Wetting

Solidification processing of metal matrixcompositesA. Mortensen and I. Jin

Fundamentals of materials processes wherein areinforcing phase is combined with bulk moltenmetal are reviewed. Capillary phenomena whichgovern the incorporation of the reinforcement intothe matrix are discussed first, with focus onthermodynamics, mechanics, and chemistry ofwetting. Engineering strategies that have beenadopted to increase reinforcement wettability arealso addressed. Relevant transport phenomena arecovered next, with focus on the infiltration processand the rheology of composite slurries. The lastsection of this review is concerned withsolidification of the matrix after it has beencombined with the reinforcement, includingnucleation and growth of the matrix, and particlepushing by the solidifying metal. This reviewplaces emphasis on the physical phenomena thatgovern the processing and the microstructuraldevelopment of cast metal matrix composites.

IMR/235

1992 The Institute of Materials and ASM International.Professor Mortensen is in the Department of MaterialsScience, Massachusetts Institute of Technology, Cam-bridge, MA, USA, and Dr Jin is with the KingstonResearch and Development Centre, Alcan InternationalLtd, Kingston, Ont., Canada.

IntroductionMore than 25 years have elapsed since Kelly andDavies,! and Cratchley2 summarised their and otherpioneering efforts on metal matrix composite mater-ials in this journal. The concept and advantages ofreinforcing a metal are now solidly established, butit is only over the past few years that reinforcedmetals have truly been considered serious contendersin the arena of engineering materials. The relativelylate blooming of these materials has several causes.One is certainly that significant advances toward theproduction of affordable reinforcing materials havebeen made over the past decade. Of at least equalimportance, however, is the stronger emphasis thathas more recently been placed on the processing ofthese materials: it is now clear to the materialscommunity that if metal matrix composites (MMCs)are to gain industrial viability, their processing mustbe rendered economical and reliable, and must betailored to produce microstructures that will optimisecritical properties of the composite. The complexityof metal matrix composite processing, and its criticalincidence on both cost and properties of the resultingmaterials, has thus led to a recent surge of scientificinterest in its underlying principles.

A large and. increasing proportion of MMCs ispresently produced by solidification processing, inwhich the metal matrix is molten before it is combined

with the particles, whiskers, or fibres that are to serveas its reinforcing phase in the final composite material.The appeal of this route compared with solid stateprocesses (such as diffusion bonding or powder metal-lurgy) results from the ready availability of bulkmolten metal, and the relative simplicity of blendingtwo phases when one is liquid.

This article is a review of the fundamental principlesof solidification processing of MMCs. It is focusedon those processes where the molten matrix is in bulkform when it is combined with the reinforcement.Thus processes in which the metal is injected into apreform of the reinforcing phase, as practiced forexample in the squeeze casting process for pistonproduction by Toyota 3 and processes in which thereinforcement is added to the metal in fine dividedform, i.e. as loose particles, whiskers, or -fibres, ofwhich the Duralcan process4,5 is an example, arecovered. Rapid solidification and spray casting pro-cesses (such as plasma spraying, or the Osprey pro-cess6), despite a recent increase in their importanceand potential, are omitted, largely for lack of space.In situ composites, obtained via directional solidifi-cation of eutectic alloys, are also excluded fromconsideration: a composite is defined here as a mater-ial that results from the artificial combination of twophases to obtain a material that can not result fromconventional alloying.

For brevity, it is assumed that the reader has somefamiliarity with the nature and general characteristicsof metal matrix composite materials and with theirprocessing. An excellent introductory reference onthese topics can be found in the recent monographby Chawla.7 Technical aspects of solidification pro-cesses for metal matrix composite fabrication, whichha ve been covered in several preceding reviews, e.g.Refs. 8-21, are not reviewed systematically.

This article is divided into three main sections,concerned with the initial, the intermediate, and thefinal steps of metal matrix composite solidificationprocesses, respectively. The first section thus addresseswetting of the reinforcement by the metal, whichgoverns that step of the process in which the liquidmatrix and the reinforcement are first combined.Fluid flow, heat transfer, and solidification phen-omena that take place in the composite materialbefore it is fully solidified are covered in the secondsection. In the third main section, final solidificationof MMCs, the last step in forming the material andits microstructure, is reviewed. The issue of interfacereactivity is omitted, despite its importance in metalmatrix composite solidification processing. This isbecause it is too system specific an issue to be treatedotherwise than with a long list of pairs of materials,their reaction products, and their kinetics. Such a listis excluded here for lack of space.

International Materials -Reviews 1992 Vol. 37 NO.3 101

102 Mortensen and Jin Solidification processing of metal matrix composites

(3)

(1)

(2)CTSA - CTSL = CTLA cos(J

where (O"SA-CTsL)O corresponds to wetting with noreaction, ACTr represents the change in interfacialenergy resulting from the replacement of the unreactedsolid/liquid interface with at least one new interfaceafter reaction, and AGr is the free energy released atthe liquid/solid/atmosphere triple contact line. Therole and evaluation of this last term is very systemdependent and still a subject of some controversy.30,31A second and more general limitation in the validity

of equation (1) stems from the fact that the mechanicsof wetting of a porous medium generally result inirreversible energy losses, a fact that is well docu-mented in hydrogeology or reservoir engineer-ing.32-38 Mechanical irreversibility in wetting isrevealed by hysteresis of drainage-imbibation curves,wherein the volume fraction liquid is plotted versus

taneous and work is required, supplied in infiltrationprocesses via pressure applied on to the molten metal.A simple calculation22-25 gives the minimum pressurerequired

where Sf is the surface area of interface per unitvolume of metal matrix .Equation (1) also shows that the critical parameter

governing the wettability of a reinforcement by ametal is the work of immersion, lti = CTSL - CTSA. lti isrelated to the contact angle (J of the metal on theceramic (measured through the liquid metal) and tothe liquid metal surface tension CTLA by the Young-Dupre equation

To achieve an improvement in wetting of thereinforcement by a molten metal, one seeks to raiseCTSA, and/or, to lower CTSL- Lowering the metal surfacetension CTLA at constant CTSA-CTSL will reduce the wettingangle (J in wetting systems, but will never cause anon-wetting system with ()> 90 to become a wettingsystem in which (J < 90 (Refs. 24, 26, 27).With the proper values for lti, equation (1)gives a

lower bound for the pressure required for infiltration,because irreversible energy losses exist in the process.Interfacial chemical reactions may result in irrevers-ible energy expenditure, which complicates the energybalance involved in deriving equation (1). O"SL is nowa function of time and, hence, of position along thecomposite. Heat generated or absorbed by the interfa-cial reaction may also influence the kinetics of wettingand infiltration which, in turn, will influence theapparent O"SL- Calculating AP y is thus complicated,though equation (1) could still be used in conjunctionwith a comprehensive model of the infiltration processand knowledge of (a) the effective value of CTSL as afunction of time and temperature (keeping in mindthat in the presence of a reaction layer at the interface,the physical meaning of 0" SL is also changed) and (b)the influence of free energy released by the reaction.Under conditions of limited interfacial reaction, it hasbeen proposed28-30 that CTSA -O"SL is given by

composite

pressure1111

000000000000000000000000000000000000000000000000000000000000000000000000000000000000000000000000000

I

Che,mistry and mechanics of wettingThermodynamics and mechanics of wettingFor the producer of a MMC, the manifestation ofpoor wettability of a reinforcement by a metal is arepulsive force exerted between the liquid metal andthe fibres or particles one seeks to combine. Asufficient - but not necessary - explanation for suchpoor wettability is that an overall increase in surfaceenergy results from combining the two phases.Wetting in solidification processing of MMCs can

be quantified by measurement of the minimum press-ure AP y that must be applied on the metal, or,equivalently, of the force that must be applied on tothe reinforcement, to combine the two phases fully atvanishingly small velocity. In the infiltration of apreform, illustrated in Fig. 1, when AP y is negativethe metal is said to 'wet' the reinforcement which itcan then infiltrate or engulf spontaneously, as doesmolten copper with a bundle of tungsten fibres. Thisis unfortunately seldom the case in metal matrixcomposite fabrication: wetting is generally unfavour-able. Positive pressure, and hence external work, mustthen be provided to create the composite.If for now infiltration is focused on and it is

assumed that it can take place reversibly (as it wouldwere the metal flowing through a bundle of parallelcapillary tubes of the reinforcing phase withoutreacting chemically with the latter), the work necess-ary for quasi static fabrication of a volume of compos-ite containing 1m2 of interfacial area is equal to thatof replacing the solid reinforcement/atmosphere sur-face, of energy CTSA, by the solid reinforcement/liquidmetal surface, of energy CTsv If CTSA > CTsv the metalcan spontaneously infiltrate and therefore wet thepreform according to basic thermodynamic analysis.Alternatively, if CTSA < CTSL, the process cannot be spon-

International Materials Reviews 1992 Vol. 37 NO.3

1 Infiltration of fibre preform by molten metal.For quasistatic infiltration, applied pressuremust oppose capillary pressure drop atinfiltration front, and constitutes therefore ameasure of preform wettability

Mortensen and Jin Solidification processing of metal matrix composites 103

pressure applied on a liquid that is infiltrating ordewetting a porous medium (examples are given inRefs. 32, 33, 37, 38). Such hysteresis occurs becauseenergy is mechanically lost in the process of wettingand dewetting the porous medium, even if there areno chemical reactions between the porous mediumand the liquid. Irreversible energy losses in infiltrationresult, among other reasons, whenever the invadingliquid encounters a constriction in the porousmedium, through which it jumps irreversibly into thepore on the other side of the constriction (these jumpsbear the name of 'Haines jumps'). Another cause forhysteresis in drainage-imbibation curves is roughnessor chemical inhomogeneity on the reinforcement sur-face, which cause other, more microscopic, jumps asthe invading fluid passes over the pore walls, andlead to a difference between the macroscopicallyapparent advancing and receding contact angles.39-41An important practical consequence of mechanicalirreversibility in infiltration is the fact that a wettingangle lower than 90 does not imply that metal willspontaneously infiltrate a packed bed of particles ora bundle of fibres.33,42The pressure given by equation (1) is therefore the

lowest pressure that may possibly drive the moltenmetal into the reinforcement preform. In reality, nosingle wetting pressure exists because of the complexmicroscopic mechanical phenomena that take placeduring wetting of a porous medium. Wetting is grad-ual, the volume fraction liquid being a step-wisecontinuous function of applied pressure, plotted indrainage-imbibation curves.43If parallel fibres touch,for example, the pressure needed to infiltrate contactlines between two fibres is calculated to be infinite,24which agrees with observations of voids at fibrecontact points24,44-48 and implies that infinite press-ure is required for perfectly full infiltration. Similarly,asperities on the reinforcement surface may preventachievement of full reinforcement contact with a non-wetting liquid.24,40 'Full infiltration' therefore oftenmeans, from a practical standpoint, infiltration of allbut perhaps a very small portion of open porosity inthe preform. Then, with AP y now defined with refer-ence to practically full infiltration, experimental evi-dence22,32 and various idealised calculations ofcapillary pressure49-51 indicate that equation (1) willonly underestimate that pressure relatively slightly(by a factor of two or so) as long as the propercontact angle and specific surface are used, and otherenergy loss mechanisms such as viscous drag areproperly modelled.25Turning now to casting processes in which the

reinforcement is added loosely to liquid metal, analy-sis is relatively simpler because the geometry of theproblem is fairly well defined if the interactionbetween neighbouring particles in the process of beingengulfed is neglected. The problem is then that of oneisolated particle at the free surface between a liquidand a gas, which is a relatively classical problem incapillarity that has recently been reviewed by severalauthors. 52-55 A conclusion from equilibrium con-figuration calculations is that for quasistaticengulfment of a smooth particle, an energy barriermust be overcome even if the contact angle is lessthan 90, in fact as long as it is finite (8) 0). This can

2 Immersion of cube in liquid. Irrespective ofwetting conditions, as long as wetting angleof liquid on cube is positive, the transitionfrom step 1 to 2 is spontaneous, while thatfrom step 4 to 5 is energetically unfavourable(equation (2), with 9> 0)

be illustrated with elementary analysis of the idealisedcase of engulfment of a cubic particle with one facehorizontal, Fig. 2. For the last face to be wetted, oneinterface (that between the particle and the gas) mustbe replaced with two interfaces (particle/liquid andliquid/gas). From equation (2), force must then beapplied on to the reinforcement to-produce the com-posite if 8> O.The very last portion of the surface ofa smooth particle that is engulfed is always horizontal,as is the top surface of the cube in Fig. 2, so the sameconclusion obtains. Because of the small dimensionsof the particles used in reinforcing metals, gravityseldom can supply sufficient force to drag them intothe liquid metal. The usual method for producingcomposites by adding discrete reinforcement particlesat the melt surface therefore makes use of motion inthe metal, which applies force on to the particles viaviscous drag along their surface to pull them into themelt (e.g. Refs. 56, 57).The process of particle engulfment in stirred liquid

metal was recently modelled by Ilegbusi andSzekely,53who added a term for viscous drag to thegeneral force balance proposed by previous research-ers, and thus created a link between wetting andprocessing parameters. An empirical drag coefficientfor flow of particles within an electromagneticallystirred turbulent fluid was used after multiplicationby the fraction of wetted particle surface. Results forspherical boron carbide particles dragged into electro-magnetically stirred magnesium are given in Fig. 3.These show the significant influence of both particlesize and contact angle on the fluid flow velocityrequired for incorporation.Limitations exist to such modelling work, of course,

if only because of the influence of cuspidal pointsgenerally present on the surface of discrete reinforce-ment materials such as particles, flakes, or short fibres,where the wetting angle 8 is not defined. The oxidelayer covering the molten metal is also' expected tobe thicker here than in infiltration processes, and thusto playa significant role in wetting. This is shown inrecent work by Stefanescu et aI.,58,59 who measuredthe centrifugal force required to immerse particles inaluminium. With increasing particle size, a suddendecrease in the force required for engulfment wasobserved, taking place at the same particle radius for



180

160 ---------- ---140

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~ 120 \\(]) \\

\

100 \\

80

60900 1000 1100 1200 1300

T,K

4 Schematic variation of wetting, angle (Jmeasured on -flat substrate with temperaturefor pure aluminium on alumina for two partialpressures of oxygen: .2:- - - total pressure",,-10-3 Pa, oxygen. partial pressure >10-6 Pa;-- total pressure 5 x 10-5 Pa, oxygen partialpressu,re 10-15Pa, Difference between the twocurves arises from presence of oxide layercovering metal at higher pressures, whichevaporates around 1150 K. From Refs. 89, 97(b)-----a

8

6

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E>- 2!:::u0..JW 0 30 60 90 . 120 150 180>0 CONTACT ANGLE, deg::::>..JI.L. 3~ o contact angle 1500::::>~ contact angle 600z 2~

o 100 200PARTICLE DIAMETER, ~m

a influence of wetting angle of liquid magnesium on particle 100 ~mdia.; b influence of particle diameter

3 Minimum fluid velocity required to engulfspherical particle of boron carbide into moltenmagnesium at vortex created by electro-magnetic stirring, based on calculations bylIegbusi and Szekely53

two oxide particles of different density. This suddendecrease was interpreted by the authors 58 as resultingfrom the particle diameter exceeding the thickness ofthe oxide skin covering the metal. Particles also havebeen found to agglomerate at the metal surface beforeincorporation,6o,61 which adds to the mechanicalcomplexity of the problem. Despite these limitations,the results from Ref. 53 show that correlation ofprocessing parameters with appropriate wetting datais possible, and agrees with empirical knowledge thatfiner particles are harder to stir into molten metal(e.g. Refs. 62,.63).

Measurement of wettabilityThe conventional experimental method used for meas-uring wettability by a metal is the sessile drop experi-ment. This consists in measuring the shape andcontact angle 8 of a drop of the liquid metal matrixat rest on a flat substrate of the ceramic reinforcementmaterial, or in some instances64,65 on fibres. Thesessile drop experiment enables measurement of bothO'LA and 8, and can therefore allow calculation of~p y via equations (1) and (2). For example, the wet-tability of carbon,64-75 of SiC,52,65.70,76-82 ofB4C,51.69.70,78.83,84 and of AI20352.71,74,75,84-94 byaluminium and its alloys has been measured and


found to be poor below about 1223 K. Reviews ofsuch wetting angle data have been published byDelannay et al.,44.83 Russell et al.,52 Nicholas,95 andNaidich,70 this last review being particularly extens-ive. A general conclusion from the large number ofwetting angle measurements that have been made todate is that'the technique is precise and reliable, andyields data amenable to scientific analysis. Examin-ation of sessile drop data relevant to systems ofinterest for metal matrix composites reveals that thesealso depend on several complicating factors:

(i) with several metals including aluminium andtin, the presence of a layer of oxide on the metaldroplet prevents it from contacting the substrateand therefore influences wetting detrimentallyduring the exp~riment. Several observationslead to this conclusion, including visual obser-vation of the drop and the influence on 8 ofmetal cleaning before melting, 84.87,88 the pres-ence of abrupt wetting/non-wetting transitiontemperatures that are relatively independent ofthe substrate but depend on the oxygen partialpressure of the atmosphere, 84,87,92,96Fig. 4, andthe strong influence on wetting angle of alloyingadditions,66,73,94 or temperature changes92,96which are known to affect the oxide coveringthe metal

(ii) wetting angle data are often time dependent.This phenomenon is attributed to chemicalreactions occurring at the metal/substrateinterface, as well as to the necessity for themetal to break thro'ugh its surface oxide layerbefore achieving intimate contact with the sub-strate. Examples of this effect can be found inRefs. 29, 66, 71, 74, 75, 77, 82, 88, 90-92, 96,98-100.


A further possible complicating effect in assessingwettability is the presence of a precursor film aheadof the spreading metal, which is reported for metallicsystems in Ref. 101 and modelled for the case oforganic and other liquids spreading on a solid sub-strate,102 or of reaction between the solid and theliquid ahead of the wetting line.29 Added compli-cations of a more mechanical nature exist as well,such as the well known fact that advancing andreceding contact angles differ as a consequence ofsurface roughness or variations in the chemical natureof the wetted substrate.39,40,92,102The dependence of sessile drop experiment data on

these epiphenomena often invalidates direct transpos-ition of sessile drop wetting angles to metal matrixcomposites solidification processes, because the velo-city of the three-phase contact line in these processesis much higher than in a sessile drop experiment,because the oxide layer covering molten aluminiumis most likely disrupted by the reinforcement as itcombines with the metal,103,104or because of micro-scopic roughness and chemical heterogeneity on thesurface of many reinforcements (e.g. commerciallyavailable polycrystalline alumina fibres105).These sev-eral limitations of the sessile drop technique haveserved as an impetus for the development of alterna-tive measurement techniques for wettability, often lessprecise and less amenable to detailed scientificinterpretation, but more directly pertinent to metalmatrix composite solidification processing.Dipping experiments have been proposed to repli-

cate more accurately the dynamic nature of wettingin composite fabrication. In hanging plate or rodexperiments, the weight of a straight-edged solidpartly immersed in the liquid metal and the shape ofthe meniscus near the solid are recorded 106to measureboth aLA and 8. Compared with the sessile dropmethod, this technique is somewhat more complicatedbut is equally precise, and has the advantage thatdynamic wetting conditions can be studied by movingthe object into and out of the liquid.A different type of dipping experiment was used by

Choh and co-workers/07-111 who studied the rate ofspreading of molten aluminium and aluminium alloyson flat discs of carbon and SiC substrates, which weresuddenly completely immersed in the molten metal.Curves of fraction of disc surface wetted by the moltenmetal as a function of time were found to have asigmoidal shape typical of thermally activatednucleation and growth reactions. Their temperaturedependence was explained in terms of possible criticalstepscin the chemical interaction between metal andsubstrate. These experiments present, however, thedisadvantage of using flat substrates instead of actualparticulate or fibrous reinforcements.Experimental procedures using particles of ceramic

material to measure wettability have recently beenestablished. In most of these,112-117molten metal isforced under pressure into a packed bed of the powderheld at the same temperature as the metal. Thedistance travelled by the metal during a time intervallit into the packed bed is measured for various appliedpressures, and from extrapolation of these data tozero infiltration distance, a minimum ('threshold')pressure for initiation of the metal movement into

the packed bed of particles is recorded as a functionof alloy composition, atmosphere, temperature, andlit. This threshold pressure does not generally equalliP y in equation (1) because (a) it is a measure ofinitiation of wetting, which can be influenced bypercolation and other preform entrance effects and(b) the infiltration distance which is recorded fre-quently includes regions of partially infiltrated com-posite. These experiments have been performed forcarbide reinforcements and matrices based on alumin-ium or zinc.112-118A practically important and scien-tifically interesting finding of these investigations isthat the infiltration distance and consequently thethreshold pressure depend on the time of pressureapplication. In particular, for a given set of experi-mental conditions, an incubation time exists, beforewhich no infiltration occurs.115-117This observationagrees with results of Choh and co-workersl07-111on the spreading of metal on dipped flat substrates.Direct measurement of the capillary pressure drop

liP y at the infiltration front, given by equation (1) forreversible infiltration, requires that data be collectedwhile the metal is being combined with a preform ofthe reinforcing phase, to replicate the dynamic andmechanical characteristics of wetting in infiltrationprocesses. Such experimental data have been gener-ated in the wake of recent studies of infiltrationprocessing. 104,119,120The measured capillary press-ures were shown to be interpretable as closelyapproximating liP y of equation (1),25and could there-fore be compared with wetting angle data via equation(2). Such comparison yielded values of the apparentwetting angle that were much closer to sessile dropwetting angles in high vacuum than to angles meas-ured in air or in lower vacuum. This, in turn, mayindicate that the oxide layer which usually coversmetals such as aluminium during static sessile dropexperiments and in measurement of threshold press-ure is washed away by the reinforcement phase duringpressure infiltration, as was first proposed by Capple-man et a1.103 This interpretation would also agreewith observations of a fine layer of alumina along thefibre/matrix interface in some cast aluminium matrixcomposites.121,122These experiments accounted forthe effect of differing metal and reinforcement initialtemperatures, which causes the wetting process totake place concomitant with intense heat flowfrom one phase to another, and possibly matrixsolidification.Measurement of liP y from infiltration experiments

is only possible when full wetting in partly infiltratedcomposites can be unambiguously identified and mea-sured during infiltration, as was the case in Ref. 104.This is generally not the case, because the geometricalcomplexity of the reinforcement preform causes wet-ting to take place gradually. To account for this,wetting in infiltration has been recently characterisedby measuring drainage-imbibation curves which givethe volume fraction metal as a function of appliedpressure, as is general practice in hydrogeology orreservoir engineering.43,123,124As seen in the section'Infiltration of preforms' below, provided there is noinfluence of infiltration velocity on these curves, thesedrainage-imbibation curves are useful in modellinginfiltration.



Engineering approachesOn the basis of wetting angle data, spontaneouswetting of a reinforcement by molten metal is foundto be promoted by. some degree of reactivity of themetal with the substrate, 11-13, 70 or by a reduction inthe tenacity of the oxide layer on metals prone tooxidation such as aluminium.13 Based on these obser-vations, much work has been done to design chemicalmeans of enhancing the generally poor wetting ofreinforcements by metals. A majority of this workcan be classified into three broad categories, namely(a) reinforcement pretreatment, (b) alloying modifi-cations of the matrix, and (c) reinforcement coating.The fibre surface energy O"SA can be raised by

changing the chemical nature of the atmospherebefore infiltration. Results have been published onthe influence of heat treatment of alumina, siliconcarbide, and graphite particles on their wettability,which improved their ease of incorporation into alu-minium melts. This was attributed to desorption ofgaseous species from the reinforcement surface duringthe heat treatment. 12,57,125-129Alloying additions to the matrix have been shown

to affect the wetting angle on, and ease of incorpor-ation of, many reinforcements. Effective additions fallin two categories:

(i) additions that promote reactions between thereinforcement and the matrix: Li in Al foralumina fibres,83,105,130,131Li for wetting ofSiC by AI,83,132,133or Mg,134Ti in AI-Sn alloys(but not pure AI) for wetting of SiC;133carbideformers in Al for wetting of carbon fibres;135,136Si in Al for wetting of carbon particles.126Theirefficiency agrees rather well with wetting angledata, which indicate that, in general, alloyingadditions that promote reactivity between themetal and the substrate lower the wettingangle68,70,83,95,137

(ii) additions to aluminium that do not promotereactions with the reinforcement, but modifythe characteristics of the oxide layer on themetal surface: Mg in Al with most reinforce-ments 13,54,60,62-64,73,94,101,113,118,133,138,139orLi in' AI13,105,131-133(with alumina or oxidecovered SiC reinforcements, these two categor-ies obviously overlap). Again, their effect corre-lates well with wetting angle data, in the sensethat the transition temperature from wetting tonon-wetting, as well as the wetting angle, aredecreased by these additions, e.g. Li, Mg, or Cain aluminium 69,71,89all elements known toaffect the oxide covering aluminium.66

These two practical methods either induce reactionsat the fibre/matrix interface, or modify the oxide layerthat usually coats molten aluminium. Reactionsbetween the reinforcement and the matrix are gener-ally undesirable because they degrade the reinforce-ment strength. A different approach based on thesame principles is to coat the fibre or particle surface.Such coatings have been extensively used, and canalso be classified broadly into the two categories usedfor alloying additions:

(i) coatings that are designed to react with thematrix. These are numerous, and include metal-lic coatings for various reinforcements in


aluminium,14,15,131,140-153oxide coatings forcarbon fibres in magnesium 154,155or siliconcarbide in aluminium 118,129 .

(ii) coatings that are designed to react with theoxide layer covering molten aluminium. Theseinclude loose coatings of K2ZrF 6,13,156-165which is a known fluxing agent for aluminium.It has been argued that most coatings thatcause reaction induced infiltration by alumin-ium with no applied pressure below 900C doso mainly because they disrupt the oxide layeron the metal.13,87Also, since it has been docu-mented that aluminium oxide is present at theinterface of carbon fibres treated by the TiBprocess in aluminium, it can be argued that thiscoating technique also acts by oxygen getteringor formation of a spinel with the oxide layeron the metal,87 Fig. 7 of Ref. 166. Similarly,MgO coatings that have been used to increasewettability of alumina particles by aluminiummay in fact be active via the oxide on themetal. 167

The dividing line that separates these two categoriesof alloying additions or coatings is thus not clear.There is nevertheless an established perception thatthere are at least two strategies for promoting wet-tability of reinforcements by metals, particularly alu-minium: to obtain wetting, one can modify the fibreor the matrix with something that will (a) react withthe metal or (b) disrupt its outer oxide layer.Many other methods have been developed to

improve wetting, which do not fall into these categor-ies, either by nature or for lack of documentationand/ or of understanding of the actual mechanism thatpromotes wetting. These include (a) various proc.essesinvolving fibre treatment by molten sodium for infil-tration of carbon or alumina fibres by aluminium ormagnesium,152,168-170(b) the TiB process involvingCVD deposition of Ti-B mixtures on carbon fibresbefore infiltration in oxygen free atmospheres byaluminium or magnesium,171-175 (c) pretreatment ofSiC by dehydrated sodium tetraborate for infiltrationby molten aluminium, 176(d) pretreatment of carbonby tetraisopropyltitanate for infiltration by moltenaluminium or magnesium,177 (e) pretreatment of B4Cby one of various alcohols or other organic solventsfor infiltration by molten aluminum at elevated tem-peratures,178 (f) dispersion of solid magnesium nitridebetween carbon fibres for infiltration by mag-nesium,179(g) the Lanxide process, reportedly usingmagnesium alloy additives, nitrogen containing,oxygen free atmospheres and non-disclosed temper-atures for infiltration by aluminium of severalreinforcements,180-184(h) the use of nitrogen in aidingwettability of SiC and Al203 particles by mag-nesium,185,186and (i) sodium tetraborate additions toaid wetting of alumina particles by AI-7Si-O-1 Mgalloys*.63 Last, the elegant experiments of Nogiet al.,187 who improved the wettability of zirconia byliquid copper and iron using a dc voltage appliedacross the metal/ceramic interface, and interpretedthis effect as resulting from the dissolution of oxygenat the zirconia/metal interface.

* All compositions are in weight per cent unless otherwise stated.


Improvements in the wetting of reinforcementsduring infiltration or incorporation of particles intoa melt have also been obtained by mechanical means.As seen above, with parallel fibres, infinite pressureis theoretically required to infiltrate the line of fibre-to-fibre contact. Calculations of capillary forcesbetween two parallel fibres188,189show that the press-ure required to infiltrate fibre contact lines decreasessignificantly with even a small separation betweentwo parallel fibres, and that forces between fibresduring infiltration are large, causing them to clusterduring infiltration. These problems were solved with'hybrid' fibre preforms, first proposed by Towata andco-workers, in which parallel fibres are individuallyheld apart during infiltration by small particles orwhiskers.46,190-199Other examples of improvementof wetting by mechanical means include the use ofsemisolid metal of increased apparent viscosity in theCompocasting process,200-210the agitation of parallelfibre bundles using alternative current in a magneticfield to ease their permeation in gravity driven infil-tration211 and vibration induced wetting improve-ments of alumina fibre preforms by aluminium.212,213

Chemistry of wettingOn a more microscopic scale, wetting can be examinedin terms of atomic bonding by comparing the initialstate (one or two free surfaces) with the final state(the interface). The results may allow one to designnovel strategies for improved wettability, ultimatelyoptimising mechanical properties of the final castmaterial as well, because an understanding of inter-facial strength can be gained from knowledge ofinterface chemistry.In the solidification processing of MMCs, one

wishes to create a bond of controlled strength betweena liquid metal and a generally non-metallic reinforce-ment. The kinetics of this bond creation process willdepend on the initial structure of each surface present,on the mutual attraction exerted by the two surfacesas they approach one another, and on the structureof the resulting final interface. Most relevant work istheoretical and, with the advent of advanced charac-terisation techniques such as high resolution micro-scopy, microstructural. A large fraction of presentresearch seeks to improve knowledge of initial andfinal surfaces and interfaces, in part because of thesomewhat static nature of sessile drop experiments(which form the bulk of available wetting data), andof the considerable experimental difficulty involvedin the experimental study of dynamic wetting.

Surface of metalsThe origin of the surface tension of a material lies inthe disruption of the bonding state for the atoms ator near to the surface or interface. In the case of ametal in contact with vacuum, this manifests itself asa local disruption of the electron gas density nearthat surface. An electric double layer, formed by apositively charged background and a negativelycharged electron cloud that extends further into space,results. Several authors have proposed calculationsof the resulting surface energy using various degreesof sophistication in their description of the electrons

and the positive ions, and have been quite successfulat predicting from first principles the surface tensionand electronic properties of several metals.214-219Ona more macroscopic level, several correlations havealso been experimentally found and modelled, relatingthe surface tension of pure liquid metals with theirheat of vaporisation and other macroscopic propertiesof the metal. 83,220-226Most metal surfaces are usually modified by the

presence of adsorbed atoms from the atmosphere orfrom within. The effect of such adsorbed elements onthe surface tension of the metal has been modelledfrom thermodynamic considerations, ranging fromGibbs isotherm treatment to more complex treat-ment, 83,95,217,219,222,223,226-228and it is well knownthat even minute quantities of alloying elements orvery small partial pressures in the vapour phase canmodify the surface composition, morphology, andenergetics to a considerable extent. Furthermore, thelevel of metal surface contamination may depend onthe process, possibly varying for aluminium betweena relatively clean surface (at elevated temperatures invacuum or because of skimming in preform pressureinfiltration, for example), to a surface covered withan oxide crust many atomic layers thick.

Reinforcement surfaceThe surface structure of oxides is of interest for tworeasons: because several reinforcements are oxides, orare coated with an oxide (such as Si02 for SiC56,79)and because the metal surface is often covered by itsoxide. From an energetic point of view, treatment ofionic solid surface energies is generally based onelectrostatics, and the surface energy can be correlatedwith the cohesive energy of the solid.217,219,225,229,230From a structural point of view, it is noted that ionicsurfaces can be electrically charged.229,231A structuralmodel of oxide surfaces by Weyl, in which the cationsat the surface are screened by the anions because ofthe much higher polarisability of the anions, is rela-tively well accepted.7o,83,232-237According to thismodel, the surface of an oxide - and a fortiori thatof many (oxidised) metals in air - consists predom-inantly of highly polarised oxygen atoms.The surface of clean graphitic carbon is complex

and depends on the orientation of the surface withrespect to the basal planes. The surface energy ofdiamond is roughly one-half of its cohesive energy,225and differs significantly from that of graphite.238 Theorientation of basal planes at the surface of carbonfibres is known to vary significantly from fibre tofibre, which results in variations in fibre chemistry(Refs. 239-241 contain recent reviews). The surface ofcarbon fibres is very active and is thus heavily con-taminated with adsorbed species, predominantlyoxygen containing atomic groups. It also is oftenmodified for improved wetting and bonding withpolymer matrices,242,243because these are used inmost applications of carbon fibres.The surface of carbides or nitrides varies with the

character of the bond (covalent or part metallic), andis for most practical purposes usually that of an oxideof one or several of their constituents (for example,the surface of SiC reinforcements generally comprisesa layer of Si02).



Interface chemistryDepending on the materials and the degree of bondingachieved, the interface can be one of two kinds:Reversible bond only, generally physical There is nochemical interaction between the two materials oneither side of the interface. There is, nonetheless, anattractive force binding the two, due to Van der Waalsforces. These are most often treated as dispersionforces. This 'physical bond' is established in a revers-ible manner, and is generally the only bonding mech-anism operating in wetting by organic solvents,aqueous solvents, or polymers.83,244,245Theoreticaltreatment of the physical bond is relatively wellestablished, and its energy can be calculated by useof the London formula or other more completetreatments.70,83,245-250The resulting bond energy isrelatively low and is temperature independent. 70Thistype of bond is reversible, since the physical act ofseparating the two surfaces after contact can be madewithout significant additional energy losses, bondstrengths being low. Reversibility of bonding at theinterface has indeed been observed in polymer matrixcomposites.242 Other types of reversible bonds ofhigher bond energy have been proposed: the hydrogenbond;251 or electrostatic attraction due to imagecharges across the interface.236,238,250With suchreversible bonding, the science of wetting and that ofadhesion are closely related, so that additional back-ground can be gained from several reviews pertainingto adhesion.251-254

Chemical bonds On the basis of sessile drop experi-ments and experience in metal matrix compositesprocessing, it has become clear that in interfacesseparating a metal from another material, a strongerchemical bond can be established, resulting fromchemical interaction across the interface. Such bondsgenerate works of adhesion (by definition, the energyliberated on their formation from two free surfaces)that are temperature dependent, are typically an orderof magnitude higher than for physical bonds andtherefore govern interfacial strength. In practicalterms, their importance translates into the observationnoted above that one of the two methods of improvingwettability in metal matrix composites solidificationprocessing is to 'add something to the fibre or thematrix that will promote reactions between the two'.These are irreversible bonds in the sense that onceestablished, they cannot be broken without also dam-aging adjoining material on either side of the interface.Evidence of the establishment of a chemical bond

after wetting of ceramics by metals has been foundin several studies of wettability by metals using thesessile drop technique, 68,78,98,99,137,232,255-261or instudies of adhesion of metals to ceramics,262-266andis reviewed by several authors. 52,70,83,98The bond isno longer constrained to remain in a surface, but maybe several atomic layers thick, or may even extendacross an interfacial region visible under an opticalmicroscope. The connection between wetting andadhesion becomes more tenuous because the interfacecannot in practice be undone once it has formed (e.g.Ref. 267). A theoretical treatment of the interface isalso more difficult for this reason. For example,thermodynamic equilibrium will in many cases not


be rigorously attained until complete transformationof at least one of the two initial phases.Models for such an interface vary in nature. Gener-

ally, it is described as sharp, i.e. as consisting of amonolayer of chemical bonds bridging dissimilarphases. Most models describe bonding across theinterface in thermodynamic terms to predict interfaceenergy.70,97,227,232,237,268-273Other models are basedon quantum mechanical calculations,274-276 or on anextension of the 'jellium' mode1.277Yet other modelsanalyse the bond qualitatively based on the bondwithin carbide reinforcements,78,98,257,273or estimateelectrostatic forces resulting from electron exchangebetween the two solids174,253,254(this contributionwas estimated by Krupp253 to be negligible; however,electrostatic attraction pressures have been foundexperimentally across the interface in composites ofaluminium reinforced with carbon fibres278).Micro-scopic investigations of interfaces are now possiblewith advanced high resolution transmission electronmicroscopy, and recent results are reviewed by Riihleand Evans.279 These include recent investigations ofmetal/oxide interfaces, for which it was shown byMader280 that oxides produced by internal oxidationdisplay a layer of oxygen atoms along their lowenergy facets with the metal matrix, in agreementwith previous work reviewed by Mader in that samereference. More detailed reviews of metal/ceramicinterface structure and chemistry can be found inRefs. 70, 83,279,281.

Influence of alloy additions and reactive wettingMost research on the structure of surfaces andinterfaces is for relatively pure metals and reinforce-ments. In practice, the metal is generally impure oralloyed, and the reinforcement surface may be oxi-dised or otherwise contaminated. When the metal isalloyed, both its surface tension and the metaljreinforcement interface energy are affected, especiallyif alloying additions tend to segregate at surfaces.Naidich 70 proposed that oxygen dissolved in metalscan form ionocovalent metal-oxygen clusters whichsegregate at ceramic/metal interfaces by coulombianattraction. This mechanism explains the strongreductions in wetting angle that result from oxygendissolution in molten metals contactingoxides.28,70,237This interpretation has recently beenconfirmed by Kritsalis et al.282The influence of alloying on wetting was recently

modelled by Li et al.28,227 assuming a regular solu-tion alloy and using statistical thermodynamic calcu-lations. Their model allows prediction of the influenceof alloying additions at low dilution, and agrees withexperimental data. 28,227,283In many practical cases, wetting is accompanied by

chemical changes that extend beyond a single metal/ceramic interface, such as dissolution of the solid inthe liquid, diffusion of the liquid into the solid, orformation of novel interfacial phases. These phen-omena introd uce a time dependence to the wettingprocess and complicate its analysis significantly. Reac-tive wetting has been addressed by severalauthors30,31,70,284from a thermodynamic point ofview. Of particular interest is the unsolved questionof the influence exerted by free energy released in


reactions between the liquid and the solid on wetting.A recent review of theory and experimental data onreactive wetting is given by Laurent. 30

Forscheimer equation297-302

f- VP = [,uVm(l- gs)K-1

+ BPmlv1 - vsl(vI - vs) (4)

where d is a characteristic length of the reinforcement(for example, the fibre diameter in fibre preforms, inwhich case Rec ~ 1 (Ref. 119)), the second term betweenbrackets can be neglected. This is most often thecase, and the Forscheimer equation then reduces toD'Arcy's law

where f is the local volumetric value of gravitational,centrifugal, or electromagnetic body forces in ~ V, Pthe liquid average pressure in ~ V, ,u the viscosity ofthe liquid metal, Vm the volume fraction metal, andgs the volume fraction of the metal that is solid; Vsand VI are, respectively, the average solid and liquidvelocities..K, the local symmetric permeability tensorof the preform and B are functions of Vm and of thevolume fraction and morphology of solidified metal.When the relevant Reynolds number is below acritical value

Transport phenomenaThis section is concerned with solidification pro-cessing steps in which the metal and the reinforcementhave been combined, and the metal matrix remainsat least partly liquid. These precede the final pro-cessing step in which solidification of the metal iscompleted. As in wetting, a distinction is madebetween cases where the reinforcement in the compos-ite constitutes a mechanically self-sustaining preform,and cases where it consists of individual particles,short fibres, or whiskers dispersed within the metalin the final composite. In the former case, the compos-ite is directly cast by infiltration of the preform,generally to its final general shape. In the latter case,the composite is generally formed by stirring thereinforcement into molten or semisolid metal to yielda free-flowing composite slurry. This slurry is sub-sequently cast into shape using processes similar tothose for unreinforced metals. These two classes ofprocesses have been combined,61,285-288 but theydiffer substantially in the transport phenomena bywhich they are governed. These are reviewed in turnin this section, beginning with infiltration processing.

(5)

(6)

where the function F describes the first drainage-imbibition curve of the preform, as described in theprevious section. If back pressure of unvented gasbuilds up, P g may be significant.188~189,314


which has been used in most recent analyses ofinfiltration processing.50,104,119,120,303-312There are,however, infiltration processes in which Rec is greaterthan one.313Assuming for simplicity that the densities of liquid

and solid metal are equal and constant, continuity ofmatter dictates

(7)

(8)

(9)

. (10)

oltfat = - V(ltfvs)o(Vm)gs-at= -V(Vmgsvs)

where ltf, Jt;" Vm are, respectively, the local volumefraction of reinforcement, pores, and metal, such thatltf+ Jt;, + Vm= 1.Drag due to flow of gas initially present in the

preform is generally ignored because of the muchlower viscosity of the gas. Provided wetting is notsignificantly influenced by the velocity of infiltration,Vm can be considered a function of the differencebetween the pressure P in the metal and that of theatmosphere Pg (if any) in the pores within ~ V

Infiltration of preformsMechanics of infiltrationThe infiltration process has been used for severaldecades to produce metallic composite materials.Early applications of the process, which were focusedon capillarity driven processing for the production ofcermets and particulate metal-metal composites, arereviewed by Lenel.289Several analyses of infiltrationwere proposed for these, in which the porous mediumwas modelled as a bundle of cylindrical tubes beinginfiltrated by the metal according to the Hagen-Poiseuille equation.289-291 This assumption, whichhas also been made in recent work by several investi-gators,116,292-294eases conceptual visualisation of theinfiltration process. However, it has no physical justi-fication, it introduces errors (for example, the effectsof inertial losseswhich arise because of tortuosity in theporous medium are unaccounted for, or erroneouslyascribed to, turbulence) and it brings no real simplifi-cation in analysis. Limitations of straight tube modelsare reviewed for example by Scheidegger295 andMorel-Seytoux,296 and practical difficulties encoun-tered in their use in materials processing are describedin Ref. 292.Generally, the reinforcing phases (fibres, particles)

are small enough that the preform can be describedas a porous continuum. Theoretical analysis is thenbased on consideration of a small volume element~ V, which contains several pores and reinforcementelements such as fibres, particles, etc. Within ~ V,temperature, matrix composition, and fraction solidare assumed to be essentially uniform, and velocitiesof the fluid and solid phases are averaged. Flow ofthe liquid metal in ~ V is generally governed by the


Immiscible flow of several liquids in porous mediahas been studied in several branches of engineeringscience, including reservoir engineering and hydro-geology.34,38,296,315,316In the particular case of uni-directional infiltration driven by a constant pressuredifference APT between the preform entrance and thegas phase, if Re


region 2

region 3

vent

infiltrationfront

region 2

infiltrationfront,

W0::::::>~ To region 50::: : 1W TE ~ .~ region 3l region 1 1 1~ Ii; i region 4

A A'

II

I region 4A'

Tm .region 3: region 1

IA

W0:::

~ To0:::Wa..2:wI-

DISTANCE

5 Schematic illustration of infiltration ofpreform by pure metal, under conditions suchthat solid metal forms at infiltration front incontact with fibres and along mould wallbecause of external cooling. Fibre preformthen contains four regions, shown in figure:region 1 is composed of fibres, solid metal,and flowing liquid metal, region 2 of fully solidmetal, region 3 of fully liquid and flowingmetal, and region 4 the uninfiltrated portionof preform. When preform or die are abovemelting point of metal, regions 1 or 2,respectively, are absent

and form 'fingers' of liquid matrix extending into theregion of semisolid matrix.120,340In the case of unidi-rectional adiabatic infiltration driven by a constantapplied pressure, it was also shown 119,311that theBoltzmann transformation is compatible with thermaland solute transport as well as fluid flow equations.With an alloy, matrix solidification during infil-

tration results in macrosegregation, featuring sig-nificant solute enrichment at the infiltrationfront. 311,339,341-343This occurs because solute is par-titioned unequally between liquid and solid duringsolidification. Since the solid phase is trapped withinthe preform while liquid metal flows further down-stream, different chemical elements concentrateupstream and downstream. With a hypoeutecticbinary alloy, solute enrichment is found downstreamin the composite, as illustrated in Fig. 6.Where solid metal has formed during infiltration

by cooling at the fibres (regions 1 and 5 in Figs. 5and 6), the final solid matrix grain size is relativelysmall, of the order of interfibre spaces with finealumina fibres and aluminium alloy matrices. This isbecause of rapid cooling by the fibres and ensuinghigh nucleation rates in the matrix. In the samecomposites, large grains are found in the region ofremelted matrix (region 3 in Figs. 5 and 6). Quantitat-ive treatment of these phenomena allows predictionof infiltration kinetics, matrix microstructure, andsolute distribution. 119,307,311,339,341-343Process para-meters can then be tailored, for example, to maximiseuniformity of microstructure and composition in thecomposite.

~ CEr~V~;HHHHHH~gion 5~ Co : l ..~~ region 31 region 1 ~ ~region 4u ,. . .

A A'DISTANCE

6 Schematic illustration of infiltration ofpreform by binary hypoeutectic alloy, underconditions such that solid metal forms atinfiltration front in contact with fibres andalong mould wall due to external cooling. Fibrepreform then contains same four regions aswith pure metal, with region 1 defined asregion where solid primary metal is present.In region 5, which is only present whenpreform temperature is significantly belowthat of eutectic, the solid metal is eutectic.When preform or die are above liquidus ofmetal, regions 1 and 5, or 2, respectively, areabsent. The two curves are schematic plots oftemperature and average matrix compositionin the composite along line A-A'. To istemperature of incoming metal, TE is eutectictemperature, Co is nominal alloy composition,CE is eutectic composition

Matrix solidification during infiltration also inter-feres with preform deformation because solid metaladds strength to the preforms. This may either preventpreform compression by locking the preform to thedie wall,104,336or prevent relaxation of compressedregions of the preform.336,339,343

Processing of metal matrix compositeslurriesWhen the reinforcement consists of isolated elementsof the reinforcing phase dispersed in the matrix, thecomposite forms a free flowing slurry if a sufficientportion of the matrix is liquid. Such composite "Slurriesare generally fabricated by stirring the reinforcementinto liquid or semisolid matrix and lend themselvesto ordinary casting processes once they are formed.The presence of a significant volume fraction of solidphase in the flowing composite modifies their behav-iour during these casting processes and creates an



added cancern because the reinfarcement distributianmust alsa be controlled.

Rheology of composite slurriesIt is well known that molten metals and alloys areNewtanian flui~s, with a viscasity in the range10- 3-10 - 2 Pa s that is independent .ofshear rate anddecreases with increasing temperature fallawing theArrhenius relation. When salid particles are dispersedin a liquid metal, twa types .ofinteractians can occur:an hydradynamic interactian between liquid and theparticle, and a non-hydradynamic interactian betweenthe particles themselves. Bath interactians praducean increase in the apparent viscasity .ofthe slurry, aneffect that has been canfirmed in several experimentalstudies.55,156,203,210,344-350These shaw that theapparent viscasities .ofvariaus campasite slurries aresignificantly higher than the unreinforced matrix allayviscasities, .often by .orders .ofmagnitude. The appar-ent viscasities .ofthe campasites furthermore decreasestrongly with increasing shear rate aver a wide rangeof shear rates, carrespanding ta nan- Newtanianpseudaplastic behaviaur. This last abservatian indi-cates that the suspended particles in metal matrixcompasite slurries interact with each ather. This givesrise ta particle agglomeratian and these aggregatesare apparently respansible for the .observed shear ratedependent behaviaur .of the slurries. Hawever, thenature .of the particle interaction itself is nat clearlyunderstaod. In the temperature range where thematrix is semisolid, the effect .of temperature ancampasite slurry viscasity is much the same as in thematrix allay melt: the viscasity decreases with increas-ing temperature predaminantly because .of thedecreasing valume fractian salid, as in the unre-inforeed partially salidified allay.55,210,305,345,349-351With campasite slurries, hawever, additianal vari-ations in apparent viscasity may result fram chemicalinteractian between the reinforcement and the matrixwhich alter the shape and valume fractian .of thereinfarcement. This effect has been shawn ta increasethe apparent viscasity .of AI-SiC particulate slurr-ies.352,353One tapic which is cantraversial in theliterature is that .ofthixatrapic behaviaur: Mada andAjersch349repart that an AI- 7Si allay reinfarced withSiC particles does nat exhibit any thixatrapic behav-iaur, whereas Maan and ca-warkers344,350 .observeda time dependent viscasity change in a similarcompasite system.The casting fluidities (in the present cantext, a

measure .of the mauld filling ability and nat theinverse .ofviscasity) .ofvariaus campasite slurries havebeen measured with the spiral test cammanly usedfar canventianal faundry allays.12,352,354,355Asexpected from viscasity data, spiral fluidity decreaseswith increasing particle valume fractian, and withdecreasing particle size far a given particle valumefractian.354,355The effect of temperature change ismore camplex. In conventional faundry allays, thecasting fluidity increases with increasing melt temper-ature, whereas the fluidity .ofAA 6061 reinfarced with15 val.-%SiC decreases with increasing temperature.This is due ta Al4C3 formatian, which increases slurryviscasity and encaurages reinfarcement particleagglameratian.352 In SiC reinfarced AI-7Si allay cam-


pasites, the fluidity increases with increasing temper-ature up to 750C, then decreases with increasingtemperature. This is because SiC particles are rela-tively stable in AI-7Si belaw 750C, while they reactat higher temperatures.With the additian of particles ta molten metal, the

viscosity increases and the casting fluidity decreases.Hawever, if the matrix allay campositian and thetype .ofparticles are tailared ta avoid interface chemi-cal reactians, compasite slurries can still be cast usingmost canventianal casting techniques in praductionfaundries.356-358 In same instances, improvements incasting quality aver unreinfarced alloys may evenresult fram the increased viscasity. An impartantexample .of this is the reduced parasity faund in diecast compasites,356,358an effect similar to that faundwith campacast or thixacast unreinfarced metals.359

Particle migrationWhen a camposite slurry is at rest, density differencesbetween matrix and reinfarcement induce settling orflaating .of the reinfarcement at a rate that dependsan lacal valume fractian and can be raughly predictedby theary.353 This results in variatians .of reinfarce-ment valume fractian within the campas-ite.12,61,352,356In .one study, campasite slurries wereheld in space abave the matrix liquidus.360 Compari-san .ofthe material processed in space with the samematerial processed similarly an the ground shawedmuch greater microscapic clustering of the reinfarce-ment in the graund pracessed material, in agreementwith other data 353which shaw that particle clusteringis present even in well stirred ground pracessedparticle reinforced aluminium. These data indicatethat same local redistributian .of the particles in theliquid matrix, an a scale much smaller than thecasting, can result from gravity induced particlematian .or fluid flaw.Often, there is matrix salidificatian cancamitant

with reinfarcement migration. If the composite slurryis stirred inta the temperature range where the matrixitself is partly salid (as in campacasting), little .or nagravity induced segregatian .of the reinfarcementoccurs even if the slurry is at rest.345,361-363This isbecause the salid matrix phase has abaut the samedensity as the liquid metal, sa it neither settles norfloats in the slurry and halds the reinfarcement inplace. Since allayed matrices salidify dendritically inmast casting pracesses, the reinfarcement particleswill nat migrate aver significant distances .once thelacal temperature falls belaw the liquidus, regardless.of whether it is pushed .or engulfed by the movingliquid/solid interface .ofthe matrix. This .observationhas been used ta madel particle segregatian withingravity cast364,365 .or centrifugally cast compasiteslurries,366,367ta obtain generally gaad agreementwith experiment. 364-371

Solidification of cast metal matrixcompositesThe influence exerted by the matrix micra structurean the an mechanical praperties .of MMCs, evenparallel ta an aligned fibre reinfarcement, has beenrepeatedly emphasised with experimental evidence


(examples can be found in Refs. 372-377). This, inturn, highlights the practical importance of solidifi-cation, which governs to a large extent' the finalcomposite microstructure. Direct transposition ofrules developed for microstructural control in thesolidification of unreinforced metals is not possiblewith MMCs, because the reinforcing phase frequentlymodifies solidification of the matrix. It is the inter-ference of the reinforcement with matrix solidificationthat is addressed in this section, for a composite thathas already been formed an4 is at rest. Nucleation ofthe matrix is dealt with, then growth of the solidphase, starting with the case where the reinforcementis fixed in space. Subsequently, effects that arise whenthe reinforcement can move in the liquid metal, inwhich case redistribution of the reinforcement mayoccur, are treated.It is assumed in what follows that the reinforcement

is chemically inert in the matrix, mainly because verylittle fundamental work has been done on the inter-action between matrix-reinforcement chemical reac-tions and matrix solidification. The' discussion isfocused on the case of an alloyed matrix, because ahighly pure metal matrix is not frequently used inpractice.

Nucleation of a reinforced metalThe solid reinforcing phase can reduce the grain sizeof the matrix significantly if it catalyses heterogeneousnucleation of the primary metal phase. This rarelyseems to occur with aluminium, since grain sizes farin excess of reinforcement diameter have frequentlybeen observed in aluminium reinforced with alu-mina 120,303,312,339,341,343,378-380carbon 381and sili-con ~arbide fibres378,382,383or particles.352,384Whenthe reinforcement does provide a propitious site fornucleation of the matrix, however, its effect on grainsize of the matrix can be quite strong. Thus, the grainsize of AI-45Cu is reduced by several orders ofmagnitude on going from SiC or Al203 fibrereinforcements to a reaction sintered porous TiCreinforcement (known to act as a heterogeneousnucleation catalyst for aluminium)385,386processedand solidified identically378 (TiC reinforced alumin-ium was also investigated by Baturinskayaet al.,387,388but though it is stated that the 'micro-structure is refined' in the conclusion of these articles,it is unclear whether any decrease in grain size wasobserved). A second example of matrix grain refine-ment due to nucleation catalysis by the reinforcementis that of hypereutectic aluminium-silicon alloys,wherein the primary phase, silicon, has been shownto nucleate preferentially on carbon, SiC, Si02, andAI20312,50,139,144,380,389-391As a consequence, thenumber of primary Si crystals per unit volume inthese alloys is increased in the composites comparedwith the unreinforced alloy.12,380Some grain refine-ment has also been found in Ti-515AI-I4Mnreinforced with less than 10 vol.- otic,TiB2 particles,which was attributed to nucleation of the primaryphase on a small fraction of the reinforcingparticles.392Braczynski393proposed that intermetallicalloy phases coat reinforcements to promote hetero-geneous nucleation of the primary phase in Cu-Pb-

Ti alloys with graphite par~icles, and AI-Cu- Ti alloyswith alumina particles. A quantitative interpretationof differential scanning calorimetry data was pre-sented to show that nucleation takes place faster inthe composite, but no reason for why the intermetallicwould precipitate before tlie primary phase was given.Grain refinement of the matrix may also result

from exchange of heat between reinforcement andmatrix during infiltration. As seen in the section'Thermal and solidification effects' above, infiltrationof a preform initially at a temperature below thematrix liquidus results in rapid solidification of aportion of the matrix during infiltration. Unless thissolid metal remelts, it will produce a fine equiaxedgrain structure in the matrix.119,120,339,342,343When the reinforcement does not induce nucleation

of the primary phase of the matrix by catalysis orheat transfer, the grain size in composite castings islikely to be somewh~t greater than that of an identicalcasting of theunreinforced matrix. This is becausethe reinforcement impedes convection of the liquidmetal, whether the reinforcement is stationary or isdiscretely distributed in the matrix. Many mechanismsresponsible for formation of fine grained, equiaxeddendritic .structures in castings depend on fluidflow.394For this reason, if there is significant convec-tion during solidification of a similar unreinforcedcasting and matrix nucleation is sluggish, an increasedpropensity for columnar dendritic solidification isexpected in composites. This effect is illustrated inthe experiments. of Cole and Bolling,395 who eluci-dated the' effect of fluid flow in ingot solidification byinserting a grid of metal wires into a mould beforecasting. The resulting castings were, in fact, lowvolume fraction wire reinforced metals, and showed agreater columnar zone than ones devoid of the wiremesh.

Growth of soiid metal with stationaryrei nforcementAs seen in the section 'Thermal and solidificationeffects' above, the time for equalisation of temperaturebetween matrix and reinforcement is at most around1ms. This is generally much less than the time formatrix solidification in casting processes used for theproduction of metal matrix composite ingots or parts.For a solidifying alloyed matrix in a cast composite,therefore, the reinforcing phase is not a heat sourceor sink of much significance, but is primarily animpermeable barrier to mass transfer.When the reinforcement is stationary and the

matrix isan alloy, it is therefore equivalent to a veryfine and narrow crucible, within which the matrixmust solidify. As in solidification of unreinforcedmetals, the simplest configuration for fundamentalstudy of solidification of composites is that where theprocess takes place at steady state, allowing simpledefinition and control of growth parameters:. temper-ature gradient G at growth rate V. Steady statesolidification requires that the reinforcing phase delin-eate tiny straight-walled crucibles, which is achievedin practice when continuous parallel fibre reinforcedmetals are solidified at steady state with the temper-ature gradient G parallel to the fibres.



7 Definition of contact angle (J between solidmetal, liquid metal, and reinforcing phase

At larger interfibre spacings, such that Ac is smallerthan the interfibre spacing by at least an order ofmagnitude, the plane-front to cellular transition isexpected to take place as in the unreinforced metal.Near the fibre, the liquid/solid interface should thenbehave essentially as it does near a grain boundaryof the unreinforced solidifying metal, Le. become aregion where instability amplitude increases fasterthan along the plane front at the onset of instability.Experimental work on directionally solidified succin-onitrile-acetone alloys illustrates the effect of 8 onmorphology of a plane front near the fibres, showinginterface curvature near the channel walls.399,400,402At values of the ratio G/V slightly below that for

breakdown of plane-front stability, the unreinforcedmatrix solidifies with a cellular morphology. Spatiallyconstrained steady state cellular growth, similar togrowth between parallel fibres along their axis, hasbeen treated theoretically by McCartney andHunt404-406for a deep single cell centred in a cylindri-cal interstice, as well as Ungar and Brown407 andTrivedi et al.402 for a two dimensional cell centredin a planar channel. McCartney and Hunt found thatnarrowing the interstice below the experimentallymeasured diameter of free growing cells increases thecell tip undercooling somewhat, although not dra-matically, for the AI- Mg-Si alloy under considerationin their work. Tip undercooling was found to decreaseby a few kelvin (or about 50% of its initial value) asthe spacing was reduced to half the spacing chosenby the cells in the unreinforced metal, while tip radiusdecreased and tip composition increased at smallinterstice radii. Overall, the cell tip undercooling wasprimarily determined by solute diffusion along thecell, as first proposed by Bower et al.,408 while theGibbs- Thomson effect increased the undercooling atsmall interstice radii by an amount that was somewhathigher than expected from proportionality of tipcurvature with interstice radius. For more drasticreductions in interstice diameter, to below 10% ofthe experimental cell spacings measured in unre-inforced metals, extrapolation of the curves points torather high undercoolings. This is in agreement withcalculations by Trivedi et al.,402 who found that thetip undercooling of a two dimensional cell is drasti-cally increased as the width of the channel is decreasedby an order of magnitude.Experimental investigations by Sekhar and Tri-

vedi399using succinonitrile indicate that steady state

Steady state solidification of metal matrixcompositesThe morphological stability of plane front alloy solidi-fication is analysed theoretically by calculating therate of growth of infinitesimal .sinusoidal pertur-bations of the plane front. 396,397In unreinforcedalloys, the front is assumed to be essentially infinitein extent, so if any perturbation wavelength shows apositive rate of growth, the front is unstable. In anarrow interstice delineated by closely spaced fibresof a composite, the spectrum of possible perturbationwavelengths has an upper limit equal to the width ofthe interstice. Calculated critical perturbation wave-lengths Ac below which an infinite plane front is stable,and above which perturbations grow, were given inRef. 398 to show that with both AI-Cu alloys andsuccinonitrile-I3 wt-O/oacetone, the wavelength formarginal stability at the breakdown of a plane frontis of the order of 100~m, equal to or larger than theinterstices left between fibres customarily used forreinforcing metals. One therefore expects, as was firstpointed out by Sekhar and Trivedi399 that fibres tendto stabilise plane front solidification when their sep-aration falls below Ac While Shangguan and Hunt400found no effect of spatial constraint on plane-frontbreakdown in succinonitrile-acetone, recent experi-ments on AI-Cu matrix composites at MIT are ingeneral agreement with this analysis.401One instance in which a reinforcement may lower

the stability of a plane front was also pointed out byTrivedi et al.,402 namely that where the interstice isa narrow channel between two wide flat plates. Inthis case, there is no upper limit to the possibledestabilising perturbation wavelengths, and becauseof curvature in the solidification front perpendicularto the reinforcement plane, the stability of a planefront is reduced by the reinforcement. However, thereare very few instances in practice where this situationwill occur in MMCs.More detailed analyses of plane-front growth take

into account the finite contact angle 8 between thesolid metal, the liquid metal, and the reinforcingphase, Fig. 7, at the jllnction line between the fibrematrix interface and the liquid/solid interface (ignor-ing roughness effects, 8 is the same angle as that usedto evaluate potency of the fibre for heterogeneousnucleation of the matrix). Unless 8=90, a strictlyplanar front can therefore not be obtained becausethe liquid/solid interface bends near the reinforcement.Ungar and Brown403have studied this problem theor-etically in two dimensions, i.e. for solidificationbetween two parallel plates, using finite elementmethods. These authors considered contact angles (}greater than 90, and found that at high e and smallinterstice widths, there is a continuous series of solu-tions for the solidification front morphology goingfrom shallow to deep cellular solidification fronts asthe temperature gradient decreases at fixed growthvelocity V past conditions for plane-front breakdownin the unreinforced metal. At smaller 8, still greaterthan 90, a pseudo-plane front is observed at high G,which breaks down, with a sudden increase in theheight of stable cells, at decreasing G. This pseudo-plane front to cellular transition takes place at higherG than that for plane-front breakdown with (}= 90.


reinforcingphase

liquid metal

solid metal


a quenched liquid metal, resulting in fine dendrites which grewunperturbed by fibres; b from sample solidified at steady state withgradient G=9100 K m-1 and growth rate R=203 ~m s-1, in squareinterstice, contorted dendritic primary dendrite arm is shown bycoring patterns; in triangular interstice, the dendritic nature of matrixis erased, and coring patterns are parallel to fibres; c G = 4500 Km-1 and R= 54 ~m s-1, longer solidification time results in non-dendritic structure in both square and triangular interstices, intriangular interstice, microsegregation is reduced, as seen by missinglow Cu coring patterns; d G= 3500 K m-1 and R= 25 ~m S-l. Atstill longer solidification times, microstructure is featureless intriangular interstices, with no coring and no second phase.

S AI-45Cu-SiC fibre composite solidified inBridgman furnace along fibre axes. Fibrediameter is 140 J.1m.From work reported inRef. 383

beyond the point where the average dendrite armspacing is of the order of the interstice radius. Ripen-ing of dendrite arms then ceases, and further coarsen-ing of the dendrite takes place entirely by dendritearm coalescence, which gradually erases the dendriticcharacter of the solid metal. If the average time tc forcoalescence of dendrite arms is of the order of, orlarger than, the total time for solidification tr, thenthe microstructure within the interstice remains den- ,dritic, with coring patterns indicating the presence ofsecondary or higher order arms, Fig. 8a and b. If, onthe other hand, tc is significantly smaller than tr,coalescence joins the dendrite arms together, and doesso early in the solidification process. The resultingfinal microstructure is then cellular in appearance,with concentric coring patterns running parallel tothe fibre surfaces, Fig. 8c and d.These changes in dendrite coarsening mechanisms

affect micro segregation in the matrix. Because themicrostructure can only coarsen up to the point wherethe scale of the matrix microstructure equals that of


cellular growth in comp


the interstices, fibres packed to a high volume fractionplace an upper limit on diffusion distances within thesolid during solidification. Therefore, solid statediffusion can reduce micro segregation to a greaterextent than in conventional unreinforced castings ofthe same alloy.383This effect is illustrated in Fig. 8dwhere at the lowest cooling rate, a fully homogenisedAI-45Cu matrix was directly solidified in thecomposite.

Unsteady solidificationEven with directional solidification at constant G andV, steady state will not be achieved if the reinforce-merit is not parallel to -the growth direction. Trivediand co-workers have studied changes in directionalsolidification morphology of succjnoriitrile basedalloys with various reinforcement geometries, includ-ing single fibres, isolated particles, two converging ordiverging fibres, particles, and a constriction followedby an enlarging channel. 399,410-412These variousgeometries caused a variety of solidification morpho-logy changes to occur, with concomitant variationsin local solidification velocity and resulting localcomposition of the solid. Observations were alleXp'lained~by analysing the influence exerted by thereinforcement on diffusion of solute in the liquid, towhich it constitutes a barrier. Provided there is noenhanced nucleation of the matrix on the reinforce-ment, the matrix alloy grows from within the inter-stices left between the reinforcement, and avoids thelatter as it grows because the reinforcement consti-tutes a barrier to solute evacuation. For this reason(and not for thermal reasons as frequently proposed),the reinforcement is quite generally found to besurrounded with the last phase(s) to solidify in thematrix alloy.12,46,48,50,367,372-374,378,381,391,411,413-419With sucinonitrile containing 0'5% impurities,412 re-peated interaction between the solute diffusive fieldand the particles caused oscillatory motion of thesolidification front, resulting in cellular growth underconditions where dendrites form if there are noparticles. Depending on particle size and volumefraction, dendrite tip splitting, particle trapping, andparticle rejection between primary dendrite arms wereobserved when dendrites encountered particles.Most of the features found with steady state experi-

ments also apply with unsteady solidification andirregular reinforcements. Several studies have beenconducted on infiltrated and remelted fibre reinforcedbinary aluminium alloys, in which samples were so-lidified in shallow temperature gradients, to induceequiaxed dendrite growth in the matrix.413,414,420,421The microstructural evolution of the matrix in thesecomposites was in agreement with observations andtheory from steady state columnar dendrite solidifi-cation studies:383at long solidification times, coarsen-ing leads to a matrix microstructure where thedendritic features are lost, and micro segregation isreduced.Changes in coarsening behavior during an iso-

thermal hold in the semisolid temperature range ofAI-7Si-0'3Mg with addition of SiC particles werestudied by Bayoumi and Suery.363It was found thatin the composite, rounding off of dendrites to form'globules' of primary aluminium is accelerated com-


pared with the unreinforced alloy, and that as theparticle volume fraction increases, the final globulediameter decreases. These observations result fromthe same coarsening phenomena founel in steady statesolidification experiments:383 as the particles areadded, the matrix cannot ripen beyond the size ofinterstices between the particles. Therefore, the finalglobule size decreases as the particle volume fractionincreases at constant particle size and, becauseincreasing the. particle volume fraction decreases thediffusion distances for coalescence, coalescence ofdendrite arms is accelerated, resulting in earlier trans-formation of dendrites of the primary phase to spheri-cal globules. Obseryations by Bryant et al.392 of areduced grain size in solidified TiB2 particle reinforcedtitanium aluminide, which were explained byenhanced nucleation, may alternatively also beexplained by similar coarsening effects.

Growth of solid metal with mobilereinforcement: particle pushingWhen a moving solid/liquid interface approaches amobile foreign particle suspended in liquid metal, theparticle can be either captured or pushed away bythe interface. If foreign particles are captured by thegrowing solid metal, little redistribution of thereinforcement will occur during solidification, andhence the particle distribution in the solidified mater-ial will be as uniform as it was in the liquid state. Onthe other hand, if particles are pushed by the solidifi-cation front, they will be redistributed, to be finallysegregated in the last pools of liquid matrix to solidify.Figures 9 and 10 show typical examples of solidifi-cation microstructures resulting from particle pushingand particle capture, respectively, in cast AI-Si alloysreinforced with SiC particles.The interaction of particles with a solidification

front has been extensively studied by various workers.Work up to - 1986 has been reviewed by Rohatgiet al.,12 and Russell et al.,52 and more recent reviewshave been presented by Stefanescu and Dhindaw,422and Rohatgi et al.139 This review therefore providesonly a brief outline of the earlier work, to analyserecent theoretical and experimental studies with afocus on those relevant to metal matrix compositesolidification.

Experimental studiesThe first systematic work on particle pushing is dueto Uhlmann et al.,423 who mixed a number of differ-ent particles into various transparent organic matrixmaterials, and carried out horizontal directionalgrowth experiments under an optical microscope.They found that, in some particle-matrix systems,the particles were not pushed at all, whereas in othersystems a critical velocity ~ existed below which theparticles were pushed and above which they werecaptured by the interface. This critical velocity wasdependent on the type of particle for a given matrixmaterial, and on the kind of matrix material for agiven particle. The effect of particle size on ~ wasnot quite straightforward: when the particle size waslarge (hundreds of micrometres in diameter), ~decreased with increasing particle size, while for small

Mortensen and Jin Solidification processing of metal ma~rix composites 117

,:,{i,' ..........'i:'iJ"",,,];,12;,f.~:flt~. };".";~I. 50 IJ.m

9 Microstructure of AI-7Si-15 vol.-%SiC com-posite solidified at cooling rate of 4 K s-'.Note that SiC particles are pushed tointerdendritic regions by AI dendrites

particles (below 15 Jlm), it was virtually independentof the particle size.A series of experimental investigations followed this

study. Cisse and Bolling424,425carried out similardirectional solidification experiments with water andsalol containing various insoluble particles, with thepurpose of examining the effect of liquid viscosity andparticle size on ~. It was found that regardless of theparticle type, ~ decreased when the particle sizeincreased, and when the liquid viscosity increased.Zubko et al.426 examined experimentally the effect ofthe thermal conductivities of the particle and matrixon particle capture, following theoretical work byChernov and Mel'nikova,427 who calculated the tem-perature field and interface shape around a particleat the solidification front. The materials used wereanthracite, glass, AI, and Zn particles in naphthalene,and Fe, Ni, and Cr particles in Zn, Bi, and Sn. It wasfound that when the thermal conductivity of theparticle was higher than that of the liquid, the particleswere captured, and that otherwise the particles werepushed.Neumann et al.428 and Omenyi and Neumann429

investigated particle pushing from a thermodynamicstandpoint using' horizontal directional solidificationexperiments with materials of known interfacial ener-gies. They found that when the net free energy changewhich occurred during the particle transfer fromliquid to solid was negative, the particles were cap-tured even at extremely low solidification rates. Con-versely, when this free energy change was positive,the particles were pushed at low velocities and cap-tured at high velocities. The critical velocities meas-ured were dependent on the particle size as well ason the value of this free energy change.The particle pushing phenomenon has. also been

observed in certain biological systems. The interactionof biological cells with growing ice crystals was stud-ied by Bronstein et al.,430 using cell suspensions ofbrewer's yeast and human red blood cells in aqueoussalt solutions. They found that the behaviour of thebiological cells at the solidification front was the sameas that of other organic or inorganic particles, i.e. the

"~

\X~~~4 ; 1'lt\(lFf~/~lI!!'..y.' " / ".' .~. " ',..

. '. ,/ ,~.," :..t'.~~~ f(~ ...~..... '. ~'J' \,.~\"M / ftI "' ..'~"."'~, . '.:?'~~I _--------t~.:.f,... .. ~ ,.,.~ \ ~10 Microstructure of AI-16Si-15vol.-%SiC

composite (cooling rate was 7 K s-'). Notethat SiC particles are not pushed by primarySi crystals

cells were pushed below and captured above somecritical velocity ~. In these experiments, ~ wasdependent on the solute concentration of the solution,decreasing with increasing solute content.The influence of the temperature gradient G on the

critical velocity was investigated by Korber andRau431 using latex sphere suspensions in distilledwater and in an aqueous solution of NaMn04' Thecritical velocity increased with increasing G.The effect of fluid convection in the melt on particle

pushing has been studied in metallic,432as well as innon-metallic,433 systems. Schvezov and Weinberg432carried out zone refining experiments with pure Pband Pb-Sb alloys which contained Fe particles. Whenthe interface was planar (pure Pb matrix) the particleswere uniformly distributed in the solid, indicatingthat the particles were not pushed by the planarinterface. When the interface was cellular (Pb-Sballoy matrix), the particles were segregated to the cellboundaries. In the experiment of Delamore et al.,433the behaviour of methyl methacrylate and carbonparticles in various organic materials was examinedunder natural and forced convection. With a planarinterface, the particles were pushed, and with a non-planar interface. the particles were deflected near thecell tips and segregated to the intercellular regions.These studies indicate that, when convection is presentin the bulk melt and if the solid/liquid interface isnon-planar, convective flow can sweep the particlesaway from cell or dendrite tips, resulting in intercellu-lar or interdendritic segregation.Recent work has focused on particle pushing in the

solidification of various MMCs. The behaviour ofSiC particles at the solid/liquid interface in variousaluminium alloys has been extensively studied usingunidirectional solidification, multidirectional solidifi-cation, and rapid solidification. 352,380,417,434-439Theparameters investigated were alloy compositio~,particle size, particle volume fraction, growth velocity,tem

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