Melt Processed Aluminum Matrix Particle Reinforced Composites

3.21 Melt Processed Aluminum Matrix Particle Reinforced Composites D. J. LLOYD and I. JIN Alcan International Limited, Kingston, ON, Canada 3.21.1 INTRODUCTION

1

3.21.2 REINFORCEMENT 3.21.2.1 The Application 3.21.2.2 The Method of Composite Manufacture 3.21.2.3 Cost

2 2 3 3

3.21.3 MATRIX±REINFORCEMENT REACTIVITY

4

3.21.4 MOLTEN METAL COMPOSITE MANUFACTURE

7

3.21.4.1 3.21.4.2 3.21.4.3 3.21.4.4 3.21.4.5 3.21.4.6

Mixing Methods Semisolid Casting Pressure Infiltration Pressureless Infiltration Spray Deposition In Situ Reaction Synthesis

7 8 9 9 10 10

3.21.5 SECONDARY FABRICATION 3.21.5.1 3.21.5.2 3.21.5.3 3.21.5.4

11

Hot Working Casting Machining Joining

11 12 14 15

3.21.6 MICROSTRUCTURES

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3.21.6.1 Reinforcement Distribution 3.21.6.2 Grain Structure

16 17

3.21.7 MECHANICAL PROPERTIES

18

3.21.8 DEFECTS

20

3.21.9 SUMMARY

21

3.21.10 REFERENCES

21

3.21.1

INTRODUCTION

with high mechanical properties restricted their use to aerospace applications. In addition, composites reinforced with continuous fibers have no secondary fabrication capability, since the fibers are heavily damaged by any forming operation, such as extrusion, rolling, or forging. This means that continuously reinforced composites must be used in essentially the same

The development of metal matrix composites (MMCs) has been one of the major innovations in materials in the last 20 years. While the early work on composites concentrated on continuous reinforcements, such as graphite and boron fibers, the prohibitive cost of continuous fibers 1

2

Melt Processed Aluminum Matrix Particle Reinforced Composites

form as they are initially fabricated, restricting them to applications involving simple shapes, such as sheet and tube. While development work is continuing on continuously reinforced MMCs, particularly for Ti/SiC composites, where several processing routes have been developed (Chou et al., 1985; Smith and Froes, 1984; see also Chapters 3.25 and 3.27, this volume), more recent effort has concentrated on discontinuous reinforcement. The availability of SiC whiskers in the mid-1960s led to a considerable effort on whisker reinforced composites, and these composites have been the subject of several reviews (Divecha et al., 1981; Nair et al., 1985). However, the high cost of whiskers together with toxicology concerns has severely restricted commercialization of MMCs, while they remain of interest for ceramic matrices. A significantly cheaper reinforcement became available with the development of short staple Al2O3 and mullite fiber (Dinwoodie et al., 1985; Stacey, 1988), and these were commercialized in aluminum matrices for selective reinforcement of engine parts. An example of such an application is the reinforcement of the ring land area of diesel engine pistons to provide wear resistance (Donomoto et al., 1983). These composites are made by pressure infiltration of a preform (see Chapter 3.20, this volume), and this can be achieved during squeeze casting to form the part itself, so it is a relatively low-cost process. However, there is still no secondary fabrication capability, and the performance of the composite is critically dependent on the quality of the preform, which itself depends on a wide range of factors such as binder content and shot content (shot is nonfiber material produced during initial fiber processing). The ready availability of a wide range of cheap ceramic powders, such as SiC and Al2O3 used for abrasives and cutting media, led to their consideration as reinforcements for metals. While particle reinforcements cannot provide the property enhancement achieved by continuous fibers (Lloyd, 1994; Clyne and Withers, 1993), they can provide sufficient improvement in some mechanical and physical properties to provide a useful increase in engineering performance. Particle reinforced MMCs have the added advantage that they can be processed using conventional equipment, and hence, minimize capital costs associated with fabrication. Because ceramic powders have very low aspect ratios (the aspect ratio is the ratio of the length of the particle to its width), typically less than 5:1, particle reinforced MMCs display quite isotropic properties, which is desirable in most engineering applications, and provides compatibility with

traditional unreinforced metals. While some properties of particle reinforced MMCs are very different to unreinforced metals, their general behavior is more akin to metals than continuous fiber reinforced composites, and this has been an important factor in their commercialization in applications such as bicycle frames, brake rotors, and other cast and wrought parts (see Chapter 3.26, this volume). In the applications to date, the main metal matrix that has been used is aluminum, primarily to minimize weight since the composites are often replacing heavier steel and cast iron presently used in the application. However, other matrices are under development, particularly magnesium, which would further reduce the weight. There are many factors which have to be considered in the successful melt processing of particle reinforced MMCs. The choice of the reinforcement is influenced by the application, the primary fabrication process, and the alloy matrix involved. Since the distribution of the reinforcement and the microstructure of the composite have important effects on the final properties, it is important to control these in the final product.

3.21.2

REINFORCEMENT

An overview of the short fibers and whiskers available for metal matrix composites has been given by Stacey (1988). Since most ceramics are available as particulates, there is a wide range of potential reinforcements for particle reinforced composites. Rohatgi and co-workers have used mica (Deonath and Rohatgi, 1980), alumina, silicon carbide, clay (Surappa and Rohatgi, 1981), zircon (Banerji et al., 1983), and graphite (Krishnan and Rohatgi, 1984), and other reinforcements, such as boron carbide (Kai et al., 1989) and titanium diboride (Roebuck and Forno, 1988), have also been examined. However, the choice of reinforcement is not as arbitrary as this list of composites might suggest, but is dictated by several factors.

3.21.2.1

The Application

If the composite is to be used in a structural application, the modulus, strength, and density of the composite will be important, which requires a high modulus and a low density reinforcement. Particle shape may be important, since angular particles can act as local stress raisers, reducing ductility. If the composite is to be used in thermal management applications,

Reinforcement

3

Table 1 Properties of SiC and Al2O3 reinforcement. Particulate

Modulus (GPa)

Density (g cm73)

CTE (K71)

Specific heat (J kg71K71)

Thermal conductivity (Wm71K71)

Poisson's ratio

SiC Al2O3

420±450 380±450

3.2 3.96

4.3 x 1076 7.0 x 1076

840 1050

10±40 (at 1100 8C) 5±10 (at 1000 8C)

0.17 0.25

the coefficient of thermal expansion and thermal conductivity are important. The coefficient of thermal expansion is generally important because it influences the strength of the composite through the development of internal stresses (Lloyd, 1994).

3.21.2.2

The Method of Composite Manufacture

There are two generic methods for composite manufacture: powder metallurgy and methods involving molten metal. In the case of powder metallurgy (see Chapter 3.25, this volume for details), the matrix alloy powder is blended with particles of the reinforcement to achieve a homogeneous mixture. To achieve this, the sizes of the metal and ceramic powders need to be carefully chosen so that agglomerates are not left after blending, and carry over into the final product. The appropriate size ratio will depend on the blending process used, but in one case a SiC:Al particle size ratio of 0.7:1 gave a more uniform reinforcement distribution than a ratio of 0.3:1 (Lewandowski et al., 1987). In powder metallurgy processing, the brittle ceramic particles are also susceptible to particle fracture, which is dependent on the particle aspect ratio and flaw density. Typically, the atomized aluminum powder particle size is in the range of 20±40 mm, and reinforcement particle sizes are 3±20 mm, with aspect ratios of 55:1. For composites processed in the molten state, there are different considerations. In some of these processes, the ceramic particles can spend considerable time in contact with the molten alloy matrix, and this can result in reaction between the two (Lloyd et al., 1990). For example, SiC is thermodynamically unstable in most molten aluminum alloys, reacting to form aluminum carbide, Al4C3, whereas it is stable in many molten magnesium alloys. On the other hand, Al2O3 is stable in most magnesium-free aluminum alloys, but unstable in magnesium alloys, reacting to form spinel, Al2MgO4. Reaction of the reinforcement can severely degrade the properties of the composite, so the reinforcement has to be chosen after considering the matrix alloy and the processing time and temperature. The reinforcement particle size is

also important because, while it is generally easier to incorporate coarser particles into the melt, large particles are more susceptible to gravity settling and can result in a heavily segregated casting (Lloyd and Chamberlain, 1988). However, finer particles increase the viscosity of the melt, making processing difficult. Most molten metal processes use ceramic particles in the 10±20 mm size range, but it is important to use a narrow size range, avoiding a large fraction of either fines or coarse particles.

3.21.2.3

Cost

A major reason for using particulates is to reduce the cost of the composite, so the reinforcement has to be readily available in the quantities, size, and shape required at low cost, i.e., around US$5 kg71. With these considerations in mind, the two reinforcements receiving the most attention are SiC and Al2O3, and some of the properties of these two reinforcements are given in Table 1. Figure 1 shows the typical morphology of SiC

Figure 1 Powders used for reinforcing composites: (a) SiC; (b) Al2O3.

4

Melt Processed Aluminum Matrix Particle Reinforced Composites ing point, reacting to form aluminum carbide, Al4C3 (Lloyd and Jin, 1988): 4Al ‡ 3SiC ! Al4 C3 ‡ 3Si

Figure 2 Decrease of the melting point with time at 675 8C for 7075/SiC/15p (after Lloyd and Dewing, 1988).

and Al2O3 powders used as reinforcement; note the clean surfaces and the absence of any very fine particles.

3.21.3

MATRIX±REINFORCEMENT REACTIVITY

The reactivity between the matrix and the reinforcement is very important in melt processed composites, since it can influence both the ease of processing and the final properties of the composite. Reaction between the reinforcement and the matrix can also result in changes in the matrix alloy metallurgy which will, in turn, influence such basic properties as the melting point of the alloy (Lloyd and Jin, 1988) and its strength. The thermodynamical stability, in aluminum and magnesium alloys, of most of the reinforcements of interest has recently been reviewed (Lloyd et al., 1990). In aluminum alloys, SiC is thermodynamically unstable above the melt-

…1†

with an increase in the silicon level of the matrix. As the reaction proceeds the silicon level increases, and the melting point of the composite decreases with time, as shown in Figure 2 for an Al±Mg±Zn matrix composite, 7075/SiC/ 15p (Lloyd and Dewing, 1988). (Throughout this chapter the Aluminum Association designation for MMCs is used: matrix/reinforcement/volume fraction followed by p, f, or w for particles, fibers, or whiskers, respectively.) The aluminum carbide reaction can be avoided by using high silicon alloys for the matrix, as shown in Figure 3 (Lloyd, 1989). So for processing routes involving long contact times between the reinforcement and the melt, high silicon aluminum alloys are preferred. In powder processing using solid-state consolidation (see Chapter 3.25, this volume), aluminum carbide formation is not a factor because silicon carbide is stable below the solidus; however, if liquid phase sintering is involved there is the potential for reaction. It is the molten metal processing routes that are particularly prone to reaction, since the particle±liquid metal contact times can be long, particularly when large-scale casting processes or remelting are involved. The spray deposition method (see Chapter 3.23, this volume for details) is the least susceptible, since in this case the contact time is only of the order of seconds. Other carbides, such as boron carbide and titanium carbide, are also thermodynamically unstable in molten aluminum, but often react in a more complex manner (Lloyd et al., 1990).

Figure 3 Silicon level required at different temperatures to prevent aluminum carbide formation (after Lloyd, 1989).

Matrix±Reinforcement Reactivity

Figure 4 Spinel crystals on the surface of extracted Al2O3 particles (after Lloyd, 1994).

Magnesium has no stable carbide, so ceramic carbides are stable in pure magnesium. However, many of the magnesium alloys of interest contain alloying elements, such as aluminum, which will form carbides, and in these magnesium alloys reaction may occur if the contact times are sufficiently long. Aluminum oxide, Al2O3, is stable in pure aluminum, but reacts with magnesium in Al± Mg alloys: 3Mg ‡ Al2 O3 ! 3MgO ‡ 2Al

…2†

and 3Mg ‡ 4Al2 O3 ! 3MgAl2 O4 ‡ 2Al

…3†

The magnesium levels in equilibrium for Equations (2) and (3) have been calculated, and show that at high magnesium levels and lower temperatures, MgO may form, while the spinel will form down to very low magnesium levels (Lloyd et al., 1990). It is not surprising, therefore, that Al2O3 is not thermodynamically stable in most aluminum alloys, since they contain magnesium. Other oxides, such as MgO, are expected to be stable. It should also be noted that, unlike SiC, which is stable below the solidus, this is not the case for Al2O3, and reaction can continue in the solid state. So solid-state processing may still result in a reinforcement reaction in this case. Figure 4 shows the spinel crystals formed on the surface of Al2O3 particles after reaction above the melting point in an Al±Mg alloy. Obviously, from Equation (3), Al2O3 will be unstable to some extent in magnesium alloys, but other oxides such as MgO and Y2O3 will be stable. While these thermodynamic considerations show a tendency for reaction to occur, it is the kinetics and the extent of reaction which are of

5

Figure 5 Change in matrix magnesium content due to spinel formation for different initial magnesium contents (after Jin and Lloyd, 1993).

practical importance. The kinetics of the aluminum carbide and spinel reactions have been considered in the literature (Lloyd and Dewing; McLeod and Gabryel, 1992), and as expected, the rate of reaction increases with increasing temperature, and the amount of reaction increases with increasing time. Minimizing the extent of particle±matrix reaction is a priority in the processing of composites. Undesirable Al4C3 can be avoided by appropriate choice of the silicon content of the alloy, which, while restricting the choice of matrix alloy, is the most convenient approach for those SiC reinforced composites involved in molten metal processing. Another approach is to oxidize the surface of the SiC, forming an outer layer of SiO2. In this case the early stages of the reaction involve reducing the SiO2, rather than dissolving the SiC. The reaction product will depend on the matrix alloy, being Al2O3 for pure aluminum, and MgO and Al2MgO4 for magnesium containing alloy (Legoux et al., 1990). This is not a very satisfactory approach, since there is still an interface reaction product, and magnesium is lost from the matrix, reducing the age hardening response (Ribes and SueÂry, 1989). For Al2O3 reinforcement, the spinel reaction can be minimized by using a low magnesium content matrix alloy. There is also some recent information that using a mixed oxide, mullite, in the crystallized, fully dense condition greatly reduces reactivity (Sritharan et al., 1990), but no detailed kinetic data have been published. Another approach is to use the reaction itself to form a barrier layer on the surface of the Al2O3 particles (Jin and Lloyd, 1993). Since reaction depends on intimate contact between the particle and the surrounding liquid, the reaction can be stopped if a

6

Melt Processed Aluminum Matrix Particle Reinforced Composites

dense interface reaction product is formed, which occurs when Al2O3 reacts in high magnesium-containing alloys. Figure 5 shows that magnesium contents of greater than 4 wt.% effectively minimizes the spinel reaction. The interface chemical reaction in alumina-reinforced composites can also be controlled by adding a surface-active element such as strontium to the matrix alloy (Lloyd et al., 1995). In this case, the surface-active element strontium effectively segregates to the matrix±particle interface and inhibits the nucleation of spinel crystals. A very unusual chemical reaction in Al2O3 reinforced composites has been recently reported. When a small amount of strontium was added to an aluminum alloy melt during composite processing, the sodium level in the melt increased to a very high value (5±10 times higher than the normal level). The results were explained by an ion exchange process at the melt±particle interface where strontium replaces sodium in Lloyd et al. (1996). This result clearly dictates another restriction to the choice of matrix alloy±reinforcement combination. The interface reaction can have several undesirable effects. Al4C3 dissolves in water, degrading the corrosion behavior of the composite, and its formation involves the release of silicon, modifying the matrix composition. Al2MgO4 is not expected to affect the corrosion behavior directly, but will modify the matrix composition. Matrix compositional modifications can in principle be allowed for by making an elemental addition, but the interface reaction products may also modify the mechanical properties of the interface (Lloyd et al., 1990; Mortensen, 1988). However, for those composites undergoing melt processing either in primary processing as in the melt mixing process, or in secondary processing such as shape casting, it is the effect that the reaction has on the viscosity of the melt, and hence on casting fluidity, which is particularly important. Casting fluidity is usually assessed by a test, such as the spiral test, which measures the distance the melt will flow before solidification occurs. So the test involves both rheological factors, such as viscosity, and solidification factors, such as latent heat, surface tension, superheat, cooling rate, and the freezing range of the alloy. For an unreinforced metal, the distance a melt will flow along a channel before solidification, L, is given by (Flemings, 1974): L ˆ frs d0 v=2h…Tm ÿ T0 †g…H ‡ CDT†

…4†

where rs is the density of solid, taken to be equal to the density of the liquid, d0 is the diameter of the channel, v is the velocity of

the liquid, h is the heat transfer coefficient, Tm is the melt temperature, T0 is the mold temperature, H is the latent heat, C is the heat capacity, and DT is the melt superheat. For composites several terms will be modified. The density, rs, will be replaced by the density of the composite, rc, rc ˆ rm …1 ÿ Vp † ‡ rp Vr

…5†

where rm is the density of the matrix, rp is the density of the particulate, and Vp is the volume fraction of particulate. The latent heat involved in solidification will be reduced because the particles are not involved in the solidification process; the effective latent heat, He, is given by He ˆ H…1 ÿ Vp †

…6†

Similarly, the effective specific heat of the composite, Ce, is given by Ce ˆ Cm …1 ÿ Wp † ‡ Cp Wp

…7†

where Cm is the specific heat of the matrix, Cp is the specific heat of the particulate, and Wp is the weight fraction of particulate. Substituting the appropriate values for the composite into Equation (4) will result in a decrease in the expected fluidity length with increasing particle content. However, of more importance is the melt velocity term, v, under the applied metallostatic head. The viscosity increases with increasing volume fraction of particles, and this will reduce the melt velocity. For nonmetallic liquids containing spherical particles, the viscosity is given by (Thomas, 1965): Zc ˆ Zm …1 ‡ 2:5Vp ‡ 10:05Vp †

…8†

where Zc is the viscosity of the composite and Zm is the viscosity of the unreinforced matrix. The non-Newtonian behavior of molten composites makes comparison with Equation (8) difficult, but it is expected to be of the right order at very high shear rates, when Newtonian behavior is approached. At low shear rates, or when reaction occurs between the reinforcement and the melt, it greatly underestimates the viscosity. If extensive reaction occurs, the viscosity can increase to essentially infinity, and the melt will not flow into the mold at all. As a result, rheological factors dominate the casting fluidity under these circumstances. This effect is demonstrated in Figure 6, where the spiral fluidities after different holding times at 750 8C and 800 8C are shown. For 750 8C the spiral fluidity remains about constant with holding time, because the extent of aluminum carbide formation is limited at this temperature

Molten Metal Composite Manufacture

7

Figure 6 Spiral fluidities of 356/SiC/15p after different holding times at (a) 750 8C and (b) 800 8C (after Lloyd, 1994).

in a 7 wt.% silicon alloy. However at 800 8C, the amount of aluminum carbide increases rapidly with time, resulting in a marked decrease in fluidity until, after 250 min, the composite will not flow into the mold. The rheological behavior of composite melts is poorly understood, but a few general statements can be made: (i) the viscosity is non-Newtonian, decreasing with increasing shear rate; (ii) the viscosity increases with increasing volume fraction of particulate; (iii) the viscosity increases with decreasing particle size; (iv) the viscosity is dependent on the history of the melt, in terms of temperature, time, and shear rate; (v) the viscosity increases with increasing reaction product at the interface. In terms of Equation (8), interface reaction will often result in an increase in the volume fraction of solid in the melt. For example, from the stoichiometry of the SiC reaction with Al4C3, and the lower density of Al4C3 compared to SiC (2.36 vs. 3.2 g cm71), a unit volume of SiC reacted with Al4C3 will result in a particulate volume increase of 60%. In addition, as pointed out by Surappa and Rohatgi (1981), the viscosity may be dependent on particle surface area, and this will increase with interface reaction because the reaction products tend to be in the form of fine crystals (Banerji et al., 1983). The non-Newtonian and thixotropic nature of composite melts indicate that the composite melts have a structure associated with them, presumably reflecting particle clustering in the melt. Particle clusters could occlude liquid within the clusters, effectively raising the solid fraction of the melt (J.A. Cornie, private communication).

It is apparent that reinforcement reactivity can influence the behavior of the composite in a variety of ways, and needs to be considered both with regard to composite processing and composite use.

3.21.4

MOLTEN METAL COMPOSITE MANUFACTURE

Early attempts to incorporate ceramic particles into metallic melts had limited success because most metals do not wet ceramic particles, and this results in rejection of the particles from the melt (Badia and Rohgati, 1969; Patton, 1972). The basic thermodynamics associated with incorporating a single particle into a melt have been considered in some detail (Neumann et al., 1973; Rohatgi, 1986; Russell et al., 1986; Rohatgi and Asthana, 1988), and show that the contact angle between molten aluminum and the ceramic particle must be less than 908 for successful incorporation. The wetting angle is usually measured by the sessile drop method, and is influenced by several variables including the heat of formation, stoichiometry, valence electron concentration, interfacial reactions, temperature, and time (Delannay et al., 1987). In general, however, molten aluminum does not wet most ceramic particles at typical casting temperatures, i.e., 5800 8C, and the molten metal methods attempt to improve this wetting behavior by supplying mechanical energy to overcome the wetting angle deficiency. 3.21.4.1

Mixing Methods

The early mixing method of Surappa and Rohatgi (1978) introduced ceramic particles

8

Melt Processed Aluminum Matrix Particle Reinforced Composites

through the sides of a vortex created in the melt with a mechanical impeller. This method is helped by the addition of surface-active elements such as 0.5 wt.% magnesium to aluminum melts, or by metal coatings to the particles (Hall, 1987). However, the process is limited to coarse ceramic particles, >50 mm, and low volume fractions, 510 vol.%, and it is difficult to avoid oxide skins being incorporated into the composite. The major breakthrough in mixing processes came with the development of Duralcan (Duralcan owned by Alcan Aluminum Limited) process by Skibo and Schuster (Alcan Aluminum Corporation, 1988). This process is applicable to conventional aluminum alloys, using uncoated ceramic particles of around 10 mm and larger, and produces reinforcement levels of up to 25 vol.%. It uses a proprietary vacuum mixing technology, and Duralcan material is now being commercially produced in batches of up to 6800 kg, and cast as pig or direct chill (DC) ingot. Two other aluminum companies have announced the development of a molten mixing method. Hydro Aluminium A.S. have discussed composites which appear to be comparable to Duralcan material (Borradaile et al., 1989), and Comalco introduced Comral, which is 6061 reinforced with a spherical, mixed oxide, Al2O3±SiO2, reinforcement (Sritharan et al., 1990), but have since stopped production. The molten metal mixing method is attractive because, in principle, it allows all the conventional metal processing routes to be used, and hence minimizes the cost. It does, however, have some problems in terms of reactivity between the reinforcement and the melt, considered earlier, and particle segregation effects, which will be considered in Section 3.21.6.1. The volume fraction of reinforcement is also limited in the mixing method because the viscosity of the melt increases with particle content and becomes non-Newtonian. The viscosity of unreinforced aluminum is around 1073 poise (1074 Pa s), but the viscosity of molten composites can be orders of magnitude higher, and shear rate dependent (Lloyd, 1990). The rheology of molten composites is poorly understood, but appears to be history dependent (Cornie et al., 1990), and the viscosity increases with increasing volume fraction and decreasing particle size. As a result, the power requirements necessary for mixing limits the amount of reinforcement that can be incorporated into the melt. Coarser particles are easier to incorporate, and it is important that the initial melt have a high level of cleanliness and a low gas content. Mixing should be carried out as quickly as possible, commensurate with achiev-

ing a good dispersion of reinforcement, since this will minimize any reaction between the particles and the melt. 3.21.4.2

Semisolid Casting

The microstructures developed by stirring a semisolid melt were investigated in the early 1970s, and have been reviewed (Cornie et al., 1990). The partially solidified, nondendritic microstructure developed has high viscosity which inhibits ceramic particle settling and floating, and can be used to retain particles in the melt (Mehrabian et al., 1975; Gibson et al., 1982), in a process often referred to as compocasting. This method has been developed commercially for magnesium alloys by the Dow Chemical Company (1984). Compocasting is restricted to longer freezing range alloys and, other than this, has the same limitations as the fully liquid mixing methods. Particles are added to a vigorously stirred melt containing about 40% solid. Initially this solid fraction consists of primary solid particles of the unreinforced alloy, which are spheroidized as a result of the imposed shear. Reinforcing particles are then added to the melt, which continues to be stirred, resulting in their incorporation. As reinforcing particles are added, the temperature of the melt should be increased to maintain the solid fractionÐprimary alloy particles plus reinforcing particlesÐat around 40%. If this is not achieved, and the volume fraction of solid becomes too high, total rejection of the reinforcing particles occurs. This is due to the fact that the semisolid melt is essentially behaving as a colloidal mixture, and the liquid content of the melt has dropped too low to accommodate the added reinforcing particles. Once all the required reinforcement has been added, stirring should be continued for a further period to achieve good particle wetting and distribution. A major problem with this method is that the stirring required for incorporation and particle entrapment also makes it very difficult to achieve a clean, low gas content melt, and the resulting composites tend to have a high porosity level. Also, the temperature control required limits the alloys which can be used, and is difficult to achieve in large melts. Nevertheless, a considerable amount of research work is currently investigating semisolid processing (Brown and Flemings, 1992). An interesting variant of the semisolid processing route is Thixomoulding, developed for magnesium matrix composites by Thixomat Inc. (Carnahan et al., 1990; Schilt et al., 1995), and shown diagrammatically in Figure 7. In this process, cold magnesium pellets, along

Molten Metal Composite Manufacture

9

Figure 7 Thixomoulding process (courtesy of Thixomat Inc.).

with reinforcing particles, are fed into a screw feeder heated into the semisolid range of the alloy, where the shear produces a semisolid composite charge of up to 65% solid. A highspeed shot system for reciprocation of the screw then forces the thixotropic slurry into a die at high velocity and pressure to form the part, then retracts the screw for the next slurry forming part of the cycle. Heating of the system is accomplished by a combination of resistance and induction heaters which bring the charge up to the precise slurry forming temperature, in the range of 550±580 8C for magnesium alloys. AZ91 alloy has been reinforced with Al2O3, SiC, and B4C using this method (Carnahan et al., 1990), and it has the attraction that composite formation and part production are combined in a single process.

3.21.4.3

Pressure Infiltration

In this process, a packed bed of ceramic particles is first evacuated and then infiltrated by a pressurized melt to form a ªmaster composite hardenerº containing about 50 vol.% of reinforcement. This hardener can then be diluted by adding it to an unreinforced melt, and the reinforcement redispersed by submerged mixing (Cornie et al., 1990; Klier, 1988). However, this dispersion is quite difficult because of the high viscosity of the master hardener, and it is difficult to obtain dispersion without incorporation of gas and oxide into the melt. Recently a bottom mixing process has been suggested, where an evacuated packed bed in the bottom of a crucible is covered with a melt, and the stirrer shears the interface between the particles and the melt, resulting in incorporation (Caron and Masounave, 1990). Combining these two processes, by replacing the ceramic bed with the pressure infiltrated composite hardener, may provide a useful way of diluting

the hardener (J.A. Cornie, private communication). The pressure infiltration route has the advantage that it is a means of producing composites with a high volume fraction of reinforcement, and since the system is evacuated, the master composite has very low porosity. The infiltration chamber can act as a mold to form a near net shape. With an additional dilution stage, the same factors come into play as with the direct mixing route, i.e., reactivity, viscosity, etc. Melt infiltration is considered in detail in Chapter 3.20, this volume.

3.21.4.4

Pressureless Infiltration

A recent molten metal process is the Lanxide Corporation PRIMEX pressureless infiltration process (Urquhart, 1991; Singer, 1985). There are two versions of this process, and these are shown diagrammatically in Figure 8. In one version a packed bed of ceramic powder is infiltrated by an Al±Mg alloy, without any applied pressure, in a nitrogen atmosphere. This requires a high temperature, and the infiltration rate is quite slow, but the efficiency of the process can be improved by incorporating magnesium powder into the ceramic bed, as in the second version shown in the Figure 8. The resulting composite, which will have a packed bed density of around 55 vol.%, can then be diluted in the appropriate matrix alloy if desired. Ceramic particles of SiC and Al2O3, with particle sizes as fine as about 1 mm, have been infiltrated in this way. Processing details of the PRIMEX route are proprietary, but it would appear to be a very competitive process for higher volume fraction composites, and can produce near net shapes directly. The process is one of reactive spontaneous infiltration, and is considered in more detail in Chapter 3.20, this volume.

10

Melt Processed Aluminum Matrix Particle Reinforced Composites

Figure 8 The PRIMEX pressureless infiltration process (courtesy of Lanxide Corporation, Newark, NJ).

3.21.4.5

Spray Deposition

The spray deposition process for unreinforced alloys was developed by Singer (1985), and commercialized by Osprey Metals (1975, 1977). It involves atomizing a melt and collecting the semisolid droplets on a substrate. As a result, the process is a hybrid rapid solidification process. The ªSpray Co-deposition Processº is a variant of the basic process, where ceramic particles are introduced into the spray and co-deposited with the alloy droplets (United Kingdom Atomic Energy Authority, 1990). The deposition rate is 6±10 kg min71, and Alcan has used the process to produce 200 kg ingots (Willis, 1988). The process has the advantage that the contact time between the melt and the reinforcing particles is brief, so reaction between the two is limited and a wider range of reinforcements are possible provided the assprayed billets are not remelted. The initial billets are typically 95±98% dense, and require a secondary fabrication step to achieve full density. As long as the alloy has a sufficient freezing range to achieve atomization at moderate superheats, any matrix composition can be used, including the advanced aerospace alloys, such as the Al±Li 8090 alloy (White et al., 1987). The particle distribution is sensitive to the spray atomization and reinforcement injection set-up. The reinforcing particles can be injected low down in the spray cone, in which case they tend to attach themselves to the surface of the spray drops, and are situated between the individual splat after solidification; alternatively, the reinforcing particles can be injected high up in the atomization cone, and with sufficient kinetic energy that they penetrate the individual drops of the atomized spray, and in this case the particles are within

the individual splats after solidification, giving a more homogeneous reinforcement distribution in the ingot. Figure 9 shows these two situations. In either case, however, there is a tendency to produce a layered type structure in large ingots, which has to be dispersed by subsequent fabrication. The spray deposition process is considered in detail in Chapter 3.23, this volume. 3.21.4.6

In Situ Reaction Synthesis

This is a rather different approach to composite manufacture than the previous molten metal methods, and one example is the XD process developed by Martin Marietta Corporation, in which ceramic particles are produced in situ in a melt (Martin Marietta Corporation, 1990). The process consists of adding to a solvent metal, such as aluminum, compounds which will react exothermally to produce the required reinforcing particles. A wide range of ceramic compounds can be formed by this process (Martin Marietta Corporation, 1987), but the two which have received most attention are TiB2 and TiC, which can be formed by the following reactions: 2B ‡ Ti ‡ Al ! TiB2 ‡ Al

…9†

and C ‡ Ti ‡ Al ! TiC ‡ Al

…10†

The particles are typically single crystal, and should have clean, unoxidized interfaces because they are formed in situ. By varying the process parameters, such as reaction temperature, the reinforcement size can be varied from approximately 0.2 to 10 mm, although the

Secondary Fabrication 3.21.5

11

SECONDARY FABRICATION

After initial manufacture, the composite has to be converted into the final part by secondary fabrication, which may involve a range of fabrication processes such as extrusion, forging, casting, machining, etc. The presence of brittle, abrasive particles in metal matrix composite means that some modification to the normal fabrication process may be required.

3.21.5.1

Figure 9 Microstructure of the spray drops with (a) SiC at the surface; (b) SiC in the interior of the drops.

material reported in the literature has particles in the 0.25±1.5 mm range. The initial composite is then used as a hardener with a pure aluminum matrix, which is then diluted in a melt of the desired matrix alloy. An alternative in situ method for producing TiC particles is to inject a carbon gas into an Al±Ti melt at a sufficient temperature for the exothermic reaction to TiC to occur (US Pat., 1989). It has the attraction of producing particles which should be inherently wet by the matrix, and therefore provide high interfacial strength. However, the fine particle size is expected to produce highly viscous melts, which may make handling and dilution difficult. For completeness it should be noted that there is a solid-state equivalent to the molten in situ method, which achieves reaction, and consolidation by means of self-propagating synthesis (Gotman et al., 1994). In this process the ignition of a lightly compacted powder blend initiates a combustion wave that propagates through the blend, leaving behind the reaction products which make up the composite. It has been used to produce about 30 vol.% of TiC and TiB2 reinforcement in aluminum. The initial composite has a high porosity, and requires highpressure consolidation to increase the density.

Hot Working

Hot working is the main fabrication route for wrought composites, regardless of the initial means of manufacture, because the high temperatures involved minimize the detrimental effect of the lower ductility exhibited by composites. The fabrication should be carried out at as high a temperature as possible, while avoiding any incipient melting, since this maximizes the ductility of the matrix, and minimizes the flow stress. A low flow stress is needed to avoid cracking the brittle reinforcing particles (Lloyd, 1991). The most difficult hot working process is rolling, since this is essentially a plane strain process where the unconstrained edges of the ingot are very susceptible to edge cracking. Edge cracking is greatly reduced if the composite is first extruded or closed die forged before rolling. The prior extrusion breaks up the as-cast structure, increasing its ductility and reducing the susceptibility to edge cracking. Extrusion and closed die forging are both constrained deformation and less susceptible to the formation of edge cracks. Extrusion is the most frequent fabrication route for wrought composites. Conventional aluminum direct extrusion uses a flat faced, shear die, and most commercial composite extrusions are produced using this type of die. However, the extrusion pressure can be reduced by using a streamlined die, which has no ªdead zoneº associated with the process, but has the disadvantage that it requires the extrusion billet to be premachined to fit the streamline profile of the die, and has a higher scrap level. Extrusion temperatures and pressures can be drastically reduced by using hydrostatic extrusion, which can be carried out at room temperature. In hydrostatic extrusion the extrusion ram does not contact the billet directly, but through a high-pressure fluid, which not only extrudes the composite, but also isolates it from the container wall, hence lowering frictional effects. The superimposed hydrostatic pressure also reduces the tendency for void formation and void growth, which improves the quality of the extrudate. The viability of the hydrostatic process

12

Melt Processed Aluminum Matrix Particle Reinforced Composites

Figure 10 Extrusion limit diagram for 6061 and 6061/Al2O3p (after Dixon and Lloyd, 1996).

for extruding a variety of composites has been demonstrated (Grow and Lewandowski, 1995). For conventional extrusion, the actual extrusion conditions required depend on the matrix alloy, volume fraction of reinforcement, and extrusion ratio, but as a guide the extrusion of 6061/Al2O3/20p can be considered, and typical conditions for this composite would be a billet temperature of 450±500 8C, with extrusion speeds of 7±15 m per minute. In conventional extrusion the main concerns are extrusion pressure, optimum extrusion speed, and die wear. At equivalent extrusion temperatures the extrusion pressure is about 10±15% higher than that of the unreinforced matrix, and the extrusion speed is limited by surface tearing due to local liquation (Brusethang et al., 1990) and the inherently lower ductility of composites. Satisfactory extrusion therefore occurs in the reduced region illustrated in an extrusion limit diagram (Jeffrey and Holcomb, 1990) (Figure 10). The high temperature limit is controlled by hot cracking of the matrix alloy while the higher flow stress of the composite results in a higher extrusion pressure and a shift in the pressure limit line. At the optimum extrusion billet temperature, the maximum extrusion speed for the composite, Smmc, is much lower than the equivalent maximum speed, Smat, for the unreinforced alloy. The extrusion process is important because, in addition to producing structural shapes, it produces a more homogeneous particle distribution, with some refinement of the particle size. Figure 11 shows the decrease in reinforcement particle size with increasing extrusion ratio in 6061/Al2O3p, with the particle size

decreasing rapidly with increasing extrusion ratio initially, but then remaining constant for extrusion ratios beyond about 20:1. Due to the high hydrostatic stress component during extrusion, there is little evidence of void formation associated with the particle fracture. While extrusion does not change the yield strength of the composite, it produces a more uniform distribution of reinforcing particles and this increases the ductility of the composite, as shown in Figure 12. One problem with extrusion is that, due to the abrasive nature of the reinforcing particles in the composite, die wear during extrusion is extensive (Dixon and Lloyd, 1996), and can limit the useful die life of a conventional steel die to as little as 100 m. For reasonable die life powder metallurgy tool steels, steel matrix TiC composites and bonded carbides can be used for simple profiles. More complex shapes may require chemical vapor deposition (CVD) coatings of TiC/TiN on conventional tool steels. 3.21.5.2

Casting

Since alloys with high castability can often be used as the composite matrix, shape casting can be used to obtain the final composite form, and the complete range of casting techniques can be used provided some process modifications are made to account for the different behavior of composite materials (Cox et al., 1993). Because of the danger of reaction between the melt and the reinforcing particles, the higher viscosity of composite melts, etc., there are several key steps which must be adhered to (most of the casting

Secondary Fabrication

13

Figure 11 Influence of extrusion ratio on reinforcement particle (after Dixon and Lloyd, 1996).

Figure 12 Influence of extrusion ratio on tensile elongation (after Dixon and Lloyd, 1996).

practices have been developed by Duralcan, but they are generally applicable): (i) Composite melts must be stirred. Since the density of the reinforcing ceramic particles is higher than molten aluminum, 3.2 g cm71 for SiC compared with 2.2 g cm71 for molten aluminum, the SiC particles will sink unless the melt is stirred. The stirring action should be just sufficient to cause visible movement on the surface, but not enough to cause vortex formation, which would entrap gas and oxide into the melt. The alternative to stirring is to use induction melting, where the natural eddy currents will keep the reinforcing particles in suspension. (ii) A critical temperature must never be exceeded. This temperature depends on the matrix alloy, and the reinforcement, but most foundry composites are Al±Si/SiC, with silicon levels of around 9 wt.%, and the upper temperature for this composite is 745 8C for normal

casting operations. If the melt suddenly becomes highly viscous it is most likely that the critical temperature has been exceeded, and reaction has occurred. (iii) Melt under an inert gas cover. Using a cover gas of dry argon or dry nitrogen minimizes oxidation of the melt. Degassing as a means of cleaning the melt can be done (Duralcan USA, 1993), and may be necessary with high levels of recycled foundry returns, but it should be avoided if possible by keeping the melt clean. Salt fluxing should not be used because it will result in dewetting of the reinforcing particles. The mechanical properties achieved in composite casting are critically dependent on the quality of the casting in terms of cleanliness and soundness. The higher viscosity of composite melts means that it is more difficult for any entrapped oxide films or gas bubbles to escape. Mold and gating configurations should be

14

Melt Processed Aluminum Matrix Particle Reinforced Composites

Figure 13 The ªInterrupted Flowº gating system (after Doutre et al., 1993).

designed to achieve as quiescent a fill as possible, and ceramic filters can be incorporated into the fill system to entrap oxides and unwanted inclusions. A ceramic filter with 10±15 pores per inch is effective without filtering out any of the reinforcing particles. Optimum gating systems have been developed to produce high-quality castings with high yields. One such system is the Interrupted Flow Gating System (Doutre et al., 1993), shown diagrammatically in Figure 13 for a bake rotor casting. The system consists of an insulating ceramic cone as the down sprue, with a plug covering a ceramic foam filter at the entrance to the mold. The pouring basin is large enough to hold the total cast weight. After filling the basin there is a delay of a few seconds before raising the plug and allowing the composite to flow into the mold. This delay is sufficient for any entrapped air to escape before the composite enters the mold. Utilizing the appropriate precautions, any of the conventional casting processes can be successfully used to cast composites, but new casting methods which are particularly suited to composites can also be developed. One example is the Lost Crucible Process (Cox et al., 1993), shown diagrammatically in Figure 14. The basic concept is to push molten composite into the mold cavity under very laminar flow by the use of a ram moving into the crucible, which is destroyed in the process. This casting method has several attractive characteristics: (i) absolute control of fill rate; (ii) slow fill rates possible due to the insulating sleeve; (iii) bottom filling to eliminate air engulfment; (iv) high initial pressure available for filter priming;

(v) potential for high casting yield; (vi) no need for coating or lubricants; (vii) can accept metal from any melting medium; (viii) mold stacking possible for increased throughput; (ix) can be used for lost foam, bonded sand, or permanent mold. Another casting method which produces low porosity shapes is Squeeze Casting, which is shown diagrammatically in Figure 15. The pressure is maintained throughout solidification, although it may not be at the same level through the complete cycle, and this pressure maintains good thermal contact and prevents pore formation. Good quality castings still require clean metal and a quiescent fill to avoid regions of premature freezing. The pressure needs to be carefully controlled to avoid segregation of both reinforcement and any low melting point eutectic, which is the last material to solidify.

3.21.5.3

Machining

Most parts are subjected to some form of machining prior to use, and again, the brittle ceramic reinforcement necessitates some special practices. By far the most cost-effective tool material for the production machining of composites is polycrystalline diamond (PCD) (Duralcan USA, 1991c); high-speed tool steels are dulled in seconds, and coated carbides only last a few minutes. Since abrasion is the main mode of tool wear, the more aggressive the cut the lower the tool wear. Tool life is also enhanced by decreased cutting speeds. Duralcan recommend a large grain size

Secondary Fabrication

Figure 14

The ªLost Crucibleº casting process (after Cox et al., 1993).

PCD tool, high feeds, and low speedsÐtypical conditions for roughing would be 51400 surface feet per minute, with a feed rate of >0.005 inch per revolution, and a depth of cut of >0.06 inches; equivalent conditions for finishing would be 52000 surface feet per minute, >0.005 inch per revolution, and >0.02 inch depth of cut. Drilling, reaming, and tapping should also use PCD tipped tools if available. The same principles of slow speed and high feed apply. Sawing can be achieved on band saws with flood cooling; for heavy sections a WC tipped blade with 53 teeth per inch running at 200± 300 surface feet per minute under medium to heavy cutting pressure is optimum, while for light sections a 60-grit industrial diamond blade gives the best performance, with a life of over a hundred hours. Industrial diamond tooling should also be used for grinding and honing.

3.21.5.4

15

Joining

Composites can be welded and brazed, but the reinforcement complicates the process (Duralcan USA, 1991a, 1991b). Because the weld pool in arc welding may contain up to 10 vol.% particulates, its fluidity is reduced, and one also has to be careful to avoid excessive reaction between the reinforcement and the molten matrix. These problems can be avoided by careful control of the welding conditions and appropriate choice of filler, and then conventional welding equipment can be used. For foundry composites, such as Al±Si/SiCp, gas tungsten arc welding (GTAW) is the preferred process for small surface defects, and gas metal arc welding (GMAW) for the repair of large defects in casting. The Al±12 wt.%Si filler alloy, ER4047, is recommended for foundry alloy

Figure 15 The ªSqueeze Castingº process (courtesy of M. Gallerneault, Queen's University, Kingston, ON).

composites to limit any aluminum carbide formation. For alumina reinforced wrought composites, such as 6061/Al2O3p, the same welding methods can be used but with the magnesiumcontaining filler wire ER5356, which reduces the tendency for particle dewetting and clumping in the weld pool region. The reduced fluidity of the weld inhibits arc penetration, and this may require a more open joint set-up. The setup also needs to take into account the reduced thermal conductivity and lower thermal expansion of composites. Conventional vacuum brazing methods can be used for composites, and will give joint strengths close to the base composite, although the elongation tends to be very low, 51%. The volume fraction of reinforcement in the joint region is much higher than in the bulk of the composite and this results in low ductilities. More development work is required to optimize the brazing process.

16

Melt Processed Aluminum Matrix Particle Reinforced Composites

Figure 16 Extents of particle settling in composites of different volume fractions of reinforcements (after Gallerneault, 1991).

3.21.6

MICROSTRUCTURES

The mechanical properties of composites are dependent on the distribution of the reinforcing particles, which is a function of the processing and fabrication routes involved. However, the other features of the microstructure which are important in the properties of unreinforced alloys, such as the grain structure, also need to be controlled in some composites.

3.21.6.1

Reinforcement Distribution

In composites processed in the molten state, the reinforcement distribution is influenced by several factors: (i) distribution in the liquid as a result of the mixing; (ii) distribution in the liquid after mixing but prior to solidification; (iii) redistribution as a result of solidification. The distribution during mixing will obviously depend on the mixing process used, and it is essential to produce as uniform a distribution as possible without any gas entrapment, since any gas bubbles will be lined with reinforcing particles. Factors affecting the clustering of particles, and the influence this can have on properties, are discussed in Chapter 3.06, this volume. After mixing and prior to solidification, the particles will segregate due to gravity (Lloyd and Chamberlain, 1988). With the relatively high volume fraction of particles and a range of particle size, the settling will be hindered:

vc ˆ vo …1 ÿ c†p

…11†

where vc is the particle velocity, vo is the Stokes' velocity, c is the particle concentration, p = 4.65 + 19.5(d/Dc) for Re50.2, p = (4.35 + 17.5 d/Dc) Re70.03 for 0.25Re51, d is the particle diameter, Dc is the container diameter, and Re is the Reynold's number, and the Stokes' velocity Vo is given by V0 ˆ 2d2 g…rp ÿ rl †=9Z

where g is the gravitational constant, rp is the density of the reinforcing particle, rl is the density of the liquid, and Z is the liquid viscosity. The settling rate will be a function of the particle density and size, and there is also the possibility that particle shape will play a role. Particles of different sizes and shapes will settle at different rates producing agglomeration, and this can produce particle clusters which settle as a cluster, rather than as individual particles. If typical values for 500-grit SiC particles in an AlSi melt are substituted in to the above equations the settling velocity is about 20 mm s71, which is slower than the typical dendrite growth velocity during solidification in most casting processes. In practice, however, extensive sedimentation occurs, indicating that the particles are settling as clusters. Figure 16 shows the variation in the particle free zone at the top of a melt for AlSi/SiCp composites with different volume fractions of reinforcement (Gallerneault, 1991), and knowing the rate of solidification in these experiments, the settling velocities for the reinforcement can be calculated. Figure 17 shows that the calculated set-

Microstructures

17

Figure 17 Calculated settling velocities in composites with different volume fractions of reinforcement (after Gallerneault, 1991).

Figure 18 Reinforcing particle distribution in a nondendritic microstructure.

tling velocities are much higher than the 20 mm s71 expected for single particles, consistent with the view that the particles are settling as clusters. It is also interesting to note that the settling behavior is independent of the matrix alloy. The third factor which influences reinforcement distribution is the solidification process itself. Reinforcing particles do not generally nucleate the primary solidifying phase (aluminum dendrites), although solidification nucleation may occur in some hypereutectic systems (Jin and Lloyd, 1990). If solidification nucleation does not occur, the reinforcing particles are rejected at the solid±liquid interface, and segregate to the interdendritic regions which solidify last. In composites made by semisolid mixing, the primary aluminum grains solidify nondendritically as a result of the imposed shear, and the reinforcing particles segregate to the inter-

stices between the spheroidal cells, as in Figure 18. If the mixing is done in the fully molten state, as in the Duralcan process, the particle distribution is influenced by the solidification rate, as shown in Figure 19, since the solidification rate controls the cell size. Thus, particle segregation due to particle rejection at the solidification front can be minimized by using a high solidification rate casting process such as twin roll casting. The composite sheet resulting from this process exhibits a higher degree of homogeneity of the reinforcing particles, and hence an improved tensile ductility (Lloyd and Jin, 1993). Secondary fabrication processing, such as extrusion or rolling, can homogenize the structure to some extent, as seen from the extrusion microstructure in Figure 20, and this increases the tensile elongation of the composite, as noted previously in Figure 11.

3.21.6.2

Grain Structure

Wrought alloy composites are solution treated and aged after fabrication, and recrystallization will usually occur during this heat treatment. Since particles of diameter larger than about 1 mm will develop an associated deformation zone sufficient to generate recrystallized nuclei (Humphreys, 1977), the reinforcing particles should produce a high density of nuclei. However, if the particles are very closely spaced, the subgrain growth necessary for recrystallization nucleation is impeded and

18

Melt Processed Aluminum Matrix Particle Reinforced Composites  1=3 D ˆ d …1 ÿ Vp †=Vp

Figure 19 Influence of solidification rate on reinforcing particle distribution: (a) slow solidification rate investment casting; (b) high solidification rate pressure die-casting (after Lloyd, 1994).

Any subsequent grain growth will be limited by Zener pinning of the particles on the grain boundaries, giving a limiting grain size of 2d/ 3Vp. For 20 vol.% of 10 mm particles, Equation (12) gives about 15 mm, and the limited grain growth gives 33 mm. Recrystallization studies on Duralcan 2014/Al2O3/20p demonstrated that the Al2O3 particles stimulate the nucleation, accelerating recrystallization and decreasing the recrystallization temperature (Ferry et al., 1991). The resulting equiaxed grain size was &15 mm, and after holding the composite for 150 h at 500 8C the grains had only grown to 17 mm. These results show that Equation (12) is appropriate for predicting both the recrystallized and limiting grain size, and the high volume fraction of reinforcing particles is very effective for stabilizing the grain size. In powder processed composites, reasonable agreement was obtained with the theory except for very coarse particles (40 mm), where the grain size was much finer than expected due to multiple nucleation occurring at each particle (Miller and Humphreys, 1990). The yield strength is dependent on grain size through the Hall±Petch relationship, and the finer grain size present in composites could contribute considerably to the strength. Aluminum alloys have a low Hall±Petch slope, and grain sizes of 10 mm and less will be required to significantly influence the strength, but in reinforced wrought magnesium alloys grain size strengthening could be very significant because the Hall±Petch slope is large (Nussbaum et al., 1989).

3.21.7

Figure 20 Particle distribution in an extrusion (after Lloyd, 1994).

recrystallization may not occur. Recrystallization is impeded when Vp/d>0.1 mm71, where Vp is the particle volume fraction and d is the particle diameter. Most commercial composites have Vp/d less than this value. The grain size of the recrystallized composite can be estimated by assuming that each reinforcing particle of diameter d acts as a nucleus for a spherical grain D, which is given by

…12†

MECHANICAL PROPERTIES

Various aspects of the mechanical properties of metal matrix composites are considered in other chapters of this volume, but it is appropriate to briefly consider how the properties are related to different aspects of the composite. Increased strength is rarely a factor in the commercial applications of particle reinforced composites, because their increased strength over the unreinforced matrix is relatively small. Of more commercial importance is the increased elastic modulus. Figure 21 shows the increase in Young's modulus with volume fraction of rein-forcement for a variety of Al±SiC and Al±Al2O3 composites. The Rule of Mixtures expression: Ec ˆ Vp Ep ‡ Vm Em

…13†

where Ec, Ep, and Em are the moduli of the

Mechanical Properties

19

Figure 22 The temperature dependence of Young's modulus (after Lloyd, 1994). Figure 21 The variation of Young's modulus with the volume fraction of reinforcement (after Lloyd, [1994).]

composite, particulate, and matrix, respectively, and Vp and Vm are the respective volume fractions, considerably overestimates the modulus. The Rule of Mixtures expression is most appropriate for continuous reinforcement and it has been modified for discontinuous reinforcement in the Halpin±Tsai equation (Halpin, 1984): Ec ˆ Em …1 ‡ 2sqVp †=…1 ÿ qVp †

…14†

where …15†

Figure 23 Aging curve at 175 8C for 6061, 6061/ Al2O3/10, and 20p.

and s is the aspect ratio of the particulate. As seen from Figure 21, the Halpin±Tsai equation gives a good representation of the results. The modulus can also be calculated using the Eshelby equivalent inclusion method (Clyne and Withers, 1993), and this approach is also in good agreement with the data. The dominant factor in controlling the modulus is the volume fraction of reinforcement, and it is relatively insensitive to the particle distribution, while variations in the type and shape of the reinforcement can be accounted for by the different expressions. The improvement in Young's modulus is retained at higher temperatures, as shown in Figure 22, where results for unreinforced matrix and composite are shown up to 500 8C. The temperature dependence of the composite modulus reflects that of the unreinforced matrix, but the larger modulus at elevated temperature is important for some higher temperature applications. While the strength increase exhibited by particle reinforced composites is relatively

small, they do respond to precipitation hardening in the heat treatable aluminum alloys. The aging kinetics may be slightly faster than in the unreinforced matrix in some composites (Lloyd, 1994 and Chapter 3.03, this volume), but the differences are small, and in practice the conventional heat treatment times can be used with little loss in properties. Figure 23 shows the aging curve at 175 8C for unreinforced 6061 and 6061/Al2O3p, where the strengthening contributed by the reinforcement is reflected in the yield stress increase at zero time, i.e., T4 temper, and the strengthening due to matrix precipitation increases with increasing time. The reinforcement strengthening increases with increasing volume fraction, and, for the same volume fraction, the yield stress increases with decrease in reinforcing particle size (Lloyd, 1994). The increased strength of composites is accompanied by a decrease in tensile elongation, as shown in Figure 24 for the same

q ˆ …Ep =Em ÿ 1†=…Ep =Em ‡ 2s†

20

Melt Processed Aluminum Matrix Particle Reinforced Composites

Figure 24 Variation in tensile elongation with aging at 175 8C for 6061, 6061/Al2O3/10, and 20p.

Figure 25 Variation of tensile elongation with the strength of the matrix for Al2O3 reinforced composites.

6061/Al2O3 p composites as in Figure 23. The ductility is a strong function of particle distribution, as shown by the influence of the extrusion ratio in Figure 11. In a homogeneous composite the ductility can be maximized by using only the minimum volume fraction of reinforcement required for the particular application, and using a high-ductility, low-strength matrix. This is demonstrated in Figure 25 where the variation in elongation with strength for different Al2O3 reinforced composites is shown.

3.21.8

DEFECTS

The mechanical properties of composites are often dictated not by the inherent properties of the composite system but by defects of various types introduced during different stages of manufacture. The issues of uniformity of reinforcement distribution have been considered in Section 3.21.6.1, and in Chapter 3.07, this volume. Coarse agglomerates of reinforcing particles, particularly if they are unwetted, will severely degrade strength and ductility. For all

References processing methods it is important to avoid the incorporation of oxide skins into the composite, since they provide planes of weakness for easy fracture. In liquid metal processing, limiting the turbulence at the surface of the melt is important for avoiding this defect. Porosity is another common defect in composites. In liquid metal processing it is usually due to gas entrained in the melt during the turbulence associated with mixing or infiltration. However, in shape castings porosity may also result from poor mold filling or solidification shrinkage typically associated with shape casting. Casting methods to minimize these defects are discussed in Section 3.21.5.2 In general, it is difficult to achieve the levels of cleanliness typical of monolithic alloys, both in terms of gas and inclusion contents, because conventional degassing and filtering technologies are not as easily applied to composites. Since remedial action is difficult, it is essential to avoid generating the potential defects at any stage in the processing.

3.21.9

SUMMARY

The fundamentals of particulate reinforced composite processing, the evolution of the processing techniques, the microstructural development, and the mechanical properties obtained in the composite materials have been reviewed. The scientific knowledge and processing techniques in each area have advanced significantly in recent years, and we are now at the stage where a range of low-volume fraction composites (>25 vol.%) are available in full commercial scale quantities both for foundry and wrought products. For high-volume fraction composites (525 vol.%), powder metallurgy and other processing techniques have been continuously evolving. However, commercialization in this area is limited only to several specific applications. At present, commercialization is progressing mainly in those applications where a higher modulus and wear resistance are used. Examples of these applications are brake rotors, cylinder liners, tire studs, bicycle frames, and drive shafts. The major barriers to the full business potential are the high cost of the composite materials and the limited capability of elevated temperature applications. Thus future research should focus on key factors which are directly related to the above problems. These may include various secondary processing technologies, recycling, nondestructive testing, and the design of new interfaces for improved wetting, matrix±reinforcement compatibility, and reactivity control.

21

ACKNOWLEDGMENTS The authors are grateful to Alcan International Limited for permission to publish, and to their colleagues at the Alcan Research and Development Centers at Kingston, Arvida, and Banbury, and at Duralcan for providing information for this chapter.

3.21.10

REFERENCES

M. K. Aghajanian, M. A. Rocazella, J. T. Burke and S. D. Keck, J. Mater. Sci., 1991, 26, 447±454. Alcan Aluminum Corporation, US Pat., 4 786 467 (1988). F. A. Badia and P. K. Rohatgi, Trans. AFS, 1969, 77, 402. A. Banerji, M. K. Surappa and P. K. Rohatgi, Metall. Trans., 1983, 14B, 273±283. J. B. Borradaile, S. Skjervold and W. Ruch, in `Proceedings of Extended Abstracts of Conference on Metal Matrix CompositesÐProperty Optimisation and Applications', Paper 10.1, The Institute of Metals, London, 1989. S. B. Brown and M. C. Flemmings (eds.), in `Proceedings of 2nd International ConferenceÐProcessing of SemiSolid Alloys and Composites', TMS, Warrendale, PA, 1992. S. Brusethang, O. Reiso and W. Ruch, in `Proceedings of Symposium±Fabrication of Particulates and Reinforced Metal Composites', eds. J. Masounave and F. G. Hamel, ASM International, Materials Park, OH, 1990, pp. 181± 186. R. D. Carnahan, R. F. Decker, N. Bradley and P. Frederick, in `Proceedings of Symposium±Fabrication of Particulates and Reinforced Metal Composites', eds. J. Masounave and F. G. Hamel, ASM International, Materials Park, OH, 1990, pp. 101±105. S. Caron and J. Masounave, in `Proceedings of SymposiumÐFabrication of Particulates and Reinforced Metal Composites', eds. J. Masounave and F. G. Hamel, , 1990, pp. 107±113. T. W. Chou, A. Kelly and A. Okura, Composites, 1985, 16, 187±206. T. W. Clyne and P. J. Withers, `An Introduction to Metal Matrix Composites', Cambridge Solid State Science Series, Cambridge University Press, London, 1993. J. A. Cornie, H.-K. Moon and M. C. Flemings, in `Proceedings of SymposiumÐFabrication of Particulates and Reinforced Metal Composites', eds. J. Masounave and F. G. Hamel, ASM International, Materials Park, OH, 1990, pp. 63±78. B. M. Cox, D. Doutre, P. Enright and R. Provencher, in `Proceedings of 2nd International Conference on Cast Metal Matrix Composites', Tuscaloosa, AL, 1993, pp. 88±105. F. Delannay, L. Froyen and A. Deruyttere, J. Mater. Sci., 1987, 22, 1±16. Deonath and P. K. Rohatgi, J. Mater. Sci., 1980, 15, 2777±2784. J. Dinwoodie, E. Moore, C. A. J. Langman and W. R. Symes, in `Proceedings of 5th International Conference on Composite Materials', San Diego, CA, ed. W. C. Harrigan, et al., The Metallurgical Society of AIME, Warrendale, PA, 1985, pp. 671±685. A. P. Divecha, S. G. Fishman and S. D. Karmarkar, J. Metals, 1981, 33, 12±17. W. Dixon and D. J. Lloyd, in `Processing, Properties and Applications of Cast Metal Matrix Composites', ed. P. K. Rohatgi, The Minerals Metals and Materials Society, PA, 1996, pp. 259±269.

22

Melt Processed Aluminum Matrix Particle Reinforced Composites

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Comprehensive Composite Materials ISBN (set): 0-08 0429939 Volume 3; (ISBN: 0-080437214); pp. 555±577