Microstructurally inhomogeneous composites: Is a homogeneous reinforcement distribution optimal?

Progress in Materials Science 71 (2015) 93–168

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Microstructurally inhomogeneous composites: Is a homogeneous reinforcement distribution optimal? L.J. Huang a,⇑, L. Geng a, H-X. Peng b,⇑ a

State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, P.O. Box 433, Harbin 150001, PR China Institute for Composites Science Innovation (InCSI), School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China b

a r t i c l e

i n f o

Article history: Received 28 July 2014 Accepted 15 December 2014 Available online 11 February 2015 Keywords: Metal matrix composites Inhomogeneous microstructure Hierarchical structures Bi-continuous Network microstructure Mechanical properties

a b s t r a c t Since the 1960s, it has been a common practice worldwide to pursue a homogeneous distribution of reinforcements within a matrix material, discontinuous metal matrix composites (DMMCs) in particular. Taking an overview of the worldwide activities in DMMC research, despite many favourable attributes such as improved specific strength, stiffness and superior wear resistance, DMMCs with a homogeneous microstructure tend to exhibit a very low room temperature damage tolerance even with a highly ductile matrix material such as aluminium. In this review, a range of uniquely multi-scale hierarchical structures have been successfully designed and fabricated by tailoring reinforcement distribution for DMMCs in order to obtain superior performance. A variety of specific microstructures that were developed in Al, Mg, Cu, Fe, Co and TiAl matrices indicate that there must be adequate plastic regions among the reinforcements to blunt or deflect cracks if one wants to toughen DMMCs. Following this path, aided by theoretical analyses, the most recent success is the design and fabrication of a network distribution of in situ reinforcing TiB whiskers (TiBw) in titanium matrix composites (TMCs), where a tailored three-dimensional (3D) quasi-continuous network microstructure displays significant improvements in mechanical properties. This resolves the brittleness surrounding TMCs fabricated by powder metallurgy. It is the large reinforcement-lean regions that remarkably improve the composite’s ductility by bearing strain, blunting the crack and decreasing the crack-propagation rate. The fracture,

⇑ Corresponding authors. E-mail addresses: [email protected] (L.J. Huang), [email protected] (H-X. Peng). http://dx.doi.org/10.1016/j.pmatsci.2015.01.002 0079-6425/Ó 2015 Elsevier Ltd. All rights reserved.

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strengthening and toughening mechanisms are comprehensively elucidated in order to further understand the advantages of such an inhomogeneous microstructure, and to justify the development of novel techniques to produce such inhomogeneous microstructures. This approach opens up a new horizon of research and applications of DMMCs and can be easily extended to general multiphase composites with enhanced physical and mechanical properties. Ó 2015 Elsevier Ltd. All rights reserved.

Nomenclature MMCs DMMCs TMCs DRA DRTMCs TiBw TiCp CNTs TEM SEM EBSD EDS XRD PM CAPAS 3D phase a phase b H–S H–T VL RoM EHS-Upper EHS-Lower FEM MFP COD YS UTS RPS GEM HAP IT ESE

metal matrix composites discontinuous metal matrix composites titanium matrix composites discontinuously reinforced aluminium matrix composites discontinuously reinforced titanium matrix composites TiB whiskers TiC particles carbon nanotubes transmission electron microscopy scanning electron microscopy Electron Backscattered Diffraction Energy Dispersive Spectrometer X-ray Diffraction powder metallurgy current-activated pressure assisted sintering three-dimensional reinforcement-rich region reinforcement-lean region Hashin–Shtrikman Halpin and Tsai local volume fraction of reinforcement rule of mixtures the elastic property of the upper H–S bounds the elastic property of the lower H–S bounds finite element model the mean free path crack opening displacement yield strength ultimate tensile strength relative particle size general effective media hydroxyapatite ice-template elastic strain energy

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Contents 1. 2.

3.

4.

5.

6.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95 Fundamental aspects of inhomogeneous microstructures in two-phase composites . . . . . . . . . . . . . . . 97 2.1. Reconsidering the matrix phase and reinforcement phase. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 97 2.2. Fundamentals behind the design of inhomogeneity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 98 Classification of inhomogeneous phases. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99 3.1. Isolated reinforcement-rich phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101 3.2. Bar and laminated/ring-like reinforcement-rich phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101 3.3. 3D network reinforcement-rich phase with isolated reinforcement-lean phase. . . . . . . . . . . . . 103 3.4. Bi-continuous reinforcement-rich phase and reinforcement-lean phase . . . . . . . . . . . . . . . . . . . 103 3.5. Theoretical justification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104 Composites with tailored inhomogeneous microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 107 4.1. Tailored microstructure with isolated reinforcement-rich phase. . . . . . . . . . . . . . . . . . . . . . . . . 107 4.2. Bar and laminated/ring-like microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111 4.2.1. Bar reinforcement-rich phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111 4.2.2. Laminated and bean-like reinforcement-lean phases. . . . . . . . . . . . . . . . . . . . . . . . . . . 113 4.2.3. Alternating macro-ring microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 118 4.3. Tailored microstructures with reinforcement-rich network . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119 4.4. Bi-continuous microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123 4.5. Bio-inspired hierarchical microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 130 Microstructural design of TMCs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133 5.1. TiBw–Ti composite with isolated TiBw-rich (dual matrix) microstructure . . . . . . . . . . . . . . . . 134 5.2. 3D continuous microstructure with dense ceramic network . . . . . . . . . . . . . . . . . . . . . . . . . . . 134 5.3. TMCs with 3D continuous reinforcement-rich network . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 138 5.3.1. Design and fabrication . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 138 5.3.2. Microstructural characteristics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 142 5.3.3. Mechanical characteristics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 144 5.3.4. Evolution of deformation and heat treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147 5.3.5. Fracture mechanism and models. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 150 5.3.6. Strengthening and toughening mechanisms. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 151 5.4. Laminated Ti–TiBw/Ti microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 156 Future outlook on tailoring inhomogeneous microstructures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 158 6.1. Research significance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 158 6.2. Outlook of future research . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 160 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 160 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 161

1. Introduction For centuries, scientists have been conducting investigations to develop materials which are ‘stronger, stiffer, lighter and hotter (capable of use at elevated temperatures)’ [1–4]. Largely led by science, over the past 50 years the investigation of an extensive spectrum of reinforcement/matrix combinations in metal matrix composites (MMCs) has generated a vast sea of literature, especially in discontinuous MMCs (DMMCs). DMMCs with a wide range of matrix materials (including aluminium (Al), magnesium (Mg), copper (Cu), titanium (Ti) and steel (Fe) among others), and ceramic reinforcements (including borides, carbides, nitrides, oxides and their mixtures) have undergone fast development. More recently, discontinuous ceramic matrix composites and intermetallic matrix composites have subsequently followed in the footsteps of DMMCs. In principle, DMMCs can combine metallic properties, such as excellent ductility, toughness, formability and good thermal and electric conductivities, and ceramic characteristics, e.g., high hardness, strength, modulus, high-temperature durability and low thermal expansion. Therefore, DMMCs are expected to exhibit higher specific strength, specific stiffness, wear resistance, thermal stability and high-temperature durability than the corresponding monolithic matrix materials. From a commercial perspective, it is worth remembering that

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discontinuous reinforcement was introduced to bring features that were not generally available with continuous reinforcement. Besides the improvements in performance for DMMCs, the relatively low cost with ease of fabrication, an ability to shape into a final component through secondary processing, such as forging, rolling, extrusion, attract more attention than continuous reinforcement [1,4]. These superiorities make them potential candidates for critical applications in the aerospace and automotive industries. Taking an overview of the worldwide activities in DMMCs, the common features are:  The majority of activities have focused on the production of a homogeneous discrete distribution of the reinforcing particles/whiskers/short fibres, which is believed to be beneficial for damage related properties, such as ductility and fracture toughness, since numerous experimental observations have shown that inhomogeneity of the spatial arrangement of particles deleteriously affects their strength, ductility, fatigue, damage and fracture behaviours [3,4]. A problem that has received much attention is particle clustering in an isolated pattern, which is an undesirable but inevitable product of traditional DMMC manufacturing processes such as stir-casting.  Despite many favourable attributes, such as improved specific strength, stiffness and superior wear resistance, DMMCs tend to exhibit a very low room-temperature damage tolerance (energy absorbing capability) even with a highly ductile matrix material such as aluminium. Consequently, despite a number of niche applications, the applications for DMMCs have not yet emerged at the rate needed to meet the high expectations and to justify the development costs [5,6]. In particular, it is unlikely that widespread usage of these materials will occur in structural applications until greater reliability is demonstrated [7,8]. That is to say, although the majority of research activities have focused on DMMCs with a homogeneous discrete distribution of reinforcement, the expected significant improvements in performances have not materialised. Their wide applications in industry are still constrained by their limited improvement in strength and their seriously decreased damage tolerance (ductility and fracture toughness). However, it can be argued that current DMMCs have relatively low properties, i.e., the potential properties of DMMCs have so far not been fully realized. In this context, an examination of the famous Hashin–Shtrikman (H–S) bounds may offer useful insights. Hashin and Shtrikman found the best possible bounds (H–S bounds) for properties, such as effective elastic modulus [9] and conductivity [10] of isotropic two-phase composites for a given phase volume fraction. With a homogeneous discrete distribution of reinforcements, composite properties are inevitably close to the H–S lower bound, as implied by the theory. This is generally true with the current DMMCs, in particular, when the reinforcement volume fraction is in the low to intermediate range [11]. That is to say, a homogeneous discrete distribution of reinforcements inevitably results in composite properties close to the H–S lower bound. This suggests that a new approach to DMMC design is needed. One possible way out of this is to look at an inhomogeneous microstructure. It is our belief that the philosophy behind these activities is that, following in the footsteps of conventional materials processing, damage tolerance of a composite material could be improved by delaying crack initiation, blunting the crack and decreasing the crack-propagation rate. One effective way is to introduce the ductile phase around the stiffer phase to constrain and blunt the crack and bear the strain. This is coincidental with the multi-scale hierarchical structures recently proposed for future metals by Lu [12]. Driven by commercial needs, in order to overcome the poor damage tolerance and to develop socalled damage tolerant MMCs, in recent years there has been much interest in the unique design opportunities afforded by the controlled distribution of reinforcements at an intermediate (or mesoscopic) scale within a given section of material, specifically in terms of graded and layered structures, as highlighted briefly by Sinclair and Gregson in their review paper on ‘Structural performance of discontinuous MMCs’ [8]. After a great deal of effort in past decades, a range of composites with a controlled inhomogeneous microstructure were successfully designed and fabricated by different routes, the named reinforcement clustering microstructure, bi-continuous microstructure, interpenetrating microstructure, laminated microstructure, network microstructure, etc., in which the reinforcing phase is inhomogeneously distributed in the micro-scale but homogeneously or still not

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homogeneously distributed in the macro-scale [10,13,14]. As expected, these microstructures exhibited significant improvements in fracture toughness, ductility and also strength. In the present review, after the fundamental considerations and classification of inhomogeneous microstructures, the research work on the specific mesoscopic structures, that were successfully developed in Al, Mg, Cu, Fe, Co and TiAl matrices, is examined in a systematic way because these types of inhomogeneity have been addressed extensively over the past decades in recognition of their effects on MMC performance. Moreover, the recent progress in bio-inspired microstructures by learning from natural multi-scale structures is also touched upon in order to illuminate microstructural design exercises. The review then moves to recent successful microstructures, including the quasi-continuous reinforcing network boundary region and isolated toughening matrix regions, which are developed in titanium matrix composites (TMCs) via the deliberate tailoring of the reinforcement spatial distribution for property improvements. Not only the critical problem of extreme brittleness surrounding titanium alloy matrix composites fabricated by powder metallurgy (PM) is resolved to exhibit a superior ductility, but also the strengthening effects at room temperature and high temperatures are remarkably enhanced to obtain much higher tensile strength and service temperature. Therefore, most of the existing research works on those mesoscopic structures are scanned to verify the necessity of microstructure design. Moreover, we will concentrate on research works undertaken in the past on the deliberate tailoring of the reinforcement spatial distribution within DMMCs for property improvement in order to further extend successful experience and enhance comprehensive understanding between the relationships of the microstructure and performance.

2. Fundamental aspects of inhomogeneous microstructures in two-phase composites 2.1. Reconsidering the matrix phase and reinforcement phase In recent years, researchers have paid more and more attention to the three-dimensional (3D) interpenetrating microstructure designed in MMCs for performance improvement. The composites with these 3D microstructures can exploit multifunctional characteristics: each phase contributing its own properties to the macroscopic properties of the composite. The development of these materials also offers opportunities for testing the theoretical understanding of composite materials in terms of the volume fraction and phase connectivity dependence of transport properties [13]. Using Newnham’s taxonomy, which is based on phase connectivity, such materials are designated 3-3 composites since both phases have connectivity in three dimensions, which is totally different from the conventional 0-3 composites with a homogeneous microstructure [15]. In these 3-3 composites, the terms ‘matrix’ and ‘reinforcement’ are ambiguous since both phases in interpenetrating two-phase composites are three-dimensionally continuous. Traditionally, in two-phase discontinuous MMC with a homogeneous microstructure, the discontinuous phase in the form of ceramic particle/whisker/short fibre phase is termed ‘reinforcement’, while the continuous metal is termed ‘matrix’. These are easily acceptable because the reinforcement phase can improve the strength, hardness or elastic modulus of the metal matrix. These definitions are still popular in polymer matrix composites. However, in the ceramic matrix composites or intermetallic matrix composites with a homogeneous or laminated microstructure, the isolated/laminated metal or other ceramic phase was added to ceramic or intermetallic compounds in order to enhance the toughness/ductility, not the strength, due to the very high strength of the ceramic or intermetallic compounds themselves [16]. It is different in that the harder ceramic/intermetallic compound phase is the ‘matrix’ phase, while the softer metal or other ceramic phase is the ‘reinforcement’ phase which is used to toughen the ceramic/intermetallic matrix phase [17,18]. Based on the above, the isolated phase is termed the ‘reinforcement’ phase, while the continuous phase as the ‘matrix’ phase. This is consistent with all the above composites with a conventional homogeneous microstructure. However, in the composites with a tailored network microstructure, such as TiCp/Ti6Al4V composites with a network microstructure [14], the continuous phase is the harder TiCp ceramic reinforcement phase. In addition, with the increasing reinforcement volume fraction in homogeneous composites, the reinforcement connectivity increases, i.e., the reinforcement

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phase becomes gradually continuous. Although the continuous phase for the above two phenomena is the harder ceramic phase, the ceramic phase is usually termed the ‘reinforcement’ phase, while the discrete metal phase is the ‘matrix’ phase, because the ceramic phase is added to strengthen the metal matrix phase. Moreover, it is still confusing to define the ‘reinforcement’ and ‘matrix’ phases for a bicontinuous microstructure in which both phases are continuous [13]. The other important view is that the phase with a higher volume fraction is termed the ‘matrix’ phase, while the phase with a lower volume fraction is the ‘reinforcement’ phase [13]. This view can resolve the above problems, even agreeing with most composites, including the conventional composites with a homogeneous microstructure and composites with tailored microstructures. However, in order to obtain significant improvements in the elastic modulus, wear resistance or high-temperature durability, SiCp ceramic phase with a volume fraction higher than 50 vol.% was deliberately added into the aluminium matrix to fabricate aluminium matrix composites [19]. In addition, 50 wt.%, 60 wt.%, 70 wt.% and 80 wt.% TiBw reinforced Ti composites (the weight fraction is similar to the volume fraction for the Ti–TiBw system due to their similar densities) were also successfully fabricated by self-propagating high-temperature synthesis, the composites also termed as TMCs [20]. In order to simplify and unify the concepts of the reinforcement and matrix phases in complicated systems, we define the phase that is the base material as the ‘matrix’ phase, while the other phase that is added to the base material for performance improvement is defined as the ‘reinforcement’ phase, which is probably used to improve either the strength or the toughness/ductility, or even the physical or chemical performances of the composites. Therefore, the reinforcement phase is possibly soft or hard, continuous or discontinuous, or high fraction or low fraction without considering the hardness, connectivity and volume fraction. Based on this viewpoint, both the individual ceramic particle phase on the micro-scale and the particle-rich composite region on the meso-scale can be viewed as ‘reinforcement’ in the composite. In addition, the ceramic particle-rich region can be viewed as a two-level ‘reinforcement’ phase, in some cases, more than two length-scale structures can be identified according to the above definition. As demonstrated by Lakes [21], many natural and artificial materials exhibit multi-scale structures, in which each structure contains a sub-structure on a lower scale which plays an individual role contributing to the overall performance. The structural hierarchy of the multi-scale structures largely determines the overall properties of the composites. Therefore, in order to reach a comprehensive understanding, the structure influence, the fracture mechanism, toughening mechanism and strengthening mechanism of the inhomogeneous microstructures must be treated hierarchically, where ‘reinforcement’ and ‘matrix’ phases are themselves a ‘composite’. Such a broad understanding will lead to novel processing techniques for the fabrication of a new class of super-composites with superior combination of properties. 2.2. Fundamentals behind the design of inhomogeneity First, it must be mentioned that, in the present work, the words ‘homogeneous’ and ‘inhomogeneous’ are used specifically in terms of the spatial distribution of the reinforcements (particles, whiskers or short fibres), i.e., their spatial position within a composite system. Other factors such as the size distribution of reinforcements can also lead to certain microstructural inhomogeneity, which will not be covered in the present paper. Inhomogeneities of spatial distribution need to be described at different levels of length scale. According to their features, the inhomogeneities can be classified as microscopic inhomogeneity and macroscopic inhomogeneity [22]. Strictly speaking, no engineering material is homogeneous at the microscopic scale, as has been revealed by many modern techniques, such as transmission electron microscopy (TEM); there are always various inhomogeneities within the micro-region (nm-um). This paper focuses on dealing with two-phase DMMCs; their microstructures can be classified into the following three types/levels: (1) Microstructurally homogeneous: Composite materials of this kind, such as conventional particle reinforced MMCs, exhibit isotropic behaviour and have a homogeneous discrete reinforcement distribution in all three dimensions. In these composites, particle contiguity is zero, while matrix contiguity is close to the maximum. (Note: according to Nan [22], these MMCs still contain microstructural inhomogeneity because two different phases are involved.)

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(2) Microstructurally inhomogeneous but macroscopically homogeneous: Composite materials of this kind also exhibit isotropic behaviour. However, in contrast to a homogeneous discrete distribution, the reinforcement distribution may be rich in some regions and lean in others in all three dimensions, for example in a dual matrix microstructure, network microstructure, bi-continuous microstructure and clustering microstructure [23,24]. In these composites, the particle contiguity is enhanced due to clustering/agglomeration. (3) Microstructurally inhomogeneous and macroscopically inhomogeneous: This kind of composite exhibits anisotropic behaviour, the reinforcement distribution may be rich in some regions and lean in the others only in a given section of the material, such as in graded or layered MMCs. It is worth noting that two sources of microstructural inhomogeneity exist. One is as an inevitable product of traditional DMMC manufacturing processes where a perfectly homogeneous distribution is sought but not attainable, and hence undesirable. Therefore, very little effort has been given to controlling this inhomogeneity except for trying to eliminate it. The other is as a result of artificially deliberate manipulation, such as the bi-continuous microstructure and the layered microstructure, which can be effectively controlled by adjusting the processing parameters, spatial distribution, and so forth. Beside a vast list of literature on numerical simulations of various forms of inhomogeneities such as finite element analyses of the effect of particle clustering (briefly described latter), one fundamental justification for the deliberate tailoring of an inhomogeneous distribution lies in the so-called phase contiguity model. In order to describe the distribution and strengthening efficiency of the reinforcement in two-phase composites, phase ‘contiguity’ was first proposed by Gurland [25,26]. The contiguity of each phase in any two-phase composites can be defined using the following parameters.

Ca ¼ Cb ¼

2Saa V ab 2Saa V þ SV

2Sbb V 2Sbb V

þ SaV b

ð1Þ ð2Þ

Ca and Cb are contiguities of a and b phases, respectively. Saa V is the surface area among a grains per ab unit volume, Sbb V is the surface area among b grains per unit volume, and SV is the surface area between a and b grains per unit volume. Therefore, in principle, both Ca and Cb are higher than zero but lower than one in any two-phase composites, i.e., 0 < Ca < 1 and 0 < Cb < 1. According to the topological transformation [27,28], a two-phase microstructure (a + b) with any grain size, grain shape and phase distribution, can be topologically transformed into a body with three basic microstructural elements (readers are referred to the relevant literature for details). Consequently, a topological parameter, continuous volume fraction, of each phase is defined by the phase contiguity and its overall volume fraction (Vf), i.e., the continuous volume fraction of a-phase, fac = CaVfa. The main conclusion of this topological exercise is that the composite property is more directly dependent on the continuous volume fraction of the reinforcing phase than its overall volume fraction and a tailored non-uniform (non-discreet) reinforcement distribution may lead to superior reinforcing efficiency.

3. Classification of inhomogeneous phases In this context of reinforcement 3D distributions, the microstructurally inhomogeneous composites effectively consist of two regions: a reinforcement-lean region (corresponding to a ‘soft’ composite phase) and a reinforcement-rich region (corresponding a ‘hard/strong’ composite phase1). According to the 3D phase connectivity or contiguity, the clustered reinforcement-rich regions would have the following four different patterns as schematically illustrated in Fig. 1 [29]. Pattern A: Isolated reinforcement-rich phase (0D), the clustered regions are separated from each other in the material, i.e., reinforcements are agglomerated in the form of isolated clusters (Fig. 1a). 1

Here the term ‘phase’ is generalized in that a phase can be a composite itself.

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Fig. 1. Schematic illustrations and representative SEM images (insets) of microstructural inhomogeneity with four different patterns of reinforcement-rich phase. (a) Pattern A: isolated, (b) Pattern B: bar/laminated/ring-like, (c) Pattern C: 3D network with isolated particle-lean phase and (d) Pattern D: 3D network with interconnected particle-lean phase forming the so-called bi-continuous microstructure. (Pattern A is similar to that used by Corbin & Wilkinson [30] and Tseng [31]. The four insets are the corresponding real microstructures of composites selected from published work that will be introduced in the following sections.)

Pattern B: Bar/laminated/ring-like reinforcement-rich phase (1D or 2D), the clustered regions are continuous in one or two dimensions (Fig. 1b). Pattern C: 3D continuous reinforcement-rich phase with isolated reinforcement-lean phase, only the reinforcement-rich phase is continuous to form a grain boundary-like network (Fig. 1c). Pattern D: Bi-continuous reinforcement-rich phase and reinforcement-lean phase. In this case, both phases are 3D continuous forming two interpenetrating 3D networks (Fig. 1d). In general, Patterns A, C and D correspond to microstructurally inhomogeneous but macroscopically homogeneous microstructures, while Pattern B corresponds to microstructurally inhomogeneous and macroscopically inhomogeneous microstructures. It is worth noting that the isolated, bar/laminated/ ring-like, network and interpenetrating characteristics are corresponding to the reinforcement-rich phase or reinforcement-lean phase, not the hard ‘ceramic’ reinforcement in itself. In the ‘soft’ reinforcement-lean phase, the local volume fraction of reinforcement (VL) can range from 0 to a particular value, such as VL1, while the VL in the ‘hard/strong’ reinforcement-rich phase (VL2) must be higher than VL1 and even up to 100%. Obviously, when VL1 = 0 and VL2 = 100%, the microstructure returns to the conventional two-phase composite microstructure. The above four patterns could be described in Newnham’s taxonomy which is based on 3D phase contiguity [15], such as Pattern A and Pattern C correspond to 0-3 type with the isolated reinforcement-rich phase and reinforcement-lean phase, respectively. Pattern B probably corresponds to the 1-3 type with a bar structure or 2-2 type with a layered structure, while

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Pattern D certainly corresponds to the 3-3 type with a bi-continuous structure, sometimes referred to as ‘interpenetrating’ structure or ‘co-continuous’ structure.2 3.1. Isolated reinforcement-rich phase Reinforcement clustering in the isolated Pattern A (Fig. 1a) observed in conventional DMMCs, such as in stir-casting MMCs [23,24], is normally believed to be detrimental to mechanical properties in relative to a homogeneous reinforcement distribution, therefore, researchers had always tried to eliminate this kind of inhomogeneity. There were no attempts to purposely manipulate such clustering from the view point of microstructural design. Deliberately controlling and producing particle clustering in DMMCs in order to investigate its effect on the mechanical properties of composites is only a recent (late 1990s) practice. Conlon and Wilkinson [32] studied the effect of an inhomogeneous distribution of second phase particles on strength and damage in Al–CuAl2 composites. When tested in compression (little or no damage), the clustered materials were significantly stronger than the homogeneous material containing the same overall CuAl2 volume fraction. They concluded that clustering can only be considered as a viable strengthening mechanism in situations for which damage can be suppressed. Moreover, when damage does occur, clustering concentrates the damage in small regions which is likely to enhance crack formation resistance and thus enhance the ductility of the material. It is necessary to mention that damage in the form of particle fracture prior to final fracture will probably results in a significant degradation of properties, such as modulus [33]. Investigations on the influence of particle clustering on the onset of damage [34] indicated that particles in the particle-rich phases are more susceptible to cracking which is a main mode of damage for the naturally aged composite materials. However, based on fracture mechanical simulation, in order to optimize the fracture toughness of DMMCs, Toda et al. [35] appear to have carried out the first practical work looking positively at the effect of reinforcement clustering and proposed a microstructurally controlled MMC with artificially agglomerated SiC whiskers. The developed materials showed much higher strength and superior crack-propagation resistance compared with the conventional composites with a homogeneous microstructure. Subsequently, Deng and Fang [36,37] designed and fabricated a dual WC–Co composite composed of the isolated WC-rich phase and continuous WC-lean phase. According to the test results, the toughness, hardness and wear resistance of the composites with the WC-rich phase were improved compared with those of the composites with homogeneous microstructures. In order to improve the toughness and ductility of TMCs with high reinforcement volume fractions, Patel and Morsi [38,39] also designed and fabricated a so-called ‘dual matrix’ TiB–Ti composite which includes isolated reinforcement-rich phase and continuous matrix phase as schematically illustrated in Fig. 2. Obviously, this microstructure belongs to the isolated reinforcement-rich case (Pattern A). 3.2. Bar and laminated/ring-like reinforcement-rich phase Reinforcement clustering occurring within a given section of material, specifically in terms of graded [40,41] and layered structures [8,42–44], is of the Pattern B type (Fig. 1b) and represents the macroscopically inhomogeneous microstructure with reinforcement clustering. The layered structure is intended to increase toughness compared to a homogeneous structure. The laminate approach has shown potential as a toughening mechanism for particulate-reinforced aluminium MMCs [42–49]. Lewandowski et al. [42,50–56] carried out a systematic investigation into laminated Al/Al–SiCp material, including the effect of lamination (layer thickness) on the impact toughness. Pandey et al. [44] also demonstrated improved toughness in the laminated particulate-reinforced aluminium composites. In addition, Pandey et al. [47] produced a ‘bean’ material consisting of large size aluminium beans in SiCp/7093 composites. This bean-like microstructure with lamellar reinforcement-lean phases can 2 In the context of phase contiguity, we prefer to use the term ‘bi-continuous’ to represent the 3-3 microstructure as the term ‘interpenetrating’ or ‘co-continuous’ may limit to cases where both phases are highly or equally connective.

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Fig. 2. Conceptual schematic diagram showing dual matrix Ti–TiBw composites [38] that belongs to Pattern A.

be approximately treated as Pattern B. The bean materials exhibited improved ‘crack growth’ toughness, and in some cases the initiation toughness was also increased. Layered MMC structures consisting of alternative lamella of reinforced and unreinforced alloy were generated using the spray-forming method [43]. Impact measurements at high strain rates have shown that the layered structure gives limited improvement in toughness; however, low strain rate testing, using Tattersall–Tappin tests, has shown that the toughness (characterized by the area below the load–displacement curves) of the layered structure is roughly twice that of comparable homogeneous MMCs, with the crack progressing in a step-like manner as it encounters successive layers [43]. Recently, Ti–TiBw/Ti composites with laminated microstructures were successfully designed and fabricated in order to further improve the ductility and toughness of TMCs [57]. In these composites, the TiBw/Ti composites with a network microstructure were employed as the reinforcing layer; this is beneficial to the mechanical properties due to the superior combination of strength and ductility [58]. It is worth noting that the prepared laminated Ti–TiBw/Ti composites with the particular laminated structure exhibited not only higher strength but also higher ductility than the pure Ti material. In addition, the laminated composites undoubtedly showed superior ductility than the monolithic network-structured TiBw/Ti composites. The increased toughness and ductility of laminated composites are believed to stem from one or more of the following factors [42–46,50,59]: (1) crack bridging, where reinforcement-lean ligaments behind the crack tip restrict crack opening in the clustered layer and force re-nucleation of the crack ahead of the crack tip; (2) crack tip plasticity, where there is more extensive plastic deformation ahead of the crack tip in the reinforcement-lean layer absorbing crack energy; (3) interfacial delamination between the clustered layer and the reinforcement-lean layer as a profound energy absorption mechanism in fibre reinforced composites [60] and conventional multi-layered materials [61]. Therefore, it is believed that there must be adequate plastic regions among the reinforcements to blunt or deflect cracks if one wants to toughen DMMCs. Likely inspired by the laminate approach, Nardone et al. [45,46,62] designed and developed the socalled microstructurally toughened particulate-reinforced Al MMCs, consisting of bar-like (not layer) SiCp/6061 rods embedded in a 6061Al matrix. Compared to conventional homogeneous SiCp/6061 composites, an increase in the notched Charpy impact energy absorption capability by as much as one order of magnitude, in a plane perpendicular to the rod-direction, was demonstrated. A similar approach was adopted by Qin and Zhang [48,49] to produce MMCs with bar clustering reinforcement showing increased toughness in the specific plane. Inspired by the concentric tree trunk growth-ring structure, Wong et al. [63] designed Mg–Al2O3/Mg composites with a concentric alternating macroring microstructure, which is composed of alternate monolithic ring Mg phases and of 1.11 vol.% Al2O3/Mg composite ring phases. It is surprising that the composites with a 3 mm ring layer thickness

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exhibited a superior combination of the tensile yield strength (YS), ultimate tensile strength (UTS) and ductility compared with the monolithic Mg and the Al2O3/Mg composites. 3.3. 3D network reinforcement-rich phase with isolated reinforcement-lean phase As early as the 1960s, using theoretical and experimental results, Hansen had confirmed that a network microstructure of hard Al2O3 particles can effectively strengthen the aluminium matrix [64]. Murphy et al. [65] generated materials with various reinforcement distributions using a number of casting processes. Particle clustering was caused by the pushing effect under different cooling rates. The ductility of the MMCs was found to decrease with the increase in the severity of the particle clustering. Using a similar technique, Kumai et al. [24] produced SiC-particle-reinforced A356 alloys with a systematically controlled dendrite cell size and particle distribution to investigate their effects on the threshold of fatigue-crack growth. Near threshold fatigue-crack growth properties were improved in the composites with a coarser cell size and an inhomogeneous particle distribution due to the enhanced roughness-induced crack closure effect [66]. In these experimental studies on cast composites, the particle distribution was deliberately controlled to a certain extent by means of different cooling rates. Due to the pushing effect by the solidification front, the regions with a higher local volume fraction form a continuous network. This form of clustering is often seen in cast composites which have undergone relatively slow cooling and is sometimes referred to as a ‘necklace’ structure [65]. It should be noted that the particle-lean phases in these composites are mostly isolated. Recently, the PM technique has also been used to fabricate composites with a network microstructure [67]. In particular, TMCs and Al MMCs with a quasi-continuous network microstructure, fabricated by PM combined with the in situ technique, exhibit a superior combination of tensile strength and ductility [58,68,69]. The superior strengthening effect is mainly attributed to the increased reinforcement connectivity in the network boundary and the strong interfacial bonding between the reinforcement and the matrix, while the superior toughening effect can be mainly attributed to the large reinforcement-lean phase which can blunt and deflect cracks, slow crack propagation and bear the strain. In addition, a continuous network microstructure was also formed in the TiBw–Ti composite system due to the excessively high local volume fraction of reinforcement in the network boundary [70]. Although the continuous network microstructure possesses very high reinforcement contiguity, it is destined to achieve inferior toughness and ductility due to the brittle continuous network boundary without an interpenetrating matrix. However, it is worth pointing out that a continuous monolithic TiC network microstructure in the TiCp/Ti composites was deliberately designed for other purposes [14,71]. In these composites, the in situ TiCp ceramic particles self-assembled to form a continuous monolithic TiC wall giving rise to a honeycomb structure. These composites exhibited not only a superior compressive strength, but also a very significant improvement in high-temperature oxidation resistance. All of the above network or ‘necklace’ microstructures are of Pattern C. It is concluded that a superior combination of strengthening effect and toughening effect needs not only high reinforcement connectivity, but also high matrix connectivity. That is to say, a balance between reinforcement connectivity and matrix connectivity is necessary for a superior combination of strength and ductility, but this balance also depends on the property contrast. 3.4. Bi-continuous reinforcement-rich phase and reinforcement-lean phase Experimental results have indicated that the composites yielded a higher energy absorbing capability due to the increased crack-propagation resistance. The composites also showed slightly improved tensile strength and elastic modulus, despite the observed low overall tensile elongation. Due to the compression applied to the fibre agglomerations during the pre-form preparation and subsequent squeeze-casting process, the fibre clusters are interconnected throughout the material [72] and therefore the microstructure is of Pattern D but with rather limited connectivity between the composite spheres.

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Wegner and Gibson [73–75] also successfully fabricated bi-continuous composites by infiltrating bronze and polymer resin into porous 420 stainless steel and 316 stainless steel foams, respectively. In the two different bronze-420 stainless steel and polymer-316 stainless steel composite systems, both the softer phase and the harder steel phase are interpenetrated and then form two separate continuous networks (i.e., bi-continuous), which result in an improvement in toughness under a special volume fraction of the ductile phase. Moreover, the initiation toughness increased with the increasing volume fraction of the softer phase. Recently, Dong et al. [76] also fabricated one Cu–Cr3C2 composite by infiltrating molten copper into the pre-formed Cr3C2 foam. Equally, the continuous Cu phase and Cr3C2 phase are continuous and then form two separate continuous networks, which result in a higher elastic property of the interpenetrating composites. Travitzky et al. [77] fabricated one dense Al2O3/TiAl composite using PM accompanied by in situ synthesis between TiO2 powders and Al powders. In the dense Al2O3/TiAl composites, not only the Al2O3 phase but also the TiAl phase was interpenetrated forming two separate Al2O3 and TiAl continuous networks (Fig. 3), which are believed to be beneficial in increasing the toughening effect of the composites. In addition, the bi-continuous microstructures involving the reinforcement-rich continuous network boundary and the interpenetrating matrix are of Pattern D, which are designed and prepared in order to increase the strengthening effect by increasing the reinforcement connectivity and retaining the superior ductility and toughness of the interpenetrating matrix. To date, bi-continuous microstructures have been successfully employed in many systems [13,78,79]. The traditional fabrication route includes pre-form preparation and the subsequent squeeze-casting process for the bi-continuous microstructure. In this route, the challenge is to fabricate a strong pre-form that is capable of withstanding the force during squeeze casting. In summary, unmistakable experimental results exist indicating that the spatial reinforcement distribution exerts significant influence on the properties of the composites. Improvements in both strength and ductility of the composites could be achieved by controlling the microstructural inhomogeneity and microstructure parameters of such inhomogeneous microstructures. 3.5. Theoretical justification It is well known that the classical rule of mixtures (RoM) model is used to estimate the elastic properties of continuous fibre reinforced composites with optimal interfacial bonding. Herein, the RoM upper bound (Voigt model) and lower bound (Reuss model) can be obtained according to the equal strain assumption and the equal stress assumption, respectively, without considering other practical factors [60].

EL ¼ V a Ea þ ð1  V a ÞEb  1 V a ð1  V a Þ þ ET ¼ Eb Ea

ð3Þ ð4Þ

Fig. 3. Secondary electron images of a polished cross section of Al2O3–TiAl composites (a), and of the Al2O3 network after the leaching of TiAl (b) [77].

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where EL and ET are the elastic modulus of the composites along the longitudinal direction and transversal direction, Ea and Eb are the elastic modulus of the reinforcement and matrix, respectively, and Va is the overall volume fraction of the reinforcement. According to the theoretical model, the elastic properties of all the composites should be located between the RoM upper and lower bounds as shown in Fig. 4 [80]. Therefore, due to the simple equations and the broad bounds, the RoM bounds are usually used to roughly predict the highest and lowest properties and monitor the experimental data. It is absolutely wrong that the experimental data of the elastic properties, including the elastic modulus and electric conductivity of composites, is located outside the RoM bounds. In addition, it is worth pointing out that the RoM is only suited to predicting elastic properties, such as modulus and electric conductivity, and is not appropriate for predicting the plastic properties of the composites, such as tensile strength and elongation. It is very easy to find superior tensile strength higher than the RoM upper bound [81,82]. In order to effectively estimate the elastic properties of an isotropic composite, such as a DMMC, Hashin and Shtrikman proposed the well-known H–S bounds as early as 1963 [80,83,84]. In the H– S theorem, the upper bound rigorously corresponds to the composites containing the ‘soft’ inclusion matrix phase encapsulated by a ‘stiffer’ reinforcement phase, while the lower bound corresponds to the composites with a ‘stiffer’ inclusion reinforcement phase encapsulated by a ‘softer’ matrix phase, such as the ceramic particle reinforced MMCs with a homogeneous microstructure. For an isotropic two-phase composite, the H–S bounds for the elastic modulus (E) can be expressed as [83]:

Ea ðEa V a þ Eb ð2  V a ÞÞ Eb V a þ Ea ð2  V a Þ Eb ðEb ð1  V a Þ þ Ea ð1 þ V a ÞÞ ¼ Ea ð1  V a Þ þ Eb ð1 þ V a Þ

EHS-Upper ¼

ð5Þ

EHS-Lower

ð6Þ

where EHS-Upper and EHS-Lower are the values of the upper and lower bounds, respectively. The EHS-Upper and EHS-Lower bounds are tighter than the RoM bounds as schematically illustrated in Fig. 5. The H–S bounds have been regarded as the best possible bounds on properties for isotropic two-phase composites, i.e., the elastic properties of all the isotropic composites without considering reinforcement distribution should be strictly located between the EHS-Upper and EHS-Lower bounds [85,86]. As illustrated in Fig. 5, the EHS-Upper bound corresponds to the microstructure in which the hard phase encapsulates the soft phase. Therefore, the hard phase is continuous while the soft phase is discrete. In contrast, the EHS-Lower bound corresponds to the microstructure in which the soft phase encapsulates the hard phase, resulting in a continuous soft phase and a discrete hard phase. That is to say, the conventional

Fig. 4. Rule of mixture bounds, black area and white area represent the reinforcement and matrix phase, respectively.

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Fig. 5. Hashin–Shtrikman bounds, black area and white area represent the reinforcement phase and matrix phase, respectively.

composites with a homogeneous reinforcement distribution can only exploit limited elastic properties close to the EHS-Lower bound. With the motivation of approximating the properties of a real composite, much effort is devoted to making the bounds even tighter and more restrictive, as in Beran’s bounds [87], Miller’s bounds [87], Torquato et al.’s bounds [88,89], Milton’s bounds [90,91], Milton and Phan-Thien’s bounds [92] and McCoy’s bounds [93], by incorporating more detailed information about the microstructure (which means more parameters in the bounding techniques/theories). The greater the number of parameters considered, the more rigorous are the bounds obtained, but the more complicated the bounds, and, therefore, it is inconvenient to calculate and apply the bounds. This also restricts the bounds to a smaller group of composites. In addition, nearly all of these composite theories/models are based on continuum mechanics approaches which are only able to predict properties as a function of the overall volume fraction of phases [10,94,95]. In contrast to the diverse patterns of reinforcement clustering investigated experimentally, nearly all the reinforcement clusterings addressed in theoretical simulations have the form of an isolated clustering region defined as Pattern A. Nonetheless, analytical analyses in this area are very limited, mainly due to the complex nature of the problem, and, therefore, alternative techniques, such as the finite element model (FEM) and the self-consistent method have been widely employed. It will be found below that, despite numerous theoretical studies, the understanding of the effect of reinforcement clustering on composite behaviour remains controversial [96–98]. Bao et al. [99] appear to have carried out the first FEM study to examine the effects of phase clustering in microstructures containing isotropic perfectly elastic and plastic phases. They found that the lowest flow stress, at fixed volume fractions of the harder phase, was associated with microstructures containing spatially homogeneous distributions of the phases. Redistributing the phases into ‘particlerich’ and ‘particle-lean’ regions with constant overall volume fractions had the effect of raising the overall flow stress. The magnitude of this effect depended on the overall volume fraction as well as the magnitude of the difference between the local volume fraction of ‘particle-rich’ regions and that of ‘particle-lean’ regions, i.e., the degree or severity of the clustering in the composites. This coincides well with Toda’s FEM results [98] and other FEM results [97,100,101] where the yield stress, strain hardening and elastic modulus were all found to increase as the severity of clustering increased. Similar results have been found in a series of papers by Wilkinson et al. [10,102–105] where a self-consistent scheme was implemented for a two-phase solid with each phase exhibiting power-law hardening. Wilkinson et al. also pointed out that an optimum degree of clustering exists at which the maximum strengthening effect is reached below the optimum value, the strengthening effect increasing with the increasing degree of clustering.

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However, Segurado et al. [100], using 3D FEM analysis, showed that the effect of particle clustering on the mechanical properties of the composite was weak. It was found that the presence of clustering greatly increased the fraction of broken particles leading to a major reduction in the composite flow stress and ductility. The discrepancy between theoretical treatments is believed to be due to the complex nature of the problem and the lack of valid model-system experimental data. Another important reason is the lack of composite theory that can properly account for the reinforcement distribution and guide the microstructural design. As a consequence, there is a lack of appreciation that a controlled inhomogeneous distribution could yield a better property combination; hence, very little effort was dedicated to controlling the inhomogeneity prior to the 1990s. Recently, in order to achieve superior DMMC performance, more and more effort has been dedicated to tailoring a better microstructure of the composites in which the reinforcement is inhomogeneously distributed in a controlled way. 4. Composites with tailored inhomogeneous microstructures In the 1990s, Hunt et al. [47,106] proposed two routes for improving DMMC toughness. The first approach, termed the intrinsic approach, controls the particle size, volume fraction, interparticle spacing, reinforcement distribution, matrix precipitates, grain size and matrix/reinforcement interface properties. Much of the open literature demonstrates that this approach has always been associated with strength loss. The other approach, termed the extrinsic approach, is an architectural/microstructural design which has been used extensively in the last two decades to improve fracture toughness [50,107,108]. Therefore, much architecture has been designed and fabricated for Al, Mg, Cu, Fe, Co and TiAl composites, as well as tailored microstructures for TMCs. The latter will be described separately in Section 5. 4.1. Tailored microstructure with isolated reinforcement-rich phase Discontinuous aluminium MMCs have attracted more and more attention due to their superior physical and mechanical properties [109,110]. Since the 1990s, limited by the performance of a homogeneous microstructure, researchers have been sedulously pursuing a superior microstructure by tailoring the reinforcement distribution in order to improve ductility and damage tolerance [3,8,10,100,111]. Guided by the phase contiguity model, Peng et al. [72,112] designed Al2O3 short-fibre agglomeration-reinforced 6061Al MMCs, in which Al2O3 short fibre agglomerations with diameters from 0.4 mm to 1 mm are uniformly distributed in the 6061Al matrix at a macroscopic level, as shown in Fig. 6a. The local volume fraction of the fibre reinforcement decreased from the outer layers (30%) to the centre region (15%) within the composite sphere. Therefore, the composite spheres possess a hard outer layer and gradient interface, which are intact, unbroken and uniformly spherical, and the fibres are randomly distributed within the fibre agglomerates. These features are beneficial to the damage tolerance of the composites. This work also overcomes the planar-random distribution of fibre reinforcement in the traditional squeeze-cast process [60]. As seen from Fig. 7, the damage tolerance of this composite with the fibre agglomeration microstructure is about 50% higher than that of the conventional MMCs with a homogeneous microstructure [112]. Interestingly, after the peak load additional energy absorption, which is positive to avoid a catastrophic fracture, is also observed. According to the crack-propagation route (Fig. 8), the improved damage tolerance can be attributed to the fibre agglomerations with a hard outer layer. For such inhomogeneous composites, fracture is a process that occurs at many length scales, from reinforcement cracking, interface failure, local voiding to large scale propagation; crack blunting can also occur by more than one mechanism. From microscale point of view, micro-cracking generally occurred through interface failure for the ex-situ composites due to weak interface bonding and strong reinforcement particles. In addition, the existence of a reinforcement-lean region at the centre of ‘composite sphere’ and the large size monolithic matrix region can compel micro-crack to coalesce along the tortuous route following the ‘strong’ composite spheres. From macro-scale point of view, the crack tends to propagate along the ‘macro-interface’ between the fibre-free Al matrix and fibre-rich composite spheres which provides significant crack

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Fig. 6. The microstructure of the composites prepared from fibre agglomerates at a lower magnification (a) and a higher magnification (b) showing the variation in the fibre volume fraction [112].

Fig. 7. Typical load–displacement curve of the conventional MMC and of the tailored MMC during three-point loading tests. (Sample N and Sample P representing a ‘V’-shaped notch parallel and vertical to the pressing direction during the squeeze-cast fabrication process, respectively. The energy absorption of I, II and III is 0.9 J, 1.19 J and 1.53 J, respectively.) [112].

Fig. 8. The crack-propagation route of the tailored MMC during fracture [112].

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deflection and absorbs much energy. Moreover, when the crack deflection is far away from the main crack direction, the crack has to pass through the composite sphere. This will require a higher stress level and lead to a higher damage tolerance by bearing strain and decreasing crack propagation rate in the centre reinforcement-lean region. On the other hand, a higher crack-propagation resistance and hence a higher damage tolerance is attained due to the ductile centre region with a lower local reinforcement volume fraction (15%). The rugged and tortuous propagation path, comprising a protruding composite ball and concave socket, effectively enhanced the energy absorption or the damage tolerance of the MMC. The tensile test results show that the UTS of the designed composite (199 MPa) is higher than that of the conventional composites (188 MPa), while the tensile fracture elongation is reduced from 7.1% to 4.9% [72]. It is worth pointing out that the elastic modulus (84 GPa) of the designed composites is also higher than that of the conventional composites (76 GPa) despite the same overall fibre volume fraction. The improvement in elastic modulus of the composites is very important and can be attributed to the increasing reinforcement contiguity due to the fibre agglomeration distribution. These improvements in UTS and elastic modulus are consistent with the work of Al MMCs with particle agglomerates [113]. In addition, Balakrishnan et al. [114] also reported that the strength of the as-sintered ceramic agglomerations increases with increasing particle contiguity, i.e. particle contact and adhesion, which can be achieved by decreasing particle sizes. Wegner and Gibson [74] also pointed out that the contact area between particles influences the properties of the composites with isolated reinforcement-rich phase. These positive conclusions further confirm that the strengthening effect of the reinforcement can be further increased by increasing its contiguity at a given total volume fraction. In order to predict and understand the short-fibre agglomeration-reinforced 6061Al MMCs, the theoretical contiguity model based on topological transformation and phase contiguity has been adopted [72]. In this phase contiguity model, not the overall volume fraction but the ‘continuous volume fraction’ of reinforcement, depending on its spatial distribution, dominates the composite properties. According to the contiguity model, the Young’s modulus and yield strength (YS) of the DMMCs increase with the increasing ‘continuous volume fraction’ of the reinforcement phase. With the aim to improve the toughness and processability of high volume fraction composites (>50%), Saha et al. [115] designed and fabricated a hierarchical SiCp-agglomeration reinforced 6061Al composite by squeeze-infiltration casting, as shown in Fig. 9. SiCp/Al composite agglomerations, with sizes of 100–500 lm, were homogeneously distributed in a 6061Al matrix as reinforcing units at a higher level, while fine SiC particles, with sizes of less than 1 lm, were homogeneously distributed in the agglomerations as reinforcement at a lower level, with a much higher local volume fraction than the overall volume fraction. The bending test results showed that not only the strength but also the damage tolerance were increased. In particular, the hierarchical composites exhibited almost twice the fracture toughness than the conventional composites due to the ductile fracture manner of the continuous matrix network.

Fig. 9. Microstructure of 50 vol.%SiCp/6061 composite with hybrid reinforcement architecture: (a) homogeneous distribution of composite agglomerate particles and (b) very good bonding between the agglomerate and matrix, although some porosity is observed within the agglomerate [115].

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Similarly, Zhou et al. [116] fabricated Al2O3/6061Al composites composed of a continuous reinforcement-free Al network phase and a discrete reinforcement-rich phase. The composites with such a controlled microstructure exhibit significant enhancements in UTS, elastic modulus, and also a slight increase in ductility compared with the conventional composites with a homogeneous microstructure at a given volume fraction. These improvements can be attributed to the reinforcement-rich phase, while the increased ductility is likely to be due to the continuous Al matrix network boundary. Recently, Roy et al. [117] also fabricated 3D Al2O3/AlSi12 composites composed of Al2O3 agglomerations, 3D continuous Al matrix with Si particles randomly distributed within them. The results showed that the internal load under tension and compression can be transferred from the softer and continuous Al matrix to the stiffer and stronger Si and Al2O3 phases. Moreover, the maximum longitudinal stress in the Al2O3 phase is almost double the external compressive stress, and it is even more than double the tensile stress. This result is likely to be due to the softer matrix being continuous while the harder reinforcement is discrete. In order to boost the fracture toughness while maintaining the superior wear resistance of WC/Co composites, Fang and Deng [36,37,118] successfully designed and fabricated double cemented (DC) carbide composites of the WC–Co system by PM, i.e., WC/Co composite granules are embedded in the continuous Co matrix (Fig. 10). These two-scale composites are described as ‘a composite within a composite’, which belongs to the aforementioned multi-scale structure (each structure contains a sub-structure at a lower scale) [21]. Moreover, the Co content of the WC/Co composite granules embedded in the DC composites can be controlled by preparing bulk WC/Co composites with different WC fractions, which are essential to the overall properties of the DC composites. In addition, the toughness of the DC composites largely depended on the width of the mean free path (MFP), which can be controlled by the granule size and the overall volume fraction of the intergranular matrix [36]. The test results showed that both the hardness and the wear resistance of the DC WC/Co composites generally increased with decreasing granule Co content and overall intergranular Co volume fraction. The improvement in hardness and wear resistance become more obvious when the Co content or overall intergranular Co volume fraction is very low. It is interesting that the fracture toughness of the DC composites remarkably increases with increase in the MFP width (Fig. 11a), which can be achieved by increasing the granule size or the overall intergranular Co volume fraction. This phenomenon is consistent with the aforementioned statement: the adequate reinforcement-lean phase is necessary to blunt or deflect the crack to improve the fracture toughness of the tailored composites. Certainly, the fracture toughness increases with the increasing Co granules in the WC/Co composite granules, which is equivalent to increasing the toughness or decreasing the brittleness of the composite granule by increasing the matrix width among individual WC particles. However, this improvement is much lower than that stemming from the increase in MFP width, due to the very high local volume fraction over 75% in the granule. Besides the toughness improvement, the wear resistance also increased with increasing granule sizes (Fig. 11b) due to greater protrusion of larger hard granules from the wear surface. In fact, it is worth pointing out that the presintering treatment is also essential to strengthen the granules by increasing WC contiguity, which is beneficial to the overall hardness, toughness and wear resistance.

Fig. 10. Microstructure of (a) DC carbide, and (b) conventional cemented carbide comprising the DC carbide granules [36].

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Fig. 11. Effect of granule size and MFP of metal matrix on (a) toughness, and (b) wear resistance of DC carbide. Fracture toughness is probably aided most strongly by increased matrix MFP while wear resistance is increased mostly by granule protrusion, due to size [36].

Finally, the most valuable result is that this DC WC/Co composite exhibits a superior combination of fracture toughness and wear resistance (a greater toughness at equivalent wear resistance, or a higher wear resistance at equivalent toughness) than the conventional WC/Co composites with a homogeneous microstructure. In addition, increasing granule size, which is viewed as increasing the overall inhomogeneity of the composites, can simultaneously improve wear resistance and fracture toughness. All the improvements can be mainly attributed to the simultaneous existence of hard and large WC/Co composite granules, and soft and wide WC-lean intergranular paths, which can inspire ductile tearing and dimple fracture, crack-bridging and crack-deflecting effects. However, the DC WC/Co composites exhibit a lower elastic modulus, which can be attributed to the overall high Co content and DC microstructure, corresponding to the H–S lower bound (i.e., soft phase encapsulates hard phase) [118]. 4.2. Bar and laminated/ring-like microstructures 4.2.1. Bar reinforcement-rich phase With the aim to overcome low fracture toughness, low fracture energy and low ductility of conventional MMCs [109,119–122], Qin and Zhang [48,49] successfully designed and fabricated a (SiCp– 6061Al)/6061Al composite, comprising 45%SiCp/6061Al composite bars as the reinforcement phase and large SiCp-lean phases as the matrix, by vacuum pressure infiltration followed by hot extrusion, as shown in Fig. 12. The SiCp–6061Al composite bars with a diameter of 560 lm can effectively bear stress, while the matrix region between neighbouring bars, with a much larger distance of about 320 lm, can effectively bear strain and blunt the crack, even bearing a partial load after the bar failure [48]. As shown in Table 1, the fracture toughness of the designed (SiCp–6061Al)/6061Al composites with a bar reinforcement clustering microstructure is increased by 35.1% compared with that of the conventional composites with a homogeneous microstructure and is even close to that of a monolithic 6000 series aluminium alloy. Moreover, the YS and UTS of the designed composites are increased by 9% and 7%, respectively, but with a similar elastic modulus when compared with the conventional composites. These results show that the tailored bar reinforcement clustering microstructure is superior to the laminated microstructure which just exhibits an improved toughening effect but a decreased strengthening effect [42,52,56]. It is worth pointing out that the composites exhibit different multi-stage fracture characteristics, which can protect the composite from catastrophic failure, while the composites with a homogeneous microstructure exhibit an abrupt fracture process, indicating lower fracture energy and catastrophic failure, as shown in Fig. 13.

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Fig. 12. Optical microstructures of the composites. (a) Designed composite perpendicular to the extruded direction; (b) designed composite parallel to the extruded direction; (c) conventional composite parallel to the extruded direction [48].

Table 1 Fracture toughness KQ and tensile properties of the two composites and unreinforced Al alloy [48]. Materials Designed SiCp–6061Al/6061Al composites Conventional SiCp/6061Al composites Unreinforced Al alloy [123]

Total SiCp Vf (%) 15 15 0

KQ (MPa/m1/2) 28.3, 28.9, 27.8 21.8, 18.4, 23.2 27–35

E (GPa)

r0.2

UTS (MPa)

e (%)

(MPa)

97.3

365

413.8

1.33

96.3

335.1

388

3.29

70

257

315

17

According to bending fracture analysis, the main crack generates first in the first SiCp/6061Al composite bars due to the high particle volume fraction, and then it is blunted by the matrix near the failure bars and has to propagate along the (SiCp–6061Al)/6061Al interface, which means that the large matrix region can absorb fracture energy by deformation and load bearing after adjacent composite bar fracture. The hierarchical fracture of the composite bars and the matrix bearing load can protect the composite from catastrophic failure, which is superior to conventional composites with a homogeneous microstructure. In addition, the large matrix region exhibiting different deformation characteristics plays an important role in improving the composite fracture toughness by delaying the onset

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Fig. 13. Load vs. crack opening displacement (COD) curves of the two composites [48].

void nucleation in the designed (SiCp–6061Al)/6061Al composites [119]. Otherwise, the main crack nucleation and growth until fracture are almost instantaneous in the composites with a small matrix region. Therefore, the superior fracture toughness and strengthening effect can be attributed to the reinforcement clustering bar, the large matrix region and the strong interfacial bonding. In addition, the improvement in fracture toughness of the designed (SiCp–6061Al)/6061Al composites may be related to the residual stress state [49,124–126]. In fact, as early as 1991, Nardone et al. [45,46,62] seemed to be the first to develop the so-called microstructurally toughened particulate-reinforced Al MMCs consisting of bar-like (not layer) clustered regions (SiCp/6061 composites) with 30 vol.% SiCp reinforcement and continuous pure Ti matrix regions. This is different in that the continuous ductile toughening regions are tubular pure Ti with different nominal outer diameters, wall thicknesses and volume fractions. After T6 heat treatment, the prepared Ti–30 vol.%SiCp/6061Al composites exhibited 20 times higher energy absorption relative to the conventional 30 vol.%SiCp/6061Al composites. The energy absorption is increased with the increasing volume fraction of the pure Ti toughening region. Moreover, the tensile fracture elongation of the toughened composite is increased with the increasing volume fraction of the toughening Ti region. The authors concluded that the main reasons for the significant improvement in energy absorption are the crack deflection and de-cohesion between the bar reinforcing region and the tubular toughening region. In addition, Nardone et al. [45] first introduced 304 stainless steel tubes of different sizes into B4C/NiAl brittle composites in order to improve the composite toughness; this is similar to the bar clustering reinforcement. The results showed a significant elongation of 15–35% and a one order of magnitude increase in toughness compared with the conventional B4C/NiAl composites. The significant improvement can be attributed to the constrained yield of the 304 stainless steel, which can prevent composite failure after the NiAl composite has cracked. Additionally, Gupta et al. [127] fabricated the hybrid (Fe + Ti)/Al composites composed of a continuous Fe mesh, discontinuous Ti particles and a 3D continuous Al matrix, which is similar to the structure of concrete, in viewing the structure and fabrication process. The microstructure observation showed the great grain refinement of the Al matrix, a relative high porosity level and nearly no interfacial reaction products. Not only were there significant improvements in the elastic modulus, YS and UTS, but also a relatively reduced decrease in ductility was achieved, all of which are superior to those of the conventional SiCp/Al composites. 4.2.2. Laminated and bean-like reinforcement-lean phases Pandey et al. [47], using PM routes, designed and fabricated ‘bean’ materials, composed of SiC/ 7093Al composites and ductile Al/Al-alloy particles with large sizes. The overall volume fraction of SiCp was kept constant at 15%, which led to 16.7% and 20% local volume fractions in the reinforcement-rich phases and ductile Al ‘bean’ regions with different volume fractions (10% and 25%), as

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Fig. 14. Optical microstructures of (a) the discontinuously reinforced aluminium matrix composites (DRA), (b) the material with 25 vol.% of small Al beans, (c) the material with 10 vol.% of large CU-50 beans, and (d) the material with 25 vol.% of large Al beans, showing the distribution of these particles [47].

shown in Fig. 14. The authors drew one conclusion: that the modulus is primarily dependent on the overall volume fraction of SiC, rather than on their distribution. This seems to be a paradox to the above conclusion that the elastic modulus increases with the increase in the reinforcement contiguity. The reason is that the SiCp volume fraction is so low, even increasing from 15% to 16.7% and 20%, that the identifiable increase in SiCp contiguity is still not formed. In addition, the YS of the composites was influenced by the pure Al particles, because the softer pure Al can cause the premature yielding of the composites. Finally, the fracture elongation of the composites decreases with the increasing volume fraction of the bean Al particles, which can be attributed to the increasing local volume fraction of SiC in the continuous composite phase. The results show that the initiation toughness of the bean composites is higher than that of the individual composite except for the bean composites with 10 vol.% pure Al particles. Moreover, the steady state toughness of the bean composites reveals a much larger improvement compared with the initiation toughness, an increase of as much as 70%. According to the toughness data, not only the high volume fraction but also the large size of the ductile Al beans is beneficial to the toughness of the bean composites. Furthermore, the composites with CU-50 Al-alloy beans exhibit a higher toughness than that with pure Al beans under the same particle size and volume fraction. According to the R-curve behaviour, the presence of beans can significantly enhance damage tolerance, particularly in the high strength condition. According to the damage and the fractography, two extrinsic toughening mechanisms may be operating in the bean composites: crack bridging by uncracked Alparticle ligaments and crack trapping by the ductile Al particles. Besides the bean microstructure, as early as the 1980s, the laminated reinforcement agglomeration microstructure seemed to be the first hierarchical architecture deliberately tailored for the performance improvement of MMCs [44,128]. In the past three decades, laminated composites have been intensively studied in various systems, including metal–metal [129], metal–composites [44], metal– ceramic [130], metal–intermetallic [131] and ceramic–ceramic [132]. In the laminated composites, the layer materials, the individual volume fractions of the layers, the layer thickness and the

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fabrication processes can be selected, designed and controlled in order to obtain the expected properties. Although the laminated composites exhibit anisotropic properties, the superior toughening effect and easy fabrication process guarantee them extensive potential, particularly for Al MMCs. In order to enhance the toughness of Al MMCs while retaining strength and stiffness compared with the conventional homogeneous composites, inspired by thin sheets of Al MMCs having significantly higher toughness than thick plates [133], Pandey et al. [44] designed and fabricated a laminated particulate-reinforced aluminium composite, comprising a 15 vol.%SiCp/7095Al composite layer (1.4 mm) and a 3003Al–Mn alloy layer (200 lm), as shown in Fig. 15, which exhibits flat and defect-free bonding between the alternate layers. As indicated in Table 2, the laminated composites exhibit a slightly low modulus, YS and UTS, but a much higher work hardening exponent compared with the individual 15 vol.%SiCp/7095Al composites with a conventional homogeneous microstructure. The slightly lower elastic modulus is corresponding to the lower overall volume fraction of SiCp reinforcement (12.9% vs. 15%). The lower UTS can be mainly attributed to the adverse testing direction of the laminated composites, which is perpendicular to the extrusion direction of the original composites, besides the low SiCp volume fraction. It is well known that the extruded materials possess anisotropic tensile properties [134]. According to analysis, the relatively low YS of the laminated composites relates to, on the one hand, the Al–Mn alloy layer being softer than the 7093 matrix alloy, which results in premature yielding of the alloy layer, and on the other hand, to the Al–Mn alloy layer in itself bearing the tensile residual stress formed in the fabrication process due to the different thermal expansion coefficients, which can aid the macro tensile stress causing the premature yielding of the alloy layer. Additionally, the tensile residual stress can also delay the onset of compression yielding, which results in the relatively higher YS and work hardening rate in compression. In fact, the early yielding of the laminated

Fig. 15. Microstructure of laminated composite containing DRA and unreinforced Al–Mn layers [44].

Table 2 Tensile and compressive properties of laminated and monolithic DRA composites in different loading and heat treatment conditions [44]. Materials

Loading and heat treatment

Young’s modulus (GPa)

YS (MPa)

UTS/compressive strength (MPa)

Strain to failure/ unlocking stain (%)

Work hardening exponent (N)

Laminated composite Laminated composite Laminated composite Monolithic DRA Monolithic DRA

Tensile (underaged)

88.0

425

546

5.3

0.102

Compression (underaged) Compression (peak-aged) Tensile (underaged) Tensile (peak-aged)

87.7

455

702

3.6

0.133

87.0

559

670

2.1

0.099

89.9 95.6

503 642

629 694

5.9 1.8

0.089 0.037

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composites can improve their work hardening exponent as an extrinsic effect, which is always beneficial in improving their damage tolerance, because there would be longer transition from the onset of yielding to the establishment of full-scale plastic deformation throughout the laminated composites. This allows the stress ahead of a crack or notch to be dispersed in a larger volume, rather than being localized to a small zone compared with the individual composites. Therefore, the damage tolerance (energy absorption) of the laminate composites is more than twice that of the individual composites, as shown in Fig. 16. Moreover, the initiation fracture toughness of the laminated composites is increased by 79% (from 19 to 34 MPa m1/2) compared with the individual composites. Additionally, the laminated composites had a tearing modulus, representing the crack resistance behaviour, more than four times higher than the individual composites according to the experimental results. According to the microstructure observations of the deformed laminate composites, the slip band representing plastic deformation was spread throughout the gauge length of the specimen, while being concentrated very close to the fracture surface for the individual composites [44,135]. That is to say, the laminated composites encouraged uniform straining, which is very useful in improving the composite ductility/toughness. Observation of the fracture surfaces showed that delamination, concentrated in the fast fracture region, occurred at the interface of the two adjacent layers not inside any of the layers. Therefore, delamination failure can encourage a significant amount of plastic deformation and energy absorption, which are beneficial to fracture toughness. Additionally, significant void growth of the primary voids, associated with the fractured SiC particle, occurred in the laminated composites, indicating that the laminated composites possess higher toughness than thick individual composites. In summary, the crack originating from the reinforcement agglomeration layer can be blunted by the toughening Al layer. This mechanism allows a desirable large plastic zone to develop before the failure conditions are reached at the crack tip, which can effectively enhance energy absorption/toughness/damage tolerance. In the context of laminated structures, the systems based on Cu and Al are largely selected due to their combination of physical properties and mechanical properties [136–140]. Guo et al. [141] designed and fabricated Cu–CuAl intermetallic–Al laminated composites for electric devices. The microstructure observation shows that the prepared laminated composites comprise three layers, Cu, Al and the intermetallic interface phase, with good bonding between any two phases. However, no delamination was observed, which indicates that the interfacial delamination is not a major failure mechanism in the Cu–CuAl intermetallic–Al laminated composites [142]. According to observation, the cracks initiated in the interface intermetallic phase and then terminated in front of ductile metal layers, indicating the key role of the ductile layers [143–145]. The tensile test results showed that the YS (147 MPa) and the UTS (243 MPa) are close to those of the commercial pure Cu [137]. The tensile fracture elongation (11%) is much larger than that of the intermetallic phases. Therefore, it can be said that the intermetallic phases were effectively toughened by the ductile Cu and Al layers.

Fig. 16. (a) Load vs. COD, and (b) K vs. Da for the laminated and monolithic DRA composites in the nominally underaged condition [44].

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Intermetallic compounds, particularly for TiAl, have attracted much attention due to their low density, high specific strength, high specific stiffness, high-temperature durability, high creep resistance and high-temperature oxidation [146]. In order to improve the low ductility and toughness and retain the high hardness of TiAl intermetallic compounds, many researchers designed alternative Ti-TiAl laminated composites [147]. Harach and Vecchio [148] successfully fabricated laminated Ti–Al3Ti composites with a well-bonded interface by reactive foil sintering in air. Subsequently, in order to retain large ductile zones for toughness improvement, they also fabricated Ti-TiAl laminated composites using thick Ti and Al foils with very high thickness, even above 1 mm, by a one-step process [149]. The test results show that the Ti–Al3Ti laminated composites exhibit much higher specific fracture toughness, partially due to their low density. Furthermore, the composites exhibit a more than one order of magnitude improvement in toughness over monolithic Al3Ti with the addition of just a 20– 40 vol.% ductile Ti phase in the form of continuous layers. The crack-initiation toughness is nearly five times that of the intermetallic compounds. The significant improvement in toughness is achieved by plastically stretching the uncracked ductile Ti layers, which can bridge the crack and shield the crack tip. Therefore, the fracture resistance and toughness are increased with increasing the ductile Ti volume fraction. Finally, taking into account a comparison of the specific toughness vs. specific elastic modulus of various structure materials, Ti–Al3Ti laminated composites may be potential candidates for structural applications requiring weight optimization. On the basis of the above experiences, Cui et al. [150–152] successfully fabricated micro-laminated TiB2/TiAl composites by roll bonding and reaction annealing followed by an effective densification process, as shown in Fig. 17. In the composites, the micro TiB2-rich layers as the reinforcement and the micron scale TiAl matrix layers are alternatively stacked to construct the named micro-laminated structure. Due to rolling deformation, Al consumption and the densification process, TiB2 particles are concentrated on the micro layer and then form the micro TiB2-rich layers, which can effectively refine the TiAl matrix grains [151]. Finally, both the unique micro-laminated structure and matrix grain refinement contribute to the significant improvement in micro-laminated TiB2/TiAl composites in elastic modulus, high-temperature tensile strength and ductility compared with conventional TiAl materials, as shown in Fig. 18 [152]. Moreover, the improvements increase with the increasing volume fraction of the added TiB2 particles. Simultaneously, laminated Ti5Si3/TiAl composites [153,154] and TiB2/NiAl composites

Fig. 17. Representative microstructures of the Ti–TiB2/Al composites. (a) Hot-pressed sample [150]; (b) hot-rolled sample [150]; (c) as-annealed sample [151]; (d) final micro-laminated TiB2/TiAl sample [151].

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Fig. 18. Tensile stress–strain curves of the micro-laminated composite sheets containing different volume fractions of TiB2 particles measured at 750 °C [152].

[155,156] were also successfully fabricated using similar processes. Moreover, the two composites also exhibit refined grain sizes by the micro-laminated reinforcement-rich layer and improved mechanical properties by the particular micro-laminated structure and the refined grains. 4.2.3. Alternating macro-ring microstructures Mg alloys represent the lowest density, the highest damping capability and modest mechanical properties among Al, Cu, Ti and Mg [157,158]. Moreover, their hardness, elastic modulus, strength and ductility can be further significantly improved by fabricating MgMMCs [159,160]. Therefore, they can be used for energy-critical structural applications, which offer significant weight savings allied with superior mechanical properties. Particularly, inspired by the fairly concentric tree-trunk growth-ring structure, Wong et al. [63] designed Mg–Al2O3/Mg composites with concentric alternating macro-ring microstructures, in order to further enhance the mechanical properties compared with the conventional MgMMCs with a homogeneous microstructure. The designed ring-structured microstructure comprises a monolithic ring Mg phase and a ring 1.11 vol.%Al2O3/Mg composite phase, which can be fabricated by the PM technique and hot extrusion deformation, as shown in Fig. 19. Fig. 20 shows the scanning electron microscope (SEM) micrographs of Mg–Al2O3/Mg ring-structured hybrid composites with different ring thicknesses. The dark rings correspond to the monolithic Mg phase, while the light grey rings correspond to Al2O3/Mg composite phase. The ring thickness decreases with increasing ring radius which is similar to tree-trunk growth rings [63,161,162]. With respect to the thinner layers before the extrusion, the thinner the rings, the higher the ring number and the more homogeneous the distribution of Al2O3 particles after extrusion [157,163]. As shown in Table 3, the micro-hardness of the ring composite phase is higher than that of the ring Mg phase, which is not affected by ring thickness [164]. It was unexpected that the tensile properties, including the YS, the UTS and the ductility, peaked for the composites with a 3 mm layer thickness, which is much higher than that of the monolithic Mg, and that the ductility would be significantly improved compared with that of the Al2O3/Mg composites. Subsequently, the tensile properties of the composites decreased with the decreasing layer thickness and the increased number of layers. The reason for the decreased tensile properties of the composites with a 4 mm layer thickness can be attributed to the irregular concentric alternating rings, which have a detrimental effect on tensile properties [165–167]. However, the increased tensile properties with increasing ring thickness from 1 mm to 3 mm can be attributed to the increasing micro-interfaces between the Mg and Al2O3/Mg composite rings. The authors also pointed out that the highest tested mechanical properties of the composites with a 3 mm layer are consistent with the estimated results of the RoM theorem. Additionally, this work

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Fig. 19. Schematic diagram showing: (a) layered arrangement (layers are of equal thickness) of Mg and Al2O3/Mg prior to extrusion and deformation of layers during extrusion and (b) deduction of tree-trunk growth-ring resemblance after extrusion [63].

appears to be the first instance of enhancing the mechanical property for MgMMCs using macro- and microstructure design. In recent work [168], the test results of the hierarchical (Al + Al2O3)/Mg nano-composites showed significant enhancement (>80%) in both tensile and compressive strengths, and the highest improvement in tensile ductility (42%) as compared with monolithic Mg, which can be mainly attributed to the grain size refinement hierarchical structure. The tensile behaviour of the hierarchical composite is remarkably affected by the particle size and volume fraction of Al2O3 nano-particles within the Al2O3-rich clustering phase [169,170]. In addition, the prismatic texture becomes stronger while the basal texture weaker by decreasing the size and volume fraction of the Al2O3–Al phase, which is related to the mechanical behaviours of the composites.

4.3. Tailored microstructures with reinforcement-rich network In order to further enhance the performance of Al MMCs, Kaveendran et al. [69,171] fabricated in situ 10 vol.%(Al3Zr + Al2O3)/2024Al composites by tailoring a controlled network microstructure [68,172]. As shown in Fig. 21, the synthesized reinforcements are distributed around large spherical

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Fig. 20. (a–d) Representative optical images of transverse cross sections of the extruded Mg–Al2O3/Mg ring-structured hybrid composites. (e) Representative SEM micrograph showing uniform distribution of nano-Al2O3 particulates in the extruded Mg– Al2O3/Mg ring-structured hybrid composites [63].

Table 3 Results of the mechanical testing of Mg and Mg–Mg/1.11 vol.%Al2O3 ring-structured hybrid composites [63]. Material

Mg Mg–Mg/1.11 vol.% Al2O3 (4 mm Mg–Mg/1.11 vol.% Al2O3 (3 mm Mg–Mg/1.11 vol.% Al2O3 (2 mm Mg–Mg/1.11 vol.% Al2O3 (1 mm Mg/1.11 vol.% Al2O3 [157]

Microhardness (Hv)

Layer) Layer) Layer) Layer)

Mg

Mg/Al2O3

40.0 ± 0.9 39.9 ± 0.8 39.9 ± 0.8 40.3 ± 1.3 41.0 ± 0.9 –

– 53.4 ± 1.5 53.4 ± 1.5 52.5 ± 1.1 52.9 ± 0.7 –

0.2%YS (MPa)

UTS (MPa)

Failure strain (%)

115.2 ± 3.4 125.0 ± 3.4 149.3 ± 5.4 112.0 ± 1.6 101.7 ± 5.1 194 ± 5

194.8 ± 10.6 195.1 ± 6.3 222.6 ± 10.1 172.6 ± 9.0 151.1 ± 7.0 250 ± 3

10.7 ± 3.3 8.3 ± 1.0 13.0 ± 0.7 5.7 ± 2.7 5.8 ± 1.9 6.9 ± 1.0

2024Al particles and then form a 3D network microstructure. Moreover, the reinforcement-rich network boundary is not totally continuous but interpenetrated by the matrix between the reinforcements. Hardness results showed that the hardness of the tailored composites with a network microstructure is increased by 35.2% compared with that of unreinforced 2024Al alloy, and even increased by 7.9% compared with the composites with a homogeneous microstructure under a given reinforcement volume fraction. As shown in Table 4, the tensile test results showed that the YS and UTS of the network structured 10 vol.%(Al3Zr + Al2O3)/2024Al composites are increased by 66.7% and 18.6% compared with those of the unreinforced alloy and increased by 4.1% and 12.5% compared with the homogeneous composites.

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Fig. 21. Microstructure of the network-structured (Al3Zrp + Al2O3np)/2024Al composite at (a) low magnification, (b) high magnification and (c) SEM showing the morphology of Al3Zr particles and the interpenetrating nature of the matrix [69].

Table 4 Comparative study of the mechanical properties of the unreinforced alloy, homogeneous and network-structured (Al2Zrp + Al2O3np)/2024Al composite [69]. Materials

Hardness (HRB)

YS (MPa)

UTS (MPa)

Elongation (%)

Elastic modulus (GPa)

Unreinforced alloy Homogeneous distribution Network distribution

42 ± 0.5 52.6 ± 0.8 56.8 ± 0.6

105 ± 3 168 ± 2 175 ± 3

220 ± 4 232 ± 2 261 ± 3

12 ± 0.3 1.33 ± 0.2 2.3 ± 0.2

74.3 ± 0.5 82.3 ± 0.3 86.4 ± 0.5

Moreover, the elastic modulus of the tailored composites is obviously higher than that of the homogeneous composites. It is interesting that the tensile fracture elongation is increased by 76.9% compared with that of the homogeneous composites. Furthermore, the elongation can be further increased by subsequent hot deformation [173]. Considering the fixed overall volume fraction of the reinforcement, the improvements can be attributed to the network distribution of the reinforcement, while the improved ductility is likely to be due to the interpenetrating microstructure and the large matrix region retained, which, according to fracture analysis, can decrease the crack-propagation speed and blunt the crack. Bhanu Prasad et al. [67] fabricated 30 vol.%SiCp/2124Al composites with a similar network microstructure by PM, consisting of isolated matrix particles and a surrounded network phase with a much higher local volume fraction of the reinforcement (even over 50%). They also fabricated a series of 30 vol.%SiCp/2124Al composites with different relative particle size (RPS) ratios of 40:1, 25:1, 18:1, 12:1, 8:1, 5:1 and 3:1. However, the tensile strength and elongation decreased with increasing RPS, i.e., Al powder sizes, which is contrary to the above conclusions. Indeed, under a given volume fraction,

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the clustering tendency increases with increasing matrix particle size, which results in the remarkably increased local volume fraction of reinforcement. This phenomenon is harmful to the mechanical properties of the composites due to the decreased matrix contiguity. Geni and Kikuchi [98] also pointed out that the non-uniform distribution of the SiC particles can lead to a large decrease in the tensile strength of SiCp/2024Al composites due to the premature local fracture generated in the particle clustering phases. In fact, these phenomena effectively demonstrate that sufficient matrix contiguity or modest local volume fraction of the hard reinforcement in the clustering phases is necessary to exploit a superior combination of mechanical properties. Prielipp’s work [174,175] also confirmed this statement, i.e., both the fracture strength and toughness of the Al2O3/Al composites with interpenetrating networks are increased by increasing the diameters of the metal matrix ligament, and, surprisingly, considering that the YS of the Al matrix is much lower than that of the Al2O3 reinforcement, are increased with increasing metal matrix content. This phenomenon effectively demonstrated that adequate amounts and large sizes of the ductile phase among the stiffer reinforcement-rich phases are necessary to exploit the strengthening and toughening effects of the reinforcement and matrix phases. Therefore, it is inevitable that the composites with a much higher local volume fraction of reinforcement fabricated by ex situ PM exhibit inferior mechanical properties. Subsequently, Slipenyuk et al. [176] fabricated 15 vol.%SiCp/Al6Cu4Mn composites by PM followed by hot extrusion. The microstructure observation showed that increasing the matrix particle size significantly increased the probability of cluster formation. The increasing reinforcement cluster led to a decrease in fabricability and mechanical properties, including YS, UTS, elongation and Young’s modulus. Similarly, loss of hot rolling fabricability for the composites with high RPS was reported by Stone and Tsakiropoulis [177]. In fact, the degradation of fabricability and the mechanical properties can be mainly attributed to the formation of the airtight reinforcement agglomerations. This phenomenon just indicates that reinforcement agglomeration phase must not be airtight but interpenetrated in a non-uniform microstructure for performance improvement. Li and Ng [178] fabricated a lightweight Al–15 wt.%Li2O composite billet with a network microstructure, in which an Li2O phase is uniformly distributed around Al particles and then forms a network microstructure at a higher hierarchy. Subsequently, the c-LiAlO2 phase was in situ synthesized in the arc-melting process and distributed around the Al matrix and then formed a nano-network microstructure at a lower level during the cooling process, i.e., two-scale network structures, which gave a Vickers’ hardness of 143. In order to overcome the poor ductility and toughness of TiAl intermetallic compounds at room temperature, Yang et al. [179] designed and fabricated in situ Ti2AlC/TiAl composites by mechanical alloying, low-energy milling and spark-plasma sintering using pure Ti powders, pure Al powders and 10 vol.% carbon nanotubes (CNTs). As shown in Fig. 22a, the in situ Ti2AlC reinforcement is distributed around the TiAl matrix particles and then forms a network microstructure due to the low-energy milling and solid state sintering processes employed. Besides the network microstructure, equiaxed ultrafine particles in the range of 200–400 nm in the TiAl matrix have been successfully formed at the optimal sintering temperature (Fig. 22b), which is related to the mechanical alloying process and the ceramic Ti2AlC network microstructure [14]. Both the reinforcement network microstructure and ultrafine matrix grains contribute to the significant improvements in the Vickers hardness of 6.12 GPa, the compressive YS of 2058 MPa and the compressive fracture strength of 2217 MPa. Inspired by the network microstructure [172,179,180], Zhou et al. [181] designed and fabricated Ti2AlN/TiAl composites with a network microstructure by reaction hot pressing using pure Al and nitrided Ti powders. The prepared Ti2AlN/TiAl composites with a network microstructure exhibited superior compressive strength of 1112.1 MPa, 958.9 MPa and 686.7 MPa at 800 °C, 900 °C and 1000 °C, respectively. The compressive strength of the network-structured composites is increased by 14.1%, 40.5% and 61.6% compared with that of TiAl alloy, and also increased by 16%, 28.1% and 33.3% compared with that of the conventional Ti2AlN/TiAl composites with a homogeneous microstructure at 800 °C, 900 °C and 1000 °C, respectively. It is worth pointing out that the strength of the network-structured composites at 900 °C is equivalent to that of the conventional composites at 800 °C. That is to say, the service temperature of the composites can be further increased by 100 °C simply by controlling the reinforcement distribution under similar reinforcement volume fractions. The improvement in the high-temperature property is attributed to the reinforcement network

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Fig. 22. SEM (a) and TEM (b) micrographs of the as-sintered Ti2AlC/TiAl composites fabricated at 950 °C [179].

distribution, which can overcome the weakening effect of the grain boundary at high temperatures as reported [182]. In addition, the increment in strength increases with increasing test temperatures from 800 °C to 1000 °C (16% to 33.3%), which is probably attributed to the fact that the continuous network can effectively bear the compressive stress when the matrix is very soft.

4.4. Bi-continuous microstructures Bi-continuous microstructure corresponds to both the reinforcement-rich phase and reinforcementlean phase being 3D continuous. Therefore, both matrix and reinforcement can contribute their own superiority to the macroscopic properties of the composites, thus endowing the multi-functional characteristics of the composites, such as higher elastic modulus, strength, ductility, electrical conductivity and thermal conductivity. Wegner and Gibson [73–75] respectively fabricated bi-continuous composites by infiltrating bronze and polymer resin into porous 420 stainless steel and 316 stainless steel for toughness improvement. Both the ductile phase and the strong steel phase are interpenetrated and then form two separate continuous networks in the two composites. Therefore, these composites were sometimes called interpenetrating or co-continuous composites in the literature [73–75]. Moreover, the steel contiguity increases while the ductile phase contiguity decreases with increasing volume fractions of the steel phase, as shown in Fig. 23. In addition, because the steel phase is integrated and coherent, this bi-continuous microstructure can be viewed as a one-scale microstructure unlike the two-scale microstructures developed [115–118]. In other words, the local volume fraction of reinforcement in the reinforcement-rich phase is 100%. The crack growth observation shows that the cracks preferred to follow the softer ductile phase and the interface between the two phases. Therefore, with the increasing volume fraction of the ductile bronze, the toughness of the bronze-420 stainless steel composites is increased [74]. The reason for this phenomenon is that more ductile phase can result in more crack bridging and unloading in the wake of a process zone by absorbing more energy, which is characterized by secondary cracking behind the crack tip and plastic deformation near the crack tip [73]. However, compared with the interpenetrated bronze phase, the polymer resin phase appears to be too weak to contribute to the mechanical properties of the bi-continuous composites [73,75]. The most valuable conclusions from these investigations are interpreted as follows. The toughening effect of the ductile phase in the bi-continuous composites increases with the increase in its volume fraction below the optimal value. Therefore, the ductile phase with very low volume fraction cannot visually contribute to the initiation toughness of the composites. However, the toughness of the composites also decreases with the increase in volume fraction of the ductile phase exceeding the optimal, due to the relatively lower strength of the ductile phase. On the other hand, the toughening effect also increases with the increase in strength of the ductile phase below the optimal value. When the

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Fig. 23. SEM micrographs of resin-316 stainless steel bi-continuous composites with different volume fractions of 316 stainless steel: (a) 90.7%, (b) 84.7%, (c) 71.8%, and (d) 65.0% [75].

strength is excessively low, the ductile phase cannot contribute to the overall mechanical properties. However, if the strength is excessively high, the ductile phase cannot contribute to the fracture toughness of the composites either due to its low ductility. In order to improve the fracture toughness of Al2O3 ceramic, Konopka et al. [16] fabricated Fe/Al2O3 composites by infiltrating molten Fe into porous Al2O3 ceramic pre-form. Both Al2O3 granules and Fe separately formed an interpenetrating network, which will probably inspire a superior combination of strength and toughness. In addition, Konopka and Ozieblo also fabricated 0 wt.%, 10 wt.%, 30 wt.% and 50 wt.% Fe/Al2O3 composites with a homogeneous microstructure by PM and high-temperature sintering [183]. In the literature, ductile Fe is termed reinforcement which is added to increase the toughness of Al2O3 base materials. Therefore, the fracture toughness of the composites increases with increasing Fe content from 0% to 30%. However, the fracture toughness no longer increases when the Fe content exceeds from 30% to 50%, which is attributed to the formation of the FeAl2O4 spinel interfacial phase. As shown in Fig. 3, Travitzky et al. [77] successfully fabricated one novel Al2O3/TiAl composites, comprising a 3D continuous Al2O3 reinforcement network and 3D continuous TiAl matrix network, by reactive infiltration of molten Al into porous Al2O3–TiO2 preform. The bi-continuous composites show a continuous Al2O3 skeleton and a continuous TiAl matrix [184]. This microstructure of the TiAl matrix composites results in fracture toughness of 7.1 MPa m1/2, bending strength of 540 MPa and Vickers hardness of 9.8 GPa. In a previous work [184], the Al2O3/TiAl composites exhibited fracture toughness of 8.6 MPa m1/2, fracture strength of 542 MPa and HV10 hardness of 565 MPa. Such superior combination of fracture toughness and the strength are related to crack deflection by the Al2O3 grains and crack bridging by the relatively ductile intermetallic matrix, as shown in Fig. 24. Certainly, the continuous intermetallic network structure inspires some plastic deformation during crack propagation, which is beneficial to fracture toughness. The damage in the form of Al2O3 particle fracture prior to final fracture (Fig. 24a) may result in a significant degradation of properties [33]. Similarly, Horvitz et al. [185] also successfully fabricated in situ Al2O3–TiAl/Ti3Al composites by the self-propagating high-temperature synthesis of compacted nano-size TiO2 and micron-size Al powders (3TiO2–7Al). The in situ Al2O3–TiAl/Ti3Al composites contain a continuous ceramic network

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Fig. 24. TiAl–Al2O3 composite with indentation crack: (a) mostly deflected by alumina grains and (b) the metal phase (mainly Ti3Al) has plastically deformed to sharp edges [77].

formed by micron Al2O3 particles and an intermetallic matrix network formed by fine sub-micron lamellar TiAl and Ti3Al phases, which are probably beneficial to the fracture toughness and high-temperature strength. Dong et al. [76] fabricated Cr3C2/Cu composites with a bi-continuous microstructure by infiltrating molten copper into a fine Cr3C2 foam pre-form as shown in Fig. 25 [186,187]. In the composites, the Cu matrix and the Cr3C2 reinforcement network are mutually interpenetrated [188]. The prepared composites exhibited a Vickers’ hardness (Hv) of 2716 MPa, which was about three times higher than that of pure Cu. The elastic modulus of the composites was increased to about 205 GPa, which is close to the theoretical upper bond value calculated by RoM. That is to say, the present improvement in the isotropic composites may be viewed as the most remarkable, and can be attributed to the finer carbide reinforcement network and the much higher contiguity of the Cr3C2 reinforcement as well as the good interfacial bonding between the Cu and Cr3C2 phases. As early as 1995, Breslin et al. [189] successfully fabricated Al2O3/Al–Si composites with a bi-continuous microstructure by a liquid phase displacement reaction. The test results showed that the prepared composites exhibited twice the thermal conductivity and a lower thermal expansion coefficient, both significant for some special applications. Due to the high contiguity of the ceramic phase with a high volume fraction of 70%, the composites exhibited a very high Young’s modulus of 215 GPa, and superior compressive strength of more than 600 MPa at 400 °C and 300 MPa at 600 °C, but relatively low ductility [189,190]. Moreover, the fracture toughness was significantly improved due to the crack blunting and crack-bridging effects. Similarly, the flexural strength and the fracture toughness of graphite can be significantly improved to 122 MPa and 1.93 MPa m1/2 from 61 MPa and 0.94 MPa m1/2, respectively, by infiltrating Al into

Fig. 25. SEM micrographs of 100% Cr2O3 green sample (a) and Cr3C2/Cu composites (b) [76].

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graphite per-forms, which is also attributed to the crack-bridging effect generated by the ductile matrix [191]. It is unusual that the bi-continuous composites exhibited a bi-linear manner including elastic and elastic–plastic behaviour, which was verified by experimental and analytical results [190]. In addition, Scherm et al. [192] reported a similar phenomenon, their Al2O3/AlSi9Cu3 composites exhibited linear behaviour under low stress and non-linear behaviour under higher stress, which is in agreement with Daehn’s observation on bi-continuous MMCs [190]. The above bi-linear behaviours can be attributed to the combination of the elastic behaviour of the continuous ceramic skeleton accompanied by the plastic behaviour of the continuous metal matrix at high stress. Certainly, the bi-continuous Al2O3/AlSi9Cu3 composites also exhibited good interfacial bonding and significant improvements in elastic modulus, tensile strength, fatigue limits and much lower thermal conductivity [192,193]. A reported bi-continuous 420 stainless/90Cu–10Sn bronze composites, fabricated by 3D printing and infiltration processes, exhibited superior strength and thermal expansion coefficients according to the experimental and FEM results [74]. The strength of the composites can be enhanced when the stainless particles become continuous, while the thermal expansion coefficient is close to the RoM upper bound due to the continuous bronze. Peng et al. [13,194,195] fabricated bi-continuous Al2O3/6061Al composites by the squeeze casting of the molten 6061Al alloy into the pre-arranged Al2O3 ceramic foams, learning from the previous work [196]. As seen from Fig. 26, it is also observed that the struts of the ceramic foams, fabricated using fine powders, are denser than those using coarse powders [197]. The composites with dense struts are expected to show high performance [198]. Furthermore, as reported in Mattern’s work [199], micro-porosity and grain size in the pore walls are influenced by changing the sintering temperatures, which subsequently affects the permeability and strength of the ceramic foams in the infiltration process. As shown in Fig. 27, not only the reinforcement phase is continuous as a result of the ceramic struts, but also the matrix is continuous due to the reticulation window and the strut pores. Moreover, the struts themselves are actually an interior Al2O3/6061Al composite phase with a much higher local reinforcement volume fraction (50–58%), while the overall composite has a much lower overall reinforcement fraction of about 5.8%. In addition, the contiguity of the matrix decreases while that of the reinforcement increases with the increase in the overall volume fraction of the reinforcement due to the decreasing window size of the ceramic foam. As expected, the bi-continuous Al2O3/6061Al composites with bi-continuous microstructures exhibit a higher modulus than the conventional composites with a homogeneous microstructure at a given volume fraction [111], which can be mainly attributed to the increased contiguity of the ceramic reinforcement in the struts. Tuchinskii deduced a specific model for the theoretical predictions of the structures consisting of two interpenetrating networks, which is expressed as follows [200]:

Ec ¼ Eb ð1  tÞ2 þ Ea t2 þ

2Ea tð1  tÞ t þ ðEa =Eb Þð1  tÞ0

ð7Þ

where Ec is the overall elastic modulus of the composites, and the parameter t is related to the volume fraction Va by

Fig. 26. Microstructure of sintered foams made with fine alumina powders (a) and coarse powders (b) [194].

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Fig. 27. The microstructure of the composite made from MA130 foam (a) and strut details at high magnification (b) [13].

V a ¼ ð3  2tÞt 2

ð8Þ

As can be seen from Fig. 28, the experimental moduli of the bi-continuous Al2O3/6061Al composites are higher than the predictions by the H–T and H–S equations, while lower than the prediction by RoM. However, the Tuchinskii model does give a good representation of the experimental results for the bi-continuous Al2O3/6061Al composites. The present results are also consistent with the prediction by finite element analysis that such topological structure arrangements offer the highest stiffness [201]. Moreover, when the reinforcement volume fraction is lower than 20%, the electrical resistivity of the bi-continuous Al2O3/6061Al composites agrees with the general effective media (GEM) equation introduced by McLachlan et al. [202]. However, the experimental results of the bi-continuous Al2O3/6061Al composites are lower than those of the GEM equation when the volume fraction exceeds 20%. This indicates that the bi-continuous Al2O3/6061Al composites with a bi-continuous microstructure have higher electrical conductivity than other composite systems, which should be attributed to the high contiguity of the matrix. Additionally, the bi-continuous Al2O3/6061Al composites should also exhibit higher strength because of the continuous reinforcement phase, and better damage tolerance and thermal conductivity due to the continuous matrix phase. Using pressureless infiltration method, Wang et al. [203] successfully fabricated Si3N4/Al–Mg composites with a bi-continuous microstructure. They concluded that the strength can be significantly

Fig. 28. The measured elastic modulus [13] of the composites (d) and the Duralcan Al-MMC (T6) () [111] compared with theoretical predications by RoM (A), Halpin–Tsai equation (B) [3], Hashin–Shtrikman equation (C) [83] and Tuchinskii equation (D) [200].

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improved by tailoring the bi-continuous microstructure when the reinforcement volume fraction is lower than 6%. However, the strength decreases with increasing the reinforcement fractions (>6%), which is attributed to cracking, voids, interface cracking, interface de-bonding and matrix damage. Recently, Binner et al. [204,205] investigated the microstructure characteristics of 20 vol.%Al2O3/Al– 10Mg composites with a bi-continuous microstructure by Electron Backscattered Diffraction (EBSD). The results showed that aluminium within each cell has the same orientation and each of the individual cells has a different orientation, although they tend towards the (0 0 1) and (1 0 1) types of orientation parallel to the infiltration direction. It is worth mentioning that San Marchi et al. [206] successfully fabricated bi-continuous 70 vol.%Al2O3/Al composites with 3D periodic architecture by infiltrating liquid Al into a pre-formed Al2O3 tower of periodic cubic symmetry on a sub-millimetre scale. The composites exhibited a moderate density of 3.4 g/cm3, but obviously low thermal expansion of 8.9  106 K1, and significantly high compressive strength of 700 MPa, which can be mainly attributed to the superior 3D periodic microstructure. On the basis of the above bi-continuous composites, Zhao et al. [207,208] successfully fabricated hybrid SiC foam-SiC particle-Al composites with overall reinforcement volume fractions higher than 53% by embedding SiC particles into SiC foam before squeeze casting. The experimental results showed that the coefficient of thermal expansion of the prepared hybrid composites is much lower than that of the conventional SiC–Al composites, which is mainly attributed to the termed double/ two-scale interpenetrating structure. This renders them with great application potential as an electronic packaging substrate material. In addition, the compressive strength can be increased to 660 MPa by employing the hybrid microstructure and ZL109 aluminium matrix. Zhao et al. [79] are confident that the bi-continuous composites possess rather high specific strength, stiffness and wear resistance, and excellent thermal and electrical conductivities because each phase is continuous and interpenetrated. Ni foam reinforced Mg matrix composites with a double continuous microstructure were fabricated by the pressureless infiltration of Mg into Ni foam. As shown in Fig. 29, the Ni foam is composed of cells and struts which form a 3D network structure at a relatively higher level. The struts of Ni foam consist of central triangular opening holes surrounded by strut walls and numerous continuous tiny cavities in the strut walls at a lower level. Therefore, the

Fig. 29. Morphologies of Ni foam reinforcement: (a) foam, in low magnification; (b) strut, in low magnification; (c) cross section of strut; (d) surface of strut [79].

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local continuous structure and the overall continuous structure are simultaneously formed at a lower and higher level after pressureless infiltration, as shown in Fig. 30. This smart microstructure is termed a double interpenetrating structure or bi-continuous structure, which is equivalent to a hierarchal architecture. Moreover, the tiny cavities in the strut walls lead to the formation of the interlocking interface which is beneficial to the mechanical properties of the composites. Therefore, it is expected that the novel Ni foam/Mg composites exhibit a superior combination of strength, ductility and damping capability. Xing et al. [78] first prepared a 3D-SiC network (3D-SiCn) using polyurethane foam and then fabricated 3D-SiCn/Cu composites by squeeze casting, as shown in Fig. 31. According to the SEM, Energy Dispersive Spectrometer (EDS) and X-ray Diffraction (XRD) results of the SiCp/Cu, Sip/Cu and 3D-SiCn/ Cu composites, the reaction between SiC and Cu did not occur in the present fabrication process under the experimental conditions, which is an improvement compared with the previous work [209,210]. However, the remnant silicon in the SiC ceramic network reacts with the Cu matrix and forms a lathy Cu3Si phase, which results in a transition layer, as shown in Fig. 31b. This fragile Cu3Si phase and cracks formed in the fabrication process must seriously degrade the mechanical properties of the composites. Therefore, it is necessary to avoid remnant silicon in the SiC ceramic network to obtain clean 3D-SiCn/Cu composites. It is reasonable to believe that the clean 3D-SiCn/Cu composites would exhibit a superior combination of electrical conductivity, thermal conductivity and mechanical properties due to the continuous SiC reinforcement and continuous Cu matrix. Before drawing an end to this section, it is worth disclosing an attempt to practically fabricate a truly bi-continuous inhomogeneous microstructure with both the reinforcement-rich phase and reinforcement-lean phase being highly 3D continuous. In contrast to the designed short-fibre agglomeration-reinforced 6061Al MMCs as described in Section 4.1 [72], Peng et al. [211] prepared a short-fibre preform containing a cellular array of fibres as shown in Fig. 32. One can easily envisage

Fig. 30. Morphologies of Ni foam/Mg interpenetrating composites: (a) low magnification image; (b) higher magnification image; (c) interface [79].

Fig. 31. SEM images of a 3D-SiC network and a Cu-matrix composite reinforced with 3D-SiC network. (a) 3D-SiC ceramic network; (b) Cu-matrix composite reinforced with 3D-SiC network [78].

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Fig. 32. SEM images of a cellular array of Al2O3 short-fibres offering 3D reticulate network [211].

the resulting microstructure of the composite after melt infiltrating with a matrix alloy, the matrix will fill the small inter-fibre space within the struts forming ‘composite struts’ that is fully 3D continuous and also fill the large space between the struts giving a fully continuous Al matrix phase. The challenge here is to prepare such cellular preforms with sufficient strength to withstand the infiltration pressure during the squeeze casting process. It is believed that such kind of microstructures will offer ultimate combination of strength and toughness.

4.5. Bio-inspired hierarchical microstructures When discussing innovatively designed microstructures, it is imperative to mention bio-inspired hierarchical microstructures. Recently, more and more researchers have realized that nature, after several billion years of stringent evolution, has provided numerous smart structures for special mechanical or functional performances [212,213]. Therefore, nature is compared to a school for scientists and engineers. We just need to reproduce these smart structures in synthetic materials in order to obtain a significant improvement in the combination of performances. However, real natural and biological structures are always significantly complex, and organized by several length scales, i.e., hierarchical structures from the macro-scale, meso-scale and micro-scale to the nano-scale and even the atom-scale [214]. They can perfectly dedicate themselves at different levels, which results in a superior combination of performances, such as the 3000-times higher toughness of mollusc shells than their individual constituents. However, it is extremely difficult to totally replicate smart natural structures. To date, researchers have just achieved imitating bio-inspired structures at relatively high twoor three-scales on the basis of comprehension at different scales. It is surprising that remarkable improvements in performances have been extensively confirmed, including mechanical properties [212,213] and functional performances [215,216]. In this review, we just focus on the improvements in mechanical properties. Therefore, due to their superior toughening and strengthening effects, we mainly review how to understand and imitate the smart structures of nacre, bone, wood and bamboo. It is unanimous that molluscs, such as nacre, are constructed from multi-scale laminated hierarchical structures, as shown in Fig. 33 [214,217], composed of about 95% hard and brittle minerals, such as calcium carbonate (calcite, aragonite), and about 5% soft organic proteins and polysaccharides, mostly located at the interface between mineral tablets [212,213]. Deville et al. [218,219] first developed layered-hybrid composites using freeze casting by imitating the smart nacre lamellar microstructure. Freeze casting can also be used to prepare bone substitutes by freezing hydroxyapatite (HAP) aqueous suspensions followed by ice sublimation and sintering [220]. As shown in Fig. 34a, the lamellar microstructure of prepared Al2O3 scaffolds is really similar

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Fig. 33. The structure of nacre over several length scales (shown for red abalone Haliotis rufescens) [214,217].

Fig. 34. SEM micrographs of the prepared Al2O3 scaffolds at a lower magnification showing lamellar structure (a) and at high magnifications showing rough surface on the lamella (b) and (c) [219].

to that of nacre [218]. Furthermore, the roughness created by the entrapped particles on the lamella surface is functionally similar to that of the inorganic walls (Fig. 34b) which is a key contributor to the final mechanical properties of nacre [219,221]. Although the thickness of the Al2O3 lamella can be decreased by increasing the cooling rate, the rough surface can play a positive role. The length scale of the prepared ceramic scaffolds, which is limited by the selected Al2O3 particle size and the processing conditions, is almost 40 times that of the smart nacre structure, which determines a discounted toughening effect compared with the real nacre structure. The prepared ice-template (IT) scaffolds can be filled with a second phase, such as epoxy or Al–Si alloy with low melting temperature, in order to obtain dense composites with bio-inspired microstructures. As shown in Fig. 35, the HAP/epoxy composites with the bio-inspired microstructure exhibit characteristics of stable crack propagation and active toughening during a three-point bending test, which is qualitatively very similar to nacre. The 45 vol.%Al2O3/Al–Si composites exhibit a general toughness of 5.5 MPa m1/2 allied with a general strength of 400 MPa. However, the toughness and the strength can be further increased to 10 MPa m1/2 and 600 MPa, respectively, by adding 0.5 wt.% Ti for interface modification. The increased toughness and strength can be attributed to the lamellar microstructure, rough surface on the lamella and good bonding between the metal and ceramic phases which result in extensive crack deflection and delamination. Subsequently, Munch et al. [222,223] successfully fabricated compacted ceramic scaffolds with 80 vol.% ceramic phase by infiltrating, heat treatment and isostatic pressing. Such scaffolds exhibit brick-and-mortar architecture, large areas with sub-micrometre polymer films between alumina blocks and ceramic bridges between bricks, as shown in Fig. 36. According to the nacre structures, the large area of sub-micrometre soft polymer phase and the special ceramic bridges are especially important to toughen and strengthen the final composites, which can control frictional sliding between ceramic bricks and the roughness of the ceramic surfaces.

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Fig. 35. Mechanical response of natural and synthetic composites. The three-point bending load–displacement data (a) and SEM micrograph (b) of HAP/epoxy composites; the bending data and SEM micrograph of nacre (c); tortuous crack paths of Al2O3/epoxy composites (d); the final mechanical properties of Al2O3/Al–Si composites (e). Scale bars indicate (b) 40 lm, (c) 1 lm, and (d) 100 lm [218].

Fig. 36. (a) Brick-and-mortar architectures with large ceramic content up to 80 vol.% are prepared through pressing of the lamellar materials and subsequent sintering. (b) These structures display large areas with sub-micrometre polymer films between alumina blocks. (c) Ceramic bridges between bricks are formed during the second sintering steps [223].

The prepared Al2O3/PMMA (polymethyl methacrylate) composites exhibit similar bending stress– strain behaviours but much higher bending strength and inelastic deformation compared with natural nacre, as shown in Fig. 37. They display nearly 300 times higher exceptional toughness than either of the monolithic Al2O3 ceramic or PMMA polymer. Moreover, fracture toughness of the prepared Al2O3/ PMMA composites is increased to above 30 MPa m1/2 allied with a tensile strength of 200 MPa, which is much higher than 10 MPa m1/2 of the previous 45 vol.%Al2O3/Al–Si composites. The significant improvement in toughness is related to the extraction of Al2O3 ceramic bricks, frictional sliding between the ceramic bricks and polymer tearing and stretching over micrometre dimensions. Certainly, the stiffness and strength of the composites can be further enhanced by infiltrating metal alloy as soft phase. However, due to the limitations of the processing parameters and the raw materials, the fraction and thickness of the soft phase and the size of the hard phase are still much higher than those of natural nacre with nano scales. Finer architectures with higher ceramic content are expected to exhibit much higher improvement in the mechanical properties. In contrast to pursuing the above sub-micrometre soft phase and the high content of the hard phase, Bonderer et al. [224] designed and fabricated layered-hybrid composites just absorbing the layered principle of nacre. In the hybrid composites, just 20 vol.% sub-micrometre-thick ceramic platelets, with a thickness comparable to that of nacre, are strongly aligned and dispersed in the 80 vol.% ductile chitosan matrix. It is interesting that the prepared 20 vol.%Al2O3/chitosan composites also exhibit significant improvements in stiffness and strength compared with the monolithic chitosan

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Fig. 37. Mechanical response and toughening mechanisms in the synthetic hybrid composites. (a) Bending stress–strain curves for the Al2O3/PMMA hybrid materials mimic those of nacre and show >1% inelastic deformation before failure. (b) These materials show exceptional toughness for crack growth, similar to natural composites, and display significant rising R-curve behaviour [222].

material or calcified tendon, while retaining the plasticity of the chitosan. Even, the tensile strength and ductility are remarkably improved compared with nacre, bone and dentin, just accompanied with a slight decrease in stiffness. These improvements can be attributed to the layered sub-micrometre ceramic platelets and strong interface bonding, which result in extensive shear deformation of the matrix before platelet extraction. Inspired by smart natural structures, Kai et al. [225] selected flake Al powders as the starting material, which was obtained by flake PM, in order to successfully fabricate 32 vol.%B4C/Al composites. The test results show that the tensile strength and ductility of the prepared flake composites are increased by 63% and 13%, respectively, compared with those of the conventional composites. By performing tensile tests on the composites with nano-scale multilayers, the nano-layered Cu/Nb composite with a nominally equal layer thickness of 40 nm exhibits an UTS of 1.55 GPa and a fracture elongation of 3.4% [226]. The improvements in tensile properties can be attributed to the nano-layered structure with a polycrystalline structure. Weinkamer and Fratzl [212,227] presented profound understanding of the hierarchical structures of bone and wood from different perspectives, which can inspire scientists to construct more composite structures with different length scales for improvement in the mechanical properties and subsequently help in understanding the relationship between structure and property. For example, the aforementioned macro-ring structure, viewed as one bio-inspired structure of Al MMCs, exhibited significant improvement in mechanical properties [48,49]. In addition, nacre-like, bone-like, wood-like and bamboo-like structures have been extensively exploited in tailoring the polymer matrix composites and MMCs for properties enhancement [228,229]. 5. Microstructural design of TMCs It is well known that discontinuously reinforced TMCs (DRTMCs) possess high specific strength, specific stiffness and high-temperature durability, which render them as potential applications in the aerospace, defence and automotive industries, among others [230,231]. Before 2000s, DRTMCs had undergone great progression, including microstructure evolution and fabrication processing. It is undoubted that in situ TiB whiskers (TiBw) and TiC particles (TiCp) are the most effective reinforcements for the Ti matrix [230,232–234], and can lead to greater improvement in stiffness, strength and creep resistance. However, some serious problems encountered by DRTMCs emerged simultaneously. As early as 2003, Panda and Ravi Chandran [70] had pointed out two problems impeding the practical application of DRTMCs. One is how to increase the volume fraction of the reinforcement in order to

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further increase the mechanical properties. The other is that the prepared DRTMCs always exhibit limited ductility, even complete brittleness, particularly for DRTMCs fabricated by PM. 5.1. TiBw–Ti composite with isolated TiBw-rich (dual matrix) microstructure Based on TiBw–Ti composite system, Patel and Morsi [38,235] designed a so-called ‘dual matrix’ microstructure, as schematically shown in Fig. 2. The composite consists of an isolated TiBw-rich composite phase and a TiBw-free titanium matrix phase (Fig. 38a), which is a typical case of Pattern A shown in Fig. 1a. In addition, highly interconnected sub-micron TiB whiskers within the TiBw-rich phase (Fig. 38b) were formed due to its high local volume fraction. As shown in Fig. 39, the macro Vickers harness of the prepared TiBw/Ti composites increases from 652 to 777 with decreasing sizes of the TiBw-rich phases from 90–120 lm to <45 lm. In addition, they predicted that the toughness and wear resistance of these special composites can be improved compared with those of DRTMCs with a conventional homogeneous microstructure. As mentioned in the literature [38], Patel and Morsi ascribed the low ductility and brittleness of DRTMCs to the TiBw interconnectivity. Probably, this has influence on the ductility; however, it is not the most important factor. It can be envisaged that the inferior properties of these special composites are related to the relatively small distance, even partially connective, between the isolated TiBwrich phase and the sintering defects, such as pores. It is worth pointing out that this work offers the tunability of the properties of Ti–TiBw composites under the same volume fraction. Especially, the authors proposed that the properties of Ti–TiBw composites could be largely controlled by varying the volume fractions of the TiBw-free Ti-matrix phase. 5.2. 3D continuous microstructure with dense ceramic network In order to improve the ductility of DRTMCs fabricated by powder metallurgy, Panda and Ravi Chandran [70] fabricated 20 vol.%TiBw/b-21S composite by selecting b-Ti alloy as the matrix due to its superior ductility. The selection of large-size b-21S powders resulted in the formation of the reinforcement-lean phase, while high volume fraction of the reinforcement led to the continuous boundary phase involving synthesized TiBw phase and residual TiB2 phase around the matrix particles. Therefore, the continuous network microstructure of TiBw/b-21S composites was formed as shown in Fig. 40, leading to inferior mechanical properties. The reason for inferior properties is that the nucleation and propagation of cracks in the continuous ceramic-like boundary become very easy due to the very high local volume fraction, which also suppresses the superior ductility of the b-Ti matrix. The failure of this work lies in ignoring the effect of reinforcement distribution. However, compared with the previous work, the ductility of the composites was finally improved by using small Ti powders and

Fig. 38. SEM micrographs of dual matrix TiBw–Ti composites at (a) low and (b) high magnification [38].

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Fig. 39. Effect of reinforcement composite particle size and powder classification on the Vickers hardness number (VHN) of current-activated pressure assisted sintering (CAPAS) processed composites [38].

Fig. 40. SEM pictures of the 20 vol.%TiBw/b-21S composite in the as-processed condition [70].

introducing b stabilizing elements [236], but this slight improvement in ductility simultaneously sacrificed the strength and modulus. In order to clarify the effect of reinforcement distribution on the mechanical properties and microstructure of DRTMCs, one interesting continuous network microstructure of TiCp/Ti6Al4V composite was successfully designed and fabricated on the basis of Hashin–Shtrikman bound theory, whose upper bound implies a ‘soft’ phase is encapsulated by a hard phase. As shown in Fig. 41, the in situ synthesized TiCp reinforcement phase is distributed around the large Ti matrix particles and forms a 3D continuous network. In addition, equiaxed a(Ti) phase instead of the typical Widmanstätten lamellar microstructure was formed in the fabricated composite, which is beneficial to the mechanical properties of the composites. The reason for this phenomenon can be attributed to

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Fig. 41. Flow chart fabricating TiCp/Ti6Al4V composite with a network microstructure and SEM micrographs of Ti6Al4V powders (a), carbon powders (b), mixture (c), the monolithic Ti6Al4V alloy (d) and TiCp/Ti6Al4V composite with a network microstructure (e) [14].

additional isotropic tensile stress or elastic strain energy (ESE) within the composite matrix generated by the stiffer ceramic network shell [237,238]. This interpretation of the effect of the reinforcement on the matrix microstructure is more profound than that in the previous work [239–244]. In addition, the continuous TiCp ceramic network can effectively refine the matrix grain size and cause the formation of an equiaxed microstructure, which result in a significant improvement in the compressive strength. Moreover, it is worth pointing out that its elastic modulus is close to the H–S upper bound due to the continuous network microstructure being able to be viewed as the microstructure with the H–S upper bound, i.e., the stiffer phase encapsulates the discrete softer phase. However, it is certain that the tensile properties are inferior due to the continuous ceramic network microstructure, as the Fe3C network formed in high carbon cast iron. It is interesting that the TiC particle reinforced Ti6Al4V composites (TiCp/Ti6Al4V) exhibit characteristics of a honeycomb with a continuous TiC network (Fig. 42). The in situ reaction synthesized TiCp formed a ceramic wall with equiaxed particle morphology [71]. As shown in Fig. 43, the oxidation scale spallation and detachment of the monolithic Ti6Al4V alloy sample are severe at 973 K for 40 h. However, the oxidation surface of the TiCp/Ti6Al4V composites with a network microstructure is still smooth without any spallation and detachment after 140 h at 973 K. The test results show that the oxidation rate of the composites is much lower than that of the monolithic Ti6Al4V alloy at 873, 973 and 1073 K. Moreover, this superiority increases with increasing test temperatures. Additionally, the activation energy (334.23 kJ/mol) of the composites is obviously

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Fig. 42. SEM micrographs of 5 vol.% TiCp/Ti6Al4V composite having undergone serious etching with increasing magnifications from (a) to (d) [71].

Ti6Al4V alloy

TiCp/Ti6Al4V

Fig. 43. Macro images of Ti6Al4V alloy and TiCp/Ti6Al4V composite oxidation surfaces at 973 K for different time periods [71].

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higher than that of the monolithic alloy (300.12 kJ/mol) by calculation. The above three phenomena sufficiently demonstrate that the oxidation resistance of the TiCp/Ti6Al4V composites is significantly improved by tailoring the 3D network microstructure. As shown in the schematic image of Fig. 44, the main reason is that the oxidation scale of the network-structured TiCp/Ti6Al4V composites is divided into small units by the network microstructure. In addition, the titanium close to TiCp network is prior to being oxidized due to the high residual stresses formed during the composite fabrication compared with the titanium in the network centre [71]. The prior-formed oxides near the TiC network boundary can effectively pin and fasten the oxidation scale. The above two factors result in no spallation and detachment of oxidation scales, not only at 973 K but also at 873 K and 1073 K [71], which can effectively obstruct the oxidation progress to protect the titanium matrix. Although little oxygen crosses the fastened oxidation scale, the self-assembled TiCp wall can arrest its progress, which in effect is similar to continuous ceramic coatings reducing oxygen ingress at high temperatures [245–247]. Therefore, not only the mechanical properties but also the chemical properties of the TiCp/Ti6Al4V composites are improved by designing the interesting network reinforcement distribution. Besides the above improvements, by designing the second phase of the MoSi2 and SiC semi-conductive materials as a continuous network distribution around discrete Si3N4 particles, not only high-temperature bending strength but also both thermal and electrical conductivities were simultaneously improved compared with those of conventional MoSi2(SiC)/Si3N4 composites with a homogeneous microstructure [248]. 5.3. TMCs with 3D continuous reinforcement-rich network

Oxidation scale

5.3.1. Design and fabrication More recently, one quasi-continuous network microstructure of DRTMCs was successfully designed in order to overturn the situation of inferior mechanical properties by absorbing the above sufficient experience. As shown in Fig. 45, according to H–S theory, the network microstructure in which the harder phase encapsulates the inner softer phase corresponds to the high bound. In addition, network microstructure is equivalent to introducing ceramic reinforcement into the grain boundary. On the one hand, this microstructure can enhance the grain-boundary strengthening effect at room temperature. On the other hand, it can overcome the grain boundary weakening effect at high temperatures. Considering the failure to achieve superior tensile properties of the continuous ceramic network microstructure and the success of the bi-continuous microstructures of the Al MMCs, the local volume fraction of ceramic phase in the network boundary region should be controlled.

Fig. 44. Schematic illustration showing the oxidation resistant mechanism of TiCp/Ti6Al4V composite with a network microstructure [71].

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+

Grain boundary strengthening

+

Room

TiBw

Phase β

temperature ductility

Phase α Phase β

DRTMCs with a quasi-continuous network microstructure Fig. 45. Design route of a quasi-continuous network microstructure for DRTMCs.

Therefore, in situ TiB whiskers were chosen as the reinforcement, because they can decrease the local reinforcement volume fraction to achieve non-continuous distribution of the reinforcement by growing into the inner matrix, which can also effectively joint the adjacent Ti matrix particles as a dowel connector. Additionally, the in situ TiB whiskers are viewed as the best reinforcement for the Ti matrix due to their high modulus, hardness, good chemical compatibility with Ti, and similar density and thermal expansion coefficient with the Ti matrix. Therefore, one quasi-continuous network microstructure in which TiBw are uniformly distributed in the network boundary is designed for DRTMCs for performance improvement. In order to obtain the designed quasi-continuous network microstructure, large spherical Ti powders as the Ti matrix and fine TiB2 powders as the B source for the TiBw reinforcement were selected. As shown in Fig. 46, the fabrication route for TiBw/Ti composites with a network microstructure includes the simplified low-energy milling and reaction hot-pressing processes. The aim of low-energy milling is not to smash large Ti powders into fine powders but to adhere fine TiB2 powders onto the surface of large Ti powders/particles. In the present process, large Ti powders instead of fine Ti powders (5–20 lm) and low-energy milling instead of high-energy milling can not only guarantee the reinforcement network distribution, but also protect the Ti matrix from absorbing oxygen (O) and hydrogen (H) which can significantly increase the brittleness of the Ti matrix.

Fig. 46. Flow chart showing the processing route together with morphologies of the raw materials and schematic illustrations of the network distribution. (a) Large Ti powder, (b) fine TiB2 powder and the blended mixture at a lower (c) and higher (d) magnification; and schematic illustrations of network distribution (d) before and (e) after reaction synthesis.

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In order to keep the interpenetrating network boundary for a superior combination of mechanical properties, the volume fraction of the reinforcement must be designed and optimized. Fig. 47a shows that half surface of every Ti particle is covered by fine TiB2 powders. For this case, all TiB2 powders have an identical opportunity to react with the Ti particle during the reactive hot-pressing process, and then the reaction between the Ti and TiB2 can be rapidly completed. This can make space for the matrix to interpenetrate through the network boundary due to TiBw reinforcement growth into the inner matrix as whisker morphology. Fig. 47b shows another case where the whole surface of each Ti particle is covered by TiB2 powders. For this case, the outside of the TiB2 powders cannot easily touch the Ti and transform it to TiB whiskers. Moreover, this sintering system covered by the whole ceramic layer is close to the ceramic system which needs much higher conditions than metal or general composite systems to be compacted. Additionally, this corresponds to the maximum addition of TiB2 powders to obtain a homogeneous distribution in the network boundary, while the first case possibly corresponds to the optimal TiBw volume fraction. Below the maximum TiB2 addition, the contiguity of the matrix decreases while that of the reinforcement increases with the increasing TiB2 addition. By deducing and calculating, the volume fraction of the TiBw for the optimal case can be expressed by the following equation:

V TiB ¼

3:4ðR þ rÞ2  r  qa R  qb þ 2ðR þ rÞ2  r  qa 3

ð9Þ

where VTiB is the volume fraction of TiBw, r and qa are the radius and density of TiB2 powders, while R and qb are those of the Ti particles. Taking the 100 lm (R) and 4.52 g/cm3 (qb ) values of the Ti particles and the 1.5 lm (r) and 4.45 g/cm3 (qa ) values of the TiB2 powders for example, the ‘optimal’ VTiB can be calculated to be approximately 5.1 vol.%. Correspondingly, the optimal volume fraction of TiBw for the 55 lm (R) of the Ti particles is 8.5 vol.%. It is easy to consider that the ‘optimal’ volume fraction and the maximum increase with decreasing matrix particle sizes. As shown in Fig. 48, firstly, TiB reinforcement is in situ synthesized in the form of whisker, which partially grew into the inner matrix. Secondly, TiBw reinforcement is distributed around the Ti particles and then forms a network microstructure, which is totally different from the conventional homogeneous distribution. Thirdly, the Ti matrix is interpenetrated through the network boundary and the TiBw reinforcement is still interpenetrated along the network boundary. Moreover, the contiguity of the Ti matrix decreases while that of the TiBw reinforcement increases with the increase in the overall TiBw volume fractions. Fourth, the network size is approximately equal to the size of the as-received Ti powders of 110 lm. Therefore, we successfully fabricated the designed TiBw/Ti composites with a quasi-continuous network microstructure. This can be mainly attributed to the following three factors: selection of raw materials with large size differences between the large spherical Ti powders and the fine TiB2 powders, the low-energy milling process protecting the large Ti powders from smashing into fine powders and solid state sintering ensuring the reaction between the Ti and the TiB2 just on the surface of the large Ti powders. Fig. 49 exhibits the stress–strain curves of all the prepared composites and the monolithic pure Ti. The milestone phenomenon is that not only the strength but also the ductility of the composites can be significantly improved by tailoring the network microstructure compared with those of the

Fig. 47. Schematic illustrations of fine TiB2 powders on the surface of large Ti6Al4V particles. (a) A half surface of Ti6Al4V covered by TiB2 and (b) the whole surface of Ti6Al4V covered by TiB2 [249].

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Fig. 48. SEM micrographs of (a) 5 vol.%, (b) 8.5 vol.% and (c) 12 vol.% TiBw/Ti composites with a network microstructure at a low magnification [58].

Fig. 49. The tensile stress–strain curves of the as-sintered TiBw/Ti composites with a network microstructure and a homogeneous microstructure and the pure Ti [58].

composites with a homogeneous microstructure. In particular, the tensile fracture elongation of 15.6% and 11.8% for the 5 vol.% and 8.5 vol.% composites can be viewed as the highest improvements to date for the as-sintered TMCs fabricated by PM. The UTS of 12 vol.% TiBw/Ti composites reaches up to 907 MPa, increased by 88%, compared with that of the pure Ti. This can also be viewed as the most effective strengthening effect for DRTMCs fabricated by PM. According to the microstructure observation, the enhanced strengthening effect can be attributed to the increased contiguity of TiBw reinforcement due to concentration on the network boundary. The increased ductility can be attributed to the interpenetrating matrix and the large TiBw-lean phase remaining, which can bear the strain, blunt the crack and decrease the crack-propagation speed. Therefore, the designed quasi-continuous

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microstructure can effectively exploit a superior combination of mechanical properties over the conventional homogeneous microstructure. 5.3.2. Microstructural characteristics After successful improvement of the mechanical properties of TiBw/Ti composite, Ti6Al4V alloy powders with average sizes of 200 lm, 110 lm and 65 lm were chosen as the matrix, due to their good mechanical properties and their wide applications in industry, in order to further improve the mechanical properties. Moreover, Ti60 alloy possessing the highest service temperature of 600 °C, which is similar to Ti-1100 (USA), IMI834 (UK) and BT36 (Russian) was used as the matrix for higher high-temperature properties, because one of the ultimate aims of DRTMCs is to enhance their service temperature and high-temperature durability on the basis of titanium alloy [250]. Fig. 50 shows the SEM micrographs of 5 vol.%TiBw/Ti6Al4V (200 lm) composites with a network microstructure at increasing magnifications and reveals different unique features. First, TiB whiskers are in situ synthesized and homogeneously distributed around the Ti6Al4V particles, and then form a network microstructure. This is similar to the primary b grain boundary structure, that is to say, we achieve introducing TiBw ceramic reinforcement into the grain boundary. Second, the ‘grain size’ is constrained into 200 lm by the ceramic network microstructure, which is simply equal to the size of the as-received Ti6Al4V particles, and the size can be further decreased by using different Ti6Al4V powders with smaller sizes such as 110 lm and 65 lm. Compared with the grain size (900 lm) of the monolithic Ti6Al4V alloy fabricated by the same raw material and sintering parameters [182], the grain size can be viewed as effectively refined. Third, due to the defined length of the TiB whiskers, the width of the network boundary is limited in a finite width (about 30 lm double TiBw length). Therefore, the overall network unit can be divided into a TiBw-rich network boundary phase and a TiBw-lean phase. That is to say, the present network microstructure can be viewed as the well-known H–S upper bound [80,83], which can exploit a superior strengthening effect. Fourth, as shown in Fig. 50b and c, TiB whiskers grew into the neighbouring Ti6Al4V particles like dowel connectors (owning to the special B27 structure [251], resulting in a strong and gradient boundary) connecting the neighbouring Ti particles, which can be viewed as a strengthened grain boundary. Fifth, as shown

Fig. 50. SEM micrographs of TiBw/Ti6Al4V composite with a network microstructure at different magnifications revealing the network and TiB whisker morphologies. (a) Network structure, (b) dowel structure and self-joining structure, (c) multibranched structure, and (d) mechanical locking, self-jointing structure and claw structure [182].

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in Fig. 50c and d, TiB whiskers with branched structures including multi-hierarchical branches and a claw-like structure are discovered for the first time. The branched structures are inferred to be formed from the reaction with polycrystalline TiB2 particles retained during the low-energy milling process [252,253]. In particular, it is extremely interesting that a single TiB whisker joins with a vertical one forming a ‘T’ junction and mechanically locks with another vertical one in a different plane, so constructing a 3D spatial TiBw structure which is beneficial for load transfer. The effective TiBw length is also obviously increased by the self-joining structure, mechanical locking structure and branched structure, which can increase the strengthening effect by increasing the reinforcement contiguity. Fig. 51 directly shows that the in-situ synthesized TiB whiskers are homogeneously distributed around Ti matrix particles and then form a 3D network microstructure. The two-scale hierarchical 3D structures, i.e. the overall 3D network and the local 3D spatial arrangement of TiBw in the network boundary (Fig. 50d), are believed to be beneficial in exploiting the strengthening effect of the TiBw reinforcement. As shown in Fig. 52, instead of the typical Widmanstätten microstructure formed in the monolithic Ti6Al4V alloy and the Ti60 alloy fabricated by the same sintering parameters [172,255], the quasi-equiaxed microstructure, including the equiaxed and platelet a phase, is formed in the matrix of the composites with a network microstructure. All the previous literature attributed the formation of the equiaxed microstructure to the reinforcement that can hinder the laminated Widmanstätten growth and provide nucleation sites to the equiaxed a phase [14]. But, it is clear that the equiaxed microstructure also formed in the large matrix region without any reinforcement, the formation of equiaxed structure is likely due to the additional isotropic tensile stress or elastic strain energy within

(b) 1

(a)

1 2

2

3

3 500μm

Fig. 51. Schematic view of the sample from the as-sintered TiBw/Ti60 composites. (a) The shape of the sample and (b) SEM micrograph of the sharp corner taken from the sample [254].

Fig. 52. The optical microscope (OM) (a) and SEM (b) [172] micrographs of 5 vol.%TiBw/Ti6Al4V composites to reveal the quasiequiaxed microstructure of Ti6Al4V matrix.

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the composite matrix generated by the strong network constraint effect when cooling from 1200 °C drags the matrix to grow into an equiaxed microstructure. The residual stress can be verified by the serious stress corrosion in the matrix near the TiBw. As shown in Fig. 53a, due to the large gaps between the loose large spherical Ti powders and the local non-uniform network distribution of the reinforcement, it is more difficult to sinter compacted composites with a network microstructure than conventional composites with a homogeneous microstructure [252,256]. Even under the optimal sintering temperature, the designed 10.2 vol.%TiBw/Ti6Al4V (200 lm) composites with the maximum volume fraction cannot be sintered to compacted composites as shown in Fig. 53b. As mentioned above, the present sintering system with the whole TiB2 ceramic layer is similar to the ceramic system which needs higher conditions than those for metal or general composite systems. Additionally, the deformation resistance is too high to compact the composite with a much higher volume fraction of reinforcement. Therefore, a modest volume fraction of reinforcement is critical to obtain the compacted composites with an interpenetrating network microstructure. Furthermore, as shown in Fig. 54a, the network boundary becomes nearly continuous and the reaction between Ti and TiB2 seems to be incomplete for the 8.5 vol.%TiBw/Ti6Al4V (200 lm) composite system, because the volume fraction of 8.5 vol.% is much higher than the optimal one of 5.1 vol.%. However, in the 12 vol.%TiBw/Ti6Al4V (65 lm) composite system, the network boundary is still interpenetrated and the reaction is complete (Fig. 54b). This phenomenon effectively demonstrates that the optimal and maximum volume fractions can be increased by decreasing the matrix particle sizes. It is also worth pointing out that the TiBw-lean phase size decreases with the decrease in the matrix particle size, which must be harmful to the ductility of the composites with a network microstructure. 5.3.3. Mechanical characteristics Fig. 55 shows the typical tensile stress–strain curves of the as-sintered composites [68]. Moreover, Table 5 summarizes the mechanical properties of various TiBw/Ti6Al4V composites with different

Fig. 53. The unmatured network microstructure of (a) 5 vol.%TiBw/Ti6Al4V composites fabricated at 1000 °C [252] and (b) 12 vol.%TiBw/Ti6Al4V composites fabricated at 1200 °C [249].

Fig. 54. SEM micrographs of (a) 8.5 vol.%TiBw/Ti6Al4V (200 lm) composites and (b) 12 vol.%TiBw/Ti6Al4V (65 lm) composites [68].

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Fig. 55. Tensile stress–strain curves of the monolithic Ti6Al4V alloy and various TiBw/Ti6Al4V composites fabricated using the same parameters (the digits in the parentheses represent the average sizes of raw Ti6Al4V powders) [68].

Table 5 Room temperature tensile properties of various TiBw/Ti6Al4V composites with different volume fractions of TiBw reinforcement and sizes of the Ti6Al4V particle [68]. Samples

Detail information of samples

VL (%)

UTS (MPa)

Elongation (%)

Elastic modulus (GPa)

Ti6Al4V V2D200 V3D200 V5D200 V8D200 V5D110 V8D110 V8D65 V12D65

Monolithic Ti6Al4V alloy 2 vol.%TiBw/Ti6Al4V (200 lm) 3.5 vol.%TiBw/Ti6Al4V (200 lm) 5 vol.%TiBw/Ti6Al4V (200 lm) 8.5 vol.%TiBw/Ti6Al4V (200 lm) 5 vol.%TiBw/Ti6Al4V (110 lm) 8.5 vol.%TiBw/Ti6Al4V (110 lm) 8.5 vol.%TiBw/Ti6Al4V (65 lm) 12 vol.%TiBw/Ti6Al4V (65 lm)

0 7.4 12.9 18.1 40.8 11.1 18.8 12.7 18.0

855 ± 7 1021 ± 5 1035 ± 5 1090 ± 10 997 ± 5 1060 ± 6 1288 ± 5 1207 ± 7 1108 ± 5

11.3 ± 1.1 9.2 ± 0.5 6.5 ± 0.5 3.6 ± 0.2 1.0 ± 0.1 5.1 ± 0.4 2.6 ± 0.3 4.6 ± 0.2 0.9 ± 0.1

112.32 ± 0.3 116.09 ± 0.3 120.79 ± 0.3 122.87 ± 0.3 131.05 ± 0.2 121.51 ± 0.2 129.58 ± 0.3 127.23 ± 0.3 136.12 ± 0.3

Note: the digits in the parentheses represent the average sizes of raw Ti6Al4V powders. VL stands for the local volume fraction of TiBw reinforcement in the network boundary, which is calculated by equation deduced in reference of [68].

network sizes and different volume fractions in order to further reveal the contribution of the network distribution and network parameters. First, the YS and the UTS of the composites with as low as 2 vol.%TiBw reinforcement increase from 700 and 855 MPa to 846 and 1021 MPa, respectively, which corresponds to increases of 20.9% and 19.4% relative to those of the monolithic Ti6Al4V alloy. Moreover, the tensile strength can be further increased to 1288 MPa, which corresponds to an increase of 50.6% by changing the Ti6Al4V particle size and TiBw volume fraction. Second, the tensile elongation of the composites can easily remain at a level higher than 3% even up to 9%. Third, the tensile strength increases while the ductility decreases with increasing the overall volume fractions under a given Ti6Al4V particle size. The increasing strength can be attributed to the increasing contiguity of TiBw reinforcement. Fourth, the ductility increases while the tensile strength decreases with the decrease in Ti6Al4V particle size under a similar overall volume fraction of TiBw reinforcement. The increasing ductility can be attributed to the increasing contiguity of the Ti matrix (the space/distance between individual reinforcements). Fifth, the ductility of the composites decreases with the decrease in the Ti6Al4V particle size under a given local volume fraction, which can be attributed to the decreased TiBw-lean phase size or termed as the space/distance between the continuous network boundaries. Sixth, the elastic modulus generally increases with the increase in the local volume fractions of the TiBw reinforcement due to the increasing contiguity of the reinforcement. Moreover, the elastic modulus of the prepared TiBw/Ti6Al4V composites with a network microstructure is close to the theoretical H–S upper bound.

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It is worth noting that the superior strengthening effects and toughening effect belong to the assintered composites without any secondary deformation. Furthermore, both the strength and the ductility of the network structured TiBw/Ti6Al4V composites fabricated by PM can be further increased by secondary deformation [257]. Additionally, not only the strengthening effect but also the toughening effect of TiBw/Ti6Al4V composites with a network microstructure can be adjusted by controlling the network parameters. The two main factors for the superior ductility of DRTMCs are the larger TiBwlean phase and the larger space/distance between the individual reinforcements. The latter can stop the micro-cracks generated in a single reinforcement from coalescing, while the former can blunt the crack, bear the strain before fracture and hence slow down the propagation of the main crack. The superior strength and elastic modulus of DRTMCs should require higher contiguity of the reinforcement and modest ductility. In a sense, the critical problem of extreme brittleness for the titanium alloy matrix composites fabricated by PM is resolved and their strengthening effect and toughening effect are further enhanced by tailoring the network distribution of the reinforcement. Interestingly, the composites with a network microstructure also exhibit a remarkable improvement in high-temperature strength, as shown in Fig. 56 [182]. Taking V5D200 for example, the tensile strength at 500 °C, 600 °C, 700 °C is increased by 36.9%, 36.1% and 27.3%, respectively, compared with that of the monolithic Ti6Al4V alloy. In particular, the tensile strength of V12D65 at 600 °C is almost equal to that of the monolithic Ti6Al4V alloy at 400 °C. That is to say, the service temperature of the composites can be increased by 200 °C compared with that of the monolithic Ti6Al4V alloy on the basis of the same tensile strength. Considering the modest volume fractions of reinforcement, the present improvement in the high-temperature strengthening effect can be viewed to be significant, which can be mainly attributed to (i) the network architecture, (ii) the 3D branched TiBw structure and (iii) a refined primary b grain with strengthened grain boundary (network boundary). In other words, the network distribution of the reinforcement can be viewed to introduce TiBw reinforcement into the grain boundary, which can overcome the weakening effect of the grain boundary at high temperatures. In addition, according to Tamirisakandala’s work [258,259], it is likely that the increased b transus due to the reinforcement introduction accelerates the improvement in the high-temperature strength. In order to further enhance the high-temperature strength and service temperature, Ti60 alloy (China) was selected as the matrix of the composites with a network microstructure. As shown in Fig. 57, the tensile strength of the as-sintered 8 vol.%TiBw/Ti60 composites is increased to 889 MPa, 721 MPa and 453 MPa from 554 MPa, 452 MPa and 303 MPa of the monolithic Ti60 alloy at 600 °C, 700 °C and 800 °C, respectively. That is to say, the tensile strength is increased by 61.1%, 57.4% and 45.5% compared with that of the monolithic Ti60 alloy [255]. Compared with the tensile strength of

Fig. 56. Comparison of high-temperature tensile strength between the TiBw/Ti6Al4V composites with a network microstructure and the monolithic Ti6Al4V alloy [182].

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147

Fig. 57. Tensile properties of the monolithic Ti60 alloy and the network-structured 5 vol.%, 8 vol.% and 12 vol.%TiBw/Ti60 composites at (a) 600 °C, (b) 700 °C and (c) 800 °C [255].

the 625 MPa and 342 MPa at 600 °C and 700 °C of 10 vol.%TiCp/TA15 composites [260,261], the network structured TiBw/Ti60 composites exhibit an obvious improvement in high-temperature tensile strength. Moreover, the strength of 721 MPa at 700 °C of the present 8 vol.%TiBw/Ti60 composites is also higher than that of 639 MPa at 650 °C of 8 vol.% (TiBw + TiCp)/Ti6242 composites [234]. Even, the present strength of the TiBw/Ti60 composites is higher than that of the TiAl intermetallic materials [262]. Moreover, the network structured TiBw/Ti60 composites possess superior formability to the TiAl intermetallic materials. Therefore, TiBw/Ti60 composites with a network microstructure have extensive application potential for high-temperature components. 5.3.4. Evolution of deformation and heat treatment It is certain that the mechanical properties of the composites with a network microstructure can be further increased by the subsequent hot deformation and heat treatment. For 5 vol.%TiBw/ Ti6Al4V(2 0 0) composites, the tensile strength can be increased to 1423 MPa from 1090 MPa by heat treatment of 870 °C/40 min/WQ + 500 °C/6 h/AC. The high-temperature tensile strength at 400 °C, 500 °C and 600 °C can be increased to 985 MPa, 848 MPa and 567 MPa by the heat treatment of 930 °C/40 min/WQ + 500 °C/6 h/AC [263]. The strength improvement is mainly attributed to the increased fraction of the transformed b phase including the martensite phase in the TiBw-lean phase, while the TiBw reinforcement has no visible change, as shown in Fig. 58. Additionally, the strength of the composites increases with the increase in the quenching temperatures due to the increase in the fraction of the transformed b phase. After extrusion deformation, the 3D equiaxed network units are extruded to the 1D column and the TiBw reinforcement is aligned along the extrusion direction, as shown in Fig. 59. Additionally, the local volume fraction of the reinforcement on the network boundary decreases due to the increasing special surface. Moreover, the whole matrix microstructure is changed to transformed b phase due to the deformation temperature of 1100 °C, higher than the b transus temperature. Finally, for one 5 vol.%TiBw/Ti6Al4V (110 lm) composite with a network microstructure, the tensile strength is increased

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Fig. 58. SEM micrographs of 5 vol.%TiBw/Ti6Al4V composite having undergone WQ at (a) 840 °C and (b) 870 °C and ageing at 500 °C for 6 h. (Inset image showing fine a phase in the transformed b phase.) [263].

(a)

(b)

Martensite

Residual stress etching

(c)

(d)

Fig. 59. SEM micrographs of the longitudinal (a and b) and cross (c and d) sections of the as-extruded 5 vol.%TiBw/Ti6Al4V composite at different magnifications; (a and c) at low magnifications, (c and d) at high magnifications [257].

to 1206 MPa, while the tensile elongation is increased to 12%, as shown in Fig. 60 [264]. The superior combination of mechanical properties can be attributed to the deformed 1D column microstructure with the aligned distribution of TiBw, decreased local volume fraction of TiBw reinforcement and the transformed b phase of the matrix. Moreover, the high-temperature tensile properties of the asextruded composites are also higher than those of the as-sintered composites [257]. As expected, the tensile strength of the as-extruded composites can be further increased by subsequent heat treatment [257]. For example, the tensile strength of the as-extruded 5 vol.%TiBw/Ti6Al4V composites can be further increased from 1206 MPa to 1364 MPa allied with 7.8% of tensile elongation after heat treatment [264]. For the 8 vol.%TiBw/Ti6Al4V composites, the tensile strength can be enhanced from 1311 MPa to 1470 MPa close to 1500 MPa. This is a significant improvement for Ti6Al4V matrix composites which can still be attributed to the tailored network microstructure composed of the TiBw-lean phase and TiBw-rich phase. Moreover, the tensile strength of TiBw/Ti60 composites can be further increased to nearly 1000 MPa and 800 MPa at 600 °C and 700 °C, respectively,

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149

Fig. 60. The tensile stress–strain curves of the as-extruded TiBw/Ti6Al4V composites before and after heat treatment [264].

by subsequent hot deformation and heat treatment according to previous experience [254,263,265] and the latest results. After rolling deformation, the 3D equiaxed network microstructure is changed to one pancake-like microstructure due to its special deformation characteristics, as shown in Fig. 61 [266,267]. The network size is enlarged along the rolling surface, while the thickness of the pancake-like microstructure decreases with the increase in rolling reductions. It is certain that the broken TiBw reinforcement increases while the local volume fraction of the reinforcement decreases with increasing rolling reductions. On the one hand, the ductility of the composites will increase due to the decrease of local volume fraction of reinforcement. On the other hand, the damage of TiBw reinforcement can result in a degradation of properties [33]. Additionally, the TiBw reinforcement is distributed not along the rolling direction but the 2D rolling surface. Therefore, the strengthening effect is slightly lower than that of hot extrusion deformation. Finally, the tensile strength of the network-structured 5 vol.%TiBw/ Ti6Al4V composites can be increased from 1090 MPa to 1330 MPa along with an elongation increase of 97%. The improvement can be mainly attributed to the deformation and heat treatment strengthening effects of the matrix. It is worth pointing out that the compressed samples with a network microstructure reveal the real deformation characteristics of hot compressive deformation [268]. The reason is that the smart network microstructure is equivalent to a marking sign at the micro local region. Therefore, the deformation characteristics can be obviously deduced by tracing the network microstructure evolution. A

Fig. 61. SEM micrographs of the as-rolling 5 vol.%TiBw/Ti6Al4V (200 lm) composites along (a) rolling surface and (b) profile surface (inset image showing schematic illustration of microstructure evolution) [266].

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similar phenomenon can be found in the superplastic tensile samples [269]. Additionally, due to the convenient condition of the inhomogeneous network microstructure, it is obviously revealed that the existence of ceramic reinforcement is beneficial to the formation of dynamic recrystallization due to the higher dislocation density near the reinforcement [268,269]. 5.3.5. Fracture mechanism and models In order to better comprehend the superior strengthening and toughening mechanisms of the network microstructure, the special fracture characteristics led by the unique network microstructure were presented. The 5 vol.%TiBw/Ti6Al4V (200 lm) composites were taken as the example. It is obvious that the main crack propagates along the network boundary, as shown in Fig. 62a. That is to say, the fracture occurred through the TiBw-rich boundary phase, which is analogous to an intercrystalline fracture. The main reason is that the large TiBw-lean region can effectively blunt the crack stemming from the TiBw-rich network boundary. This can encourage all the TiBw reinforcement to bear the stress and strengthen the composites with a modest local volume fraction of reinforcement. This certainly enhances the strengthening effect on the basis of the grain boundary strengthening effect by introducing ceramic reinforcement into the ‘grain boundary’, which is consistent with the superior strengthening effect even for the network-structured composites with a low 2 vol.% TiBw reinforcement. As shown in Fig. 62b, due to the large 3D network structure and the ‘intercrystalline fracture’, the fracture surface of the composites with a network microstructure exhibits a rather rough characteristic indicating torturous crack propagation, which can more effectively toughen the composites than plain crack propagation. The above two phenomena effectively demonstrate that the continuous stronger network boundary dominates the mechanical behaviour of the network-structured composites. As shown in Fig. 62c, TiBw reinforcement is fractured accompanied by matrix dimples and tearing ridge lines. The fracture of TiBw is similar to that in the conventional homogenous composites: the micro-crack originates from reinforcement cracking since the in situ composites have strong interface between reinforcement and matrix. This means that the strengthening effect of the TiBw has been fully utilized owing to the strong interfacial bonding. The dimples and tearing ridge lines are beneficial to

TiBw Ti6Al4V

Fig. 62. The secondary electron micrographs of crack propagation and fracture surfaces of the 5 vol.%TiBw/Ti6Al4V composite. (a) Crack propagation at a low magnification. (b) Mirror images of fracture surfaces at a low magnification. (c) Fracture surface at a higher magnification [68].

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the strengthening effect of the TiBw reinforcement and the toughening effect of the Ti matrix, which verifies the necessity for the interpenetrating matrix structure. It is different in that the composites with higher volume fractures exhibit TiBw fractures without matrix dimples and tearing ride lines similar to ceramic fractures [249], which cannot explore the strengthening effect of the TiBw reinforcement and the toughening effect of the matrix. Therefore, the composites with higher local volume fractions exhibit inferior tensile elongations. Fig. 63 shows the magnified images of region ‘A’ marked in Fig. 62a, presenting the detailed fracture characteristics of the tensile sample. A number of micro voids were found in the boundary phase away from the main crack, which indicates that crack propagation progressed via micro-void coalescence. However, the micro voids did not quickly grow and propagate but were blunted by the TiBw-lean phase, which is beneficial to further extending the strengthening and toughening effects of the network microstructure. Crack branching and secondary crack propagation near the main crack path are also popular. These can absorb more fracture energy and slow the crack-propagation speed [270], which effectively exploits the strengthening and toughening effects of the network microstructure combined with the formation of micro voids. In addition, at the beginning of tensile deformation, the plastic deformation of the TiBw-lean phase can be effectively obstructed by the stronger TiBw-rich network boundary, which can correspondingly enhance the yield strength and elastic modulus of the network-structured composites. In the TiBw-lean matrix phase, extensive slip bands, representing plastic deformation, are clearly observed, as shown in Fig. 63b, which demonstrate the high strain-bearing ability of the TiBw-lean phase [271]. As shown in Fig. 63c, many micro-cracks formed by TiBw breakages are also observed far away from the main crack. In general, reinforcement particle fracture prior to the ductile matrix would result in a degradation of mechanical properties. However, the micro-cracks were blunted immediately preventing further propagation; this encourages the fractured TiBw segments to further bear stress and strengthen the composites, eventually resulting in the multiple fractures of dowel-like TiBw as shown in Fig. 63c and d. In addition, the small cracks perpendicular to the network boundary are immediately blunted by the soft TiBw-lean phase, as shown in Fig. 63c. These phenomena indicate that the network microstructure composed of the large isolated TiBw-lean phase and the interpenetrating network boundary can efficiently exploit the strengthening effect of the TiBw reinforcement and the toughening effect of the Ti matrix. Fig. 63d also shows micro-crack branches, multiple fractures, crack kinks and micro-cracks, which can be attributed to the 3D distribution, 3D branch structures and the dowel-like structure of the TiBw reinforcement. Therefore, the special fracture characteristics are consistent with the superior combination of the mechanical properties of the composites with the network microstructure [270]. Fig. 64 shows the typical fractographs of the 12 vol.%TiBw/Ti6Al4V (65) composites possessing the highest strength at high temperatures. It is interesting that the main crack still propagated along the network boundary and then formed an intercrystalline fracture accompanying the TiBw fracture. That is to say, the strengthening effect of the TiBw reinforcement is utilized at high temperatures due to the strong in situ bonding between the TiBw reinforcement and the Ti matrix. However, with increasing test temperatures, the strength of the TiBw-lean phase decreases, which decreases the necessary energy (i.e., the resistance to crack propagation) for the transcrystalline fracture. It is possible that the energy of the transcrystalline fracture is lower than that of the intercrystalline fracture at boundary junction regions at high temperatures. Therefore, the fraction of the transcrystalline fracture increases with the increasing test temperatures, as shown in Fig. 64, which certainly decreases the strengthening effect of the composites. This is consistent with the sharp decrease in strength and the sharp increase in ductility at 700 °C for the TiBw/Ti6Al4V composites [182].

5.3.6. Strengthening and toughening mechanisms From a macro view, the present network microstructure can be treated as the stiffer TiBw-rich phase encapsulating the softer TiBw-lean phase, which corresponds to the H–S upper bound as shown in Fig. 5. Therefore, the network microstructure can exploit a higher strengthening effect than the conventional homogeneous microstructure, which can be attributed to the higher contiguity of reinforcement-rich phase. Moreover, the whisker morphology, 3D spatial branched structure, mechanical

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Fig. 63. The secondary electron micrographs of fracture and crack propagation of the 5 vol.%TiBw/Ti6Al4V composite. (a) Magnified micrograph of the boxed A region in Fig. 62a. (b) High magnification micrograph of the boxed B region in (a) showing plastic deformation or slip bands in the TiBw-lean phase. (c) High magnification micrograph of the boxed C region in (a) showing TiBw cracking and crack blunting. (d) Micrograph of the boxed D region in (a) showing the detailed secondary crackpropagation path [68].

locking structure and self-joining structure of the TiB reinforcement can further enhance the contiguity of the TiBw reinforcement and then enhance the strengthening effect. For the as-extruded TiBw/Ti6Al4V composites, the adjustable Halpin and Tsai (H–T) theorem is used to effectively estimate their elastic properties. The expressions are listed as follows [60,243]:

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Fig. 64. High-temperature tensile fractographs of 12 vol.%TiBw/Ti6Al4V composite with a network microstructure at different temperatures: (a) 500 °C, (b) 600 °C, and (c) 700 °C [182].

 1 þ 2 dl gL V a EC;L ¼ Eb 1  gL V a E E gL ¼ a;L  l b Ea;L þ 2 d Eb 1 þ 2gT V a Eb EC;T ¼ 1  gT V a E  Eb gT ¼ a;T Ea;T þ 2Eb

ð10Þ ð11Þ ð12Þ ð13Þ

where EC,L and EC,T are the elastic modulus of the 1D composites along the longitudinal and transversal direction, respectively; Ea,L and Ea,T are the elastic modulus of the whisker along the longitudinal and transversal direction; l and d are the length and diameter of the whiskers. As shown in Fig. 65, the elastic modulus of the as-sintered composites with the network microstructure is close to the H–S upper bound due to the network microstructure. After hot extrusion, the as-extruded TiBw/Ti6Al4V composites exhibit an elastic modulus very close to the H–T upper bound along the extrusion direction, which is much higher than that of the as-sintered composites due to the alignment distribution of the TiBw reinforcement. It is certain that the as-extruded composites become anisotropic, and the elastic modulus along the transversal direction is lower than that of the as-sintered isotropic composites. From a micro view, in the network boundary, TiBw reinforcement bears higher stress than the local matrix around the TiBw due to the stress concentration generated by the dislocation pile up or prior bearing stress due to the high contiguity, as shown in Fig. 66a, when loading to the TiBw/Ti6Al4V composites. From a meso view, the stronger TiBw-rich network boundary preferentially bears higher stress than the softer TiBw-lean phase due to the continuous network microstructure with higher local volume fractions. Therefore, the TiBw whiskers and the network boundary fractured before the matrix around the TiBw and the TiBw-lean phase, respectively. The micro-crack of the TiBw and the small crack of the network boundary can be easily blunted by the matrix around the TiBw and the TiBw-lean phase, respectively, when the local volume fraction is lower than the optimal and the TiBw-lean phase

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Fig. 65. The H–S upper and lower, the H–T upper and lower bounds as a function of volume fraction and the elastic modulus comparisons of the as-sintered and as-extruded 5 vol.%TiBw/Ti6Al4V composites [257].

Fig. 66. Schematic illustrations of the stress distribution (a) before and (b) after the origin micro-crack and (c) the stress distribution at the crack tip; (d) the distance of the adjacent reinforcements, D: the size of the plastic deformation at the crack tip) [257].

size is greater than the optimal, i.e. the distance (d) of the adjacent reinforcements (TiB whiskers or network boundaries) is larger than twice the size (D) of the plastic deformation region at the crack tip (d > 2D). The stress at the front of the crack tip can be expressed as follows:

K

I ffi ryy ¼ pffiffiffiffiffiffiffiffi 2p r

ð14Þ

where KI is the fracture toughness of the reinforcement, r is the distance to the crack tip in the matrix. Taking the yield strength (rys) of the matrix into Eq. (14) replacing ryy, the largest distance (D) of the plastic deformation can be expressed as follows:

D¼

1 2p



KI

rys

2 ð15Þ

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The above two cases can effectively exploit the strengthening effect of the reinforcement and the toughening effect of the matrix, and result in a superior combination of mechanical properties. However, if the local volume fraction is higher than the optimal, or the TiBw-lean size is smaller than the optimal, the micro-crack or crack can quickly propagate under a low external load due to the high local stress at the crack tip, which would remarkably decrease the strengthening effect and toughening effect of the composites, even resulting in a brittle fracture. Additionally, the blunting crack allows the already fractured TiBw segments or the fractured network boundary to continue to bear load until further fracturing occurs, which leads to multiple fractures of the same TiB whisker and micro-void coalescence fractures and can probably prevent the properties from being degraded. These are consistent with a superior strengthening effect of the reinforcement and a superior toughening effect of the matrix. This model also suits other inhomogeneous microstructures including the laminated microstructure and bi-continuous microstructure. As in the schematic illustration in Fig. 67a, a crack perpendicular to the network boundary cannot propagate through the softer matrix phase because of crack tip blunting (crack A), unless the TiBwlean phase is very small. Therefore, the crack can only propagate by micro-void coalescence under increased load (cracks C and D to crack B), which are beneficial to extending their strengthening and toughening effects. The propagation routes (path I or path II) of crack B depend on their absorbing energy themselves. According to Eq. (16) of the effective reinforcement strength [272], the absorbing energy equation can be deduced as Eq. (17):

sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi pEI GIII rF ¼ 2ð1  m2I ÞdI sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 3 pEI GIII pEI GIII 2  d ¼ AL  ðdI Þ2 Q ¼ rF  S  L ¼ AL 2ð1  m2I Þ  dI I 2ð1  m2I Þ

ð16Þ ð17Þ

where rF is the effective reinforcement strength, E is elastic modulus, G is the critical strain energy release rate for the dynamic propagation of a crack into the softer TiBw-lean phase, m is Poisson’s ratio, d is the reinforcement size along the crack direction, S is the crack surface area, L is the open crack length and A is a the parameter related to surface and length. Therefore, the absorbing energy of crack propagation along paths I (QI) and II (QII) can be expressed as Eqs. (18) and (19), respectively:

sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 3 pEI GII  ð3aÞ2 2ð1  m2I Þ

Q I ¼ AL  B

ð18Þ

where B is the coefficient of the crack kinking difficulty.

sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 3 pEII GIIII  ð2aÞ2 2ð1  m2II Þ

Q II ¼ AL

ð19Þ

Fig. 67. Schematic illustrations of the composites with a continuous stronger phase (a) and a continuous flexible phase (b).

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Therefore, the absorbing energy of crack propagation mainly depends on the elastic modulus E, critical strain energy release rate G, Poisson ratio m and the particle size. Moreover, the difference between GI and GII sharply increases with increasing local volume fractions. Therefore, when the local volume fraction is very low, due to similar E and m, it is possible that QII is lower than QI at the local region with high B, which can lead to partial transcrystalline fractures. This is consistent with the local transcrystalline fractures of the 1.7 vol.%TiBw/Ti6Al4V composites [249]. Additionally, EII and mII swiftly decrease with increasing temperatures. It is possible that QII is lower than QI at very high temperatures, such as 700 °C [182]. This also results in transcrystalline fractures in the local region with high B. In general, QII is always higher than QI for composites with a network microstructure due to large differences between GI and GII, which results in crack propagation along the network boundary (path I). If the crack B propagates along path II, the transcrystalline fracture will need much more energy due to the plastic deformation of the TiBw-lean phase when the TiBw-lean phase is sufficiently large. Therefore, crack B has to propagate along path I due to the lower energy needed. However, for composites with a very small TiBw-lean phase, i.e. d < 2D, crack B can easily propagate along path II not path I due to the high local stress at the crack tip, which just consumes low fracture energy. Therefore, the large size of the TiBw-lean phase is necessary to exploit the strengthening and toughening effects of the composites. Additionally, if the local volume fraction is too high to stop microcrack coalescence (d < 2D), crack B can easily propagate along path I just consuming very low energy, which results in a brittle fracture and degrades the tensile properties of the composites. It is worth pointing out that more energy or higher load is needed when crack B kinks along path I, which can increase the strengthening and toughening effects of the composites compared with the composites with a homogeneous microstructure. For comparison, the continuous softer phase can first bear stress and strain as shown in Fig. 67b, which discounts the strengthening effect of the reinforcement. Therefore, both the continuous reinforcement phase and the continuous matrix phase are necessary to exploit a superior combination of the strengthening effect and the toughening effect due to the continuous phase being able to dominate the behaviour of the composites. Additionally, the branched structure and dowel-like structure of the TiB whiskers certainly play a positive role in strengthening the composite by increasing reinforcement contiguity. In summary, fracture failure of the TiBw/Ti6Al4V composites with a network microstructure mainly occurs at three length scales that are describe as following: (1) At the TiBw reinforcement level, micro-cracks stem from TiBw without interface failure and pull-out due to strong interface bonding of the in situ composites, which can effectively strengthen the composites by prohibiting dislocation movement (Figs. 62 and 63). (2) At reinforcement-rich region level, the blunting effect of matrix around TiBw reinforcement compels the fractured TiBw segments to further bear load and then further strengthen the composites, which results in multiple fracture of TiB whiskers (Fig. 63). Finally, the micro-cracks coalesce to crack when increased load is applied. (3) At macro network structure level, the internal stress state means that the main crack path has to follow along network boundary and become very tortuous. It is worth noting that the above tortuous crack propagation path bears similarity to the fracture process of the ex-situ composites consisting of an isolated reinforcement-rich phase and continuous matrix phase, such as the short-fibre agglomeration-reinforced 6061Al MMCs (Fig. 8), that has been described in Section 4.1. 5.4. Laminated Ti–TiBw/Ti microstructure On the basis of the above network-structured TMCs, laminated Ti–TiBw/Ti composites composed of alternate pure Ti layers and network-structured TiBw/Ti composite layers were also successfully designed and fabricated in order to further enhance the mechanical properties [57]. Considering the brittleness of the graded TiBw/Ti composites with high volume fractions [273,274], the TiBw/Ti composite layer with just low 5 vol.%TiBw reinforcement is introduced in order to obtain superior ductility. As shown in Fig. 68, the thickness of every layer is about 400 lm and the total volume

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Fig. 68. SEM and OM micrographs of laminated Ti–TiBw/Ti composites: (a) low magnification; (b) high magnification; (c) the pure Ti layer (OM); (d) TiBw/Ti composite layer (OM); and (e) TiBw-rich zone [57].

fraction of the reinforcement is 2.5%. The prepared laminated Ti–TiBw/Ti composites showed a uniform layered structure and a well bonded interface without porosity and cracks. In the pure Ti layer, the matrix near the TiBw/Ti composite layer exhibits fine equiaxed a grains while the inner matrix of the pure layer far away from the composite layer presents lath-like a grains. Additionally, equiaxed a grains are formed in the matrix of the TiBw/Ti composite layer. In other words, all of the in situ TiBw, the network microstructure of the composite layer and the macro laminated microstructure together modify the matrix microstructure characteristics and are beneficial to the overall mechanical properties. As shown in Fig. 69, the prepared Ti–5 vol.%TiBw/Ti composites exhibit 617 MPa tensile strength and 20.5% tensile elongation, which are obviously higher than those (546 MPa and 17.5%) of the pure Ti fabricated using the same raw material and sintering parameters. Although the tensile strength is lower than that of the monolithic 5 vol.%TiBw/Ti composites, the point is that the tensile elongation is higher than not only that of the 5 vol.%TiBw/Ti composites but also that of pure Ti. Combined with fracture surfaces and microstructure observation, several possible reasons are listed as follows. First, the modified matrix microstructure, involving the equiaxed a grains and grain refinement, is positive to the ductility of the laminated composites. Second, the softer pure Ti layer can blunt cracks originating from the stronger composite layer and constrain crack propagation. Third, the stronger composite layer can delay the premature local necking of the softer pure Ti layer. Finally, the modest ductility of the network-structured composite layer is beneficial to overall elongation. In the overall structure, it is believed that the softer pure Ti layer can tolerate a large amount of strain and obstruct crack propagation, while the stronger TiBw/Ti composite layer can bear stress and dominate the tensile behaviour. According to the fracture observation and schematic illustrations of crack initiation and propagation, the cracks first form in the TiB whiskers, and then propagate along the network boundary in the composite layer. The cracks were blunted and arrested when they reached the pure Ti layer. Under higher load, the crack with a size equal to the layer thickness can cause local plastic necking and local shear bands in the pure Ti layer, and then develop until the final failure of the overall laminated Ti– TiBw/Ti composites.

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Fig. 69. The tensile stress–strain curves of pure Ti, 5 vol.%TiBw/Ti composites, the laminated Ti–TiBw/Ti composites [57].

The future work will focus on optimizing the layer parameters including the layer thickness, the layer thickness ratio, the volume fraction of the reinforcement and the matrix type. Therefore, the laminated composites will exhibit much higher mechanical properties, and more interesting results will be presented in future. 6. Future outlook on tailoring inhomogeneous microstructures 6.1. Research significance It is unmistakable that the overall performances, such as strength, stiffness, ductility, toughness, wear resistance and the thermal and electrical conductivities, of discontinuous composites can be significantly enhanced by tailoring the 3D distribution of the reinforcements, and optimizing the structure parameters. For example, 3D bi-continuous composites can simultaneously exhibit the high strengthening effect and high toughening effect. According to the phase contiguity model, for discontinuous composite systems, the elastic modulus increases with the increasing contiguity of the hard reinforcing phase, and the higher the ratio of the reinforcement to matrix modulus (Eb/Ea), the greater the increment in stiffness with increasing contiguity of the reinforcement phase at a fixed overall volume fraction. Moreover, the structures at multiple length scales can be effective in exploiting superior combinations of properties. However, there have been relatively few attempts to create the real multiscale architectures. The difficulty lies in achieving the desired contiguity and spatial distribution of the individual phases, especially on a smaller scale. However, bi-continuous microstructures offer the possibility of developing materials with truly multifunctional properties, i.e., each phase contributing its own characteristics to the macroscopic properties of the composite. For instance, one phase might provide high strength or wear resistance while the other contributes a different property such as electrical conductivity. Research seeking for the optimal structure and structural parameters can help us to systematically understand the relationship between microstructure and performance. Furthermore, understanding the relationship can again guide microstructural design to obtain further enhancements in composite properties. It is established that tailoring the microstructures at different levels to meet specific engineering applications or to significantly improve the combination of properties will create a new class of multiphase materials. The microstructures and performance characteristics of the composites with deliberately tailored reinforcement distributions have been reviewed in the context of microstructural design. The primary conclusions drawn are as follows:

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1. Besides the type, morphology and volume fraction of the reinforcement and the interface bonding between the matrix and the reinforcement, the spatial distribution of the reinforcement seems to have a more profound influence on the overall performances by controlling the contiguities of the matrix and the reinforcement, the dimensions of the reinforcement-lean and reinforcement-rich phases and the distance between the adjacent reinforcement-rich phases. It has been demonstrated that the mechanical, physical and even chemical properties of the composites can be further improved by tailoring the spatial distribution of the reinforcements rather than pursuing a conventional homogeneous composite microstructure. 2. In general, increasing the reinforcement contiguity can exert a superior strengthening effect, which results in improvement in the stiffness, elastic modulus, strength and hardness. However, increasing the matrix contiguity can inspire a superior toughening effect, leading to enhancement of the matrix dominated properties, such as toughness, ductility and deformability. Therefore, a bi-continuous microstructure with the optimal contiguities of the reinforcement and matrix phases possesses a superior combination of mechanical properties. Moreover, in order to enhance toughness or ductility, it is necessary to retain an adequate size matrix phase (reinforcement-lean phase) and ensure a strong interface in order to blunt the crack or form a bridging crack stemming from the adjacent reinforcement phase (reinforcement or reinforcement-rich phase). Crack deflection, depending on the size, morphology and strength of the reinforcement phase, certainly plays a modest toughening effect. Combined with strong interface bonding, delamination, which mainly occurs in laminated composites, can also toughen the composites by absorbing fracture energy. However, delamination, formed due to the weak interface, almost has no toughening effect. 3. The critical problem, i.e., extreme brittleness surrounding the discontinuously reinforced Ti matrix composites (DRTMCs) fabricated by PM, is resolved by tailoring a 3D quasi-continuous network architecture. The large reinforcement-lean phase can remarkably improve the composite ductility by blunting the crack, bearing the strain and slowing down the crack-propagation rate. In addition, the strengthening effect is further enhanced by introducing ceramic reinforcement into the grain boundary phase to form a network distribution, which also overcomes the weakening effect of the grain boundary at high temperatures. The improvement in the strengthening effect is also related to the high reinforcement contiguity within the TiBw-rich network phase, which can dominate the mechanical behaviour of DRTMCs. Certainly, in the overall DRTMCs, the dual 3D architectures give superior strengthening and toughening effects respectively at two different levels. 4. For titanium alloy matrix composites, the subsequent heat treatment exerts significant influence on the matrix microstructure which determines the overall performance of the composites, but almost has no effect on the ceramic reinforcements. The subsequent hot deformation can significantly enhance not only the ductility of the composites by increasing matrix contiguity and decreasing matrix grain size, but also the strength along the deformation direction through work hardening of the matrix and aligned distribution of the reinforcement. 5. Although the work described here is focused on DMMCs, the fundamental principle in tailoring phase distribution can be applied to any two-phase composites: ceramic matrix composites, intermetallic matrix composites and polymer matrix composites. They can exhibit a superior combination of properties by tailoring particular inhomogeneous microstructures rather than pursuing a homogeneous distribution pattern, such as interpenetrating microstructures, bi-continuous microstructures, quasi-continuous network microstructures and laminated microstructures. 6. Through understanding nature’s hierarchical structures, bio-inspired microstructures exhibit a remarkable combination of mechanical properties, such as strength and toughness. However, it is critical to replicate bio-inspired microstructures with more than one length scale structures, which is not a trivial task facing composite researchers. 7. It is worth pointing out that reproducibility of mechanical properties is an issue for inhomogeneous composites. Because the properties of the composites with inhomogeneity of reinforcement might vary more widely than those of the homogenous composites. In addition, the inhomogeneous architecture designed for improvement of properties may be changed or, at least to some extent, destroyed by the secondary processing. Therefore, some of the advantages of a tailored inhomogeneous composite may lost in a final processed component.

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6.2. Outlook of future research Having justified the effort in deliberately tailoring the microstructure of discontinuous composites, several relevant research directions are proposed as follows: 1. The inhomogeneous microstructure should be gradually and systematically analysed using numerical methods, such as the FEM. That is to say, the tailored multi-scale structures should be analysed at different length scales, such as treating the reinforcement-rich phase as one integrated reinforcement unit at a higher level. Understanding the stress and strain characteristics, fracture and strengthening mechanisms of the reinforcement phases at different levels can obviously help identify optimal microstructures. 2. The optimal structures, such as bi-continuous structure or network structure, can be deduced with targeted performances under the identical matrix, reinforcement and their corresponding volume fractions. Furthermore, under a constant volume fraction of reinforcement, the structure types can be separately constructed and optimized for different systems of the matrix and reinforcement, such as Ti–TiBw, Al–SiCp systems, as designed and formulated by the FE model in Zuo’s work [275]. 3. For a given system, by assigning the fundamental parameters of the matrix and reinforcement into the FE model, the structure parameters, including the sizes of the individual reinforcements and the reinforcement-rich phase, the overall volume fraction of the reinforcement and the reinforcement-rich phase, the distance between the adjacent reinforcement-rich phases can be optimized for different applications. For example, high reinforcement contiguity may correspond to the highest strengthening effect, while high matrix contiguity may correspond to the highest toughening effect. 4. New fabrication processes should be developed to prepare the optimized structures with tailored spatial distributions of the reinforcement. In addition, the structures with multi-scale hierarchies can be designed and fabricated down to nano-scale using new fabrication processes in order to significantly enhance the performances of the composites, which are closer to natural structures. It will be interesting to construct a hierarchical reinforcement scaffold using 3D printing techniques, followed by infiltrating the matrix into the scaffold in order to prepare the composites with tailored microstructures. Furthermore, the composites with tailored microstructures can be directly fabricated by one-step 3D printing techniques with a multiple feed matrix and reinforcement raw materials. 5. The composites exhibiting superior performances should be extensively exploited to replace conventional materials which can in turn lead to further progress in fundamental theories and fabrication techniques of the composites. The final aim is to explore different structures with different parameters for different systems in order to satisfy different performance requirements or the optimal combination of properties. In this context, it is fitting to end this review with a quote from Milton [276]: ‘It is hard to look into the future, but undoubtedly it will become increasingly possible to produce ‘‘designer composites’’, where the microstructure has been tailored to produce desirable properties. Obviously, the better understanding of the link between the microstructure and the macroscopic properties will be essential in the endeavor’. Acknowledgements LJH thanks the China Scholarship Council for supporting his one year exchange (2009-2010) in the Advanced Composites Centre for Innovation and Science (ACCIS) at the University of Bristol (UK) under the supervision of HXP where some of the research work reported here were carried out. HXP is under The Recruitment Program of Global Experts and is thankful to Professor Zhongyun Fan of Brunel University (UK) and Professor Julian Evans of University College London (UK) for their inspiration in some of the early works in this area. Financial supports from the High Technology Research and Development Program of China (863) under Grant No. 2013AA031202, the National Natural Science Foundation of China (NSFC) under Grant Nos. 51101042, 51271064, 51228102 and

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51471063 as well as the Fundamental Research Funds for the Central Universities under Grant No. HIT.BRETIII.201401 are acknowledged. The authors are indebted to those who kindly authorised permission to use or reproduce figures and drawings where corresponding references are duly indicated as well as the anonymous referees for their invaluable and constructive comments that helped to shape this work. References [1] Clyne TW, Withers PJ. An introduction to metal matrix composites. 1st ed. Cambridge: Cambridge University Press; 1995. p. 1–60. [2] Tjong SC. Recent progress in the development and properties of novel metal matrix nanocomposites reinforced with carbon nanotubes and graphene nanosheets. Mater Sci Eng R 2013;74(10):281–350. [3] Lloyd DJ. Particle reinforced aluminium and magnesium matrix composites. Int Mater Rev 1994;39(1):1–23. [4] Tjong SC, Ma ZY. Microstructural and mechanical characteristics of in-situ metal matrix composites. Mater Sci Eng R 2000;29:49–113. 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