The Introduction of Fine SiC Particles Into a Molten Al Alloy Matrix: Application to Composite Material Casting
Vnašanje finih SiC delcev v talino iz Al zlitine: uporaba pri litju kompozitnih materialov
Kevorkijan V.1, zasebni raziskovalec A kategorije B. Suštaršič, IMT, Ljubljana
The immersion of fine SiC particulates into a Al aiioys is one of the major factors hampering the commercial development of particie reinforced composites prepared by liguid metallurgy techniques. Smaller particles are generally more difficult to wet and disperse than larger particles because of their inherently greater surface area On the other hand, smaller particles generally give MMCs with superior mechanical properties I he challenge is to control the reaction between the matrix and reinforcement particles to achieve improved vvetting so that good distribution and interfacial bonding are obtamed. This paper will revievv the actual technological challenge of fine particulate reinforced MMC fabrication, and will describe the most promising new processes.
Key words: vvetting, interfacial reactions, model of vvetting, activation complex theory
Uvajanje finih SiC delcev v talino iz Al zlitine je tehnološko zahtevno zaradi njihove slabe omočljivosti m izrazite kemijske reaktivnosti. Hkrati je uvajanje finih SiC delcev predpogoj za izboljšanje mehanskih lastnosti kompozitov in njihovo večjo konkurenčnost na trgu. V delu so opisani dosedanji rezultati na tem področju in predstavljene nove smeri razvoja.
Ključne besede: omakanje, reakcije na mejah, model omakanja, teorija aktivnega kompleksa
Introduction
The interest in discontinuous reinforcement of alu-minium alloy relates primarily to producing improve-ments in strength and stiffness combined with re-duction in density. Improved wear resistance also has an important bearing on potential applications.
The ability to produce vvhiskers, particles and platelets at lovver costs has stimulated the search for low cost routes permitting the discontinuous reinforcement of aluminium. As standard shaping meth-ods such as forging, rolling and extrusion can be em-ployed on discontinuous reinforced Al MMCs, their production and fabrication is more readily integrated into existing metal processing facilities than of con-tinuously reinforced MMCs.
Reinforcement with SiC has undoubtedly received the greatest attention. Its attractions include: relativen low cost and ready availability, high modulus and strength and density only slightly higherthan Al.
Developments in MMC have to date been led by the needs of the aerospace and defance industries, but there is increasing interest in their application in the automotive and other high-volume manufactur-ing sectors. In the leisure sector (for example, in sports goods) there is also great activity. There is a
' Dr. Varužan KEVORKIJAN Borova vas 4 62000 Maribor
perception that, in the near term at least, automotive applications will drive the vvide-scale exploitation of MMC because of the high-volume requirements of this sector. Particular applications are envisaged in piston crovvns, gudgeon pins and connecting rods where the higher-temperature strength allovvs im-provements in engine design and greater efficiency. For example, several million MMC piston crovvns are currenty being manufactured each year for use in Toyota diesel engines. Other clearly identified applications for MMC are limited, but recent announce-ments have included forgings for helicopter parts, instrument racks for aircraft, bicycle grames and car-brake components.
Hovvever, MMCs generally have not yet achieved vvidespead industrial application. This situation is at-tributable to a number of factors including high cost, lack of commercially acceptable means of production, quality assurance of MMC products and the mechanical properties not reaching the theoretically predicted values. The latter point is paticularly rele-vant to MMCs produced by foundry processes. This class concerns ali the techniques that elaborate MMCs by stirring an alloy above the liquidus, or be-tvveen the solidus and the liquidus, vvhile proceeding to the addition of the reinforcement particulates. These techniques correspond to the most inexpen-sive processes for producing MMCs. It is likely that in the next decade most of the MMC volume produced in the vvorld will be done using one of them.
A common problem with MMCs is excessive reac-tion between the matrix and reinforcement. This re-action often results in inferior properties in the com-posite compared with the matrix alloy. Some interfacial reaction is required, however, to achieve vvetting of the particles by the melt which, in turn, is necessary to disperse the particles uniformly in the melt. Thus the challenge is to control the interfacial reactions to obtain MMCs with a uniform dispersion of reinforement particles and acceptable interface strength, vvhile keeping the contamination in gases and inclusions at the lovvest level. These require-ments, which may appear to be easily achievable, are in reality very demaning. The problems in achiev-ing ali these requirements is that surface energies of most liquid metals are very important, and their re-sulting high surface tension penalize the vvetting of intruding ceramic particulates. This results in partic-ulate rejection by the melt, due to non-wetting condi-tions. Note that smaller particles are generally more difficult to wet and disperse than larger particles be-cause of their inherently greater surface area. Moreover, once vvetted, smaller particles have a greater propensity to react vvith molten Al. This latter effect is attributed to the larger surface area available vvith the finer SiC particles and consequently a high-er level of surface reaction.
Several authors12 have confirmed that there is no difference in strength betvveen the PM and melted discontinuous MMCs if the comparison is made betvveen composites that are comparable; in other vvords if the bonding betvveen the reinforcement and the matrix is the same, and vvith the same size, dis-tribution and volume fraction of the reinforcement. In order to meet these requirements, it is necessary to find a way of introducing a large amount (20-30 wt%) of fine SiC particles (less than 10 pm) into an Al alloy melt vvithout the rejection of reinforcement and with-out excessive oxidation or other unvvanted chemical reactions. Note that typical results achieved by vor-tex method are, for example, 2 wt% 53 um silicon carbide in Al-Si alloy3. Vortex method is probably the simplest and most instinctive method that can be used for the fabrication of MMCs. The method con-sists of vigorously stirring an alloy maintained over its liquidus, vvhile adding reinforcement in the vortex4 7. More promising results are obtained by compocast-ing - a process consisting of vigorously stirring a se-mi-solid alloy, vvhile adding the particulates to the surface8"11. Composites containing 29 wt% 14 |.im SiC in Al-Si alloy have been successfully fabricated using this technique12. Hovvever, before commercial application can be developed, problems of exces-sive porosities1317, high level of inclusions1819 and tight temeprature control for larger batch size, must be solved.
It can be noticed that a large amount of fine particulates are difficult to incorporate vvith these tech-niques, vvithout at least partial rejection of particulates and chemical reaction (dissolution) of SiC vvith molten Al. Finer than 10 |am particulates also seems to be a lovver limit in particulate dimension under vvhich complete rejection and dissolution become
the rule. Hence, the possibility of using such a processes for the commercial fabrication of structur-al MMCs containing much finer particulates, at high-er volumic fraction, remains questionnable. In order to achieve this end, the new fabrication processes must be developed.
This paper will revievv the challenge of mass pro-ductiing fine particulate reinforced metal matrix composites and will focus on the most promising new processes and composite designes.
Theoretical background
VVetting and bonding of SiC particulates to Al and Al alloys
VVetting and bonding are of practical importance in discontinuously reinforced composites technology. first in determing the effectivness vvith vvhich SiC particulates and matrix can be combined in foundry methods of composite preparation and second in de-termining the particulates/matrix bond strength. Improvements in vvetting betvveen the matrix and the reinforcing phase in MMCs are an important goal of surface engineering.
The vvetting properties of ceramics by liquid metals are governed by a number of variables, including heat of formation, stoichiometrx, valence electron concentration in the ceramic phase, interfacial chemical reactions, temperature, and contact tirne15. The vvork of adhesion betvveen a ceramic and a melt de-creases vvith increasing heat of formation of car-bides. The high energy of formation of a stable carbide implies strong interatomic bonds and correspondingly weak vvith metals. This leads to a high interfacial energy and a small vvork of immer-sion, resulting in poor vvetting.
High valence electron concentration generally implies lovver stability of carbides and improved vvetta-bility betvveen ceramics and metals. High temperature and long contact times promote melt-ceramic wettability due to reactions at the melt/ceramic interface, resulting in reduced contact angle.
The most common method of assessing vvettabili-ty in tvvo-phase systems is the sessile drop experi-ment, in vvhcih a drop of the lovver meltig-point con-stituent lies on a plane substrate of the other constituent. The angle of contact, ©, of the drop is often taken as a mesure of wettability. For ideal vvetting, a liquid drop must spread completely over the solid surface, that is the contact angle becomes ze-ro, although a contact angle of less than 90° is ade-quate to cause effective vvetting. Similarly, complete non-wettability is defined by a contact angle of 180°; hovvever angles greater than 90° are practically suf-ficient to cause devvetting.
The different interfacial forces acting betvveen sol-id-liquid, solid-vapour and liquid-vapour-phases are represented by vectors 7sh ysy and -y,v respectively. At equilibrium conditions Thomas Young16 equated the vectors in the horizontal direction to give the follovv-ing fundamental relation:
7sv = 7si + 7iv ' COSM	(1)
Since surface tension of solids in vapour(7sv) can-not be measured easily, Dupre17 suggested the work of adhesion of a solid to a liquid as given by:
Ws, = 7sv + 7iv - 7si	(2)
where, Wsl is defined as the energy required to separate an solid-liquid interface into free surfaces of solid and liquid. Combinig Equations (1) and (2)
gives:
Wsl = 7lv(1 + cosH)	(3)
Thus Wsl can be determined vvithout direct knovvl-edge of the interfacial energy. Note that it is a com-bination of (■) and 7,v rather than (-) alone that gives a measure of wetting (bonding). Thus, a reduced val-ue of 0 does not necessarily indicate improved bonding since it may occur in association with a low-ering of 7|V.
The relative strength of bonding can be assessed by comparing Wsl with 27lv and 27sv, these being a rough reflection of the atomic bond strengths in liq-uid and solid respectively. A high Wsl value is a pre-requisite but not a garantie for good mechanical stength in the bond. In general, the mechanical strengths will depend on the vveakest path in the interface, vvhich may exist in one of the pathes close to the interface or in a layer of reaction products rather than in the interface as such18.
A revievv of the vvetting of SiC by liquid metals has been published by several authors18 23. As a rule, the wettability of covalent carbides such as SiC and B4C follovvs the same dependance on the nature of metal as the wettability of carbon. The magnitude of the chemical interaction is vveaker than for carbon, due to the larger strength of the bonds in these carbides. Wetting was usually not observed below about 900°C and the contact angle decreased (in several steps) when the temperature increased. It was shovved24 that the Al/SiC system exhibits a non-vvet-ting behaviour up to 1223 K, vvhere a sharp transition to vvetting of SiC by liquid aluminium occurs. This kind of transition is typically found in aluminium/non-metallic refractories systems (carbon, Al203, TiB2, etc.). It has been attributed to the disappearance of the effect of Al203 layer on liquid aluminium at about 1223 K under high vacuum2425 or/and Si02 layer on the surface of SiC particles18 2425. Note that very of-ten, the production of Al/SiC composites involves a double AI&AI203/Si02&SiC interface, since the Silicon carbide reinforcement employed is oxidized nat-urally or on purpose. Penetration of aluminium trough this diffusion barrier is facilitated either by in-creasing the temperature, or by the addition of alloy-ing elements26. At lovver temperatures, this layer pre-vents a true metal/substrate interface from developing20'24'27. Being more electropositive than aluminium, these elements probably substitute for aluminium in the oxide layer, bringing about a vveak-ening of the film27. These elements thus enhance vvetting even though they form vveaker bonds with SiC than aluminium does. The same effect might al-
so explain the enhancement of vvetting brought about by adding magnesium in aluminium23.
Therefore, keeping in view of the theories of surface energy and chemical theory of vvetting28 follovv-ing techniques were developed to improve vvetting of non-metallic solids by liquid metals: use of metal coatings4'2930, addition of elements in the liquid metal31"36, heat-treatment of ceramic particles37 and use of ultrasonic38 and other mixing techniques39"41.
The most widely used technique to promote vvetta-bility is to add a vvetting agent, usually in the form of an alloying element in the melt. A mixture of an al-loying element in a liquid metal can promote the lat-ter's wettability with a solid surface by three mecha-nisms: reducing the surface tension of the liquid, decreasing the solid-liquid interfacial energy and chemical reactions at the solid-liquid interface. An early patent32 describes additions of alloying elements (vvetting agents) such as lithium, magnesium, silicon and calcium to the liquid metals to promote wettability of non-metallic particles vvith the melts. It vvas observed that additions of small pieces of magnesium to the surface of the melts along vvith disper-soids vvere more effective in dispersing these particles than the čase vvhen magnesium was already present in the melt. This is probably because magnesium added to liquid aluminium initially melts and spreads on the melt surface, thereby vvetting the dis-persoids42.
The enhanced vvetting of ceramic materials result-ing from the addition of magnesium to an aluminium alloy is well documented13 42 45. Several mechanisms are generally discussed vvhen the role of magnesium is considered. Authors23 claim that the addition of an alloying element can enhance the vvetting of a solid surfacein three ways, namely (i) by reducing the surface tension of the alloy, (ii) by decreasing the solid-liquid interfacial energy, and (iii) by promoting a chemical reaction at the solid-liquid interface. Other investigators13 state that magnesium is effective in reducing the surface tension of the melt and induc-ing interfacial reactions. Finaly, authors43 who found magnesium to significantly improve wettability, be-lieve that the promotion of interfacial reactions is the most active mechanism for enhancing the vvetting of a solid ceramic surface vvith a molten aluminium al-loy.
According to19, it vvas found that he vvetting process of SiC by aluminium alloy initially had an incubation period that vvas decreased by the addition of alloying elements. Hence, authors19 proposed that the vvetting process could be explained by assuming that the vvetting rate is propotional to the number of vvetting nuclei and to the ratio of unvvetted area. Thay also found that the vvetting rate seems to be rate determined by the dissociation of SiC because the same value of activation energy for the vvetting process vvas obtained from the four systems: pure aluminium, Al-Ti, Al-V, and Al-Zr. Generally, the value of the reaction rate constant increased and that of an incubation period decreased vvith the addition of alloying elements giving high carbon solubility19.
Two other factors that are of importance in the vvet-ting experiment are temperature and atmosphere since these can have an important influence on the various interfaces. Atmospheric impurities can markedly depress the liquid surface energy, -ylv, by adsorbtion. However, in Al-SiC system this effect is small18. Probably more significant effects are: the influence of temperature and atmosphere on the breakdovvn of the passive oxide film on the free SiC surface and the effect of temperature on the breakdovvn of this film at SiC/AI interface. Concequently, several patented methods of fluxing and degassing composite melts has been developed vvhich uses a rotating impeller-like device to both stir the bath and inject a blend of purge gases46"48. Purge gases can be either inert (argon or nitrogen) or reactive (chlo-rine, fluorine and Freon 12). The standard methods for degassing aluminium are also very effective49 and are widely used in foundry processing of MMCs. Fluxes can also be used to minimize oxide formation (and to remove suspended nonmetallics). For this reason, melts containing magnesium are often pro-tected by the use of salts that form liquid layers, most often of magnesium chloride, on the melt surface. These fluxes, termed covering fluxes, must be peri-odically removed and replaced. Carbon, graphite and boron povvder also effectively retard oxidation vvhen applied to the melt surface.
Dissolution of SiC particulates in Al alloy matrix
A common problem with metal matrix composites is excessive reaction betvveen the matrix and rein-forcement. It is well knovvn50'51 that SiC is thermody-namically unstable in molten Al and reacts to form AI4C3 and Si according to the reaction:
4AI + 3SiC = AI4C3 + 3Si	(4)
The AI4C3 forms at the interface vvhile the Si dis-solves in the Al matrix. The formation of AI4C3 gen-erally leads to degradation of the mechanical prop-erties of the reinforcement and the composite52. The SiC-AI reaction is irreversible and once AI4C3 formation is initiated at higher temperatures, lovvering the melt temperature will only increase the viscosity and decrease the fluidity of the melt due to the bridging effect of the AI4C3 and silicon particle formation on the SiC particles53. Hence this reaction should be avoided. Note that only the control of the kinetics of the reaction trough the use of thermodynamics or protective coatings vvhich reduce reaction rate can prevent the degradation of matrix and composite properties.
From equation (4) it is seen that, in addition to AI4C3, Si is produced, and if the Si content is suffi-ciently high the reaction will tend to go to the left, ie. the SiC vvill be stable in the matrix and not react. The required Si levels at different temperatures can be calculated54'55, and at typical casting temperatures 10 to 12 wt% Si is required for complete thermody-namic stability. Hovvever, in practice it is the kinetics of aluminium carbide formation vvhich has to be cosidered. In practice melt holding times can be
many hours, depending on the foundry, and for very long times the equilibrium Si levels vvill be needed. Hovvever, for holding times of 2 to 4 hours, vvhich is viable for most casting operations, much lovver Si levels are sufficient to prevent extensive AI4C3 formation. For example, vvith a common hypoeutectic Al-Si-Mg casting alloy, vvhich contains 6-7 wt% Si, 4 hours at 1023 K produces barely detectable AI4C3, and even at 1073 K the rate of AI4C3 formation is quite slovv. Note that this is not valid at higher processing temperatures. At 1400K, for example, the re-quired Si level varied betvveen 15 to 25 wt%56 vvhich is completely unconvenient for practice.
So the Si level required to limit AI4C3 formation vvill depend on the reactivity of SiC phase and the par-ticular foundry practice. The finer SiC particles, the higher the melt temperatures and the longer the holding times involved, the higher the required Si level slould be applied. A good practice is to min-imise superheat and holding times as much as pos-sible - vvhich is normal, good foundry practice.
The reaction betvveen SiC and Al has been studied in detail by several authors55 61. They observed that the reaction proceeds in tvvo steps. In a first stage. the silicon carbide dissolves in the liquid metal:
SiC(s) => (Si), + (C),	(5)
Secondly, the AI4C3 compound precipitates grovv-ing in random islands on the SiC surface. More de-tailed, the interfacial reactions in the Al/SiC composite vvith molten metal manufacturings are likely to include: (a) the chemical reaction (dissolution) of SiC vvith molten Al, (b) the diffusion of silicon and carbon atoms away from the SiC surface into the bulk molten Al pool, (c) the formation of AI4C3, (d) upon cooling, further precipitation of AI4C3 and (e) the so-lidification of the matrix. It seems that the rate con-troling step in the overall interfacial reaction is the chemical dissolution of SiC in molten aluminium62. The rate of the SiC dissolution may be expressed as Equation (6):
[Si] = 0.8 + 2.06t at 1253 K	(6)
vvhere [Si] is the average silicon in wt% in aluminium matrix and t is the composite manufacturing tirne in hours. Applying that kinetic model developed by62, the rate of molten aluminium attack on SiC was de-termined to be:
In k (jam/min) = 6.36 - 7180/T	(7)
vvhere k is the rate of SiC dissolution in molten aluminium and T is the processing temperature for the composite.
Several modifications of the surface chemistry of SiC particulates have been also examined in order to overcome the problems of chemical reaction betvveen the dispersoids and the matrix. The surface of the dispersoids is coated by a refractory materials vvhich is non-reactive both vvith dispersoid and the matrix63. This prevents chemical reaction and simul-taneously improves vvetting vvith the matrix.
However, these types of coatings are very expen-sive. The metallic coatings to the surface of SiC par-ticulates such as nickel64 and coper65 are also found to be effective. In these cases the coatings dissolve in the matrix alloy during synthesis to give precipita-tion of brittle intermetalics adjacent to the SiC dis-persoids. In ali cases the morphology of the coating structure. thickness of the coating and nature of the coating namely adhesion and bonding vvith the dis-persoid surface, are likely to play a significant role.
An easily scaled-up method for the preparation of protective surface layer to the SiC dispersoids vvas suggested by several authors24'376668. Note that Silicon carbide particles usually have a vitreous surface layer of Si02. Therefore, the initial interfacial reaction is betvveen Si02 and Al and, hence, a thick surface layer of Si02 can serve as a barrier for the undesir-able reaction betvveen SiC and Al. The Si02 layer on SiC can easily be thickened by heating in air. It is es-timated that heating in air at 973 K for 1 hour, in-creases the thickness of the oxide layer to betvveen 30 and 50 nm from the native thickness of betvveen 2 and 4 nm. As predicted, the use of SiC preheated at 973 K in air vvas found to reduce the reaction betvveen Al and SiC and also to improve the vvetting ac-cording to the several possible chemical reactions at the interface AI'Si0266 67. Si02 can also react vvith Mg vvhich is usually present in an Al alloy66 67. For low ox-idation levels (<4 wt%), the reaction of Si02 and Al leads to the formation of spinel, MgAI204. The reaction is rapid and completed during fabrication of the composites as indicated by the fact that no residual Si02 is observed in the reaction zone.
For high oxidation levels (=16 wt% Si02), the trans-formation of Si02 results into a continuous layer sur-rounding the SiC particles and is also complete during fabrication. The interface is polycrystalline and constituted mainly of MgAI204 vvith some Mg2Si and chanels of Al.
The oxidation of SiC particles performed in order to improve its wettability and, in the same tirne, to remote AI4C3 formation during composite process-ing, seems to be the key factor in the proprietary commercial foundry proces for MMCs produc-tion37.
Other sol-gel coating techniques based on MgO70, Al2037°, Zr0270, K2ZrF671 are also applied. None of these techniques, hovvever, can overcome the problem of rejection at higher particulate content in Al alloy melt. These methods are, also, too costly for most commercial application.
Carbon coatings produced by pyrolysis of phenolic resin or high yield polymers72 vvhich are frequently used in ceramic matrix composites are recently cosidered as a nevv promising way for the low cost large scale production of MMCs73.
Effect of dispersoids size and aspect ratio on the composite mechanical properties
Strength
A very strong dependence of strength of discontin-uously reinforced MMCs on particle size vvas ob-
served74. This relationship is approximately linear on a semi-logarithmic plot: bend strength versus rein-forcement diameter. As the particulate diameter de-creases from 800 to 6 |im, bend strength increases 4 to 6 fold. More precisely, the measured bedning stregths in Al/SiC MMCs vvith 100, 10 and 3 (im particulate diameter vvere 300, 500 and 600 MPa re-spectively. Finally, the extrapolated bending strength for SiC reinforcement vvith particle size only 1 (im vvas found to be around 700 MPa74. The authors74 are also found the similar strength increasing effect in MMCs vvith submicrometre SiC particles.
The observed increase in strength, as a function of particulate diameter may be a consequence of in-creased particle strength at smaller sizes due to sta-tistical (e.g. VVeibulI) arguments. An alternative ex-planation relates to the interaction of the dislocation fields surrounding the reinforcement particles. Strengthening of discontinuous MMCs has been shovvn to arise from the generation of a high dislocation density around the filler during cooldovvn from processing75-76.
Toughness
Again, as the particulate diameter decreases the composite toughness significantly increases. The increase in toughness is approximately linear on a se-mi - logarithmic plot (=10 MPam 1/2 in Al/SiC composites vvith 100 |im particulate diameter and s15 MPam 1/2 in samples vvith SiC particulate diameter around 5 |am)74.
Stiffness
Generally, there appears to be no effect of particulate diameter on the elastic modulus of either composite system74.
Wear
Ali of the discontinuously reinforced composites exibited improved vvear resistance vvith increasing particulate diameter74. As the particulate diameter increases from 8 to 800 (im, vvear resistance increases 5 to 6 fold. Hovvever, note that further increasing of SiC particle size over 65 (im there is no significan effect on vvear resistance74.
Aspect ratio and mechanical properties
For the optimum properties in the particulate composites, high aspect ratio and uniform particle distri-bution is important if conventional shear - lag composite strengthening is operating. Hovvever, the high aspect ratio needs to be achieved at fine particle sizes if particle fracture during composite fabrication is to be avoided. Most of the particulate composites developed so far utilize particles vvith aspect ratio of less than two, and particle sizes in the 10 to 20 (im size range because higher aspect ratio are only available in much coarser povvders.
While vvhiskers, according to its high aspect ratio, give the highest properties of ali the discontinuous reinforcement, the high cost of vvhiskers and their po-tential health hazards have resulted in the major ef-fort being concentrated on particulate reinforcement.
Whisker composites have approximately the same yield strength, a higher ultimate strength, and a lovver strain to failure than do paticulate reinforced composites in the extrusion direction. Hovvever, it's im-portant to note that vvhisker composites have lovver properties in the directions perpendicular to the ex-trusion.
An alternatively approach in using of particulates vvith high aspect ratio is based on SiC platelets pro-duced by heating a porous alpha silicon carbide pre-cursor composition comprising silicon and carbon in intimate contact to a temperature of from 2373 to 2773 K in a non-reactive atmosphere. By controling the thickness, the material can be tailored to have different aspect ratio.
According to77, silicon carbide platelets of three different sizes fabricated into 6061 aluminium povvder metallurgy compacts containing 25 vol% of the silicon carbide platelets have shovved improved properties vvith each reduction in platelet size. The SiC platelets (at 15 vol%) vvere also incorporated into an AI-356 aluminium alloy using a propietary molten metal mixing method - aspect of vvhich have been discussed previously77. Again, the tensile properties of the platelet reinforced composites vvere found su-perior to those of the particulate reinforced. Because of the improvement in properties of MMCs vvith re-duced platelet diameter, the development of this product was extended to include smaller thickness (=0.5 |.im) vvhich resulted in aspect ratio of about 10 vvith particles in the 5 to 10 |am size range.
The initial results of the incorporation of SiC platelets in aluminium alloys suggest that by limiting the size of platelets and by improving their dispersion and alignment, it should be possible to improve the composite properties significantly.
Foundry processing of Al/SiC composites vvith fine dispersoids: the state-of-the art and future trends
This class concerns ali techniques that elaborate MMCs by stirring an alloy above the liquidus (vortex method), or betvveen the solidus and the liquidus (compocasting), vvhile proceeding to the addition of the reinforcement dispersoids.
Mixing techniques generally used for the introduc-ing and homogeneously dispersing a discontinuous phase in a melt are:
•	Addition of particles to a vigorously agitated fully or partially molten alloy49,
•	Injection of discontinuous phase into the melt vvith an injection gun49,
•	Dispersion of pellets or briqurttes, formed by com-pressing povvders of base alloys and the ceramic phase, into a mildly agitated melt49,
■ Addition of povvders to an ultrasonically irradiated melt. The pressure gradients caused by cavitation phenomena promote homogeneous mixing of ce-ramics in metallic melts,
•	Addition of povvders to an electromagnetically stirred melt. The turbulent flovv caused by electro-magnetic stirring is used to obtain a uniform suspen-sion,
■ Centrifugal dispersion of particles in a melt, • Addition of ceramic phase to an accessory metallic melt (e.g. Si) vvhich wets previously (or in situ) sur-face engineered ceramic dispersoids, and can be. after that, successfully "diluted" by the basic metallic matrix (i.e. magnesium reach Al alloy) to the required final composition73.
In ali the above techniques, external force is used to transfer a nonvvettable ceramic phase into a melt and to create a homogeneous suspension of the ceramic in the melt.
A broad range of SiC particulates size (10 to 120 um) and amount of dispersoids (3 to 20 vol%) vvhich have been successfully incorporated into Al alloy matrix by foundry procedures is vvell document-ed49 78. Hovvever, these reports mainly discusse the introduction of relatively large SiC dispersoids (vvith particle size range betvveen 15 and 60 um). It can be noticed that there is only a fevv existing reports about the immersion of fine (less than 10 um in size) SiC particulates into Al melt using above described foundry routes1173'79. As mentioned earlier, finer than 10 f.im particulates also seems to be a lovver reason-able limit in particulate dimension under vvhich al-most complete rejection becomes the rule. Hance, these routes can be used mainly for the fabrication of vvear resistant composites, for vvhich larger particulates results in better performance. Hovvever, the possibility of using such a processes for the com-mercial fabrication of structural MMCs containing much finer particulates, at higher volumic fraction, re-mains questionnable.
The exception is the last of the above listed tech-niques, based on the combination of two compatible metallic melts. In the first of them, ceramic particulates can be successfully dispergated according to the chemical reaction betvveen the melt and SiC particulates vvith previously (or in situ) surface engineered layer. Once dispergated, ceramic particulates will not be rejected during the introduction of the second melt, if the process performs by carefully vvetting control73.
By follovving this procedure, Al/SiC composites vvith 10 - 20 vol% of SiC particulates in the size range less than 10 um are successfully prepared. This process is also advantageous in the fabrication of composites vvith particularly difficult to deal vvith rein-forcements, such as vvhiskers or platelets73. According to author's proposal, chemically treated fine SiC particles (or other morpohologies vvith high aspect ratio) vvere first dispersed into a Si melt (at ap-prox. 1773 K in vacuum). The wettability of SiC by Si(i, was enhanced by chemical reaction betvveen a carbon layer previously deposited on the surface of the SiC particles and the matrix. Carbon layer vvas produced by pyrolysis of phenolic resin or high carbon yield polymer source. When a sufficient portion of SiC particles vvas incorporated into Si matrix, an Al alloy vvas carefully added at a controled rate under vigorous stirring conditions and in a protective atmosphere in order to fit the final matrix composition.
The main advantages of this MMC preparation technique are:
•	The concentration of both matrix constituents - Al and Si can be tirne programmable, and,
•	The temperature of the intermediately melt can be effectivelx regulated, especially during the addition of AI-alloy into previously formed suspension of SiC particulates in Si melt.
In this way, it is possible to regulate: ■ The vvetting kinetics of SiC particulate reinforce-ments in the metallic matrix, and,
•	The kinetics of unvvanted interfaciai reaction vvhich leads to the formation of AI4C4 (Reaction 4). Note that the rate of Reaction 4 is expressed as:
v = k1[AI]4[SiC]3
At the beginning of the alloying process, the concentration of Al is practically zero vvhich results in a lovv AI4C3 formation rate. Concequently, the process-ing temperature can be kept sufficiently high in order to maintain the wettability of ceramic particulates vvith Al in the vvetting region73.
Hovvever, this method has also some limitations. In practice, it's difficult to select a pair of compatible metallic matrixes vvith ali requested performances. Another important unconvenience is the high pro-cessing temperature caused by high melting point of silicon.
The commercial importance of this method should be carefully evaluated in future, based on the im-provement of the mechanical properties od MMCs caused by the introducion of finer SiC paticulates.
Conclusions
Foundry processes, generally used for the prepa-ration of MMCs, concerns ali the techniques that elaborate MMCs by stirring an alloy above the liq-uidus (Vortex method, Duralcan technology) or be-tvveen the soiidus and the liquidus (Compocasting).
In spite of many existing problems, it is likely that in the next decade most of the MMC volume pro-duced in the vvorld vvill be done using one of them.
Discontinuously reinforced Al/SiC composites can be generally classified in wear resistant grade, for vvhich larger particulates (20 um < d < 60 um) at medium volumic fraction (10-15 vol%) results in better performance, and structural MMCs containig much finer particulates (< 10 |.im), at higher volumic ratio (20-25 vol%).
It can be noticed that a current foundry processes are used mainly for the fabrication of wear resistant composites. Hovvever, the possibility of using such a processes for the commercial fabrication of structural MMCs stili remains questionnable. Regarding poor wettability of fine ceramic particulates vvith molten metals, finer than 10 |om particulates seems to be a lovver limit in particulate dimension vvhich could be successfully introduced by existing foundry techniques. Also, the incorporation of large amount of SiC particulates (> 20-25 vol%) becomes difficult or even unpossible.
In order to overcome these problems, the new foundry processes vvhich enable the routinely intro-
duction of larger amount of finer SiC particles into metallic melt, must be developed.
One possible solution, based on the combination two compatible metallic melts, is presented in this re-vievv and theoreticaly evaluated.
References
1	R. J. Arsenault, S. B. Wu, in Proceedings of the International Symposium on Advances in Čast Reinforced Metal Composites, Chicago, September 1988, edited by S. G. Fishman and A. K. Dhingra (ASM International, Ohio, 1988), 231
2	R. J. Arsenault, S. B. Wu, Scripta Metallurgica, 22,1988, 767
3	P. K. Rohatgi, B. C. Pai, S. C. Panda, J. Mater. Sci., 14, 1979, 2277
4	F. A. Badia, P. K. Rohatgi, Trans. AFS, 77, 1969, 402
5	F. A. Badia, Trans. AFS, 77, 1971, 347
6	M. K. Surappa, P. K. Rohatgi, J. Mater. Sci., 16, 1981, 983
7	M. K. Surappa, P. K. Rohatgi, Met. Tech., 5, 1978, 358
8	R. Mehrabian, R. G. Riek, M. C. Flemings, Metal. Trans., 5A,1974, 1899
9	B. F. Quigley, G. J. Abbaschian, R. VVunderliand, R. Mehrabian, Metal. Trans., 13A, 1982, 93
10	C. G. Levi, G. J. Abbaschian, R. Mehrabian, Metal. Trans., 9A, 1978, 697
11	F. M. Hosking, F. FolgarPotillo, R. VVunderlinand, R. Mehrabian, J. Mater. Sci., 17, 1982, 477
12	C. Miliere, M. Suery, J. Mater. Sci. Technol., 4, 1988, 41
13	J. M. McCoy, C. Jones, F. E. Wawner, SAMPE Quarterly, 19, 1988, 2, 37
14	S. Ray, in Proceedings of the International Symposium on Advances in Čast Reinforced Metal Composites, Chicago, September 1988, edited by S. G. Fishman and A. K. Dhingra (ASM International, Ohio, 1988), 77
15	J. V.Naidich, Prog. Surf. Membr. Sci.. 14, 1981, 353
16	T. Young, Trans. Roy. Soc., 95, 1805, 65
17	A. Dupre, Theorie Mechanique de la chaleur, (Gauthier Villars, Pariš, 1869)
18	R. VVarren, C. H. Andersson, Composites, 15,1984,101
19	T. Choh, T. Oki, Mater. Sci. Technol., 3, 1987, 1
20	V. Laurent, D. Chatain, N. Eustathopoulos, J. Mater. Sci., 22, 1987, 244
21	J. P. Rocher, J. M. Ouenisset, R. Naslain, J. Mater. Sci. Lett., 4, 1985, 1527
22	F. Delannay, L. Froyen, A. Deruyttere, J. Mater. Sci., 22, 1987, 1
23	A. Banerji, P. K. Rohatgi, W. Reif, Metallvviss. Technik, 38, 1984, 656
24	V. Laurent, D. Chatain, N. Eustathopoulos, X. Dumant, in Proceedings of the International symposium on Advances in Čast Reinforced Metal Composites, Chicago, September 1988, edited by S. G. Fishman and A. K. Dhingra (ASM International, Ohio, 1988, 27
25	J. G. Legoux, L. Salvo, G. L'Esperance, M. Suery, in Proceedings of the International Conference on Fabrication of Particulates Reinforced Metal Composites, Montreal, September, 1990, edited by J. Masounave and F. G. Hamel) ASM International, Ohio, 1990), 31
26	W. Kohler, Aluminium, 51, 1975, 443
27	C. R. Manning, T. B. Gurganus, J. Amer. Ceram. Soc., 52, 1969, 115
28	V. Naidich, N. Chuvashov, J. Mater. Sci., 18,1983, 2071
29	B. C. Pai, P. K. Rohatgi, Mater. Sci. Eng.,21, 1975, 161
30	H. Tokisue, G. J. Abbaschian, Mater. Sci. Eng., 34, 1978, 75
31	S. K. Rhee, J. Amer. Cer. Soc., 54, 1971, 332
32	G. Imich, US Patent No.2, 793 949, 1957
33	J. V. Naidich, V. S. Zhuravlev, G. V. Chuprina, L. V. Strashinskaya, Soviet Povvder Metallurgy and Metal Ceramlcs, 12, 1973, 895
34	Hitachi Ltd., Jpn. Kokal Tokyo koho 8073, 839, 1980
35	B. N. Keshavaram, A. Banerji, M. K. Surappa, P. K. Rohatgi, J. Mater. Sel. Letters, 1, 1982, 29
36	T. P. Murali, M. K. Surappa, P. K. Rohatgi, Met. Trans., 13B.1982, 485
37	M. D. Skibo, D. M. Schuster, US Patent No.4, 865, 806
38	Y. Babaskin, Russian Castings Production, 1972, 328
39	R. Mehrabian, A. Sato, M. C. Flemings, Light Metals, 2, 1975, 177
40	R. Mehrabian, M. C. Flemings, New Trends In Materials Processing, (ASM Publication, Metals Park, Ohio, 1976), 98
41	H. V. VValter, G. Ziegler, Eur. Space Agency, Special Publication ESA SP/35, 1978
42	S. Das, T. K. Dan, S. V. Prasad, P. K. Rohatgi, J. Mater. Sci. Lett., 5, 1986, 562
43	S. Y. Oh, J. A. Cornie, K. C. Russell, Ceram. Engng. Sci. Proc., 8, 1987, 912
44J. A. Cornie, A. Mortensen, M. C. Flemings, in Proceedings of the Sixth International Conference on Composite Materials and the Second European Conference on Composite Materials (ICCM and ECCM), London, UK, Vol.2, edited by F. L. Matthevvs, N. C. R. Buskell, J. M. Hodgkinson and J. Morton (Elsevier Applied Science, London, 1987, p.2.297
45	T. Choh, R. Kammel, T. Oki, Z Metallkde, 78,1987, 286
46	D. O. Kennedy, Advances Materials&Processes, 6, 1991, 42
47	M. D. Skibo, D. M. Schuster, US Patent No.4, 786,467, 1988
48	M. D. Skibo, D. M. Schuster, US Patent No.4. 759,995, 1988
49	Metal Handbook, Ninth Edition, Vol.15 (Metals Park, Ohio, 1988
50	M. Skibo, P. L. Moriš, D. J. Lloyd, in Proceedings of the International Symposium on Advances in Čast Reinforced Metal Composites, Chicago, September 1988, edited by S. G. Fishman and A. K. Dhingra (ASM International, Ohio, 1988), 257
51	D. J. Lloyd, I. Jin, Met. Trans., 19A, 1988, 3107
52	A. Mortensen, in Proceedings of the 9th Riso International Symposium on Metallurgy and Materials Science, Roskilde, September 1988, edited by S. Anderson, H. Lilholt and O. Pedersen, (Riso National Laboratory, Rosklide, 1988), 144
53	A. M. Samuel, H. Liu, F. H. Samuel, Compos. Sci. and Techn.. 49, 1993, 1
54	D. J. Lloyd, Compos. Sci. and Techn., 35, 1989, 159
55	J. C. Viala, P. Fortier, J. Bouix, J. Mater. Sci.. 25, 1990, 1842
56	J. C. Viala, P. Fortier, C. Bernard, J. Bouix, Developments in the Science and Technology of Composite Materials (A. E. C. M., Bordeaux, 1985)
57	T. Yano, S. Kato, T. Iseki, J. Am. Ceram. Soc.. 75.1992, 580
58	S. D. Peteves, P. Tambuyser, P. Helbach, M. Audier, V. Laurent, D. Chatain, J. Mater. Sci., 25, 1990, 3765
59	T. Iseki, T. Kameda, T. Maruyama, J. Mater. Sci.. 1984. 1692
60	R. J. Arsenault, C. S. Pande, Scripta Metali., 18. 1984. 1131
61	D. J. Lee, M. D. Vaudin, C. A. Handvverker, U. R. Kattner. Mat. Res. Soc. Symp. Proc., 120, 1988, 357
62	R. Y. Lin, K. Kanniceedings, on the International Conference on Interfaces in Metal-Ceramics Composites, Anaheim (CA). February, 1990. edited by R. Y. Lin, R. J. Arsenault, G. P. Martins and S. G. Fishman (TMS, VVarrendale, 1990), p.153
63	R. K. Everett, R. J. Arsenault, Metal Matrix Composites: Processing and Interfaces, (Academic Press, Boston, 1991)
64	C. Y. Liue, J. W. VVang, Y. M. Peng. H. J. Chen, J. H. Shen, C. A. Hung, MRL Buli. Res. De v., 4, 1990, 31
65	ibid., 4, 1990, 35
66	J. G. Legoux, G. L'Esperance, L. Salvo, M. Suery, in Proceedings of the International Conference on Fabrication of Particuiates Reinforced Metal Composites, Montreal, September 1990, edited by J. Masounave, F. G. Hamel, (ASM, Ohio), p.31
67	T. Sritharan, K. Xia, J. Haethcock, J. Mihelich. in Proceedings of the International Conference on Metal and Ceramic Matrix Composites: Processing. Modeling and Mechanical Behavior, Anaheim (CA), February 1990, edited by R. B. Bhagat. A. H. Clauer, P. Kumar and A. M. Ritter, (TMS, VVarrendale, 1990), p.13
68	H. Ribes. M. Suery, Scripta Metali., 23, 1989, 705
69	N. I. A. Lattef, A. R. I. Khedar, S. K. Goel. J. Mater. Sci. Lett., 4, 1985, 385
70B. Kindl, Y. H. Teng, Y. L. Liu, Composites, 25,1994. 671 71J. P. Rocher, J. M. Ouenisset, R. Naslain. J. Mater. Sci..
24, 1989, 2697 72K. Sugihara, M. Yamamoto, T. Kida, M. Fukazavva. US. Patent No. 4, 929,472, 1990
73	V. M. Kevorkijan, B. Šuštaršič, Paper no. SIV-25-95 on the 97th Annuai Meeting of the American Cer. Soc.. Cincinnati, May 1995
74	J. T. Burke, M. K. Aghajanian, M. A. Rocazella, Proc. Int. SAMPE Symp., 34, 1989, 2440
75	D. McDanels, Met. Trans., 16A, 1985, 1105
76	M. Vogelsang, R. J. Arsenault, R. M. Fisher. Met. Trans., 17A, 1986, 379
77	M. K. Jain, D. J. Lloyd, S. P. Tremblay, in Proceedings of the International Conference on Metal and Ceramic Matrix Composites: Processing, Modeling and Mechanical Behavior, Anaheim (CA), February 1990. edited by R. B. Bhagat, A. H. Clauer, P. Kumar and A. M. Ritter (TMS, VVarrendale, 1990), p.549
78	P. Rohatgi, Adv. Mater. Proc., 137, 1991, 39
79	R. Chen, G. Zhang, Composite Science and Technology, 47, 1993, 51