CommunicationsExperimental Observations of CrystalGrowth in Alloys Rapidly Quenchedfrom the Liquid State

JOHN WOOD AND IAN SARE

A recent paper by Shingu and Ozaki l compared solid-ification rates in splat cooling calculated from a heatflow balance with crystal growth rates obtained fromclassical molecular kinetic equations. Under certainwell defined conditions they were able to estimate thekinetic undercooling in the liquid at the solid-liquidinterface, and they subsequently derived a limitingcondition where the interface will advance at a greatervelocity than that required for crystal growth, thus en-abling the solid to remain in an `amorphous' state. Itis not our purpose to contend the basic assumptions ofthe model described in the above paper concerning thenature of the moving solid-liquid interface, though it isunsettling not to have a rigorous definition for such aconcept. It is our aim to present some metallographicobservations of splat-cooled alloys which show that theassumption of one-dimensional heat flow made byShingu and Ozaki, 1 and by Ruhl2 in an earlier theoret-ical analysis, is not often borne out in practice in re-gions that are subject to structural analysis by trans-mission electron microscopy. In both of these theoret-ical studies, growth of the solid was assumed to occurperpendicular to the foil-substrate interface, but wehave found crystal growth, and hence heat extraction,in most cases takes place parallel to this interface inunthinned electron-transparent areas of splats pro-duced by the 'gun' 3 technique.

The present work is derived from structural studiesof transformable 4 and non transformable 5 steels aftersplat quenching in a high-temperature, inert atmo-sphere `gun' device. Initial work on austenitic steels s

showed a large variation in structure within thin areasof a splat, and these were ascribed to the differentcooling conditions present.

In all iron base alloys featureless grains were com-monly observed in the thinnest regions of unthinnedfoils examined by transmission electron microscopy.The longitudinal grain boundaries generally lay par-allel to each other, and there was always a commonboundary in the center of such areas, where grainswhich had grown from each side have impinged (Fig. 1).Jones 7 has collated the various observations of suchelongated grains in other alloy systems and quotes asuggestion of Furrer that such grains grow in liquidspreading outwards from areas where solidificationis already in progress. This implies that the nucleat-ing point for growth of the solid must lie somewhere

JOHN WOOD is Foseco Research Fellow at the Department of Metal-lurgy and Materials Science, University of Cambridge, Pembroke Street,Cambridge, England. IAN SARE, formerly at the Department of Metal-lurgy and Materials Science, University of Cambridge, Cambridge,England, is now with the CSIRO Division of Tribophysics, Universityof Melbourne, Parkville, Victoria 3052, Australia.

Manuscript submitted May 20, 1975.

within the thin area, and that growth would occur radi-ally from it. The absence of any such growth centermeans that this suggestion is probably not correct, andindeed, it cannot account for the observed impingementboundary in the center of thin areas. We believe it ismore logical to attribute such observations to crystalgrowth from the edges of such a thin area to the center,implying that heat extraction occurs parallel to the foilplane, and not perpendicular to it, into the substrate.

The growth of grains perpendicular to the substratedoes occur, however, and is readily apparent fromscanning electron micrographs. Fig. 2, which illus-trates such a structure, is taken from the top side ofa pure nickel splat (i.e. the surface in contact with theatmosphere after solidification has ceased). The ap-pearance of a fine-grained structure with the principlegrowth direction perpendicular to the foil plane indi-cates that heat flow has occurred in accordance withthe two theoretical treatments. Unfortunately suchareas are relatively thick and are seldom observable

Fig. 1—Transmission electron micrograph from unthinnedarea of a Fe-4Mo-0.2C splat with elongated grains of S fer-rite sharing common boundary A-A.

Fig. 2—Scanning electron micrograph from top surface of anickel splat showing fine equiaxed grain structure.

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by transmission electron microscopy unless furtherthinning is carried out (Fig. 3). Since splats producedby the 'gun' technique are highly porous (Fig. 4) dueto their multidrop constitution, 'edge' effects cannotbe ignored and indeed they have provided the areasthat yielded anomalous results with respect to meta-stable and microcrystalline phases. 8

The other principle type of crystalline growth mor-phology observed in splat-cooled material is that of

branched dendritic structures (Fig. 5). In unthinnedfoils these are invariably found in larger electron-transparent areas than are the elongated grains ofFig. 1. The measurement of secondary arms in theseareas is a standard method of estimating cooling rates, 9

and the fact that secondary dendrite arms lie in the foil

plane means that the primary dendrites from which theyemanate must also lie in this plane. Since the growthdirection of primary dendrites is largely determinedby the prevailing direction of heat extraction, theirpresence in the plane of a splat foil means that theprinciple heat flow direction also lies in this plane.

The reason why the heat flow conditions in theseareas are contrary to those assumed in the theoreticalanalyses became clear from examination of the under-side of splats by scanning electron microscopy. 6 Be-tween 60 and 90 pct of this surface replicates theroughened copper substrate (which in this study wasroughened in two orthogonal directions with 400 gradeemery paper), but in the remaining areas the splatflakes show no visible signs of contact with the sub-strate (Fig. 6), and have been described previously 6

as `lift-off' areas. It is not possible unambiguously

Fig. 3—Equiaxed grains of austenite in a splat of AISI MI toolsteel after further ion beam thinning. (Transmission electronmicrograph).

Fig. 5—Transmission electron micrograph from large un-thinned area of a Fe-20 pct Cr-25 pct Ni austenitic steel withprimary and secondary dendrite arms in the foil plane.

Fig. 4—Low magnification scanning electron micrograph ofan Ag-Sn alloy showing porosity and multidrop nature of splat(top surface).

Fig. 6—Underside of splat showing `lift-off' areas. Dendriteshave grown in the plane of the foil in this Fe-20 pet Cr-25pct Ni splat. (Scanning electron micrograph).

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to determine how these are formed, but there are atleast three possible causes, viz, a) gas entrapment,b) thermal contraction during solidification, and c)shearing of droplets over substrate surface irregu-larities.

In these areas two extremes of heat flow can be con-sidered: i) heat removal into the gas from the top andbottom surfaces; ii) heat removal to the nearest con-tact point on the substrate. Denoting these as class i)and class ii) respectively, it can be seen that the for-mer gives relatively low values for h (the heat trans-fer coefficient) and d (the thickness across which heatflows), and that the latter leads to higher values for hand d. It is worth noting that Shingu and Ozaki' use alarge value for h and a small value for d when consid-ering the limiting condition for noncrystalline growth.The only areas that conform to this situation are thosevery small regions that protrude from the edge of foils.In general, areas such as Fig. 1 are observed, and ifd is <5 µm in these instances, the structure cannot beadequately analyzed, since it is an analogous situationto a deep well on the splat surface, and any criticaltilting will be sabotaged by the high walls on eitherside. In class ii) cooling, d is dependent upon the di-ameter of the `lift-off' regions, and hence the areatransparent to electrons will give an indication of themagnitude of d.

The observation by transmission electron micros-copy of elongated grains and dendritic structures inthe foil plane of unthinned splats and the observationof `lift-off' areas by scanning electron microscopy,indicate that class ii) is the predominant type of cool-ing in thin regions. Using the results of Ruhl z for ironquenched onto a copper substrate, it can be demon-strated that for a given dh between class i) and classii) there will be a critical diameter below which classii) will give a faster cooling rate than class i). Sinceit is postulated that in all cases cooling is Newtonian,it is necessary to extend Ruhl's calculations for thepresent case. In practice the breakdown to segregatedstructures in austenitic steels (Fig. 5) occurs at diam-eters greater than 20 µm, and from this it may be de-duced that Mh will be of the order of two to three ordersof magnitude. Since scanning electron micrographs ofthis austenitic steel indicate that the majority of thesplat consists of unbranched structures like those inFig. 2, the area of Fig. 5 must have cooled at a slowerrate than the bulk foil.

In conclusion, whilst the theoretical treatments ofone dimensional heat flow are useful, they often do notapply to very thin areas of foils because such regionsneither solidify nor cool in contact with the substrate.This is unfortunate as the majority of reported obser-vations are taken from such areas. Thus it is neces-sary to unambiguously determine heat flow conditionsfor a particular area to determine whether these theo-retical treatments are applicable.

We would like to acknowledge the support of theScience Research Council for an equipment grant andthe maintenance of one author (JW) and CSIRO formaintenance of the other (IS). Our gratitude is due toProfessor R. W. K. Honeycombe for encouraging super-vision and for providing laboratory facilities.

1. P. H. Shingu and R. Ozaki: Met. Trans. A, 1975, vol. 6A, pp. 33-37.2. R. C. Ruhl: Mater. Sci. Eng., 1967, vol. 1, pp. 313-20.

3. P. Duwez: Trans. ASM, 1967, vol. 60, pp. 607-33.4. I. R. Sare: Ph.D. Dissertation, 1975, University of Cambridge.5. J. V. Wood: Ph.D. Dissertation, 1974, University of Cambridge.6. J. V. Wood and R. W. K. Honeycombe: J. Mater. Sei., 1974, vol. 9, pp. 1183-88.

7. H. Jones: Rep. Progr. Phys., 1973, vol. 36, pp. 1425-97.8. H. A. Davies and J. B. Hull: J. Mater. Sci., 1974, vol. 19, pp. 707-17.9. H. Matyja, B. C. Giessen, and N. J. Grant: J. Inst. Metals, 1968, vol. 96, pp.

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Observation of Transient Stress Statesin Eutectic Composite Creep

G. GARMONG

In a previous paper' the author proposed a new cal-culational approach to the problem of transient stressstates produced during composite creep deformation,a possibility first recognized by deSilva. 2 While thereis no doubt from the mathematics that such transientsshould exist, their experimental verification is difficult.Since the in situ creep parameters of the two phasescannot be individually measured, it is generally notpossible to tell whether the externally measured com-posite strain reflects transient or nontransient behav-ior. In specific circumstances, however, transientsdue directly to the mechanical interaction of the phasesmay be observed and related to the analytical approach.It is the purpose of this note to describe the results ofsome experiments in which the strain-time creep rec-ord at constant load provides evidence for the stresstransients present in the individual phases of a eutec-tic composite.

If a composite creeps with the matrix plastic and thereinforcement elastic, the reinforcement assumes anincreasing share and the matrix a decreasing share ofthe load as time and strain progress, due to the iso-strain nature of the deformation. Eventually the matrixstress decreases to zero or to a stress at which no fur -ther plastic straining can occur. A mathematical illus-tration of the nature of this stress relaxation of the ma-trix and buildup of reinforcement stress may be foundin the development leading to Eq. [8] of Ref. 1. In thatexample the matrix stress approaches zero as strainincreases because matrix strain rate depends only onstress. In other cases the matrix stress need not fallto zero, but only to a point below which no furthercreep can occur (the creep yield stress). In the caseof matrix stress falling to zero, the elastic reinforce-ment eventually carries all of the load imposed on thecomposite, while in the other case the two phases bearproportions of the load that cannot be determined fromsimple elasticity considerations only. In either eventthe composite strain rate eventually becomes zero.

Specimens of the Ni -45.5 wt pct W eutectic, preparedas described in a prior paper 3 and given a solutionizingheat treatment of 1323 K for 1 h followed by a waterquench, were deformed in constant-applied-load tension

G. GARMONG is a Member of Technical Staff, Rockwell Internatio-nal Science Center, Thousand Oaks, CA 91360.

Manuscript submitted February 28, 1975.

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