Why Al-B₄C Metal Matrix Composites? A Review

Abstract

The Al-B4C metal matrix composite (MMC) is characterized by its ability to absorb neutrons which makes it the most suitable shielding material for nuclear reactors. The present work was performed on two series of Al-B4C metal matrix composites made using a powder injection apparatus. In one series, commercially pure aluminum (A5) served as the matrix. For the second set, 6063 alloy was used. In all cases the volume fraction of B4C reinforcement particles (grit size 400 mesh, purity 99.5%) was approximately 15%. The volume fraction of the injected B4C particles was determined using a computer driven image analyzer. Measured amounts of Ti, Zr, and Ti + Zr, were added to the molten composites of both series. Microstructural characterization was carried out employing a field emission scanning electron microscope operating at 20 kV and equipped with an electron dispersive x-ray spectroscopic system (EDS). The same technique was applied to characterize the fracture behavior of the tested composites. Mechanical properties of these composites were investigated using impact testing, and ambient and high temperature tensile testing methods. Almost 1000 impact and tensile samples were tested following different heat treatments. The obtained results from these investigations are reported in this Chapter.

1. Introduction

The Al-B₄C metal matrix composite (MMC) is characterized by its high thermal conductivity and its ability to absorb neutrons which makes it a suitable shielding material [1]. Increasing the concentration of B₄C (>30%) increases the composite strength as well as its neutron absorption capacity. Roy et al. [2] suggested the use of 7xxx alloys as base material for the MMC due to its low density and its hardening ability caused by heat treatment which would contribute to the strength of the MMC. The use of 2124 Al alloy composites reinforced with B₄C particulates has been proposed by Öksüz and Oskay [3]. The authors claim that the volumetric wear rates of the 2124 Al alloy and its composites are increased with increase in the applied load. Singla et al. [4] proposed the use of molten technique for the production of Al-B₄C MMC. The authors studied an MMC made of Al-7075 alloy as the matrix and B₄C 32 μm particulate as the reinforcement agent. Mohan and Kennedy [5] investigated the machinability of Al-(7 and 14) wt.% Si alloys reinforced with B₄C. The MMCs were developed using the stir casting technique. Their results show that the composite reinforced with B₄C with a particle size of 100 nanometers has better mechanical properties and wear behavior compared to those reinforced with 24-micron or 6-micron sized particulates. Vaidya et al. [6] found that the strength of B₄C particle reinforced Al 6061 composite was significantly greater than the unreinforced alloy.

Drilling experiments were conducted by Kumar et al. [7] on 6061 alloy-15%B₄C (220 μm particulate diameter) using a vertical machine with High Speed Steel drills of 6 mm, 9 mm and 12 mm diameter under dry drilling conditions. It was found that speed, design of the experiment and drill diameter have a marked influence on the Over-Cut (half the difference of the diameter of the hole produced to the tool diameter). Topcu [8] and Manjunatha et al. [9] used the powder atomization technique to produce Al-5% B₄C and Al-15%B₄C MMCs. The authors reported that the wear resistance increased in proportion to the amount of the boron carbide reinforced. Tribo-surface characteristics of two aluminum metal matrix composites (Al-MMC) of compositions Al–13 vol%B₄C and Al–13 vol%SiC sliding against a commercial phenolic brake pad under dry conditions were investigated by Shorowordi et al. [10, 11, 12]. The friction coefficient was found to decrease slightly at high contact pressure.

The wear rate and friction coefficient of Al–B₄C was lower than that of Al–SiC. Several studies on friction behavior involving Al-MMC friction against ferrous materials revealed that during sliding, a layer, termed as mechanically mixed layer (MML), was formed on the worn surface of the Al-MMC [13, 14, 15, 16, 17]. Such layer, however, was not found to form on unreinforced aluminum. Several researchers [18, 19, 20, 21, 22] studied the production of Al-11%B₄C using stir melt technique. The 6061 alloy was the matrix to which B₄C particles were added. Prior to addition, the B₄C particles were preheated along with K₂TiF₆ halide salt. The resulting composite was found to have improved mechanical properties compared to the base alloy. Uthayakumar et al. [23] performed a study on the wear performance of Al–5%SiC–5%B₄C hybrid composites under dry sliding conditions using a pin on disc tribometer method. The main conclusion was that the hybrid composites can retain the wear resistance properties up to 60 N load and sliding speed ranges of 1–4 m/s.

Comparison of microstructural and mechanical properties of Al–10 vol% TiC, Al–10 vol% B₄C and Al–5 vol% TiC–5 vol% B₄C composites prepared by casting techniques was made by Mazaheri et al. [24]. The results show that the wear behavior of Al-B₄C MMC is the best among the three composites studied. The wettability of B₄C particulates was investigated by Toptan et al. [25]. They found that addition of Ti leads to formation of thin layers (80–180 nm in thickness) of Ti-C and Ti-B around the B₄C particulates which would solve the wettability issue. Similar observation on titanium as one of the reactive metals that can be used to increase wettability in Al-B₄C system was reported by other researchers [26, 27, 28, 29, 30, 31]. According to Wang et al. [32] and Yang et al. [33], the stress distribution within a particle-reinforced composite subjected to external loading is non-uniform. Nanostructured Al–B₄C composite sheets were processed by accumulative roll bonding (ARB), and the effect of the number of ARB cycles on the distribution of the B₄C particles in the Al matrix was evaluated by Yazdani and Salahinejad [34] who noted an improvement in the reinforcement distribution by increasing the ARB cycles.

The present chapter summarizes the work that was carried out by the present authors using two types of Al-B₄C composites: (i) a mechanically alloyed composite supplied by Ceradyne Canada ULC, a 3 M Company, Chicoutimi, Québec, Canada, and (ii) an in-house made composite using powder injection at the Université du Québec a Chicoutimi [35, 36, 37, 38, 39, 40, 41, 42, 43, 44].

2. Experimental procedure

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.

Figure 1.

A schematic diagram of the powder injector used in the present work.

Alloy Code	Composition (Ti, Zr, Sc in wt%)
1A	Al-15v/oB4C
2A	Al-15v/oB4C + 0.45%Ti
3A	Al-15v/oB4C + 0.45%Ti + 0.25%Zr
4A	Al-15v/oB4C + 0.45%Ti + 0.15%Sc
5A	Al-15v/oB4C + 0.45%Ti + 0.15%Sc + 0.25%Zr
1B	6063-15v/oB4C
2B	6063-15v/oB4C + 0.45%Ti
3B	6063-15v/oB4C + 0.45%Ti + 0.25%Zr
4B	6063-15v/oB4C + 0.45%Ti + 0.15%Sc
5B	6063-15v/oB4C + 0.45%Ti + 0.15%Sc + 0.25%Zr

Table 1.

Codes and compositions of the MMCs used in this study.

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.

Figure 2.

(a) L-shaped mold, (b) book-type mold, (c) L-shaped casting, (d) book-mold casting.

Figure 3.

(a, c) Optical micrographs of Al-7%Si alloy for (a) block casting, 60 μm; (c) L-shaped casting, 30 μm; (b, d) Macrographs showing grain size in (b) block casting; (d) L-shaped casting.

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.

Figure 4.

Hot rolled sheets.

Figure 5.

Typical sample for room and high temperature tensile testing (dimensions are in mm).

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. Results and discussion

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).

Figure 6.

Secondary electron micrographs showing: (a) a uniform distribution of B₄C in matrix of base alloy using in-house powder injecting technique, (b) segregation of B₄C (white circle).

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 7.

SE images from composite B showing (a) regular B₄C particles protected by layers of Zr–Ti rich particles, (II) irregular forms of B₄C showing partial reaction with matrix forming AlBC compound; (b) B₄C particle surrounded by two layers of Zr–Ti rich particles; and (c) B₄C particle showing progress of cumulative reaction towards 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].

Figure 8.

Element distribution: (a) backscattered electron image-note formation of several particles around a B₄C particle, white arrows, (b) boron, (c) alumium, (d) carbon, (e) titanium, (f) zirconium.

Figure 9.

B₄C-matrix interactions: (a) formation of Al-C, (b) formation of Zr rich phases-see Figure 7(c).

Figure 10.

EDS spectra obtained from Figure 9(b) confirming the B4C-matrix interaction and the dependence of the composition of outcome on its position with respect to B4C particles.

4. Mechanical properties

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).

Figure 11.

Total absorbed energies of the present composites: (a) SHT, (b) aging at 200°C, (c) aging at 300°C, (d) aging at 400°C.

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.

Figure 12.

Fracture characteristics of the present composites: (a) stacking faults-SHT, (b) cracks- aging at 200°C/10 h, (c) B₄C/matrix coherency, (d) vertical section beneath (c) confirming particle/matrix bonding-note the severe reaction around some of the B₄C particles - circled areas.

Figure 13.

(a) Fracture surface of composite 5B aged at 400°C/10 h, (b) EDS spectrum corresponding to white circle in (a).

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.

Figure 14.

Stress–strain diagrams corresponding to: (a) Al/B₄C 5A composite - SHT, (b) Al/B₄C composite aged at 200°C/ 10 h, (c) 5B 6063/B₄C composite - SHT, (d) 6063/B₄C composite aged at 200°C/ 10 h.

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.

Figure 15.

Fracture surface of tensile tested samples of 5A composite: (a,b) SHT, (c,d) aging at 200 C/10h.

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].

Figure 16.

(a) Stress–strain curves obtained from 5A composite, (b) stress–strain curves obtained from 5B composite, (c) % elongation as a function of testing temperatire (composite 1 is 5A, composite 2 is 5B).

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]:

$$\stackrel{.}{\epsilon }exp\left(\frac{Q_a}{RT}\right)=A{\sigma }^n=Z$$E1

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:

$${\sigma }^n=Bexp\left(\frac{Q_a}{RT}\right)$$E2

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

$$\frac{\partial ln\sigma }{\partial \left(\raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$T$}\right.\right)}=\frac{Q_a}{nR}$$E3

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

Figure 17.

Relationship of flow stress vs 1/T (T = in kelvin degree).

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].

Figure 18.

Fracture surface of tensile tested samples of 5A composite: (a) 250°C, (b) 350°C, (c) 450 C, (d) an enlarged portion of (c) showing the crack, (e) EDS corresponding to (c) revealing reflections due to Al, Zr, Ti elements.

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).

Figure 19.

Fracture behavior of 5B composite tested at: (a) 250°C, (b) 350°C, (c) 450°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 n_a (> 5) as well as the activation energy Q_a. 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 σ_o. In this case, the deformation process is related to an effective stress, σ_e = (σ -σ_o) not the applied stress σ. Therefore, equations 1–2 can be rewritten to take into consideration σ_o as follows

$$\stackrel{.}{\epsilon }exp\left(Q_t/RT\right)=Z=A^{\prime }\frac{Gb}{kT}{\left(\frac{\sigma -{\sigma }_0}{G}\right)}^{n_t}$$E4

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

Figure 20.

True stress–strain curves at different strain rates at (a) 300°C, (b) 400°C, (c) 500°C, (d) strain rate 10⁻³ s⁻¹.

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 n_a 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 n_a and Q_a threshold stress in the composite materials.

Figure 21.

Strain rate and stress relationship in the temperature range 300–500°C.

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].

Figure 22.

Fracture surface of samples tested at 300°C: (a) 10⁻² s⁻¹, (b) 10⁻⁴ s⁻¹.

5. Summary

The results obtained from the present investigations revealed that the powder injection technique used in our study proved to be effective in producing composites with a uniform distribution of B₄C particles throughout the matrix (commercial aluminum or 6063 alloy). The combined addition of Zr and Ti improved the possibility of increasing the number of B₄C particles in the matrix by improving the particulate wettability. The precipitating Al₃(Zr_1–xTi_x) particles decorating the B₄C particles were found to grow into the surrounding matrix. Precipitation of Mg₂Si in 6063/B₄C was more effective in controlling the composite toughness than Al₃Zr in the under-aging conditions. Overaging occurred at 400°C for prolonged aging times (i.e. 10 h), resulting in a significant improvement in the composite toughness regardless the type of the matrix. Cracks were always initiated at the particle/matrix interfaces and propagated either through the B₄C particles or along the protecting Al₃(Zr_1–xTi_x) layer. No particle debonding was observed regardless the type of matrix or the testing method. Formation of the Zr/Ti rich layers surrounding the B₄C particles strengthen their adhesion to the surrounding matrix. Increasing the testing temperature leads to rapid decrease in the composite strength in an exponential pattern which appeared in the gradual fracturing of the reinforcement B₄C particles. The plots of flow stress as a function of testing temperature are linear with a fitting factor of 0.955. The value of n_t ~ 5 and Q_t of 130 kJ mol⁻¹ along with subgrain formation may conclude that dislocation climb is the main controlling process. Similar observation was made in pure Al with dispersed particles. The pct elongation to failure reached a maximum value at intermediate value of Z, which can determine the optimum conditions for the composite formability.