High modulus Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite: microstructural

characteristics and micromechanics-based analysis

Sajjad Amirkhanlou1, Shouxun Ji1*, Yijie Zhang1, Douglas Watson2, Zhongyun Fan1

1Brunel Centre for Advanced Solidification Technology (BCAST), Brunel University

London, Uxbridge, Middlesex UB8 3PH, United Kingdom

2Engineering Centre, Jaguar Cars Ltd, Abbey Road, Coventry CV34 4LF, United Kingdom

*Corresponding author, Tel: +44-1895-266663; Fax: +44-1895-269758; Email:

Shouxun.Ji@brunel.ac.uk

Abstract

As an intrinsic materials property and an important criterion in structural design, the Young’s

modulus of cast aluminium alloys can be significantly increased through adding Si and Cu

elements as well as in-situ forming Mg2Si and TiB2 particles in the alloys to make castings

with complex geometry. The microstructural evaluation and mechanical properties of Al-Si-

Mg alloy and Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite were examined by X-ray

diffractometer (XRD), optical microscopy (OM), scanning and high resolution transmission

electron microscopes (SEM and HRTEM), ultrasonic pulse technique and tensile test. The

results revealed that the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite could provide the

Young’s modulus more than 94 GPa and the yield strength up to 235 MPa by forming the α-

Al (Cu, Mg), Si, Mg2Si and TiB2 phases in the microstructure. Micromechanics-based models

were also employed to explain important factors in Young’s modulus and yield strength. The

theoretical calculation confirmed that the contribution of thermal mismatch, Orowan, elastic

mismatch, load bearing and grain boundary strengthening mechanisms to the yield strength

are 65.5, 38.3, 26.2, 6.7 and 1.4 MPa, respectively. The Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid

composite showed a homogeneous distribution of strengthening phases throughout the

aluminium matrix, and a microstructure without any detrimental phases. The interfaces

between the aluminium matrix and the TiB2 particles are clean, smooth and well-bonded. TiB2

particles tend to have hexagonal and cubic prism shapes with chamfered edges and corners

due to the specific growth rate of facets.

Keywords: A. Metal matrix composites; B. Liquid-solid reactions; C. Elasticity; C.

Microstructure; C. Mechanical properties; D. Transmission electron microscopy, TEM.

1- Introduction

Weight reduction through applying aluminium structural components in aerospace and

automobile industries is one of the most promising ways to decrease energy and fuel

consumption [1-3]. These structural components, in particular shaped castings, are usually

designed on the criteria of either yield strength (YS) or Young’s modulus (E) of metallic

materials. When the yield strength is the design criterion, aluminium alloys with much higher

strength than pure aluminium are commercially available and these can be selected for

industrial applications[3]. However, when the modulus is the design criterion, there are

limited options for the aluminium alloys with significantly increased Young’s modulus

compared with that of conventional aluminium alloys with 70 GPa modulus [4, 5].

Furthermore, almost all the strengthening mechanisms which result in a significantly

improvement in yield strength have no positive effect on the Young’s modulus since the

modulus is predominantly governed by the nature of interatomic bonding [6-8]. This has

limited the applications of aluminium alloys in the shaped castings and components that

require high modulus to achieve further weight reduction.

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Although the Young’s modulus of aluminium alloys cannot be significantly increased by

mechanical working and heat treatment processes [9], it can be remarkably improved by

changing the atomic structure of aluminium alloys to form strong interatomic bonding, and/or

introducing high modulus phases in aluminium matrix. Generally, high modulus phases can

be introduced to the aluminium matrix through alloying elements and/or ceramic particles to

form aluminium matrix composites (AMCs) [10, 11]. Li, Be, Si and Cu have been found to

be able to increase the Young’s modulus of aluminium alloys [12-14]. However, Al-Li and

Al-Be alloys are brittle, expensive, toxic and there are difficulties in making shaped castings

with complex geometry [15, 16]. Among the available elements, Si and Cu are the favourite

candidates to enhance the modulus of aluminium alloys because the presence of Si and Cu

can form solid solution and second phases with increased Young’s modulus. More

importantly, the aluminium alloys with Si and Cu are capable of casting complex shapes.

However, the increases in Young’s modulus are usually less than 15% by adding Si and Cu

[17]. In order to further improve the Young’s modulus, aluminium alloys need to be

reinforced by adding high modulus phases such as Al2O3 (E=400 GPa), SiC (E=480 GPa),

TiB2 (E =560 GPa) into aluminium matrix. The capability of making complex shaped

castings of these materials depends on the introduction methods, processing methods, volume

fraction, size and distribution of high modulus phases. When considering the massive

production, the casting approach is favourable because of the low cost and high productivity

[18].

One of the most attractive ways to achieve desirable mechanical properties is to use in-situ

precipitations of reinforcement phase since the reactions taking place within the molten metal

can generate strong bonding at interfaces [19]. Among all the high modulus reinforcement,

TiB2 is one of the popular candidates due to its ultra-high Young’s modulus and the capability

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of forming thermodynamically stable phase. In addition, other high modulus phases such as

Mg2Si, having Young’s modulus of 120 GPa, can be considered as a suitable phase to

improve the modulus of aluminium alloys [20, 21]. Aluminium matrix composites reinforced

by either TiB2 or Mg2Si particles have been investigated, however, most of the existing

studies are concentrated on processing methods [22], wear resistances [23], high temperature

properties [24] and thermodynamic evaluations [25]. Limited attempts have been performed

for Young’s modulus improvement, which has different emphasis in terms of the phase

formation, reinforcement size and distribution, and particle/matrix interface.

Therefore, the aim of the present work is to investigate the capability of fabricating high

modulus aluminium-based alloys. In comparison with the popular gravity die cast Al-Si-Mg

(A356) alloy, the Si and Cu elements were increased and Mg2Si and TiB2 particles were

introduced to improve the modulus of aluminium alloys. The microstructures and mechanical

properties of the experimental materials, in particular the interface between the reinforced

phases and the aluminium matrix were studied by optical, scanning electron and high

resolution transmission electron microscopes. Moreover, the Young’s modulus and yield

strength of the experimental materials were theoretically analysed by micromechanics-based

models, and the calculated data were compared with the experimental results. The discussion

is focused on the understanding of strengthening mechanism and the contribution of different

mechanisms on the mechanical properties of the experimental materials.

2- Experimental

The Al-9Si-1Mg-0.7Cu/TiB2-Mg2Si (Al-Si-Mg-Cu/TiB2-Mg2Si hereafter) hybrid composite

was produced by adding the mixture of KBF4 and K2TiF6 into the molten Al-7Si-0.3Mg alloy

with the addition of Si, Mg and Cu through the exothermic reaction. Two types of salts were

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added into the molten aluminium alloy in the atomic ratio in accordance with Ti/2B ratio.

The salts were maintained at 850 °C for 30 min to complete the reactions. Then the melt was

degassed by a rotatory machine at 500 rpm for 3 min, followed by casting tensile samples in a

standard ASTM B108 mould. The alloy was reinforced with 8.9 vol.% TiB2 and 1.6 vol.%

Mg2Si particles. The Al-7Si-0.3Mg alloy (Al-Si-Mg hereafter) with a composition

specification of ENAC 42100 was also fabricated for comparison. The Al-Si-Mg alloy was

chosen for comparison because ENAC 42100 has been broadly used in industry and was the

target materials to be replaced by the developed materials. Both the Al-Si-Mg-Cu/Mg2Si-TiB2

hybrid material and the Al-Si-Mg alloy were solution-treated at 525 °C for 9 hours and

immediately quenched into cold water to achieve T4 condition. The chemical compositions

of the experimental materials were analysed using a Perkin-Elmer Optima 5300 dual view

ICP-AES and the results are given in Table 1.

The microstructures were characterized by optical microscopy (OM, Zeiss, Jena, Germany),

high-resolution transmission electron microscopy (HR-TEM, JEOL 2100F, JEOL Ltd.

Tokyo, Japan) operated at 200 kV and field-emission scanning electron microscopy (SEM,

SUPRA 35VP, Carl-Zeiss Company, Jena, Germany) operated at 15 kV and equipped with an

energy dispersive X-ray spectroscopy (EDS, EDAX International Company). Thin foils

required for TEM were mechanically ground and punched into 3 mm discs with an average

thickness of less than 100 μm. The discs were subsequently ion beam thinned using a Gatan

precision ion polishing system (PIPS) at 5.0 kV and at an incident angle of 4°. The X-ray

patterns of the specimens were recorded by a D8 advanced Bruker X-ray diffractometer

(Bruker Corporation, Billerica, Massachusetts, United States) with CuKα radiation in the

range of 20-90° using a step size of 0.05o and a counting time of 1s per step. Consequently,

XRD patterns were analysed using X’Pert High Score software. Clemex Vision Image

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analysis software was utilized to determine different attributes of the reinforcements and

grains. These included the volume fraction, average grain and particle sizes, and the aspect

ratio of reinforcement particles. The average size of particles and grains was defined as the

diameter of a circle having the same area as that of the particle and was calculated using Eq.

(1). The aspect ratio of particles was measured as the ratio of the length of its longest feret to

the length of its shortest feret as illustrated by Eq. (2). The average diameter and aspect ratio

were taken from at least 500 measurements.

Average diameter = √ 4(Area )π

(1)

Aspect ratio = Lengthof longest feretLengthof shortest feret (2)

Young’s modulus of the specimens was measured by a dynamic method namely ultrasonic

pulse technique according to ASTM E1875-13 standard. Parallel smooth surfaces in the

specimens were shaped by grinding and polishing before measuring the transversal and

longitudinal wave velocities. The wave velocities of the specimens were measured using a

38DL PLUS Ultrasonic Thickness Gage (Olympus Industrial) and the Poisson’s coefficient (

ϑ ) and Young’s modulus (E) were calculated according to the following equations [26]:

Poisso n' s ratio(ϑ )=1−2(V T /V L)

2

2−2(V T /V L)2 (3)

Elastic modulus(E , Gpa)=V L

2 ρ(1+ϑ )(1−2ϑ )1−ϑ

(4)

where V T is transverse velocity, V L is longitudinal velocity and ρ is density.

The tensile test were conducted according to the ASTM E8/E8M standard using an Instron

5500 Universal Electromechanical Testing Systems equipped with Bluehill software and a 50

kN load cell. The tensile test was carried out at a nominal strain rate of 1.6×10-1s-1 and at

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ambient temperature (~25 °C). The total elongation of specimens was measured from the

difference in the gauge length before and after testing via an extensometer. To have accurate

results, three tensile experiments were conducted on each specimen and the average was

taken as the result.

3- Results and discussion

3-1- Microstructural evaluation

X-ray diffraction patterns of the Al-Si-Mg alloy and Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid

composite are shown in Figure 1. It is clear that the main phases of the hybrid composite are

α-Al, β-Si, Mg2Si and TiB2, while the Al-Si-Mg alloy consists of only α-Al and β-Si phases.

The characteristic peaks of the Al-Si-Mg-Cu/Mg2Si-TiB2 composite represent high intensities

and are well matched with JCPDS cards (Mg2Si, No. 00-035-0773 and TiB2, No. 01-075-

0967). The reactions between the molten aluminium and the K2TiF6 and KBF4 salts lead to

the formation of TiB2 phases. These reactions have been reported as follows [27]:

3 K 2Ti F6+13 Al→ 3Ti Al3+3 KAl F4+ K3 Al F6 (5)

2 KB F4+3 Al → Al B2+2 KAl F4 (6)

Al B2+Ti Al3 →Ti B2+4 Al (7)

Gibbs free energy of TiB2 is more negative than that of other potential phases such as Al3Ti

and AlB2 during solidification temperature of the hybrid composite [28], which

thermodynamically resulted in the formation of TiB2 phase. In order to achieve the maximum

potential of particulate reinforced aluminium matrix composites, it is important to avoid

forming any undesirable phases, which are usually observed in cast aluminium matrix

composites. The phases such as Al3Ti have been identified to be detrimental on the

mechanical properties of cast Al/TiB2 composites due to their brittle nature [29]. The XRD

pattern of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite in Figure 1 shows no hint of

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undesired phases in the microstructure, which confirms that the materials manufacturing has

been properly controlled at the reaction conditions with appropriate Ti/B ratio. Figure 1b and

c show that the α-Al peaks are shifted toward the lower angles in the Al-Si-Mg-Cu/Mg2Si-

TiB2 hybrid composite compared with that of the Al-Si-Mg alloy. According to Bragg's

diffraction law, nλ=2 d sinθ, where d=a/√h2+k2+l2 for cubic crystals, when the diffraction

angle (θ) of specific plane (h ,k , l) decreases, it means that lattice constant (a) has been

increased. Therefore, the displacement of the α-Al peaks (θ) is due to solutionising more

alloying elements in the α-Al lattice. This implies that the α-Al lattice expands as Cu and/or

Mg atoms diffuse therein, forming the α-Al (Cu, Mg) solid solution with a structure similar to

that of α-Al but with a slightly larger lattice constant.

Figures 2 and 3 show the optical and SEM micrographs of the Al-Si-Mg alloy and Al-Si-Mg-

Cu/Mg2Si-TiB2 hybrid composite in two magnifications. The figures show that the typical

microstructure of the Al-Si-Mg alloy consists of the needle/plate shaped eutectic β-Si phase

surrounded by α-Al dendrites. Optical micrographs represent eutectic β-Si phase, whereas

back-scattered SEM images reveal the size and distribution of the TiB2 phases, as shown in

Figure 3c and d. It is obvious that homogeneous distribution of the reinforcement, in

macroscopic scale, has been successfully achieved in the hybrid composite. Although there

are many individual particles in the aluminium matrix, TiB2 particles are partially segregated

in the eutectic β-Si regions. This phenomenon is attributed to the particle pushing mechanism

during solidification [30]. When aluminium alloys melt solidifies, the TiB2 and β-Si particles

are pushed away by the solidification front of solid α-Al dendrites and segregated within the

inter-dendritic and eutectic regions. However, according to Figure 3d, many of the fine TiB2

particles are engulfed by the α-Al phase due to their excellent wettability. The individual TiB2

particle forms strong interfaces with the aluminium matrix and can significantly resist against

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larger stresses during deformation. Figure 2 also shows a decrease in the average grain size

with the addition of reinforcement to the matrix. The average size of α-Al grains is 25 µm in

Al-Si-Mg alloy and 18 µm for in Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite.

EDS microanalysis of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite, as shown in Figure 4,

confirms the formation of Si, TiB2 and Mg2Si phases. High magnification SEM micrographs

of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite are shown in Figure 5, representing the

TiB2 phase in the morphology of clusters and/or individual particles. Different shapes of TiB2

particles can be seen in the micrographs. In general, the TiB2 particles exhibit hexagonal

prism in Al/TiB2 composites, which has been reported in ref. [31, 32]. However, apart from

the hexagonal TiB2 particles, it is clear that some of the reinforcement phases in Figure 5 tend

to have a cubic prism or other polyhedron prisms. Based on the Wuff’s theorem [33], TiB2

crystals tend to form the minimum surface energy in the equilibrium state. Close-packed

planes always show low surface energy. On the other hand, based on Bravais-Friedel law, the

number of a specific crystal plane increases with its compactness [34]. Abdel-Hamid et al.

[35] reported that the elementary growth layers for TiB2 are {0001}, {10 1 0 }, and {10 11}

and close-packed planes run as follows: {0001 }> {10 10 }>{10 11}. As a result, microstructural

observation of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite, as shown in Figure 5 and 6, is

validated by theoretical models that the crystal morphology of TiB2 tends to be hexagonal

prisms with {0001} basal and {10 1 0 } prism planes. A careful examination by TEM

represented in Figure 5c and 6 confirms that almost all of the TiB2 particles have chamfered

edges and corners. This means that the growth rate of the facets is not capable of obtaining

perfect hexagonal TiB2 particles. It has been reported by Sun et al. [32] that the growth rate

of the faces, edges and the corners should maintain a certain relationship in order to achieve

complete shape of the hexagonal TiB2 particles. Otherwise, higher-index {10 1 1}, {1 213 }

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and/or {1120 } planes form at the edges and corners of the TiB2 particles, as shown in Figure

7. SEM and TEM observations confirm that the sizes of the TiB2 particles are from 25 nm to

3 µm and most of the particles are in nanoscale, and thereby, the average particle size is 385

nm. A typical interface between the aluminium matrix and the TiB2 particle is shown in

Figure 7, which confirms the clear and well-bonded interface is formed between the

reinforcement phase and the matrix. The projection axis is [155]Al /¿[1213]TiB2.

Indeed, in order to achieve maximum potential of reinforced materials with desirable

mechanical properties, several important factors are required [36, 37]: (a) homogeneous

distribution of reinforcement in the matrix without particle free zones, (b) clear and well-

bonded interface between the reinforcement and the matrix without any undesirable reactions

between them, and (c) lack of porosity in the created materials. According to the

microstructural observations in the present study, it is achievable for the

Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite without any noticeable porosity and

detrimental phases.

3-2- Mechanical properties

Young’s modulus can be determined statically (tension, compression, torsion or flexure), or

dynamically through the study of vibrating bars, or by measuring the velocity of sound

passing through the materials. However, it has been stated [38][1] that dynamic methods have

advantages over the traditional static methods. The dynamic methods don’t alter stresses on

the materials above the elastic limit, which eliminate complex creep effects or elastic

hysteresis. Therefore, elastic constants calculated by dynamic methods are more accurate in

comparison to static techniques. Young’s modulus of the specimens measured by ultrasonic

pulse technique is shown in Table 2. The Young’s modulus of the Al-Si-Mg-Cu/Mg2Si-TiB2

10

hybrid composite is 94.7 GPa, significantly higher than that of the Al-Si-Mg alloy with 74

GPa modulus. The increase is 28 % and the significant improvement can be attributed to the

presence of high modulus Mg2Si (120 GPa) and TiB2 (560 GPa) phases in the microstructure

and the formation of α-Al (Cu, Mg) phase, as shown in Figure 1. In the elastic deformation

stage, the applied loads are shared by the α-Al (Cu, Mg) and high modulus Mg2Si and TiB2

phases, resulting noteworthy improvement of the modulus.

The Young’s modulus of composite materials has been studied by means of several

theoretical models to deal with the morphological arrangement of reinforcement [39]. Among

them, Halpin-Tsai model [19] is the frequently discussed ones for particulate metal matrix

composites. In the Halpin-Tsai model, in predicting the Young’s modulus of the composites,

besides the volume fraction and Young’s modulus of the particles, the aspect ratio of the

reinforcement is also taken into consideration:

E=Em( 1+2ξηV p

1−ηV p) (8)

η=(E p/Em−1)/(E p/Em+2ξ) (9)

where Ep and Em are the Young’s modulus of the particles and the matrix, respectively. V p is

the volume fraction of the particles (V p (Mg2 Si)=0.016and V p (TiB2)=0.089), ξ refers to the aspect

ratio of the particulate reinforcement. According to the microstructural studies, an aspect ratio

of 1.5 can be considered for both Mg2Si and TiB2 particles. Predicted value based on the

Halpin-Tsai model for the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite is 93 GPa, which is in

good agreement with the experimental value of 94.7 GPa.

Figure 8 and Table 3 demonstrate the tensile properties of the Al-Si-Mg alloy and Al-Si-Mg-

Cu/Mg2Si-TiB2 hybrid composite. As seen, a great enhancement in the yield strength (YS)

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and the ultimate tensile strength (UTS) is observed in the alloy reinforced by Mg2Si and TiB2.

The yield strength is increased from 104 MPa to 235 MPa. The strength increment is

partially due to the increase of Si and Cu elements in the alloy. Analysis of this feature is

beyond the scope of the present article, but it is well-presented in literature [40-42] that the

addition of Si and Cu to the aluminium alloys increases the strength mainly by precipitation

strengthening mechanism and the formation of intermetallic phases [43, 44]. In the present

investigation, effect of Mg2Si and TiB2 reinforcement particles on the strength is the main

consideration. The interfaces between the aluminium matrix and the reinforcement of Mg2Si

and TiB2 phase are incoherent since the structure and lattice constants of the reinforcement

are different from that of the aluminium matrix. Therefore, the strengthening for the hybrid

composite can be divided into (a) dislocation strengthening due to thermal mismatch, (b)

elastic mismatch, (c) dislocation-nanoparticle interactions by the Orowan process, (d) load

transfer from matrix to the particles, and (e) grain boundaries strengthening mechanism

according to Hall-Petch equation [51].

(a) Dislocation strengthening due to thermal mismatch (TM): dislocation density generated

by the difference in thermal expansion coefficient (CTE) between the aluminium matrix and

the particles causes an increase in yield strength [45]. The CTE of the aluminium matrix (

23 ×10−6 K−1) is different from that of the TiB2 (7.8 ×10−6 K−1) and Mg2Si (7.5 ×10−6 K−1)

particles. Assuming that dislocations are homogeneously generated throughout the matrix and

that all thermal stresses are relieved by the generation of dislocations, an increase in the yield

strength due to thermal mismatch (∆ σTM) can be given by the following equation [46]:

∆ σTM=αGb√ ρTM (10)

where α is the average value of dislocation strengthening efficiency (∼ 1 [47]), G is shear

modulus (~ 25.4 GPa for aluminium matrix), b is the Burgers vector (=0.286 nm [48, 49])

12

and ρ is dislocation density. Dislocation density induced by thermal mismatch can be

calculated by [50]:

❑TM=12 ∆ Tb [ V p(Mg2 Si)∆C (Mg2 Si)

(1−V p (Mg2 Si))d p (Mg2 Si)

+V p(TiB2 )

∆C (TiB2)

(1−V p (TiB2))d p (TiB2)] (11)

where V p is the volume fraction of particles, d p is the average diameter of the particles (385

nm for TiB2 and 8 µ for Mg2Si), ∆ T is the difference between the processing temperature

(525 ºC) and test temperature (25 ºC), ∆ C is the difference in CTE between the particles and

the aluminium matrix. The calculated ∆ σTM for the hybrid composite is 65.6 MPa.

(b) Dislocation strengthening due to elastic mismatch (EM): Eq. (12) estimates the density of

dislocation generated by modulus mismatch. The stress contribution, ∆ σEM , can be expressed

by Eq. (13) [51]:

ρEM=8 εb [V p (Mg2 Si)

d p(Mg2 Si)+

V p (TiB2)

d p(TiB2)] (12)

∆ σEM=αGb√ ρEM (13)

where is the yield strain (0.2%) and ❑EM is the density of dislocations caused by elastic

mismatch. Elastic mismatch strengthening is calculated to be 26.2 MPa for the Al-Si-Mg-

Cu/Mg2Si-TiB2 hybrid composite.

(c) Orowan strengthening: Orowan mechanism corresponds to the interaction of the

nanoparticles and dislocations, in which nanoparticles pin the motion of dislocations,

resulting in significant improvement of the yield strength. In general, the particle size should

be smaller than 1 µm to significantly contribute in Orowan strengthening mechanism. As the

average Mg2Si particle size is 8 µm in the present work, the contribution of Mg2Si particles to

the Orowan mechanism is negligible. ∆ σOrowan is the contribution of Orowan strengthening

13

from the presence of TiB2 nanoparticles and can be calculated using Orowan-Ashby equation

[52, 53]:

∆ σorowan=0.4 MGb ln(√2/3d p(TiB2)

b )π √2/3d p (TiB2)

(√π /4V p(TiB2 )−1)√1−ϑ

(14)

where M is the Taylor factor (=3.06 for aluminium matrix [47]) and ϑ is the Poisson’s ratio

(=0.3). The contribution of Orowan strengthening mechanism to the final yield strength is

38.3 MPa.

(d) Grain boundary strengthening (∆ σGB): Hall-Petch equation describes the effect of grain

size on the yield strength of materials, as shown in following well-known equation [54]:

∆ σGB=k y (DG)−12 (15)

where DGis average grain size, k y is constant and typically equals to 40 MPa √ μm for

aluminum alloys [55, 56]. Therefore, the contribution of grain boundary strengthening is 1.4

MPa for the cast Al-Si-Mg-Cu/Mg2Si-TiB2 nanocomposite.

(e) Load-bearing effect: The contribution of load-bearing in increasing of the yield strength is

expressed by [51, 55]:

∆ σL=0.5 [V p (Mg2 Si)+V p (TiB2)] σ m (16)

where vp and σ m are the volume fraction of particles and matrix yield strength, respectively.

∆ σL is 6.7 MPa for the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite.

The total yield strength can be calculated by the Clyne models referred as arithmetic

summation (Eq. 10) and quadratic summation (Eq. 11) [51]:

σ Arith.=σ m+∆ σ TM+∆ σ EM +∆ σ Load+∆ σGB+∆ σOrowan (17)

σ Quad.=σ m+√(∆ σTM )2+(∆ σEM )2+(∆ σ Load)2+(∆ σ GB)

2+(∆ σOrowan)2 (18)

14

The contribution of each strengthening mechanism on the yield strength, the experimental

results as well as the yield strength obtained by arithmetic summation and Clyne models are

summarised in Table 4. Thermal mismatch and Orowan strengthening mechanisms provide

the significant effect to enhance the yield strength in the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid

composite. These two strengthening mechanisms prove the importance of (1) TiB2

nanoparticles introduced into the matrix and (2) dislocation density formed in the matrix due

to thermal mismatch during heat treatment process.

3-3- Fractography

Figure 9 shows the SEM fractographs after tensile test for the Al-Si-Mg alloy and the Al-Si-

Mg-Cu/Mg2Si-TiB2 hybrid composite. As shown in Figure 9a and b, large dimples in the

fractured surface of the Al-Si-Mg alloy indicate the existence of beneficent plastic

deformation. The hybrid composite shows a quasi-cleavage fracture, characterised by the

cleavage fracture of Si and small and shallow dimples of aluminium phase. The clusters of

TiB2 particles are also observed in the fracture surface (Figure 9c). The clustered TiB2

particles in the aluminium matrix can lead to a much higher crack formation, which can

ultimately reduce the ductility. However, TiB2 particles identified at high magnification

micrographs in Figures 9d and e confirm that there is no cracking or significant debonding

between the TiB2 particles and the aluminium matrix during fracture. These observations

provide the firm evidence of strong bonding between the individual TiB2 particle and the

aluminium matrix.

For complement of fracture analysis, the cross section of the fractured tensile specimens is

shown in Figure 10. The cracks are initiated from the Si eutectic phase in the Al-Si-Mg alloy,

while the cracks are nucleated from the clustered TiB2 particles in the hybrid composite.

15

However, it is clear that the cracks tend to propagate through eutectic Si phase for both the

Al-Si-Mg alloy and the hybrid nanocomposite. In general, based on fractographic analysis,

extensive matrix plastic flow around the TiB2 particles can be observed, showing strong

interfacial bonding between the TiB2 particles and the aluminium matrix. The reported data in

literature [57, 58] show that the ductility usually is less than 0.5 % when the Young’s

modulus is more than 85 GPa for particulate reinforced aluminium matrix composites.

However, the materials developed in the present study can provide reasonable ductility of 1.2

% when the Young’s modulus reaches to 94.7 GPa. This is reasonably good to satisfy

industrial requirement. Of course, the improvement of ductility in the developed materials is

important for industrial application and it can be investigated in future.

4- Conclusions

(1) The Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite can provide the Young’s modulus

over 94 GPa and the yield strength up to 235 MPa by the formation of α-Al (Cu, Mg), Si,

Mg2Si and TiB2 phases in the microstructure.

(2) The theoretical calculation confirms that the contribution of thermal mismatch, Orowan,

elastic mismatch, load bearing and grain boundary strengthening mechanisms to the yield

strength are 65.5, 38.3, 26.2, 6.7 and 1.4 MPa, respectively. The calculated yield strength

is in good agreement with the experimental results.

(3) The microstructure of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite contains no

detrimental phases for property enhancement. The strengthening phases are

homogeneously distributed throughout the aluminium matrix. TiB2 particles are in the

range from 25 nm to 3 µm and the average particle size is 385 nm. The interfaces

between the aluminium matrix and the TiB2 particles are clean, smooth and well-bonded.

16

TiB2 particles show hexagonal and cubic prism shapes with chamfered edges and corners

due to the specific growth rate of the facets.

(4) The hybrid composite shows a quasi-cleavage fracture, characterised by the cleavage

fracture of the Si phase and small shallow dimples of the aluminium phase.

Acknowledgment

Financial support from Jaguar Range Rover (JLR) [grant number R33232] is gratefully

acknowledged.

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Table 1. Chemical composition of the experimental materials identified by ICP-EDS (wt.%).

Elements Si Cu Mg Fe Ti Mn Zn B Al

Al-Si-Mg alloy 7.13 ˂0.01 0.28 0.08 0.11 ˂0.01 0.01 ˂0.01 Bal.

Al-Si-Mg-Cu/Mg2Si-

TiB2 hybrid composite9.11 0.73 1.00 0.2 7.74 0.02 0.03 3.49 Bal.

Table 2. Density, velocity parameters and Young’s modulus of the experimental materials

measured by Archimedes’ principle and ultrasonic method.

PropertiesDensity

(g/cm3)

Longitudinal

velocity

(mm/µs)

Shear

velocity

(mm/µs)

Poisson'

s ratio

Young’s

modulus

(GPa)

Al-Si-Mg alloy 2.651 6.34 3.25 0.3 74.02

Al-Si-Mg-Cu/Mg2Si-

TiB2 hybrid composite2.805 6.80 3.59 0.3 94.72

Table 3. Mechanical properties of the Al-Si-Mg alloy and the Al-Si-Mg-Cu/Mg2Si-TiB2

hybrid nanocomposite.

Property 0.2% Yield Strength (MPa) Ultimate tensile strength (MPa) Elongation (%)Young’s

modulus (GPa)

Al-Si-Mg alloy 104 222 10.7 74

Al-Si-Mg-Cu/Mg2Si-TiB2

hybrid composite235 305 1.2 94

25

26

Table 4. Micromechanics-based analysis of Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid nanocomposite.

Experimental results

/ModelsEquations Values

Grain boundary (∆ σGB): ∆ σGB=k y (DG)−12 1.4 MPa

Thermal mismatch (∆ σTM

)

∆ σTM=αGb√ ρTM

❑TM=12 ∆ Tb [ V p ( Mg2 Si ) ∆ C( Mg2 Si )

(1−V p ( Mg2 Si ) )d p ( Mg2 Si )+

V p ( TiB2 ) ∆ C( TiB2 )

(1−V p (TiB2) ) dp ( TiB2 ) ]65.6 MPa

Elastic mismatch (∆ σEM )

∆ σEM=αGb√ ρEM

ρEM=8 εb [V p (Mg2 Si)

d p(Mg2 Si)+

V p (TiB2)

d p(TiB2)] 26.2 MPa

Orowan (∆ σOrowan)∆ σorowan=

0.4 MGb ln(√2/3 d p(TiB2)

b )π √2/3d p (TiB2)

(√π /4V p(TiB2 )−1)√1−ϑ

38.3 MPa

Load-bearing (∆ σLoad) ∆ σL=0.5 [V p (Mg2 Si)+V p (TiB2)] σ m 6.7 MPa

Arithmetic summation

yield strength (σ Arith.)σ Arith.=σ m+∆ σ TM+∆ σ EM +∆ σ Load+∆ σGB+∆ σOrowan 242 MPa

Quadratic summation yield

strength (σ Quad.)σ Quad.=σ m+√(∆ σTM )2+(∆ σEM )2+(∆ σ Load)

2+(∆ σ GB)2+(∆ σOrowan)

2185 MPa

Experimental yield

strength (σ Ex .)- 235 MPa

Halpin-Tsai modulusE=Ep( 1+2 ξηV p

1−η V p)

η=(E p/Em−1)/(E p/Em+2 ξ)

93 GPa

Experimental Young’s - 94.7 GPa

27

modulus (σ Ex .)

28

Figure 1. X-Ray diffraction patterns of the Al-Si-Mg alloy and the Al-Si-Mg-Cu/Mg2Si-TiB2

hybrid composite.

Figure 2. Optical micrographs of (a and b) the Al-Si-Mg alloy and (c and d) the Al-Si-Mg-

Cu/Mg2Si-TiB2 hybrid composite.

29

Figure 3. Backscattered SEM micrographs of (a and b) the Al-Si-Mg alloy and (c and d) the

Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite.

30

Figure 4. (a) Backscattered SEM micrograph showing the morphology of identified phases in

the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite, (b), (c) and (d) EDS results corresponding to

the phase at spots 1, 2 and 3 shown in (a), respectively.

(a)

31

Figure 5. TEM and high magnification TEM images of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid

composite showing (a) a TiB2 cluster, and (b) and (c) different shapes of TiB2 particles.

32

Figure 6. Bright-field TEM micrographs of the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid

nanocomposite showing the different shapes of TiB2 nanoparticles, (a) cuboid and (b) cubic

with chamfered corners and (c) hexagonal TiB2 nanoparticles. The corresponding Fast

Fourier Transform (FFT) patterns are given as the insets at the top-right corners.

33

Figure 7. (a) High resolution TEM image representing Al/TiB2 interface being viewed in

[155]Al /¿[12 13]TiB2 projection axis, (b) the corresponding selected area diffraction pattern

(SEDP).

34

Figure 8. Engineering stress-strain curves of the Al-Si-Mg alloy and the Al-Si-Mg-Cu/Mg2Si-

TiB2 hybrid composite.

35

Figure 9. SEM micrographs of fractured surface for (a and b) the Al-Si-Mg alloy and (c-e)

the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite.

36

Figure 10. Optical micrographs showing the cross section perpendicular to the fractured

surface, (a) the Al-Si-Mg alloy and (b) the Al-Si-Mg-Cu/Mg2Si-TiB2 hybrid composite.

37