Microstructure and mechanical properties of a high volume fraction SiC particle reinforced AlCu4MgAg squeeze casting

Materials Science and Engineering A269 (1999) 175 – 185 www.elsevier.com/locate/msea

Microstructure and mechanical properties of a high volume fraction SiC particle reinforced AlCu4MgAg squeeze casting S. Long a, O. Beffort a,*, C. Cayron b, C. Bonjour c a

Swiss Federal Laboratories for Materials Testing and Research, CH-3602 Thun, Switzerland Interdepartmental Centre of Electron Microscopy, EPFL, CH-1015 Lausanne, Switzerland c Valais Engineering School, Swiss Engineering College of West Switzerland, CH-1950 Sion, Switzerland b

Received 21 December 1998; accepted 8 March 1999

Abstract The microstructure and mechanical properties of a high volume fraction SiC particle reinforced AlCu4MgAg alloy produced by squeeze pressurised infiltration of dense packed particle preforms were characterised. It was found that the addition of 60 vol.% 12 mm SiC particles eliminates the processing related intrinsic macrosegregation of the matrix alloy, induces coarse precipitates in as-cast matrix, but does not impede matrix grain growth. The SiC – Al interfacial chemical reaction is suppressed to a limited extent. In T6 condition, the interfacial intermetallics are largely dissolved and the matrix is decorated with fine and dense precipitates of u% and Q phases. The addition of SiC particles accelerates the age-hardening response of the matrix alloy. Mechanically, the stiffness, hardness, flexural strength, fatigue strength and abrasive wear resistance are substantially improved, together with significantly reduced fracture toughness and ductility in comparison with the matrix alloy. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Metal matrix composite; Microstructure; Mechanical properties; Squeeze infiltration

1. Introduction Addition of ceramic particles into light alloys provides an attractive technique to economically produce particulate metal matrix composites (PMMCs) with improved mechanical performance at room and elevated temperatures and modified physical properties attractive to electronic and structural applications [1]. These property modifications are expected to be more pronounced when the volume fraction of the ceramic phases is high. Among the liquid metal composite synthesising techniques capable of producing near net-shape composite structures at low cost in a one-shot processing, squeeze casting is characterised by its high pressure desirable for composite quality. Therefore, this process has been * Corresponding author. Tel.: +41-33-228-3041; fax: + 41-33-2284490. E-mail address: [email protected] (O. Beffort)

intensively investigated in fabrication of composite castings via liquid metal infiltration of short fibres and continuous multifilament preforms, and via casting the re-molten composite ingots containing low volume fraction ceramic particles produced by other liquid metal composite techniques. The production of high volume fraction particulate metal matrix composites (HVPMMCs) via matrix melt infiltration of ceramic preforms with structural and property satisfactions presents an engineering challenge to the squeeze casting technique. Up to now, cast HVPMMCs produced by reactive infiltration [2,3], gas pressurised infiltration with or without the aid of vacuum [4,5] have been reported. As reviewed by Ejiofor and Reddy [6], little investigation has been reported on HVPMMCs fabricated by squeeze infiltration of particle preforms except for the most recent studies on the effects of processing conditions on microstructure and mechanical properties of Al 6061/SiC particles fabricated by squeeze infiltration of sintered or oxidised SiC

0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 9 ) 0 0 1 6 3 - X

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

176

Table 1 Chemical composition of SiC particles and matrix alloy (in wt.%) Particles [11]

SiC

Norton F500HD

99.75

0.07

0.12

Matrix alloy AlCu4MgAg

Cu 4.29

Mg 1.01

Ag 0.62

a

Free C

Free SiO2

Free Si

APSa (mm)

Fe2O3

Al2O3

CaO

0.04

0.015

0.003

0.002

11.8–13.8

Si B0.005

Mn 0.016

Fe 0.038

Ti 0.003

Cr 0.001

APS, average particle size.

particle preforms [7,8], where the expected improvement in mechanical performance was impeded by the employment of low preform preheat. A full comprehension of the microstructure features and mechanical properties of squeeze cast HVPMMCs produced under optimal conditions remains to be achieved to evaluate the potential and limits of squeeze infiltration for fabrication of HVPMMCs. Hence, the present study has been focused on systematic characterisation of the microstructure and mechanical properties of an AlCu4MgAg alloy reinforced with a high volume fraction of SiC particles fabricated by direct squeeze infiltration of tap-packed SiC particle preforms under processing conditions optimised previously for chopped fibre MMC squeeze castings [9].

2. Experimental description

2.1. Component materials Norton 500HD SiC particles were used as reinforcement. The composition, size distribution and morphology of the particles are given in Table 1 and Fig. 1, respectively. A high purity age-hardenable AlCu4MgAg alloy [10] was used as matrix. The chemistry of the alloy is specified in Table 1.

condition. The rest of the composite plate was homogenised at 480°C/2 h + 500°C/2 h, followed by water quenching and room temperature exposure for more than 100 h for natural ageing to T4 condition. Then, the specimen was artificially aged at 170°C to reach its peak-age hardened condition (T6). The age-hardening response of the composite was determined by hardness measurements (HB30). Microstructurally, the composite specimens in both as-cast and T6 condition were characterised with light and transmission electron microscopy (TEM) to examine the infiltration quality, matrix and interfacial microstructures. The morphologies of the fracture surfaces of specimens subjected to mechanical tests were characterised by scanning electron microscopy (SEM) to examine the failure behaviour. Mechanically, Young’s modulus and flexural strength were determined by four-point bending tests on 3 mm× 4 mm × 45 mm specimens in as-cast and T6 conditions. The mode I fracture toughness (KIC) was measured by three-point bending test on 2 mm×2.5 mm× 30 mm pre-notched specimens in both as-cast and T6 conditions. The details of the testing methods are specified in Ref. [12]. The fatigue strength was tested under completely reversed bending (R= −1)

2.2. Squeeze infiltration processing The SiC particles were tap-packed into a steel housing to form an 8 mm ×90 mm × 100 mm particle preform with  60% volume fraction. Before infiltration, the preform was preheated to 700°C in N2 atmosphere, the matrix alloy was superheated to 750°C and the squeeze die cavity and ram were preheated to 300°C. During casting, the preform and the steel housing were transferred into the die, immediately followed by melt pouring and squeeze pressurisation at a ram speed of 5 mm/s until the pressure reached 130 MPa. Then, the pressure was maintained at 130 MPa until the complete solidification of the casting.

2.3. Material characterisation Part of the composite plate was examined in as-cast

Fig. 1. Morphology of F500HD Norton high density (Green Alpha Silicon Carbide) SiC particles of round shape and an average dimension of  12 mm.

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

177

Fig. 2. Micrographs of the 60% SiCF500HD/AlCu4MgAg composite, showing particle distribution and the matrix-rich channels in the preform entry region.

Fig. 3. Solidification structure of the monolithic AlCu4MgAg squeeze casting. (a) macrosegregation; (b) interdendritic intermetallics.

condition on 2 mm× 5 mm× 60 mm specimens in T6 condition. The abrasive wear resistance of the composite in both as-cast and T6 condition was examined by the Miller Slurry Abrasion Test under G75-89 ASTM standard with the following parameters: wear surface area, 12 mm ×24.5 mm; loading force (F), 22 N; relative moving rate, 2880 r/h ×0.5 m; and the HP7 slurry consisting of 50 wt.% water+50 wt.% F22 Al2O3 particles of 45–125 mm diameter.

3. Results and discussion

3.1. Microstructure 3.1.1. Non-infiltration porosity and particle distribution As indicated by all the following micrographs, no non-infiltration pores can be observed, suggesting that a high quality infiltration has been achieved although the lack of the aid of vacuum, capillarity and nitrogen gas entrapment are expected to induce fine scale porosity at particle contacting sites. The SiC particles are found uniformly distributed in the matrix throughout the casting due to their

dense packing nature in the preform. However, some brighter channels of 0.1–0.3 mm width were observed in the region close to the preform entry, where melt starts penetrating into the SiC particle preform. These channels are characterised by a low local particle volume fraction (Vp) of 25–40%, as shown in Fig. 2. The formation of these matrix enriched channels can be attributed to the further condensation of the tap-packed particles in the preform during infiltration. According to the infiltration hydrodynamics [13], the infiltration pressure generated by the forced flow of viscous melt at the pre-selected speed can be higher than the condensation force applied during tap-preforming, which results in a relative movement of the particles and the consequent formation of the matrixrich channels.

3.1.2. As-cast microstructures The solidification structure of the monolithic AlCu4MgAg casting produced with identical processing parameters is given for comparison in Fig. 3 and the solidification microstructure of the composite is given in Fig. 4.

178

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

As the figures show, the solidification structure of the monolithic casting is characterised by its dendritic grain structure decorated by intermetallic eutectic phases on the interdendritic boundaries. Macrosegregation of the alloying elements Cu, Mg and Ag was observed at they casting centre where the melt solidification finally completes. In contrast, the composite possesses an equiaxial grain structure with an average grain size of  0.5–1 mm throughout the casting plate. Neither macro nor grain boundary segregation can be observed, except for the preferential precipitation of the intermetallic phases on the SiC/Al interface, the common distribution feature of the intermetallics in cast MMCs. Previously, the addition of ceramic particles was reported to reduce matrix grain size in the composites containing particles of low volume fraction [14] and this structural modification was attributed to the particles impairing the growth of the solidifying matrix. In contrast, the present observation indicates that, with the employed particle size, particle volume fraction and processing conditions, the solidifying Al primary phase was permitted to grow through the particle interspaces to form a grain size two orders of magnitude larger than that of the SiC particles. The contradiction of the previous observation with the present observation can be attributed to the differences in the composite systems

and in the nature of the composite process, which significantly alter the nucleation and growth kinetics of the matrix alloy during solidification. The change of the matrix grain structure from dendritic in the monolithic casting to equiaxial in the composite is due to the use of the steel housing that encapsulates the SiC particle preform and isolates the composite plate from the chilling effect of the cold die wall during casting. Subjected to the isolation, the solidification of the composite relies entirely on nucleation within the encapsulated region, instead of being affected by the chilling of the casting/die interface as the monolithic casting does. The as-cast matrix microstructure is characterised by the presence of coarse precipitates of mainly u%-Al2Cu and Q-Al5Cu2Mg8Si6 phases as Fig. 4c shows. The phase identification of the precipitates was detailed in Ref. [15]. The appearance of the coarse precipitates that are absent in the as-cast monolithic casting of the matrix alloy (Fig. 3b) can be associated with the mismatch of the coefficients of thermal expansion (CTE) between SiC and matrix alloy and the appearance of Si in the matrix. During the cooling-down post-solidification, the thermal stresses and the resultant high density dislocations in the Al primary phase catalyse the formation of the precipitates at elevated temperature.

Fig. 4. Micrographs of the as-cast SiCF500HD/AlCu4MgAg composite. (a) Solidification structure; (b) matrix structure in as-cast condition; (c) coarse Q and u% precipitates in the matrix.

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

179

Fig. 5. TEM micrographs of as-cast SiCF500HD/AlCu4MgAg composite, showing the preferential deposition of intermetallics on the SiC/Al interface (a and b) and a Al4C3 crystallite on the interface (c).

3.1.3. Interfacial microstructure Fig. 5 gives the morphology of the as-cast interface, showing that the intermetallic phases, mainly u-Al2Cu, Q-Al5Cu2Mg8Si6, b-Mg2Si and Si phases that usually appear on the grain boundaries and concentrate in the final solidification regions in monolithic castings of corresponding compositions, preferentially deposit on the particle/matrix interface. Except for the above intermetallics, the SiC/Al interface is distinctive and free of Al4C3 under low magnifications, as shown in Fig. 4cFig. 5b. Observations under higher magnification reveal the existence of fine Al4C3 crystallites in some interfacial regions (Fig. 5c) indicating that the SiC/Al reaction has proceeded to a very limited extent during casting. Considering the low initial concentration of Si in the matrix alloy, the prevalent appearance of Si-containing intermetallic phases is apparently the result of the chemical reactions in the SiC/AlCu4MgAg system during squeeze casting. For the SiC/AlCu4MgAg system, the interfacial chemical reaction 3SiC(s)+4Al(l)=Al4C3(s) +3Si(l)

(1)

is able to produce Si. The possible existence of a thin SiO2 film on the SiC particle surface after particle

processing [16] and/or preform preheating may also contribute to the common presence of Si through the following reactions. 3SiO2(s)+ 4Al(l)=2Al2O3(s)+ 3Si(l)

(2)

SiO2(s)+ 2Mg(l)=2MgO(s)+Si(l)

(3)

2SiO2(s)+ Mg(l)+ 2Al(l)= MgAl2O4(s)+ 2Si(l)

(4)

However, no notable interfacial oxides were observed in the present TEM study, suggesting that the Si in the intermetallics mainly originates from the SiC/Al reaction. SiC/Al composites are notorious for their intensive interfacial reaction. The presence of large Al4C3 phases on the interface results in severe degradation of the mechanical performance and the corrosion resistance of the composites. To suppress the reaction, high Si concentration in the Al matrix alloy is recommended [17] at the expense of inducing a substantial amount of Si phase into the composite causing material embrittlement and possible strength loss. The formation of a very limited amount of fine Al4C3 in the squeeze cast composite indicates that the squeeze infiltration processing is able to suppress the SiC/Al reaction to a great extent because the high pressure used accelerates infiltration and subsequent solidification, thus significantly shortening the melt–particle contacting time to

180

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

Fig. 8. Age-hardening response of SiCF500HD/AlCu4MgAg and its matrix alloy during 170°C artificial age-hardening. Fig. 6. TEM micrographs showing the interface of the T6 SiCF500HD/AlCu4MgAg composite.

less than 20 s. Therefore, squeeze casting, as compared with the other liquid route composite processing, offers a widened freedom in employment of different particle/ matrix chemistries suitable to prescribed property requirements.

3.1.4. Effect of heat treatment on microstructure In the T6 condition, the as-cast interfacial intermetallics have been largely dissolved after homogenisation, as shown in Fig. 6, except for the occasionally observable u and Q phases smaller than 200 nm on the SiC/Al interface. Although solid state SiC/Al reaction is thermodynamically possible during homogenisation [18], the present TEM investigation shows no obvious difference in the size and amount of the interfacial Al4C3 in the as-cast and T6 conditions.

Fig. 7 shows the microstructure of the matrix in the T6 condition. As the figures indicate, the coarse precipitates in the matrix and the large interfacial intermetallics observed in the as-cast condition were replaced with fine and dense matrix precipitates of u%-Al2Cu, Q, QC and QP phases. The later two types of precipitates are identified as the precursors of the Q phase with different composition and lattice constants. The details of the related phase identification and their formation mechanism are given in Ref. [15].

3.2. Age-hardening response of the composite The AlCu4MgAg alloy is a typical age-hardenable Al alloy. The age-hardening response of the composite at 170°C after homogenisation and  100 h room temperature natural ageing is given in Fig. 8 in comparison with that of the monolithic casting.

Fig. 7. TEM micrographs of the matrix precipitation state in T6 SiCF500HD/AlCu4MgAg composite. (a) Bright field image of matrix showing the fine rod shaped Q-phase along with u% with selected area diffraction pattern in [100]a-Al axis (inset); (b) dark field image of u% phase near the [110]a on u% spot, showing the presence of u% phase precipitates (bright phase) in the matrix.

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

181

Table 2 Mechanical properties of SiCF500HD/AlCu4MgAg and AlCu4MgAg castings

Composite As-cast T6

Monolithic matrix alloy As-cast T6

E (GPa)

s (MPa)

HB(30)

KIC (MPa m1/2)

Fatigue s (MPa)

200.09 0.4 198.49 0.3

673.29 50 703.59 60

273 360

10.039 0.02 9.590.16

n.a. 200

E (GPa)

s* (MPa)

s*0.2 (MPa)

d* (%)

HB(30)

KIC (MPa m1/2)

Fatigue s** (MPa)

70 72

256 456

164 402

5.7 3.6

97 170

21.5 24

n.a. 5100

* Tensile properties measured on A5 round bar specimens with 35 mm parallel gauge length. ** Room temperature flexural fatigue strength for T7-tempered Al201 casting alloy (4.0–5.0% Cu, 0.15–0.50% Mg, 0.15–30% Fe, 0.2–0.5% Mn, 0.15–0.35% Ti) in R = −1 condition [20].

As the curves indicate, the precipitation kinetics of the monolithic alloy is characterised by a hardness decrease in its quenched state and a subsequent hardness increase after 4 days of natural ageing, corresponding to the formation of the GP zones. During 170°C artificial ageing, due to the dissolution of the GP zones, a slight hardness reversion appears during the initial holding stage. After an incubation time of  0.2 h, the hardness increases slowly until its peak hardness is reached after 18 h. For the composite, the hardness is increased after homogenisation, as well as during natural ageing. During artificial ageing, the hardness increases steadily until the peak-aged condition is reached after 10 h and remains at the peak-aged condition for a longer duration. This observation suggests that, the addition of 60% SiC particles accelerates the process of precipitation. Again, this precipitation enhancement can be attributed to the CTE mismatch between the SiC particles and the Al matrix and the resultant dense dislocations formed during water quenching, on which the precipitates are susceptible to nucleate [19].

3.3.2. Flexural strength The flexural strength of the composite is about 670 and 700 MPa in as-cast and T6 conditions, respectively. However, as the loading-time curves in Fig. 9 show, regardless of the matrix heat treatment states the composite starts to yield at a loading level of 350 MPa, a stress close to the yielding strength of the age hardened monolithic matrix in Table 2, although it shows no macrodeformation upon failure. The lowered yielding stress of the composite is consistent with the early onset of matrix plastic deformation at the pole positions of the contacting particles [21]. The significantly increased failure strength can be attributed to the constrained matrix plastic deformation and the associated work hardening [22] that does not prevailingly exist in the composites containing a low volume fraction of particles.

3.3. Mechanical properties The mechanical properties of the composite in both as-cast and T6 conditions are summarised in Table 2. As a reference, the mechanical properties of the monolithic AlCu4MgAg alloy are also shown in the table.

3.3.1. Young’s modulus As the data indicate, the Young’s modulus (E) of the composite possesses a value of  200 GPa, indicating a substantial increase in comparison with the matrix alloy, but 10% lower than the prediction of the Rule of Mixture. Like the matrix alloy, the Young’s modulus of the composite is insensitive to the matrix precipitation state.

Fig. 9. Three-point bending loading curves of SiCF500HD/ AlCu4MgAg in its as-cast and T6 conditions. Note the early appearance of yielding.

182

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

The composite fracture toughness is about 40–50% of the value of the monolithic alloy, indicating that the addition of the high volume fraction of SiC particles significantly weakens the crack arresting capability of the matrix alloy via its plastic deformation. However, the value of the fracture toughness is two times higher than the 2–4 MPa m1/2 for monolithic polycrystalline SiC ceramic, indicating that the presence of the matrix– particle interfaces and the matrix between the particles significantly raises the energy dissipation during crack propagation, leading to a significant improvement in fracture toughness.

Fig. 10. Fractograph of as-cast SiCF500HD/AlCu4MgAg failed at 655 MPa flexural stress.

Fig. 11. Stress-cyclic life relationship of T6 SiCF500HD/AlCu4MgAg composite and T7 monolithic Al201 alloy under R= − 1 flexural loading condition at room temperature.

3.3.4. Fatigue beha6iour The fatigue behaviour under completely reversed three-point bending (R= −1) is plotted in Fig. 11. As the figure shows, the fatigue limit of the composite is 5200 MPa. When the loading stress was raised from 200 to ] 230 MPa, the specimens failed at a dramatically reduced cyclic life. As compared with the 100 MPa fatigue strength of the over-aged (T7) A201 alloy at 107 cycles [20], the addition of 60% SiC particles improves the fatigue strength significantly. Fig. 12 shows the fractograph of the transition region between the fracture of the initial fatigue damage zone (right-hand side of the image) and the fracture produced by the subsequent static failure of the specimen at the end of its fatigue endurance (left-hand side of the image). As the morphology indicates, the fatigue failure proceeds by damage accumulation in the matrix during sub-critical crack propagation. The subsequent overcritical fast crack propagation proceeds via particle breakage and matrix plastic failure and leads to the failure of the specimen.

The bending fracture surfaces of the composite in both as-cast and T6 conditions present the same features: mainly consisting of SiC particle cleavages and matrix ductile plastic dimples, in addition to the occasional interfacial failure, as selectively shown in Fig. 10. The morphological features indicate that a strong SiC– matrix interface has been established during casting and the failure of the composites is dominated by SiC particle breakage and, thus, the addition of the SiC particles diminishes the ductility of the composite.

3.3.3. Fracture toughness The fracture toughness of the composite in as-cast and T6 conditions is 10 and 9.5 MPa m1/2, respectively. A fracture morphology is produced similar to that of the bending fracture surfaces in Fig. 10.

Fig. 12. Fracture morphology of T6 SiCF500HD/AlCu4MgAg composite specimen subjected to 5.6 × 104 cycles under 230 MPa (R= − 1) flexural loading.

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

183

be associated with the mechanical behaviour of the matrix in the composite with a strong interface. When the loading stress is low, the matrix remains elastic and no fatigue damage can be generated. However, when the loading level is high enough (\200 MPa) to induce plastic deformation in the matrix, fatigue damage initiates and accumulates quickly until the concentrated stress induced by the fatigue cracks is able to break the SiC particles in front of the fatigue crack tips, leading to the catastrophic failure of the material at a significantly reduced cyclic life.

Fig. 13. The abrasive wear resistance of SiCF500HD/AlCu4MgAg in as-cast and T6 conditions.

Fig. 14. A comparison of the abrasive wear resistance of SiCF500HD/AlCu4MgAg with that of Duralcan 6061/Al2O3/20p in T6 condition.

The significant improvement of the fatigue endurance of the composite under B200 MPa flexural loading can

3.3.5. Abrasi6e wear resistance The abrasive wear resistance of the composite in terms of weight loss is given in Fig. 13, and compared with that of A6061 reinforced with 20% Al2O3 particles (Duralcan 6061/Al2O3/20p) tested under identical conditions in Fig. 14. As the figures indicate, the addition of 60% SiC particles significantly improves the abrasive wear resistance, especially in the peak aged condition. The common surface morphologies of the tested specimens in Fig. 15 indicate that, despite the presence of the large area of worn matrix in Duralcan 6061/ Al2O3/20p, both composites share the same wearing mechanism during the abrasive test: reinforcement particle wearing via direct contacting with the abrasive Al2O3 particles, matrix cutting by Al2O3 particles, and falling-out of the protruding reinforcement particles when the worn matrix becomes unable to support them. Regulated by the above mechanism, the addition of the high volume fraction of SiC particles will naturally reduce the direct exposure of the matrix phase to the abrasive Al2O3 particles, leading to a substantial improvement in the abrasive wear resistance and the morphological difference shown in Fig. 15. 3.3.6. Comparison of mechanical properties Finally, as a comparison, the properties of the cast HVPMMCs with similar SiC particles fabricated by different infiltration techniques are given in Table 3. As the data indicate, with the optimal processing conditions, squeeze pressurised infiltration of densely packed

Fig. 15. The surface morphology of SiCF500HD/AlCu4MgAg composite subjected to abrasive wear test (a), in comparison with Duralcan 6061/Al2O3/20p in T6 condition (b).

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185

184

Table 3 A comparison of similar HVPMMCs fabricated by different infiltration techniques Material

SiC (%)

Matrix alloy

Process route

r (g/cm3)

SiC/6061 [8] AlSiC™ [23]

45 63

A6061 AlSiMg

n.a. 3.01

500–565 369

150–165 220

n.a. n.a.

APIC [24]

55

A295

2.97

502

183

–

Primex [25]

70

AlSiMg

3.0

n.a.

265

150

10

Keramal [26]

50

AlCu4TiMg

3.0

]540

\150

n.a.

n.a.

SICF500HD/ AlCu4MgAg

60

AlCu4MgAg

Squeeze casting Pressurised infiltration 10 MPa gas pressure Reactive infiltration 10 MPa gas pressure Squeeze casting

3.0

\670

200

\200

10.4

SiC particles is able to produce composites with satisfactory microstructure and mechanical performance superior to those fabricated by the other infiltration techniques.

4. Conclusions In the present work, the microstructure and mechanical properties of a squeeze cast composite consisting of 60 vol.% F500HD SiC particles and an AlCu4MgAg matrix alloy were microscopically and mechanically characterised. Microstructurally, the particle preform is well infiltrated and the particles are uniformly distributed in the composite casting except for the formation of matrix-rich channels at the preform entry due to further condensation of the tap-packed particle preform. In the as-cast condition, the composite possesses a large equiaxial solidification structure and is free of macroscopic and grain boundary segregation. The intermetallics preferentially deposit at the SiC/Al interface and coarse precipitates form in the matrix apart from the interface. Fine Al4C3 crystallites are observable in some interfacial regions. In the T6 condition, the matrix is enriched with fine precipitates of u% and Q phases and the SiC/Al interface is decorated with small intermetallics and fine Al4C3 crystallites. The addition of SiC particles is found to accelerate the precipitation and to reduce the time for peak-age hardening during artificial age-hardening. Mechanically, the addition of 60 vol.% SiC particles significantly improves the stiffness, hardness, flexural strength, fatigue strength and abrasive wear resistance compared with the monolithic matrix alloy. However, the addition of SiC particles decreases the fracture toughness and dramatically reduces the ductility. The mechanical properties of the composite are dominated by the addition of SiC particles and relatively insensitive to the matrix precipitation state.

s (MPa)

E (GPa)

s (R=−1) (MPa) (107 cyc.)

KIC (MPa m1/2)

n.a. n.a. 9.4–10.4

References [1] D.J. Lloyd, Int. Mater. Rev. 39 (1994) 1 – 23. [2] A.W. Urquhart, Mater. Sci. Eng. 144A (1991) 75 – 82. [3] J.A. Cornie, H.-K. Moon, M.C. Flemings, in: J. Masounave, F.G. Hamel (Eds.), Fabrication of Particulate Reinforced Metal Composites, ASM International, Materials Park, OH, 1990, pp. 63 – 78. [4] M. Kobashi, T. Choh, J. Mater. Sci. 32 (1997) 6283–6289. [5] E. Candan, H.V. Atkinson, H. Jones, Key Eng. Mater. 127-131 (1997) 463 – 470. [6] J.U. Ejiofor, R.G. Reddy, JOM 49 (11) (1997) 31 –37. [7] B.K. Hwu, S.J. Lin, M.T. Jahn, Mater. Sci. Eng. 207A (1996) 135 – 141. [8] Y. Cui, L. Geng, C. Yao, X. Li, in: Proceedings of ICCM-11, Vol. III, Gold Coast, Australia, 14 – 18 July 1997, pp. 215–220. [9] S. Long, H.M. Flower, O. Beffort, in: J. Beech, H. Jones (Eds.), Solidification Processing 1997, Proc. of the 4th Decennial Int. Conf. on Solidification Processing, 7 – 10 July, 1997, Sheffield, UK, pp. 92 – 96. [10] O. Beffort, C. Solenthaler, P.J. Uggowitzer, M.O. Speidel, Mater. Sci. Eng. 191A (1995) 121 – 134. [11] Product data sheet of NORTON AS, NORTON-Lillesand, Norway, 1997. [12] J. Ku¨bler, Ceram. Eng. Sci. Proc. 18 (4) (1997) 155–162. [13] S. Long, Z. Zhang, H.M. Flower, Acta Metall. Mater. 43 (1995) 3499 – 3509. [14] M. Ferry, P. Munroe, A. Crosky, T. Chandra, in: N. Hansen et al. (Eds.), Metal Matrix Composites — Processing, Microstructure and Properties, 12th Risø Int. Symp. on Metallurgy and Materials Science, Risø National Laboratory, Roskilde, Denmark, 1991, pp. 337 – 342. [15] C. Cayron, P.A. Buffat, O. Beffort, S. Long, 14th International Congress on Electron Microscopy (ICEM-14), Cancun, Mexico, 31 Aug. – 4 Sept. 1998. [16] H. Ribes, M. Suery, G. L’esperance, Metall. Trans. 21A (1990) 2489 – 2496. [17] J.C. Lee, S.B. Park, C.S. Oh, H.I. Lee, Acta Mater. 44 (1998) 2635 – 2643. [18] B. Noble, A.J. Trowsdale, S.J. Harris, J. Mater. Sci. 32 (1997) 5969 – 5978. [19] P.B. Prangnell, W.M. Stobbs, in: N. Hansen et al. (Eds.), Metal Matrix Composites — Processing, Microstructure and Properties, 12th Risø Int. Symp. on Metallurgy and Materials Science, Risø National Laboratory, Roskilde, Denmark, 1991, pp. 603–610. [20] Properties of Cast Aluminium Alloys, ASM Handbook Vol. 2 — Properties and Selection: Nonferrous Alloys and Special-

S. Long et al. / Materials Science and Engineering A269 (1999) 175–185 Purpose Materials, ASM International, Materials Park, OH, 1992, 152 – 153. [21] D.F. Watt, X.Q. Xu, D.J. Lloyd, Acta Mater. 42 (1996) 789 – 799. [22] Z. Dong, Y. Wu, J. Mater. Sci. 31 (1996) 4401–4405. [23] Data sheet of PCC-AlSiC™, AFT Advanced Forming Technology, USA, 1997.

.

185

[24] J.A. Cornie, Mater. Technol. 10 (3/4) (1995) 43 – 58. [25] Data sheet of composites fabricated by LANXIDE™, PRIMEX™ and PRIMEX CAST™, Lanxide Corporation, USA, 1997. [26] Data sheet of Keramal Aluminium Matrix Composites, AMAGForschung und Entwirklung, Austria, 1995.