Why Al-B₄C Metal Matrix Composites? A Review

2.1 Composite preparation

Reinforcement powder (grit size 400 and 95.4% purity) additions of 15 vol. % were made using a powder injection apparatus (Figure 1). The B₄C particulate was injected into molten Al. Fe, Ti and Zr additions were introduced into the molten bath, using Al-25%Fe, Al-10% Ti and Al-15% Zr master alloys, respectively, whereas Mg and Si were added as pure elements. Chemical compositions of the investigated composites are listed in Table 1. A general view of the powder injection set-up showing a schematic of the injection system is shown in Figure 1. It consists of the following components:

1. a fluidizer tube

2. a carrier tube and a quartz nozzle

3. resistance heating coils

4. an adjustable two-dimensional movable stand

5. a melting unit with resistance heating

6. an impeller (stirrer) with adjustable rotation speed

7. flow diversion baffles.

In order to ensure uniform distribution of the B₄C particulates, the molten composite melt was stirred vigorously (300 rpm) at 730 ± 5°C. Thereafter, the molten composite was poured in two different metallic molds preheated at 450°C, as shown in Figure 2: an L-shaped mold (3.5x 3.8 x 30.5 cm) which was used for microstructure characterization, and a book-type mold (4 x 17 x 34 cm). In order to determine the solidification rate obtained from each mold, trials were made using Al-7%Si. Figure 3 depicts the dendrite arm spacing (DAS) and grain size corresponding to each mold. The castings made using the book-mold were hot rolled into slabs of 1-3 mm thickness, depending on the type of test carried out.

2.2 Microstructural investigation

Samples for microstructural characterization were prepared from the L-shaped mold casting in the as cast condition using 5A composite. The volume fraction and average size of the B₄C particles was measured using Clemex image analyzer. Fracture surfaces were examined of samples sectioned from both tensile- and impact tested bars. The samples were examined using Hitachi S-7000 and Hitachi SU-8000 FE-SEM microscopes equipped with EDS facilities at McGill University, Montreal.

2.3 Mechanical testing

Charpy impact testing was carried out on un-notched test specimens (10x 10 x 55 mm). The samples were sectioned from the L-shaped mold castings and heat treated in an electrical air forced furnace. An instrumented Charpy impact testing machine, equipped with a data acquisition unit was employed to measure the load, total absorbed energy (E_t) to fracture. The mean values of 6 impact-tested samples for each composite/condition were reported.

Slabs (25x 20x 400 mm) were prepared from the book mold castings. Prior to rolling using a four cylinder mill, the slabs were annealed at 500°C for 16 h. The last two passes were carried out at room temperature to straighten the rolled slabs (sheets) – see Figure 4. Figure 5 shows the dimensions of samples prepared from the rolled sheets and used for room and high temperature tensile testing. Tensile samples (matrix is aluminum) were solutionised at 620°C for 24 h. In spite of the fact that pure aluminum normally is not heat treatable, it could benefit from the precipitation of Zr-rich particles during aging. The 6063/B₄C/15p composite samples were solutionized at 540°C to minimize surface oxidation (MgO). After solution heat treatment, the tensile bars were quenched in warm water (60° C), followed by aging for 10 h at 200, 300 and 400° C, and then air cooling. Room temperature testing was carried out using an MTS Servohydraulic mechanical testing machine at a strain rate of 4 x 10⁻⁴/s.

High temperature testing was done at strain rate of 5 x 10⁻⁴/s in a temperature range 25–500°C. In all cases, tensile properties were measured: ultimate tensile strength (UTS), the 0.2% offset yield strength (YS) and percentage elongation (%El). For each working condition, at least five specimens were tested and mean values were reported (SD ±5%). Microstructure and fracture behavior of selected samples were examined using optical microscopy and Field Emission Scanning Electron Microscopy (FESEM) techniques.

3.1 Microstructural characterization (as cast condition)

The main function of the addition of Zr and Ti, is to protect the B₄C particles from reacting with the molten Al [45, 46, 47, 48]. Figure 6(a) depicts the microstructure of a specimen sectioned from the L-shaped castings, revealing a uniform distribution of B₄C particles throughout the matrix. From such micrographs, the volume fraction of B₄C particles was determined (~15 vol.%). According to the Al–Ti binary diagram [49], at 730°C an amount of 0.5 wt-%Ti could be added to the molten composite. From the reported findings of Tahiri et al. [50, 51, 52, 53, 54] it was reported that increasing the concentration of Ti in the molten alloy above 0.5 wt.% will increase the temperature of the molten alloy. As a result, fluidity of the composite will be markedly reduced caused by the segregation of B₄C particles as exhibited in Figure 6(b).

FESEM examination of 5A composite treated with Ti revealed that in addition to B₄C particles, possible precipitation of several intermetallics mainly, TiB₂, TiC and traces of AlB₂₄C₄, Al₄C₃, Al₃BC and AlB₁₂, along with the primary intermetallic phases TiAl, Ti₃Al and TiAl₃ could also occur. It is expected that the formation of these phases in layers would lead to an improvement in the adhesion between the matrix and the B₄C reinforcement [21, 22, 23]. Figure 7 displays electron micrographs of alloy B, containing Ti and Zr. As can be seen in Figure 7(a), the B₄C particles are surrounded by several layers of Zr–Ti rich phases (area marked I). Area marked II shows B₄C particles that have partially reacted with the matrix due to formation of the layer of AlBC existing in the matrix. In area III, fine B₄C particles are found be transformed completely into AlBC compounds. Figure 7(b) reveals a B₄C particle surrounded by a thin layer of Ti-rich phase followed by several layers of Zr–Ti rich phases. Some of these Zr–Ti rich phase particles are seen to grow into the aluminum matrix.

Figure 8 is produced from 5A composite remelted for multiple times at 730°C. Figure 8(a) shows the distribution of the B₄C particles. The B, Al and C distribution are presented in Figures 8(b),8(c),8(d), respectively. Another point to be considered is that Ti covers the entire surface of the B₄C particle (similar to C and B) whereas Zr is limited to the layer decorating the B₄C particle. It is inferred from Figure 9 that the layers surrounding the B₄C particles are a mixture of Al-Ti, and Al-C-Ti compounds. The EDS spectra obtained from Figure 9(b) are presented in Figure 10. Based on these EDSs, areas near the B₄C particles could be made of Al-B-Zr compound whereas those away are probably Al-Ti-Zr compound [55].

4.1 Impact testing

Total energies (Et) produced from the ten studied composites in the solutionized condition are shown in Figure 11(a). Apparently the absorbed energy of the composite depends on what matrix is used and the volume fraction of the undissolved intermetallics. Following aging at 200°C for 10 h (Figure 11(b)), the precipitation of Zr-and Sc containing phases [55, 56, 57, 58] led to significant decrease in the values of Et which may be attributed to precipitation of Mg₂Si phase particles during aging in particular in B-series. It is inferred from Figure 11(b) that precipitation of Zr and/or Sc phases has an insignificant effect on the absorbed energy in series A-composites. According to Fuller et al. [59] aging 6063 alloy at 300°C would result in alloy softening due to coarsening of Mg₂Si phases particles. Simultaneous precipitation of Zr-rich phase may lead to balancing the composite toughness to some extent (Figure 11(c)). Aging at higher temperatures i.e. 400°C for 10 h resulted in coarsening of all types of precipitated phases causing important improvement in the composite toughness, regardless of the type of the matrix used, as exhibited in Figure 11(d).

Fracture mechanism of Al-2%Cu composite was investigated by Miserez [60]. The study showed that the fracture may occur in two stages: (i) particle fracture leading to void nucleation in the matrix, and (ii) voids nucleated in the matrix in areas of high stress concentrations. The blue arrow in Figure 12(a) shows that the crack is propagating through the protecting layer surrounded by stacking faults (white arrows). On the left hand side of the micrograph several stacking faults appear in the form of steps (white arrows). Two distinctive types of cracks were observed in Figure 12(b): cracks that took place at the interior of the B₄C particles and continued through the protecting layers i.e. intergranular, or those occurring at the B₄C/matrix interfaces (black arrow). No particle debonding was observed due to the existence of the protecting layers as displayed in Figure 12(c). The microstructure beneath the fracture surface (vertical section-loading direction) shown in Figure 12(d) demonstrates the coherency between the B₄C particles and the surrounding matrix. Aging the composite at 400°C led to marked coarsening of the Al₃Zr phase particles as seen in Figure 13(a), which explains the improvement in the composite toughness in Figure 11(d). The EDS spectrum in Figure 13(b) corresponds to the circled area in Figure 13(a) revealing strong reflections from Al and Zr elements.

4.2 Tensile testing

4.2.1 Room temperature testing

The stress–strain curves of two aluminum matrix composites in the SHT condition and after aging at 200°C/h are shown in Figure 14. The main observation to be made is the slow working hardening rate illustrated by low work hardening and the slow increase in the composite UTS (Figure 14(a)). As a consequence of aging at 200°C/10 h, the UTS increased by approximately 80 MPa which may be attributed to the precipitation of Al₃Zr phase particles [61]. Considering the solutionizing treatment of 5B composite the maximum attainable strength is about 280 MPa- Figure 14(b). Using a heat treatable matrix i.e. 6063 alloy, the composite revealed significant improvement in both the UTS levels as well as work hardening rate as displayed in Figure 14(c) which may be caused by the precipitation of Mg₂Si phase particles during the storing period prior to testing (~10 minutes at room temperature). As expected, aging at 200°C/10 h resulted in increasing the composite strength from 280 MPa to 500 MPa, Figure 14(d), which may be interpreted in terms of simultaneous precipitation of both Mg₂Si and Al₃Zr phase particles.

Following the solution heat treatment of composite 5A, the fracture surface is characterized by the formation of deep dimple network as demonstrated in Figure 15(a). Some of these dimples revealed the presence of deformation bands (arrowed) due to composite ductility. The marking seen on the surface of the B₄C particles in Figure 15(b) may be caused by gradual fracture of the particles, maintaining at the same time their coherency with the aluminum matrix. Precipitation of Al₃Zr phase particles during aging at 200°C/10 h is clearly seen in Figure 15(c). Due to reduction in the composite ductility, some of the B₄C particles were cracked as shown by the white arrows in the same figure. In the case of 6063 alloy matrix, with the significant increase in the composite UTS level following aging at 200 C/10 h (500 MPa), cracks are seen to initiate and propagate through the Zr-Ti protecting layer as demonstrated in Figure 15(d) – see blue arrow. No B₄C particle debonding is observed to take place under axial loading.

4.2.2 High temperature testing

The 5A and 5B composites (Table 1) were tested in the temperature range of 25–500°C and their corresponding stress–strain curves are displayed in Figure 16. Composite 5B showed a slightly higher strength compared to composite 5A. It should be mentioned here that samples of composite 5B were tested in the T4 condition which involves natural aging. Increasing the testing temperature up to 450°C resulted in significant increase in the composite pct. Elongation to failure. Aging at further higher temperature would lead to precipitation of Al₃Zr which would result in reducing the composite ductility, as shown in Figure 16(c) [6, 62, 63, 64, 65].

For aging at temperatures higher than 0.5 of the melting temperature (Tm), there is a similarity between creep and hot deformation. Under this condition, the relation between the measured parameters can be expressed using power law relationships as described by Eq. 1 [66, 67, 68, 69, 70]:

ε.expQaRT=Aσn=ZE1

where ε = strain rate, σ = flow stress, n = stress exponent, Qa = activation energy, R = gas constant, T = absolute temperature, A is a constant and Z = Zener-Hollomon parameter. At a constantε, σn can expressed as:

σn=BexpQaRTE2

where B = constant. Differentiation of Eq. 2 coupled with (1/T), gives Qa as:

∂lnσ∂1T=QanRE3

Applying these equations, the plot of lnσ vs 1/T, will give a straight line with a slope of (Qa/nR) as shown in Figure 17 [66, 67, 68].

The fracture surface of composite 5A tested at 25°C was characterized by the presence of deformation bands covering the internal surface of the dimples as shown previously. Figure 18(a) exhibits the fracture surface of 5A composite tested at 250°C revealing multiple contour-type markings (white arrow) due to the high ductility ~15%. Testing at 350°C resulted in major increase in the deformation bands in the form of steps (blue arrows) as shown in Figure 18(b). The thin white arrows point to cracked B4C particles. The black arrow in Figure 18(c) - 5A composite pulled to fracture at 450°C- indicates the presence of a long crack within the protective layer. In addition, the 5A composite exhibited elongated dimples as depicted in Figure 18(c). Figure 18(d) is an enlarged portion of the crack in Figure 18(c).From the associated EDS spectrum in Figure 18(e), the possibility of precipitation of a large amount of Zr-rich particles, which would explain the reduction in the composite ductility when tested at this temperature. Fractographic observations made by Zhang et al. [69] on 6092/(B₄C)p indicated the possibility of several interfacial bonding characteristics such as good bonding (extruded composites) and weak bonding (hot isostatic pressing). The fracture surfaces of composites would also show a mixture of ductile and brittle types of fracture [70].

The fracture surface of 5B composite pulled to fracture at 250°C revealed that in addition to the deep dimples, some stacking faults could also be seen in the fracture surface (Figure 19(a)-white arrows). As in the case of 5A composite, at 350°C, the fracture surface exhibited a well-defined dimple structure as a result of the increase in the composite % elongation to fracture, Figure 19(b)-black arrows point to the thickness of the protection layer. Due to the strong particle/matrix interface adhesion, some of the B₄C particles have been cracked at their interior. In this case, the crack was initiated at the particle/matrix interface and propagated through the particle. When the sample was tested at 450°C, the fracture surface revealed the formation of very large and deep dimple network as shown in Figure 19(c).

4.2.3 Effect of strain rate

The main characterized parameters of hot deformation or creep behavior of commercial Al alloys and Al-based composites are by high values of na (> 5) as well as the activation energy Qa. These values are higher than those for solute diffusion [71, 72, 73, 74, 75]. This behavior can be explained in terms of interaction of dislocations with the dispersed strengthening particles resulting in a threshold stress σo. In this case, the deformation process is related to an effective stress, σe = (σ -σo) not the applied stress σ. Therefore, equations 1-2 can be rewritten to take into consideration σo as follows

ε.expQt/RT=Z=A′GbkTσ−σ0Gnt

where: A = constant, k = Boltzmann’s constant, b = magnitude of Burgers vector, G = shear modulus and Qt = true activation energy.

The true stress–strain curves obtained from testing the 5A composite tested at 300°C (a), 400°C (b) and 500°C (c), are respectively depicted in Figure 20. These curves can be divided into three stages; strain hardening where the stress increases with strain until reaches a steady state. Stage 2 represents maximum stress, followed stage 3 where necking takes place leading to failure. Generally speaking, with increasing the strain rate would result in an increase in the flow stress. The effect of testing temperature on the behavior of the stress- strain curves at a constant strain rate of 10⁻³ s⁻¹ is displayed in Figure 20(d). Increasing the testing temperature led to an increase in the composite ductility at 500°C and higher strain rates higher than 10⁻² s⁻¹. The ductility was decreased in temperature range of 350° C–450°C.

The relationship between the strain rate, ε̇ and stress σ, at a constant temperature, is governed by plotting ε̇ vs. σ applying a log log scale (Figure 21) for different testing temperatures. The results reported in Figure 20 may suggest that the data points at each testing temperature fall on a straight line with a constant na that increases from 5.8 at 500°C to ~7 at 350–450°C, thereafter to 10.4 at 300°C. The high values of na are close to those obtained for commercial aluminum alloys [76, 77, 78, 79, 80] and metal matrix composites. As mentioned before, dislocations -second phase particles interaction would lead to high values of na and Qa threshold stress in the composite materials.

Figure 22 (a) shows the fracture surface of the samples tested at 300°C (strain rate of 10⁻² s⁻¹), consisting of a mixture of small dimples and intragranular fracture. Figure 22 (b) is the fracture surface at strain rate of 10⁻⁴ s⁻¹ at 300°C, exhibiting larger dimples with precipitation of Al₃Zr in particles at their interiors (circled areas) [81].