Fracture toughness and wear resistance of aluminum-boron particulate composites cast using metallic and non-metallic chills

Materials and Design 23 Ž2002. 41᎐50

Fracture toughness and wear resistance of aluminum-boron particulate composites cast using metallic and non-metallic chills Joel HemanthU Department of Mechanical Engineering, Siddaganga Institute of Technology (S.I.T.), Tumkur, 572-103 Karnataka, India Received 20 February 2001; accepted 14 May 2001

Abstract This paper describes the fabrication and testing of Al-boron particulate composites cast in sand moulds containing metallic Žcopper, steel and cast iron. and non-metallic Žsilicon carbide. chills, respectively. The size of the boron particulates is between 30 and 100 ␮ m. The boron contents are 3, 6 and 9% by weight. The superior mechanical properties of the castings, particularly their ultimate tensile strength ŽUTS., hardness, wear resistance and fracture toughness, are discussed in relation to their microstructure. 䊚 2002 Elsevier Science Ltd. All rights reserved. Keywords: Chill; MMC; Boron; Fracture toughness; Wear; UTS

1. Introduction Reinforcing an Al alloy with boron particulates yields a material that displays physical and mechanical properties of both the metal matrix and the boron. For example, the toughness and formability of Al can be combined with the strength of the boron particles. On a weight adjusted basis, many Al-base composite materials can outperform cast iron, Al, Mg and virtually any other reinforced metal or alloy in a wide variety of applications. Hence, it seems probable that metal matrix composites ŽMMCs. will replace conventional materials in many commercial and industrial applications in the near future w1᎐5x.

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Tel.: q91-816-282-990; fax: q91-816-282-876. E-mail address: [email protected] ŽJ. Hemanth..

Fabrication of the discontinuously-reinforced Albased MMCs can be achieved by standard metallurgical processing methods such as powder metallurgy, direct casting, rolling, forging and extrusion, and the products can be shaped, machined and drilled using conventional facilities. Thus, they can be made available in quantities suitable for automobile applications. In general, the primary disadvantages of some MMCs for automobile applications are their low hardness and inadequate or poor fracture toughness and fatigue performance compared to those of the constituent matrix material w6,7x. The superior properties offered by particulate-reinforced Al based MMCs make these materials attractive for applications in the automobile, industrial, aerospace, defense and leisure industries. The wear resistance of unreinforced Al alloys is relatively poor. Hence, their applications as structural and automobile parts are often limited. Fortunately,

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the strength, fracture toughness and the wear resistance can be improved remarkably by the incorporation of hard particulates such as boron.

perature gradient in the casting during solidification w27x. 2.3. Volumetric heat capacity (VHC) of chills

2. Literature review 2.1. Freezing of Al alloys It is well known that Al alloy-based composites freeze over a wide range of temperatures and are difficult to feed into a mould cavity during solidification w8x. The dispersed porosity caused by the pastry type of solidification of long freezing range alloy castings can be effectively reduced by the use of chills. Chills extract heat from the melt more rapidly and promote directional solidification. Therefore, they are widely used by foundrymen in the production of quality castings w9᎐15x. With the increase in the demand for quality composites, it has become essential to produce Al-based composite free from micro-porosity. Al-based composite castings are prone to defects in the form of microshrinkage or dispersed porosity, which can be minimized by the judicious location of chills w16x. In spite of the increased applications of chills on the Al alloy founding, there are currently no data available on the action of chills on the mechanical properties of Al-boron composites although there has been some research done on the influence of chills on the solidification and soundness of long freezing range alloy castings w17᎐25x. Furthermore, Redemske et al. w26x did point out from their investigation that chilling has an effect on the structure, properties and soundness of Al alloy castings. 2.2. Applications of chills The casting process is the most flexible, versatile and usually the simplest and least expensive among several feasible manufacturing processes. The primary and most important aspect to be considered in a casting process is the solidification shrinkage of the casting because it contributes substantially to the problems faced during the feeding of castings. In the case of sand moulds, pasty zones pose severe problems, the solution to which lies in establishing a steep temperature gradient during solidification. This can be done by incorporating heat sinks in the form of chills located at pre-determined places in the mould. This is the most popular and simplest way of establishing a steep tem-

The ability of the chill to extract heat from the molten metal during freezing of the casting is dependant on the chill size and the properties of the chill material. The capacity of the chill to absorb heat from the casting is taken as a measure of its efficiency. The VHC of the chill, which takes into account the volume, specific heat and density of the chill material, has been identified as an important factor in evaluating the efficiency of the chill. By definition: VHCs V = Cp = ␳ Where V s Volume of the chill, Cp s specific heat of the chill material and ␳ s density of the chill material. It is obvious that an increase in any one of the factors, V, Cp or ␳ increases the value of VHC, thus enhancing the chilling. On placing a chill in the mould, heat transfer takes place through the metal-chill interface during the initial stages of solidification. Subsequently, an air gap forms between the casting and the chill face due to dissimilar thermal behavior of the chill and the surface of the thin layer of solidified metal. This air gap results in a temporary delay in the heat transfer. In the final stages of solidification, heat transfer takes place again, unobstructed when there is direct contact between the castings and chill surface. From the existing literature, it is evident that no investigation has been carried out so far regarding the influence of metallic and non-metallic chills on the mechanical properties of AlrBoron particulate composites. The present investigation is intended to fill the void.

3. Experimental procedure 3.1. Chemical composition of MMCs The chemical composition of the Al alloy used in the matrix material is given in Table 1. In this investigation, boron particles of 3, 6 and 9 wt.% are dispersed in the matrix. The density of boron is 2.3 grcm3, melting point 2300⬚C, Youngs modulus 440 GPa and its crystal structure is ortho-rhombic. The size of the boron varies from 30 to 100 ␮ m.

Table 1 Chemical composition of matrix alloy Elements Composition wt.%

Zn 3.01

Mg 3.00

Si Trace

Cu Nil

Fe 0.001

Al 93.989

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Fig. 1. Sand mould for casting Al-boron MMCs with chill.

3.2. Preparation of test specimens After melting the matrix material in a furnace at approximately 750⬚C in an inert atmosphere, boron particulates pre-heated to 600⬚C were introduced evenly into the molten metal alloy by means of special feeding attachments. Meanwhile, the molten MMC was well agitated by means of a mechanical impeller rotating at 760 rev.rmin to create a vortex. The melt was next poured into a sand mould which allowed for different types of chills to be attached to it at one end, as shown in Fig. 1. The moulds produced are plate shaped ingots of dimensions 225 = 150 = 25 mm. The former were prepared using silica sand with 5% bentonite as binder and 5% moisture according to American Foundrymen Society ŽAFS. standards, and were dried in an air furnace. Ingots were cast employing different chills Žmade of copper, steel, cast iron and silicon carbide, respectively. in order to study the effect of VHC of the chill on ultimate tensile strength ŽUTS., hardness, fracture toughness, wear behavior and microstructure of the composite. Unless otherwise stated, all the chills were 170-mm long, 35-mm high and 25-mm thick. The tensometer specimens for the strength tests were prepared according to AFS standards. The specimens were taken from various locations in the casting, namely, at the chill end and 75, 150 and 225 mm, respectively, from the chill end, the latter case being at the riser end. The longitudinal axes of these specimens were parallel to the longitudinal axis of the chill during casting. Specimens for wear testing were selected only at the chill end of casting. The dimension of each specimen are 50 = 60 = 20 mm, with the largest face parallel to the adjacent chill surface. All the specimens were heat

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treated by aging before wear testing. Aging was done at a fixed temperature of 180⬚C for duration of 1, 2, 3, 4 and 5 h, respectively. The purpose of this aging was to promote nucleation and to allow the Al matrix to undergo precipitation, a time-dependent process. For Al alloys, an aging temperature of 180⬚C and an aging duration of 4 h are ideal for the material to reach maximum values in its mechanical properties. Such heat treatment accelerates the precipitation process during which the dislocation proliferates near the particulatermatrix interface. These high-density dislocations act as heterogeneous nucleation sites for further precipitation, a process, which enhances the hardness and wear resistance of the MMC. To study the effect of the rate of chilling on the properties of the MMCs, the procedure was repeated with chills of the same shape and size made of steel, cast iron and silicon carbide, respectively. 3.3. Microstructural examination Microscopic examinations were conducted on all the specimens using SEM as well as a Neophot-21 metallurgical microscope. Various etchants were tried but dilute Keller’s reagent proved to be the best and was therefore used. Photomicrographs of all the specimens were taken to study their micro-constituents and the distribution of the boron particulates. SEM photographs were also taken of all the worn surfaces after testing to study the wear mechanism. 3.4. Tensile test Tensile tests were performed using an Instron tensile testing machine on AFS standard tensometer specimens. Each test result reported in this paper is the average obtained from at least three test specimens taken from the same location in the mould and cast under identical conditions. 3.5. Fracture toughness test Fracture toughness test were performed using a closed-loop Instron servo-hydraulic material testing system in accordance with ASTM E 399-1990 standards. The method of testing involved the three-point bend testing of notched Žchevron type. specimens, which had been pre-cracked by fatigue. The fatigue test specimens were smooth and rectangular in section, their machined surfaces ground using 600 grit silicon carbide and mechanically polished to remove scratches and machining marks. Fully reversed push᎐pull, total strain-performed, tension᎐compression Ž R s y1. fatigue tests were performed. The tests were performed in a controlled laboratory air environment Žtemperature 26⬚C, relative humidity 56%.. An axial 12.5-mm

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Fig. 2. Schematic diagram of abrasive wear tester.

mouth chip gauge extensometer was attached to the gauge section of the test specimen to measure the total strain. The stressrstrain hysterisis loops were recorded on an X-Y recorder. From the load, the stress intensity factor K 1C Žwhich is a measure of the fracture toughness of the material. was calculated using equations, which have been established on the basis of elastic strain analysis. The validity of this method depends on the establishment of a steep crack condition at the tip of the crack in a specimen of adequate size. All these conditions were fulfilled in this experiment. 3.6. Wear test procedure Dry wear tests were conducted at room temperature according to ASTM standards using the spindle-time wear tester manufactured by Riken-Ogoshi and Co., Korea, shown schematically in Fig. 2. The surface of each test specimen, which was adjacent to the chill during casting, was polished using 600-grit emery paper before being subjected to the wear test. Oil-quenched SCM4 Žequivalent to AISI 4140. steel was used as the material for the rotating disc. An electronic balance capable of weighing to an accuracy of 10y5 g was used to measure weight loss in the specimens. Wear test were performed at a final load of 14.8 Kgf, a fixed sliding distance of 400 m, and at four sliding speeds of 0.52, 1.15, 1.98 and 3.64 mrs, respectively. Hardness tests were performed using a Rockwell hardness tester.

4. Results and discussion 4.1. Heat treatment In ceramic-reinforced MMCs, there is generally a big difference between the mechanical properties of the dispersoid and those of the matrix. This results in incoherence or high density of dislocations near the interference between the dispersoid and the matrix. Precipitation reactions are accelerated because incoherence and the high density of dislocations act as heterogeneous nucleation sites for precipitation w28᎐30x. Therefore, if all other factors are kept constant, the aging rate of a composite is generally faster than that of the matrix alloy. After solution treatment, optimum aging conditions can be determined by observing the hardness of the MMCs cast with copper chills for different aging duration, as tabulated in Table 2. It is known that the optimum aging conditions are Table 2 Effect of aging duration on hardness for MMCs cast using copper chill and aged at 180⬚C Aging time Žh.

Hardness ŽRockwell B scale. 3 wt.% boron

6 wt.% boron

9 wt.% boron

1 2 3 4 5

76 78 79 81 77

80 81 83 89 81

79 80 83 85 82

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strongly dependent upon the amount of dispersoid present w28x. From Table 2, it can be seen that for each MMC, as the aging time increases, the hardness of the MMCs increases to a peak value and then drops again. As boron content is increased, there is a tendency for the peak aging time to be reduced because dispersoids provide more nucleation sites for precipitation. As expected, for any fixed aging temperature and duration, increasing the boron content causes the hardness of the MMC to increase since the boron particulates are so much harder than the Al matrix.

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Fig. 4. Microstructure of Al-boron MMC Ž6 wt.% boron. cast using copper chill.

4.2. Microstructure A microstructural study of test specimens was carried out using a VG-9000 scanner. The microstructural examination and fractographical analysis revealed good distribution of boron particles throughout the matrix. The microstructure of the chilled Alrboron composites containing 3 and 6 wt.% boron cast using copper chills are shown in Figs. 3 and 4, respectively. These photomicrographs show that the boron particles are of nearly uniform size and are uniformly dispersed in the aluminum matrix. However, microstructural studies reveal that, Mg migrated to the grain boundaries. This migration of alloying elements into the grain boundaries leaving behind the dispersoids in the grains result in a higher concentration of boron within the grains, which may be one of the main reasons for the increase in strength and soundness of the composite developed, as will be described below.

the specimen from the chill end of the casting for MMCs containing 6 wt.% boron. For a particular chill, UTS can be seen to decrease with distance from the chill up to a distance of approximately 150 mm, beyond Table 3 UTS of Al-boron composites with various boron contents cast using different chills Žchill thickness, 25 mm. Type of chill

UTS,MPa Žat the chill end. 3 wt.% boron

Copper Steel Cast iron Silicon carbide

116 114 102 97

6 wt.% boron 134 128 125 119

9 wt.% boron 129 122 119 112

4.3. Ultimate tensile strength (UTS) Table 3 tabulates the UTS of specimens taken from the chill end of the casting for composite cast using different types of chills. It is evident from the table that for a particular chill, the UTS of the composite increases as boron content is increased up to 6 wt.% by weight, beyond which it drops again. There is therefore, no advantage in reinforcing the Al matrix with boron contents above 6 wt.% as far as UTS is concerned. Fig. 5 shows a plot of UTS vs. distance of the location of

Fig. 3. Microstructure of Al-boron MMC Ž3 wt.% boron. cast using copper chill.

Fig. 5. Plot of UTS of MMC Ž6 wt.% boron. vs. distance from chill end for various types of chills.

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4.4. Fracture toughness

Fig. 6. Plot of UTS of MMC Ž6 wt.% boron. vs. VHC for various types of chills.

which it increases again. In both Table 3 and Fig. 5, it can be seen that if the location of a specimen is kept constant, of the four types of chills used, the copper one produces specimens of the highest UTS, followed by steel, cast-iron and SiC chills, in that order. This shows that as heat conductivity of the chill increases so does the UTS of the specimen. In addition to these, more specimens were later tested which were cast using chills of varying thickness. If the UTS of every specimen tested is plotted against VCH as shown in Fig. 6, all the points fall within a narrow linear band. It can be inferred that the effect of increasing the VHC of the chill is to increase the UTS of the MMC. In other words, if all other factors were kept constant, the faster the heat extraction from the molten MMC during casting, the higher would be the UTS of the cast MMC. The relationship is roughly linear.

The manner in which stress response varies with the number of cycles and the plastic-strain amplitude is an important feature of the low-cycle fatigue process. The cyclic stress required for pre-cracking the specimen provides useful information pertaining to the mechanical stability of the intrinsic microstructural features during reverse plastic straining. This and the ability of the material to distribute the plastic strain over the entire bulk of the material are the two key factors governing the cyclic response of the material w31x. The experimental results of the fracture toughness tests done on specimens taken from the chill end for composites cast using different types of chills are shown in Table 4. They are represented graphically in Fig. 7. From these results, it can be seen that changing the type of chill does have a profound effect on the fracture toughness of the material. In fact, increasing the rate of chilling by increasing the VHC of the chill tends to result in an increase in the fracture toughness of the material. It can be seen that if all other factors are kept constant, copper chilled castings invariably have the highest fracture toughness followed by steel, cast iron and silicon carbide chilled castings, in that order. This agrees with the deductions made by the author in an earlier paper w32x that increasing the rate of chilling tends to increase in fracture toughness of the material. Another factor affecting the fracture toughness of a specimen is its boron content. The fracture toughness is seen to increase up to approximately 6 wt.%, beyond which the fracture toughness drops again. There is therefore, no advantage in adding more than 6 wt.% of boron to the matrix material since it would cause the MMC to become more brittle. Fig. 8a,b shows the SEM photographs of fracture surfaces of typical specimens cast using copper chills and Fig. 9a,b shows those of typical specimen casts using steel chills. Of all the various types of chills used in this research, these two types provide the highest rate of heat extraction during casting. The photographs reveal a ductile mode of fracture, accompanied by isolated micro-cracks in the matrix as shown in Fig. 8a,b. Large areas of the fracture surface were covered with a bimodal distribution

Table 4 Fracture toughness of Al-boron composites with various boron contents cast using different chills, Žchill thickness 25 mm. Type of chill

Copper Steel Cast iron Silicon carbide

Fracture toughness Ž K1C ., Mpa 6m Žat the chill end. 3 wt.% boron

6 wt.% boron

9 wt.% boron

10.82 9.21 9.02 6.71

12.87 10.76 10.31 7.42

11.62 10.61 9.96 6.81

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formed mainly at the particle-matrix interface. However, growth of the voids was limited by the completing and synergistic influences of the brittleness of the reinforcing boron particles and the cyclic ductility of the matrix material. In contrast, specimens cast without chills displayed low strength with fracture essentially normal to the stress axis. The mechanisms, which control the variation of fracture toughness of chilled Al-boron composites, are dependant on both microstructure and strain range. The possible micro-mechanisms controlling the fracture behavior during cyclic loading are ascribed to the following synergistic influences: 1. load transfer between the soft Al matrix and the hard brittle boron particulate reinforcement; 2. hardening arising from constrained plastic flow and tri-axiability in the Al matrix due to the presence of the brittle boron reinforcements. As a direct result of the particles resisting the plastic flow of the matrix, especially in chilled composites, an internal stress or back stress is created; and 3. residual stresses generated in the Al matrix arising from the mismatch in thermal expansion coefficients between the soft matrix and the hard reinforcement particulates. Fig. 7. Plot of fracture toughness of MMC vs. boron content for various types of chills.

of dimples, indicative of ductile rupture as shown in Fig. 9a,b. A dimple is a half void. The voids were

During cyclic deformation it seems possible that the mismatch that exists between the brittle reinforcing particle and the ductile matrix favors concentration of stress at and near the particle-matrix interface, causing

Fig. 8. Ža. Fractograph of Al-boron chilled composite Ž3 wt.% boron. cast using copper chill. Žb. Fractograph of Al-boron chilled composite Ž6 wt.% boron. cast using copper chill.

Fig. 9. Ža. Fractograph of Al-boron chilled composite Ž3 wt.% boron. cast using steel chill. Žb. Fractograph of Al-boron chilled composite Ž6 wt.% boron. cast using steel chill.

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the matrix in the immediate vicinity to fail permanently or the particle to separate from the matrix. In addition, the improvement in fracture toughness when chills are employed during casting can also be attributed to the refining of the grain structure of the matrix. 4.5. Wear beha¨ ior Fig. 10 shows the effect of sliding speed of the wear of Al-boron MMCs containing 3% by weight of boron tested under a final load of 14.8 kgf with a sliding distance of 400 m at sliding speeds of 0.52, 1.15, 1.98 and 3.64 mrs, respectively. All these MMCs have been aged for 4 h at 180⬚C, which are optimum aging conditions. Each curve plotted shows the results obtained for MMCs cast using a particular type of chill, as indicated in the legend in Fig. 10. Although only four sliding speeds were used, there are seven points for each curve plotted because the MS Excel package used for plotting the graphs interpolated three additional points for curve fitting. 4.5.1. Low sliding speed As the sliding speed increases from 0.52 to 1.98 mrs, the wear resistance of the composite improves remarkably. The major wear mechanisms of the chilled com-

Fig. 10. Effect of sliding speed on wear of Al-6 wt.% boron MMCs cast withvarious chills.

posite are abrasive and adhesive wear at sliding speeds of up to 1.15 mrs, during which the MMCs is worn by the frictional force on the wear surface w33x. Above the sliding speed of 1.15 mrs, however, the major wear mechanism changes to melt wear because of the rise in temperature on the localized wear surface, resulting in the MMCs being worn out less rapidly. The high thermal stability of the MMCs during this stage is due to the presence of boron particulate, and is further enhanced by chilling. The above wear behavior can be explained based on microstructural analysis as follows. Fig. 11 shows typical photograph of the wear surfaces of Al-boron chilled composite containing 3 and 6% boron, chilled using a copper chill at a sliding speeds of 0.52 and 1.98 and 3.64 mrs tested under a final load of 14.5 kgf and with a sliding distance of 400 m. At a low sliding speed of 0.52 mrs, it was found that abrasive wear was dominant in the ploughing and grooving as indicated by the arrows in Fig. 11a,b. The number of stripes on the wear surface seems to progressively increase with sliding speed because of the increase in the abrasive action between the wear specimen and the counter material due to the increase in frictional force on the wear surface. The surface damage in the MMC with 3 wt.% boron is more severe than that in the MMC with 6 wt.% boron. At low sliding speed of 0.52 mrs, the MMC with 6% boron has grooves formed by the shearing action of the friction on the wear surface. In fact, even the wear surface of the MMC with 3 wt.% boron contains a number of wear grooves. It was also perceived that even at such a low sliding speed, frictional forces on the MMCs are large enough to fracture and pull out dispersoids from the matrix and these particulate abrade the wear surface of the composite. 4.5.2. Intermediate sliding speed The wear surfaces of the MMCs tested at an intermediate sliding speed of 1.98 mrs are as shown in Fig. 11c,d. The wear debris and distorted surface of the MMC with 3 wt.% boron are distinct than those in MMC with 6 wt.% boron. From the wear surface, it can be seen that wear debris and distorted surface play an important role in the wear, which increases with their formation and growth. The distorted surface and debris are found at the locally fractured area of the matrix alloy, as indicated by the arrows. These localized fractures are caused by highly localized frictional forces between the non-uniform surface of the counter material and the defective areas on the wear surface. The growth of the damage section is aggravated by the fracture of the wear surface. The wear surface of the MMC with 3 wt.% boron shows that abrasive and adhesive wear are the dominant wear mechanisms at this intermediate sliding speed. By contrast, the wear

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Fig. 11. Wear surfaces of Al-boron MMCs cast with copper chill under the following conditions: Ža. sliding speed of 0.52 mrs for MMC containing 3 wt.% boron; Žb. sliding speed of 0.52 mrs for MMC containing 6 wt.% boron; Žc. sliding speed of 1.98 mrs for MMC containing 3 wt.% boron; Žd. sliding speed of 1.98 mrs for MMC containing 6 wt.% boron; Že. sliding speed of 3.68 mrs for MMC containing 3 wt.% boron; and Žf. sliding speed of 3.68 mrs for MMC containing 6 wt.% boron.

surface of the MMC with 6 wt.% boron is completely different in that on this surface, abrasive wear is hardly seen. This is due to the reduction in frictional forces on the wear surface. Fig. 11d shows a boron particle being pulled out and being stuck on the wear surface, as indicated by the arrow. 4.5.3. High sliding speed Above the sliding speed of 1.98 mrs, the wear of the composite increases severely, indicating the significant reduction in wear resistance. This change in wear behaviour of the composite with sliding speed shows that the major wear mechanism of the composites developed by chilling are strongly dependent upon the sliding speed. At a high sliding speed of 3.68 mrs, adhesive and slip phenomena also appear. The removal of the material seems to be accelerated by fractures of boron particulates and the matrix which might be due to high frictional force on the wear surface. Wear surfaces of the 3 wt.% boron MMC and 6 wt.% boron MMC

tested at a sliding speed of 3.68 mrs are shown in Fig. 11e,f, respectively. Localized melted areas owing to rise in temperature are seen in both the composites. As shown in Fig. 11e, wear of the 3 wt.% boron chilled composite seems to start by localized melting of the surface and proceed by delaminations from the matrix in which there is severe plastic deformation. In the 6 wt.% boron chilled composite shown in Fig. 11f, some wave like wear patterns which might be related to melt and slip are also seen. These photomicrographs illustrate that the major wear mechanisms at high sliding speeds are melt and slip wear of the matrix due to the elevated temperatures.

5. Conclusions In the Al-boron composites tested, both UTS and fracture toughness of the chilled composites were found to increase as the content of boron particulates was increased up to approximately 6% by weight. Further

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addition of boron particulates only serves to reduce these two mechanical properties. There is therefore, no advantage in increasing indefinitely the boron content in such MMCs. If all other factors were kept constant, the faster the heat extraction from the molten MMC during casting, the higher would be the UTS and the fracture toughness of the castings. Fracture analysis of the MMCs cast using copper and steel chills showed mixed mode fracture with isolated micro-cracks and a bimodal distribution of dimples on the fracture surface. In contrast, fracture analysis of the MMCs cast without chills revealed ductile failure with separation of the boron particles from the matrix. The wear resistance of boron reinforced Al-based MMCs is highly dependent on the rate of chilling during casting. Wear resistance improves as the rate of chilling is increased. Wear resistance also improves as sliding speed increases from low to intermediate, beyond which, wear resistance is reduced again. Hardness and wear resistance both increase as boron content is increased. For each MMC, as the aging time increases, the hardness of MMCs increases to a peak value and then drops again. As the boron content is increased, there is a tendency for the peak aging time to be reduced because dispersoids provide more nucleation sights for precipitation. For any fixed aging temperature and duration, increasing the boron content causes the hardness of the MMC to increase. References w1x Razaghian A, Yu D, Chandra T. Composites Sc Tech 1998;58:293.

w2x Wu SQ, Wang ans HZ, Tjong SC. Composite Sc Tech 1996;56:1261. w3x Tjong SC, Wu SQ, Liao HC. Composite Sc Tech 1997;57:1551. w4x Harrigan WC. Trans AIME 1983;33:169 Warrendale, Pennsylvania. w5x Friend CM. J Mater Sci 1987;22:3005. w6x C.R. Crowe, R.A. Gray and D.F. Haryson, Metall Soc AIME, Warrendale, Pennsylvania, USA, 1985, p. 843. w7x J.J. Lewandowski, C. Liu and W.H. Hunt, Metall. Soc. of AIME, Warrendale, Pennsylvania, USA, 1985, p. 163. w8x Reddy GP, Pal PK. Br Foundryman 1976;69:265. w9x J. Hemanth, SME year book, Singapore, EM 1994; 94r154:31. w10x Hemanth J. J Mater Sci 1995;30:4986. w11x Hemanth J. Wear 1996;192:134. w12x Hemanth J. Trans AFS 1999;5:769. w13x Hemanth J. J Mater Sci 1998;33:23. w14x Hemanth J. J Mater Eng Proc Tech ASM Int 1999;8:417. w15x Hemanth J. Mater Design 1998;19:269. w16x Radhakrishna K, Seshan S. Trans AFS 1981;89:158. w17x Walther WD. J Inst Mater 1954;62:219. w18x Ruddle RW. J Inst Mater 1950;77:37. w19x Flinn RA. Trans AFS 1959;67:385. w20x Couture A. Trans AFS 1966;74:164. w21x Chamberlin MV. Trans AFS 1946;54:648. w22x Seshan S, Seshadri MR. Br Foundryman 1968;61:339. w23x Berry JT. Trans AFS 1970;78:421. w24x Prabhakar KV, Reddy GP. Trans Trans AFS 1979;87:377. w25x Seshan S, Radhakrishna K. Trans AFS 1984;92:234. w26x Redemske J. Trans AFS 1977;85:609. w27x Hemanth J. Mater Design 2000;21:1. w28x Towle DJ, Friend CM. Scripta Metall 1992;26:437. w29x Murthy KSS, Ramachandra A. Br Foundryman 1964;52:451. w30x Murthy KSS, Seshadri MR, Ramachandra A. Metalen 1964;19:367. w31x E.A. Starke, G. Lutjering, J Mater Proc Tech, 1979;4:205. w32x Hemanth J. Mater Sc Tech 1999;15:201. w33x Prasad SV, Mecklenburg KR. Wear 1993;162:47.