 2007, vol. 20 (6),pp. 319–24.

71. A.K.M.B. Rashid and J. Campbell:Metall. Mater. Trans. A, 2004,vol. 35A (7), pp. 2063–71.

72. G.P. Faubert, D.J. Moore, and K.B. Rundman: AFS Trans., 1990,vol. 98, pp. 831–45.

73. M. Hillert and H. Steinhauser: Jernkontorets Ann., 1960, vol. 144,pp. 520–22.


a hypothesis for cast iron micro structures

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