Progress on discontinuously reinforced titanium matrix composites

Accepted Manuscript Progress on discontinuously reinforced titanium matrix composites Yang Jiao, Lujun Huang, Lin Geng PII:

S0925-8388(18)32601-X

DOI:

10.1016/j.jallcom.2018.07.100

Reference:

JALCOM 46816

To appear in:

Journal of Alloys and Compounds

Received Date: 16 May 2018 Revised Date:

7 July 2018

Accepted Date: 9 July 2018

Please cite this article as: Y. Jiao, L. Huang, L. Geng, Progress on discontinuously reinforced titanium matrix composites, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.07.100. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT

Progress on discontinuously reinforced titanium matrix composites Yang Jiao, Lujun Huang*, Lin Geng

RI PT

State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, P.O. Box 433, Harbin, PR China

Abstract

Discontinuously reinforced titanium matrix composites (DRTMCs) are attractive

SC

materials for the application of aeronautics and astronautics, automotive and military fields, due to their low density, heat-resistance, and excellent room- and

M AN U

high-temperature properties. Recently, numerous literatures have been published about the DRTMCs, it is essential to summarize their research progress, in order to further comprehend their development and guide the future work. This work includes several significant aspects of DRTMCs: recent development in the reinforcements,

TE D

microstructure characteristics, mechanical properties as well as modeling and calculation. The present review attempts to reveal the existing problems and propose further research directions for the aforementioned topics. In order to further improve

EP

the overall properties and extend the applications of DRTMCs, bringing the hierarchical structures of biological materials into the composites is an effective

AC C

method. Furthermore, the numerical methods should be performed to study the composites fundamentally and obtain microscopic information of the composites.

Keywords: discontinuously reinforced titanium matrix composites; reinforcements; microstructure characteristics; overall performances; modeling and calculation *Corresponding author: E-mail address: [email protected] 1

ACCEPTED MANUSCRIPT

Contents 1. Introduction .............................................................................................................. 3 2. Reinforcements ......................................................................................................... 8 3. Microstructure characteristics.............................................................................. 12

RI PT

4. Overall performances ............................................................................................ 19 4.1 Tensile properties ........................................................................................... 20 4.2 Wear and Fatigue behavior............................................................................. 24

SC

4.3 Oxidation resistance ....................................................................................... 27 4.4 Creep resistance ............................................................................................. 33

M AN U

5. Modeling and Calculation ..................................................................................... 36 5.1 Analytical model ............................................................................................ 37 5.2 First-principles calculation............................................................................. 39 6. Summary and Future outlook............................................................................... 41

TE D

Acknowledgement ...................................................................................................... 41

AC C

EP

References ................................................................................................................... 41

2

ACCEPTED MANUSCRIPT

1. Introduction Numerous studies manifest that compared with Ti alloys, discontinuously reinforced titanium matrix composites (DRTMCs) reinforced with ceramic particles or whiskers possess similar density (4.5 g/cm3), however, the service temperature of

RI PT

DRTMCs can be enhanced by 200 oC, even reaches 600-800 oC. Compared with TiAl, NiAl and Ti2AlNb alloys, the DRTMCs have higher plasticity, and advantages of formability and weldability. Especially, the weight reduction can reach 40% when the

SC

high-temperature alloys or steels were replaced by DRTMCs. Therefore, DRTMCs have been widely regarded as promising structural materials in aerospace, automotive

M AN U

and military fields [1, 2]. Compared with the DRTMCs fabricated by ex-situ methods, the in-situ DRTMCs showed the advantages of strong bonding between matrix and reinforcement, little contamination, ease of processing and low cost [3]. The in-situ synthesis techniques include mechanical alloying (MA) [4], powder metallurgy (PM)

TE D

[5] and self-propagation high-temperature synthesis (SHS) [6]. In addition, casting exhibited flexibility and large scale production [7, 8]. However, some defects, such as pores and materials waste, presented during the melting process, which could be solved

EP

by cost-intensive process including many steps of hot deformation (such as forge and

AC C

extrusion) and heat treatments (such as annealing and solution). In recent years, based on the conventional powder metallurgy process, some prevailing methods including reaction hot pressing (RHP), spark plasma sintering (SPS) and hot isostatic pressing (HIP) emerged. Taking the RHP method as an example to illustrate the fabrication process of DRTMCs, as shown in Fig. 1. The first step is to mix the Ti matrix and reinforcement raw powders by low- or high-energy ball milling. Afterwards, the blended mixture was sintered in a mould under high temperature and pressure. The in-situ synthesis combined with RHP method possesses microstructural control, near 3

ACCEPTED MANUSCRIPT net shape processing and minimal materials waste, particularly its ability to render composites with densification and excellent mechanical properties [1]. Compared with the conventional hot pressing, SPS is an optional approach. It can increase the density during hot compaction and avoid the grain growth, due to the shorter sintering time [9].

M AN U

SC

11] and RHP combined with stacking powders [12].

RI PT

Recent approaches fabricating the DRTMCs contain additive manufacturing (AM) [10,

TE D

Fig. 1 A schematic illustration of powder metallurgy process [1, 9]. (a) Ti mixed with ceramic powders; (b) Mixed powders were low/high-energy ball milled under Argon atmosphere; (c) Blended powders were sintered under temperature and pressure via reaction hot pressing (RHP).

In-situ synthesized TiB whiskers (TiBw) and TiC particles (TiCp) are commonly

EP

used to reinforce Ti alloys for fabricating TMCs with superior properties [3, 13]. The

AC C

addition of slight TiB or TiC can increase the strength and wear performance of composites, on account of their high modulus and hardness [13]. Due to their density and coefficient of thermal expansion (CTE) are similar with those of Ti, the residual stresses at ceramics/Ti interfaces can be reduced or even eliminated [14]. In addition, good chemical and thermodynamic stability of TiB can bring in the formation of good interfacial bonding between Ti and TiB [14]. TiB/Ti composite was also considered as biomedical materials, owing to the biocompatibility of B element [15]. TiC possesses high oxidation resistance [16] and low friction coefficient [17]. Except for TiB and 4

ACCEPTED MANUSCRIPT TiC, a variety of ceramic reinforcements including carbon nanotubes (CNTs) [18], graphene [18], TiN [19], B4C [20] and Ti5Si3 [21] are used for strengthening Ti alloys. For conventional metal matrix composites (MMCs), the strength can be improved by increasing fractions of reinforcements, however, the elongation of MMCs will decrease

RI PT

sharply. In the past decades, many efforts have been paid to fabricating the DRTMCs with a homogeneous distribution of reinforcements. However, the homogeneous microstructures were proved to not only limit the improvement of strength, but also

SC

lead to a remarkable decrease in ductility for the composites without secondary processing [22]. In contrast, controlling the inhomogeneous distribution of

M AN U

reinforcements to simultaneously increase strength and ductility is an innovative approach. For example, Huang et al. [13] reported for the first time that by tailoring inhomogeneous distribution of TiBw around Ti6Al4V alloys (Fig. 2a), unprecedented increase of room-temperature tensile properties can be achieved. The feasibility of such

TE D

a strategy in other titanium matrix composites was also validated [1]. Additionally, laminated structure [23], fiber-like arrays of nanoparticles [24] and ring structure [25] as the typical inhomogeneous architectures were presented in Fig. 2b-2d. Recently, the

EP

biological materials have inspired enormous interests among researchers, due to their

AC C

hierarchical structures and outstanding mechanical properties [26, 27]. Therefore, bio-inspired structures as a popular idea were proposed in designing and fabricating alloys and composites, in order to surmount the trade-off between strength and ductility [28]. Wang et. al simultaneously [29] enhanced the strength and ductility of stainless steels, due to the hierarchically heterogeneous microstructure with multiple length scales. An unprecedented tensile ductility has been achieved from the Mo alloys with a hierarchical microstructure [30]. Through designing the DRTMCs with an inventive two-level hierarchical microstructure and micro and nano reinforcements, good 5

ACCEPTED MANUSCRIPT mechanical properties were displayed in recent work [21]. In this sense, hierarchy and multi-scale structures will become one of the future perspectives of further improving

M AN U

SC

RI PT

overall properties of composites.

Fig. 2 Schematic illustrations and representative SEM images of different architectures with inhomogeneous distribution of reinforcements. (a) Network structure [1]; (b) Laminated structure [23]; (c) Fiber-like arrays of nanoparticles [24]; (d) Ring structure [25].

Due to good overall properties, the discontinuously reinforced titanium matrix

TE D

composites had been used in military, automotive, industrial and biomedical fields [31]. For example, Dynamet Technology had produced valves, connecting rods and piston pins made of titanium matrix composites, as shown in Fig. 3 [31]. Compared

EP

with the aluminum matrix composites and high-strength steels, the titanium matrix composites as connecting rods had higher strength and lower weight, respectively.

AC C

Hockey skate blades composed of TiC/Ti6Al4V composite displayed lighter weight than their steel counterparts by 40% and higher fracture resistance [31].

Fig. 3 (a) A schematic illustration of engine valve with Ti6Al4V core and TMC clad layer. (b) The forged automotive connecting rods made of TiC/Ti6Al6V2Sn composite [31]. 6

ACCEPTED MANUSCRIPT During the past several decades, researchers made concentration on the studies of fabrication methods, reinforcements distribution adjustment and improvement in room-temperature properties of Ti matrix composites (TMCs). For example, Shin et al. prepared the TiC/Ti6Al4V composite by laser direct deposition and investigated the

RI PT

effect of melting degree of TiC reinforcements on microstructure (Fig. 4) [11]. Wang et al. designed and fabricated the Ti matrix composites with a laminate-network microstructure [12]. In recent years, in order to increase strength without decreasing

SC

ductility drastically by adjusting the microstructures of composites, one should consider to bring the multi-scale and hierarchical microstructures into the composites

M AN U

without any treatments. Furthermore, the high-temperature properties, such as oxidation resistance and creep resistance are important for broadening the application of DRTMCs at high-temperature fields [32, 33]. Both analytical and numerical models can be used to investigate the strengthening and toughening mechanisms of

TE D

DRTMCs. In the finer scale, first-principles calculation can be employed to obtain the

AC C

EP

atomistic information of DRTMCs.

Fig. 4 (a) Laser deposited TiC/Ti6Al4V composite cylinder. (b) Morphologies of TiC in the deposited composite [11].

To date, a large amount of studies have reported the microstructures and properties

of DRTMCs. The present review emphasizes on the recent progress of reinforcements, microstructure characteristics, overall properties as well as modeling and calculation of DRTMCs. In addition, the unsolved problems and future opportunities in the present research are provided. 7

ACCEPTED MANUSCRIPT

2. Reinforcements Among the literatures studying the reinforcements of DRTMCs, in-situ TiBw [5, 13] and TiCp [16] have been unanimously considered as the best reinforcements, due to their desirable properties descripted in section 1. As shown in Fig. 5 [34], the

RI PT

cross-section of the in-situ TiB is characterized by (100), (101) and (10 1 ) crystal planes with a hexagonal shape. The growth direction of TiB is the [010] direction, therefore, the morphology of whisker or short fiber forms easily. The orientation

SC

relationship between TiB and α-Ti matrix was identified as: (101)TiB//(0001)α-Ti and

M AN U

[010]TiB//[11 2 0]α-Ti. TiC exhibited ellipse shape, which is related to the fabrication process and C content. In generally, the morphology of TiC is dendritic in the liquid-solid reaction (Fig. 4b). TiC reinforcements can be generated by the reactions of Ti and various carbon materials, such as carbon nanotube, graphite and carbon fiber. In addition, hybrid reinforcements, such as TiBw and TiCp [2], TiBw, TiCp and La2O3

TE D

[35] were used to achieve higher mechanical properties. However, excess ceramic reinforcements in the composites led to high strength and low elongation. That is to say,

EP

constrained by the strength-ductility trade-off. The high-temperature oxidation resistance of TMCs reinforced by TiBw is inferior, due to the evaporation of B2O3 [32].

AC C

For the composites reinforced by TiCp, element diffusion in the Ti/TiC interface and the formation of non-stoichiometric TiC layer can affect the properties of composites. Most of rare earth oxides (REOs) are poisonous and their density is much higher than that of Ti alloys. The purpose of fabricating DRTMCs is to develop the tougher, stronger and lighter structural materials. Therefore, other reinforcements should be developed to improve the properties and extend the applications of DRTMCs.

8

RI PT

ACCEPTED MANUSCRIPT

Fig. 5 TEM results of the (TiB+TiC)/Ti6Al4V composite. (a) TEM image of different phases and selected area electron diffraction (SAED) pattern of TiC; (b) TEM image of the cross section of TiB; (c) SAED pattern of two phases shown in (b) [34].

SC

For weight savings in aerospace and automotive industries, carbon nanotubes (CNTs) are attractive from the view of their superior mechanical properties, such as low

M AN U

density (~1.7-2.0 g/cm3) and high Young’s modulus (~1TPa). CNTs can be classified into single walls, double walls and multi walls, as shown in Fig. 6a [36]. TEM image of large bundles and clustering of the multi-wall CNTs (MWCNTs) were presented in Fig. 6b. HRTEM image (Fig. 6c) showed an individual MWCNT with an inter-layer

TE D

spacing of 0.380 nm [37]. CNTs show a promising contribution in enhancing the mechanical properties of composites with lightweight in recent studies [37, 38]. Graphene is a two-dimensional carbon material and shows similar mechanical

EP

properties with CNTs. Moreover, graphene has large specific area and high aspect ratio. Hence, graphene is also a good substitution for conventional ceramics as reinforcement

AC C

for metal matrix composites. The strength of Ti alloys could be effectively increased without much sacrifice to ductility when introducing slight graphene nanoplatelets [39]. However, for the CNTs and graphene as reinforcements in TMCs, some existing problems are as follows [40, 41]. On the one hand, the dispersion of CNTs or graphene is difficult and slight fractions of reinforcements can limit the enhancement of performances. On the other hand, CNTs or graphene react with Ti easily and the formation of TiC can decrease the strengthening effect. The achievement of CNTs or graphene are difficult, due to the complicated fabrication process and high cost. 9

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 6 (a) Single-, double-, and multi-walled CNTs [36]; (b) TEM image of the multi-walled CNTs (MWCNTs) [37]; (c) HRTEM image and selected area electron diffraction of the MWCNT [37].

M AN U

For aeronautical, marine and chemical industries, Ti5Si3 is a promising candidate as the reinforcement in TMCs, due to its high melting point (2130 °C), low density (4.26 g/cm3), retaining strength up to 1200 °C, especially good oxidation and creep resistance at elevated temperatures [42]. Moreover, the coefficient of thermal

TE D

expansion of Ti5Si3 is also close to that of Ti alloys [43]. Therefore, Ti matrix composites reinforced with Ti5Si3 particles are potential materials for high-temperature applications. Ti5Si3 is also potentially used in wear and corrosion field, due to its

EP

inherent high hardness, unique chemical composition and strong covalent bonds [44]. Previous literatures indicated that the Ti5Si3 reinforcements were commonly observed

AC C

in the TiAl matrix composites [45, 46]. While, limited systematic work is available on Ti5Si3

reinforced

Ti

matrix.

Huang

et

al.

designed

and

sintered

the

1vol.%(Ti5Si3+Ti2C)/Ti composite, whose room-temperature strength and ductility were drastically enhanced, compared with that of as-sintered pure Ti alloy [47]. Owing to its complex hexagonal structure (as shown in Fig. 7a), Ti5Si3 possesses considerably low fracture toughness at room temperature. The formation of coarse Ti5Si3 particles in the conventional cast Ti alloys deteriorated the ductility of alloys. However, the 10

ACCEPTED MANUSCRIPT strength and ductility of TiBw/Ti6Al4V composite was improved by adding nano-Ti5Si3 particles [48] (Fig. 7b and 7c), which were obtained via adjusting the

M AN U

SC

RI PT

fabrication process and tailoring the microstructure of composite.

Fig. 7 (a) A schematic illustration of Ti5Si3 crystalline structure; (b) TEM image and energy dispersive spectrum (EDS) result of Ti5Si3 [48]; (c) The selected area electron diffraction pattern of Ti5Si3 in (b) [48].

TE D

In conclusion, the properties of reinforcements commonly used and investigated in DRTMCs are summarized in Table 1. TiB and TiC reinforcements are beneficial to improving strength and modulus of the Ti matrix composites. However, excess

EP

reinforcements will decrease the tensile properties significantly. Carbon nanotubes

AC C

(CNTs) and graphene are suitable reinforcements for the sake of reducing weight and enhancing hardness. Adjusting the distribution and orientation of CNTs or graphene without alteration of constituents is an effective way in enhancing strengthening efficiency or balancing strength and ductility of composites. Ti5Si3 particles are promising reinforcements for DRTMCs applying to high-temperature fields, while a rapid reduction of ductility caused by coarse Ti5Si3 particles should be eliminated.

11

ACCEPTED MANUSCRIPT Table 1 Properties of reinforcements commonly used in discontinuously reinforced Ti matrix composites Density (g⋅cm-3)

Elastic modulus (GPa)

Coefficient of thermal expansion (×10-6 K-1)

Ti5Si3 Graphene CNTs TiB TiC TiN SiC TiB2 B 4C La2O3

2130 ~3852 ~3379 2200 3160 3290 2697 2980 2447 2217

4.32 2.00 1.70-2.00 4.50 4.99 3.97 3.19 4.52 2.51 6.51

225 ~1000 ~1000 425-480 440 420 430 500 445 --

9.70 --8.60 6.52-7.15 (25-500 oC) 8.30 4.63 (25-500 oC) 4.60-8.10 4.78 (25-500 oC) 5.86-12.01 (100-1000 oC )

SC

M AN U

3. Microstructure characteristics

RI PT

Reinforcements

Melting point (oC)

In order to optimize the overall properties of DRTMCs, the unique structure design by controlling the distribution of reinforcements at different scales attracts much interest in recent years. A series of composites with a controlled inhomogeneous

TE D

microstructure (Fig. 2) were successfully designed and fabricated through different routes. In those systems, the reinforcements were inhomogeneously distributed in the micro-scale, but homogeneously or still inhomogeneously distributed in the

EP

macro-scale [1]. As expected, under the same conditions (without heat treatments and

AC C

plastic deformation), these microstructures in the fabricated composites led to remarkable improvement in strength, ductility and fracture toughness, compared with conventional metal matrix composites with homogenous distribution of reinforcements [1].

Functionally graded materials (FGMs) are inhomogeneous composite materials, and the constituents change continuously or step-wise as a function of position [49]. Zhang et al. [49] made a four-layer TiB-Ti system functionally graded material via spark plasma sintering (SPS) method. The mechanical properties of different area in 12

ACCEPTED MANUSCRIPT the materials can be adjusted by the fractions of TiB whiskers. Wang et al. [50] adopted the laser melting deposition (LMD) method and synthesized the functionally graded TiCp/Ti6Al4V composite successfully. In the composite, the volume fraction of TiC particles changed gradually from 0% to 50%. Fig. 8 illustrated the macrostructure

RI PT

of the functionally graded TiCp/Ti6Al4V composite [50]. The height of the as-deposited composite was approximately 37.0 mm and no cracks were observed. The blocky TiC particles distributed inhomogeneously in the Ti matrix. Moreover, their

SC

quantity increased with increasing the volume fraction from 0% to 50%. The interfaces were smooth and exhibited good bonding between individual layers. Although the

M AN U

strength of the composite increased, the low ductility with high fraction of TiC was

TE D

presented.

Fig. 8 Morphologies of functionally graded TiCp/Ti6Al4V composite. (a) Macrostructure of the

EP

graded composite; (b) 5vol%TiCp; (c) 10vol%TiCp; (d) 15vol%TiCp; (e) 20vol%TiCp; (f) 30vol%TiCp; (g) 40vol%TiCp; (h) 50vol%TiCp [50].

AC C

Recently, Zou et al. [34] designed and fabricated the (TiBnw+TiCp)/Ti6Al4V composite with a three-dimensional network-woven microstructure by SPS (Fig. 9a and 9b). The as-sintered composites exhibited superior tensile properties, compared with those cost-intensive products such as the heat treated Ti6Al4V or forged Ti10V2Fe3Al alloys followed by heat treatments. The in-situ formed TiB nanowires with high aspect ratios improved the tensile strength and the architectural design resulted in crack deflection that contributed to retain good tensile ductility of the as-sintered

composites.

Furthermore,

Cong 13

et

al.

[10]

reported

the 3D

ACCEPTED MANUSCRIPT quasi-continuous network (3DQCN) microstructure (Fig. 9c and 9d) in TiB-Ti system fabricated by laser deposition-additive manufacturing (LD-AM) process, which has advantages of small substrate deformation, alterable powder mixture ratio and controllable cooling rate. The formation mechanism of the network structure was as

RI PT

follows. Firstly, TiB formed by the reaction of Ti and B was completely melted into molten material. Secondly, during the solidification process, TiB nucleated and grew into long whiskers, providing heterogeneous nucleation sites for the liquid Ti. Then,

SC

the high-temperature β-Ti nuclei was generated. TiB whiskers were pushed into the liquid ahead of the nuclei-liquid interface, due to the growth of β-Ti nuclei. Thirdly,

M AN U

TiB whiskers accumulated at Ti grain boundaries, forming the 3DQCN microstructure in the previous molten pool area. The strength was increased and the ductility of

AC C

EP

TE D

composite was decreased significantly, compared with Ti alloys.

Fig. 9 SEM micrographs of the (TiBnw+TiCp)/Ti6Al4V (a and b) composite [34] and TiB/Ti

composite (c and d) [10]. (a) and (c) 3D presentation of the composites; (b) and (d) Magnified view of the composites.

SEM micrographs of the MWCNTs/Ti composites fabricated by SPS method were exhibited in Fig. 10 [51]. Few pores were presented, and no TiC particles were observed in the composites. The multi-wall carbon nanotubes (MWCNTs) mainly 14

ACCEPTED MANUSCRIPT distributed at the grain boundaries of the Ti matrix, forming the analogous network structure. With increasing the weight fraction of MWCNTs, the grain boundaries and grain size of Ti matrix became clearer and smaller, respectively. The Zener pinning effect induced by MWCNTs could hinder the grain growth and grain boundary sliding

M AN U

SC

RI PT

of matrix, therefore, the strength of the composites was increased.

Fig. 10 SEM images of the MWCNTs/Ti composites with different MWCNTs fractions: (a) 0.2 wt.%; (b) 1.0 wt.%(the inset shows a magnified image of the CNTs) [51].

In order to further improve the ductility of discontinuously reinforced titanium matrix composites without secondary processing, laminated architecture was

TE D

performed in the composites. A large amount of studies had been carried out on the laminated composites fabricated by sintering or diffusion bonding. On the basis of the above-mentioned network structure (Fig. 2a), laminated Ti-(TiBw/Ti) composite

EP

consisting of alternate pure Ti layers and network-structured TiBw/Ti composite layers (Fig. 2b) were designed and fabricated [23]. The tensile strength (617 MPa) was higher

AC C

than that of pure Ti (546 MPa), however, lower than that of network-structured TiBw/Ti composite (720 MPa). The elongation was about 20.5%, which was much higher than those of TiBw/Ti composite (7.5%) and pure Ti alloy (17.5%) under the same condition [52]. Sometimes, the large grains or weak interface bonding in the laminated composites caused unsatisfactory performance. Thus, further processing should be introduced to enhance the interface bonding in the laminated DRTMCs, so as to improve the strength of the composites. 15

ACCEPTED MANUSCRIPT Natural biomaterials, such as shells, bones and nacres reveal low weight, high strength and toughness, due to their structure changes with hierarchical level (from the macroscale to nanoscale in length-scale) [26, 27]. In nacre [28], at the millimeter scale, the shell contains two lamellar system. One is a hard outer layer with large calcite

RI PT

crystals; the other is a softer inner layer with nacre, as shown in Fig. 11a and 11b. The microscale structure of nacre consists of a 3D brick and mortar wall. The bricks are densely packed layers of microscopic aragonite polygonal tablets, which are about 5-8

SC

µm in diameter and 0.5 µm in thickness. While, the tablets are held together by organic materials layers (20-30 nm thick, Fig. 11c-e). Jiang et al. fabricated biomimetic

M AN U

Al2O3/Al composite with nanolaminated structure by imitating nacre structure, and achieved a well-balanced strength and ductility [28]. Zhang et al. investigated the bio-inspired and artificial wing-like structures, which may be used in advanced sensors and solar cells [53]. Therefore, learning from natural biological structure may be an

TE D

effective strategy to obtain a significant improvement in the combination of strength and ductility for TMCs. However, few literatures reported that DRTMCs have hierarchical structures design and outstanding mechanical properties without any

EP

treatments. It is difficult to replicate the hierarchical structures of biomaterials

AC C

completely in the laboratories.

Fig. 11 The multi-scale structure of nacre: (a) inside view of the shell; (b) cross section of a red 16

ACCEPTED MANUSCRIPT abalone shell; (c) a schematic illustration of the brick wall microstructure; (d) OM of the tiling in the tablets; (e) SEM of a fracture surface [26].

Combining the one-level network structure (Fig. 2a) with hierarchical structure design, Jiao et al. [21] conceived and fabricated the (Ti5Si3+TiBw)/Ti6Al4V composite

RI PT

with a two-level hierarchical microstructure by RHP. Owing to the unique distribution of two reinforcements, the composite possess controllability. For instance, the volume fraction of different reinforcements could be adjusted to control the mechanical

SC

properties of composites. The morphology and fraction of Ti5Si3 reinforcement could be adjusted by sintering parameters, TiBw fraction and heat treatments, owing to the

M AN U

unique formation mechanism of Ti5Si3 [21, 54]. Ti5Si3 particles could dissolve into the Ti matrix after solution at 1200 oC and precipitate after aging at 600 oC [54]. The as-sintered composite exhibited superior room- and high-temperature properties, such as high tensile strength and creep resistance. Fig. 12 showed SEM micrographs of the (Ti5Si3+TiBw)/Ti6Al4V

composite

[54].

On

the

whole,

TiBw

TE D

as-sintered

reinforcement formed the regular first-level network structure (FLNS, Fig. 12a), due to the in-situ reaction between Ti and TiB2. The FLNS could contribute to the

EP

enhancement of strength by refining the matrix size. The matrix morphology transferred from Widmanstätten microstructure (in as-sintered Ti6Al4V alloy) to short

AC C

lamella microstructure and the α colony matrix was refined. From the view of lower level, the Ti5Si3 particles precipitated in the β-Ti phases (distributed around the α-Ti phases), and formed the second-level network structure (SLNS, Fig. 12b and 12c), due to high-temperature solid solution of in-situ Ti5Si3 and precipitation of Ti5Si3 during cooling process. The SLNS was beneficial to improving the ductility of composite via enhance the deformation compatibility, compared with the as-sintered TiBw/Ti6Al4V composites with one-level network structure. Since the two reinforcements located at different scales, the strength and ductility of as-sintered composite with a hierarchical 17

ACCEPTED MANUSCRIPT structure were improved simultaneously, compared with

the conventional

hybrid-reinforced casted and sintered composites without heat treatments and plastic

M AN U

SC

RI PT

deformation [2, 55].

Fig. 12 SEM micrographs of the (Ti5Si3+TiBw)/Ti6Al4V composite. (a) The first-level network structure; (b) and (c) The second-level network structure [54].

Recently, Wang et al. fabricated a novel Ti6Al4V-(TiBw/Ti6Al4V) composite

TE D

with two-level laminate-network structure (Fig. 13) and enhanced the impact toughness, compared with the TiBw/Ti6Al4V composite with one-level network

EP

structure under the same condition [12]. The increased toughness was attributed to the laminate structure, microstructure refinement and interface crack deflection. From the

AC C

view of macro-level, the novel composites were composed of Ti6Al4V alloy layers and network structured TiBw/Ti6Al4V composite layers. From the view of micro-level,

the

TiBw/Ti6Al4V composite

layer

possesses

refined

matrix

microstructure. Layer thickness and TiBw fraction could be adjusted by controlling the weight of raw powders.

18

RI PT

ACCEPTED MANUSCRIPT

SC

Fig. 13 Morphology of the two-level laminate-network structured Ti6Al4V-(TiBw/Ti6Al4V) composite. (a) Low magnification; (b) High magnification; (c) Macrograph [12].

M AN U

To summarize, the inhomogeneity of specific arrangement of reinforcements was beneficial for the mechanical properties of TMCs. Moreover, imitating the structures of natural biomaterials is another possible route in developing composites with superior mechanical properties, particularly in the aspects of improving the strength

TE D

without sacrificing the ductility. The hierarchical structures can defeat the strength-ductility trade-off, which recently proposed for structural materials by Ritchie [56] and Lu [57].

EP

4. Overall performances

Ti matrix composites were used as structural components under different

AC C

conditions, and the aim of developing TMCs is to further improve the strength, modulus and service temperature of Ti alloys. Therefore, the performances including tensile properties, wear, fatigue, oxidation and creep resistance are the important criteria to examine the success of the designed and fabricated composites. The mechanical properties of DRTMCs mainly depend upon three factors: 1) the composition or microstructure of matrix; 2) the distribution, morphology, fraction and size of reinforcement; 3) the interface between matrix and reinforcement. At room 19

ACCEPTED MANUSCRIPT temperature, the addition of ceramic phases to Ti alloys increases strength, but at the expense of decreasing ductility. At high temperature, the strength, oxidation and creep resistances of Ti alloys were considerably improved [16, 58]. 4.1 Tensile properties

RI PT

Owing to the high specific strength, good biocompatibility and outstanding corrosion resistance, Ti and Ti-based alloy are extensively used in engineering fields. For example, in most aircraft engines, Ti alloy parts account for 20-30% of engine

SC

weight [59, 60]. Ti6Al4V is a prevalent Ti alloy, which has α+β dual-phase and has

M AN U

been widely applied for aerospace, marine and biomedical implants [59]. Due to most Ti matrix composites fabricated by in-situ methods, the interface bonding between Ti matrix and ceramic reinforcement was strong [3]. Therefore, according to the aforementioned influence factors of mechanical properties, taking pure Ti (Fig. 14a) and Ti6Al4V (Fig. 14b) matrix composites as an example, to investigate the

TE D

room-temperature tensile properties of composites reinforced with different ceramics (such as TiB, TiC, Ti5Si3 and CNTs). Moreover, the composites were fabricated by different routes, including reaction hot pressing (RHP), spark plasma sintering (SPS),

EP

powder metallurgy (PM) and casting.

AC C

The addition of excess ceramic reinforcements resulted in a significant decrease in the tensile ductility. For example, the elongation of the Ti matrix composite reinforced with 8.15vol.%TiB, 1.26vol.%TiC and 0.59vol.%La2O3 decreased from 10.2% to 3.3% [35]. For the functional gradient TiC-Ti composite fabricated by laser powder deposition, the tensile strength changed slightly with TiC addition, but the ductility decreased rapidly [61]. However, the hardness and wear resistance of Ti-TiC composite increased, due to the strengthening effect of hard TiC phase and the good interface bonding between TiC and Ti [61]. When compared with the DRTMCs fabricated by 20

ACCEPTED MANUSCRIPT casting, PM and SPS methods without any processing [62-64], the tensile properties of the as-sintered composites fabricated via RHP were superior. As compared with Ti composite reinforced with carbon nanotubes [65], the 10wt.%W/Ti composite had a higher tensile strength and elongation (920 MPa and 14%) [66]. Through designing

RI PT

novel structure for DRTMCs, Huang et al. improved the mechanical properties of the as-sintered composites without any secondary processing [67]. For example, as shown in Fig. 14a, the mechanical properties of network-structured composite increased

SC

significantly (842.3 MPa and 11.8%), compared with the 8.5vol.%TiBw/Ti composite with a conventional homogeneous structure (687.9 MPa and 2.9%). Compared with

M AN U

pure Ti (17.8%), the strength of the 5vol.%TiBw/Ti composite increased from 482.4 MPa to 753.8 MPa, and the elongation remained at 15.6%. Unprecedentedly, the 1vol.%(Ti5Si3+Ti2C)/Ti composite presented a superior elongation of 28.9%, which was attributed to the tailored network microstructure and the refined matrix grain [47].

TE D

Similarly, Zou et al. [68] fabricated the network-woven structured Ti matrix nanocomposite reinforced with TiB nanowires, and achieved high tensile strength (808 MPa) and sufficient tensile ductility (20.3%). The composites were merely

EP

fabricated under low-energy short mixing process and a fast sintering technique. As

AC C

shown in Fig. 14a, the above results can infer that the network structure design can improve strength of the fabricated composites without decreasing ductility largely. Owing to the increasingly demanding lightweight applications, it is desirable that

the tensile strengths of DRTMCs can be further improved while maintaining sufficient (>6%) or good (>10%) tensile ductility [34]. Fig. 14b compares the tensile properties reported for Ti6Al4V matrix composites reinforced with different ceramics and fabricated by different routes. It can be seen obviously that the strengths of most composites are enhanced, while the elongations are under 6% [2, 5, 69-75]. Recently, 21

ACCEPTED MANUSCRIPT some literatures improved the strength with slight decrease in the ductility. For example, the tensile strength of the as-sintered (4vol.%Ti5Si3+1vol.%TiBw)/Ti6Al4V composite with two-level network structure reached 1150 MPa and the elongation remained at 6.4% [21]. Through introducing TiB nanowires and TiC particles, the tensile strength of the

RI PT

network-woven structured composite increased to 1267 MPa, while an average strain-to-fracture of 6.1% was retained [34]. For the composite with one-level network structure, the tensile strength of the as-sintered (Ti5Si3+Ti3SiC2+TiC)/Ti6Al4V

SC

composite were enhanced while maintaining sufficient ductility (8.5%) [76], and the elongation of the as-sintered 3vol.%TiBw/Ti6Al4V composite reached 6.5% [55]. The

M AN U

aforementioned superior ductility is attributed to low fractions and unique distribution of reinforcements. However, the strengthening and toughening mechanisms of

EP

TE D

DRTMCs reinforced by reinforcements with low fractions remain elusive.

AC C

Fig. 14 Room-temperature tensile properties of the Ti matrix (a) and Ti6Al4V matrix (b) composites reinforced with different ceramics and fabricated by different routes.

The

near-α

Ti

alloys,

(Ti-5.8Al-3.4Zr-4.0Sn-0.4Mo-0.4Nb-0.4Si-0.06C),

such TA15

as

Ti60

(Ti-6Al-2Zr-1Mo-1V),

VT18U (Ti-6.8Al-4Zr-2.5Sn-1Nb-0.7Mo-0.15Si) alloys, are usually applied in high-temperature fields, since the alloys possess good comprehensive mechanical properties [77-80]. Therefore, adopting them as matrix materials can improve the high-temperature properties of composites. Furthermore, the tensile properties of the 22

ACCEPTED MANUSCRIPT composites fabricated through single-step can be significantly enhanced via subsequent heat treatments and plastic deformation [77, 81]. Zhang et al. reported that the room-temperature tensile strength and elongation of casted TiB/near-α Ti composite increases from 1054 MPa and 0.9% to 1342 MPa and 5.7% [81]. The

RI PT

room-temperature tensile strength and elongation of the as-cast TiB/VT18U composite reached 1240 MPa and 7%, moreover, the tensile strength reached 740 MPa at 600 °C and 500 MPa at 700 °C after forging followed by heat treatments [77].

SC

The as-extruded TiBw/Ti60 composite exhibited excellent high temperature tensile strength, which were close to 1000 MPa at 600 °C and 800 MPa at 700 °C,

M AN U

respectively [79]. For the as-sintered TiBw/TA15 composite after solution and aging treatments, the tensile strength was increased to around 800 MPa at 600 °C and 500 MPa at 700 °C, respectively [80].

Recent studies indicated that the fabrication of DRTMCs need to meet three

TE D

conditions to obtain excellent tensile properties: i) reinforcements with low fraction and nano-scaled level. Slight reinforcements can result in improvement of strength without decreasing ductility drastically. Decreasing the sizes of ceramic phases to

EP

sub-micron or nanoscale level can lead to substantial improvement in mechanical

AC C

performances of TMCs [24, 45, 68]; (ii) unique distribution of reinforcements. Controlling and optimizing each hierarchical microstructure in the composites can further enhance the properties of DRTMCs. However, for the metal matrix composites, it is important to maintain the stability of mechanical properties, and investigate the strengthening and toughening mechanisms thoroughly; (iii) subsequent deformation or heat treatments. Compared with the as-sintered or casted composites, hot or cold working and heat treatments can lead to significant improvement in both tensile strength and ductility. 23

ACCEPTED MANUSCRIPT 4.2 Wear and Fatigue behavior In terms of TMCs applied to aerospace and automotive industries, wear resistance is one of the important properties. Several researchers have reported the

can overcome the poor wear resistance of Ti alloys [82, 83].

RI PT

proper addition of hard ceramic particles, such as TiC, TiB and B4C, with Ti alloys

Choi et al. [82] investigated the wear behavior of in-situ (TiB+TiC)/Ti composites and measured the coefficient of friction (COF) of pure Ti and composites

SC

under an applied load of 1 N, as shown in Fig. 15. The COF values in the steady-state region are similar for all composites with different reinforcement fractions. However,

M AN U

the length of the run-in period increased with increasing fractions, which indicated that the reinforcements are beneficial for enhancing the wear resistance. In addition,

AC C

EP

TE D

the length of the run-in period increased with increasing the size of reinforcements.

Fig. 15 The COF of Ti and (TiB+TiC)/Ti composites with different reinforcement fractions under an applied load of 1N [82].

Recent studies demonstrate that the wear resistance of Ti6Al4V alloy can be remarkably improved by tungsten inert gas (TIG) cladded TiBw/Ti6Al4V composite coating [84]. Fig. 16 displayed the COF of Ti6Al4V alloy (0.48) and composite coating (0.36) after dry sliding wear test. Furthermore, the worn surfaces of alloy and 24

ACCEPTED MANUSCRIPT composite coating were presented [84]. Obvious parallel grooves and wear debris were generated on the surface of alloy. The dominant worn mechanism was microploughing. However, for the composite coating, the size of grooves was

M AN U

SC

mode dominated the worn mechanism of composite coating.

RI PT

decreased by hard TiB reinforcements and refined matrix grains. The abrasive wear

Fig. 16 The COF of Ti6Al4V and TiBw/Ti6Al4V composite coating (a). SEM images of worn surfaces after dry sliding wear tests of Ti6Al4V alloy (b) and TiBw/Ti6Al4V composite coating (c)

TE D

under a load of 0.49 N [84].

Thermal oxidation (TO) is a simple and low cost method, which has been adopted to improve the tribological properties of Ti alloys under high loading

EP

conditions [17]. For example, Dalili et al. [17] enhanced the wear resistance of TiC/Ti6Al4V composite at 2-10 N through TO method. Owing to the hard oxide scale

AC C

and diffused zone, not only the load-bearing capacity of TMCs was improved, but also the abrasive effect of TiC particles and the extensive plastic deformation of Ti matrix was blocked.

The fatigue resistance of TMCs can be influenced by intrinsic and extrinsic

factors. Intrinsic factors include volume fraction of reinforcements and matrix microstructure. Extrinsic parameters contain stress ratio, mean stress, temperature and environment [3]. In generally, TMCs exhibit higher fatigue resistance, when compared with Ti alloys, owing to the better mechanical properties of composites. 25

ACCEPTED MANUSCRIPT Quast et al. [85] examined the microstructure and high-cycle sonic fatigue behavior of Ti-TiB functionally graded materials. The regions with high TiB fractions showed brittle fracture characteristics, in contrast, the Ti-rich layers showed ductile failure characteristics. Chen et al. reported the fatigue S-N curves of the Ti6Al4V alloys with

RI PT

B addition at 455 oC, as shown in Fig. 17 [86]. Under the same conditions, the Ti6Al4V alloys with 0.1wt.%B addition exhibited longer average fatigue lives than those of Ti6Al4V alloys. The results were related to the refined grain size and the

SC

addition of strong TiB reinforcement. As shown in Fig. 18a, a large amount of cracks presented in the α phases, while cracks were rarely observed within the fine β phases.

M AN U

For the Ti alloys with B element, TiB cracks were evident within the deformed gauge sections (Fig. 18b and 18c), which inferred that TiB can bear higher load and hinder

EP

TE D

dislocation motion [86].

AC C

Fig. 17 Maximum applied stress versus cycles to failure curves where runout specimens were marked with an arrow [86].

26

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 18 BSE micrographs of (a) Ti6Al4V alloy failed after 44,153 cycles at a maximum stress of 350 MPa; (b) Ti6Al4V0.1B alloy failed after 990,554 cycles at a maximum stress of 400 MPa; (c)

M AN U

Ti6Al4V1B alloy failed after 32,718 cycles at a maximum stress of 350 MPa [86].

However, reinforcements may have no effect or even deteriorate the fatigue properties of TMCs. Majumdar et al. [87] studied the fatigue behavior of in-situ TiBp/Ti35Nb5.7Ta7.2Zr (β-type, TNTZ) composite and found that TiB particles have

TE D

no substantial influence on the yield strength and the fatigue strength of TNTZ alloy. Xie et al. [88] investigated the effect of reinforcements on rolling contact fatigue (RCF) behavior of the (TiB+TiC)/Ti6Al4V composite. The RCF life of the composite

EP

was decreased, owing to more stress concentration distributed around the reinforcements and accelerated flaking and fracture. Hence, decreasing the stress

AC C

concentration is important for improving the fatigue resistance of TMCs. 4.3 Oxidation resistance Factors that affecting the oxidation behavior of TMCs can be classified into two

categories, namely the intrinsic characteristics and extrinsic conditions. Intrinsic characteristics are more material-related, such as the type, volume fraction and distribution of reinforcements as well as matrix composition. Extrinsic conditions are mainly referred to testing methods, which include continuous in thermal gravity 27

ACCEPTED MANUSCRIPT analyzer (TGA) [89, 90] or discontinuous in air furnace [58, 91]. For the oxidation behavior of TiBw reinforced Ti matrix composite, Zhang et al. [89] reported that in-situ TiBw reinforcement deteriorated the oxidation via forming evaporable B2O3 at 750-950 oC. Moreover, Hu et al. [32] found that oxidation rate of

RI PT

the TiBw/Ti60 composites is higher than that of the Ti60 alloy, and increases with increasing the TiBw fractions. TiBw bringing in more phase boundaries deteriorated the oxidation of composites. It can be concluded from the aforementioned results that

SC

the oxidation resistance of DRTMCs reinforced by TiBw will be decreased, compared with Ti alloys. However, the oxidation resistance of the TiBw-reinforced composites

M AN U

were remarkably enhanced by compacted TiAl3 coating, which was formed by hot-dipping method [92]. Fig. 19 showed obviously that the oxidation rate of the TiAl3-coated composites significantly reduces, compared with the TiBw/Ti6Al4V composites at 700 oC and 800 oC [92]. The partial spallation of oxide scales on the

TE D

uncoated composites was observed after 30 h at 700 oC and 20 h at 800 oC, respectively. While, no spallation was occurred for the TiAl3-coated composites after oxidation tests for 100 h at 700 oC and 800 oC. The mass gain of the coated

EP

composites was 0.19 mg/cm2 and 1.04 mg/cm2 for 100 h at 700 oC and 800 oC, and

AC C

that of the uncoated composites was 7.21 mg/cm2 and 22.17 mg/cm2, respectively.

Fig. 19 The cycle oxidation kinetic curves of the TiAl3-coated and uncoated TiBw/Ti6Al4V composites at (a) 700 oC; (b) 800 oC [92].

Fig. 20 displayed the corresponding SEM images of the oxidized surface and 28

ACCEPTED MANUSCRIPT cross-section on the TiAl3-coated TiBw/Ti6Al4V composites at 800 oC [92]. The network structured oxide scales were formed after oxidation for 100 h. No spallation of the oxide scale and the oxidation product of Al2O3 with larger size were observed. The oxidized coating kept good adhesion to the composite substrate. Four different

RI PT

layers were formed on the composite substrate. The outmost oxide scale consist of Al2O3 (~33 µm in thickness) could effectively hinder O element into the coating. Furthermore, the interface reaction zone (about 15 µm) including TiAl2 and TiAl layer

SC

was generated, due to the outward diffusion of Ti atom in the composite substrate. Therefore, the TiAl3 coating was consumed by the oxidation reaction in the outmost

AC C

EP

TE D

temperature oxidation process.

M AN U

layer and interface reaction between the coating and the substrate during high

Fig. 20 The surface morphology (a), the cross-section morphology and EDS line scanning (b and c) of the oxidized surface on TiAl3 coating at 800 oC for 100 h [92].

Many literatures reported that TiC and/or REOs can improve the oxidation resistance of TMCs. For example, Qin et al. [91] investigated the oxidation behavior of TiC/Ti composite. The introduction of TiC reinforcement could effectively decrease the oxidation rate and enhance the oxidation resistance of the composite 29

ACCEPTED MANUSCRIPT through refining the TiO2 grains. Zhang et al. [90] also reported the oxidation resistance of in-situ TiCp/Ti6Al composites was improved at 600 oC and 700 oC. Yang et al. [58] compared the oxidation resistance of (TiC+TiB+Nd2O3)/Ti composite with that of TiC/Ti composite at 600 oC. The multiple-reinforced composite exhibited

RI PT

much lower oxidation rate and higher activation energy, which suggested much higher oxidation resistance. The TiCp/Ti6Al4V composite with one-level network structure displayed a superior oxidation resistance in the temperature range of 600-800 oC, due

SC

to the wall-like structure of TiCp formed around the Ti6Al4V particles [16]. Recently, Wei [93] reported that the oxidation resistance of Ti6Al4V alloy and composites

M AN U

reinforced individually with TiCp or TiBw was remarkably improved by the network architecture of TiBw and TiCp reinforcements. Moreover, the interaction between cyclic oxidation and structural evolution could be adjusted by the thermal-stress induced spallation of oxide scales. Fig. 21 displayed the cyclic oxidation kinetics plots

TE D

at 600 oC, 700 oC and 800 oC [93]. In the three composites, the volume fraction of reinforcement is same (5vol.%). At lower temperature of 600 oC, all the mass-gain kinetics approximatively followed parabolic principles, inferring that protective oxide

EP

scales were formed and macro-scale spallation did not occur during cyclic oxidation.

AC C

With increasing temperature to 700 oC, the kinetics of Ti6Al4V alloy changed into a quasi-linear trend. The spallation of the oxide scales on Ti6Al4V alloy was severe. Although the masses of all composites showed a parabolic-like trend from 70 to 100 h, the

oxide

scales

maintained

adhesive

on

the

alloy

substrate

in

the

(TiCp+TiBw)/Ti6Al4V composite. Among all the oxidized materials, hybrid reinforced composite showed the least mass-gain, which was 33.9%, 18.9% and 10.5% lower than Ti6Al4V alloy, TiBw/Ti6Al4V and TiCp/Ti6Al4V composites, respectively. When the temperature increased to 800 oC, the curves of all the four materials 30

ACCEPTED MANUSCRIPT followed a linear law. During the oxidation process, the oxide scales detached for the four materials, where (TiCp+TiBw)/Ti6Al4V composite displayed the least amount of spallation. The above-mentioned comparison illustrated that the oxidation resistance of Ti6Al4V alloy and composites with TiBw or TiCp is improved by the combination

TE D

M AN U

SC

RI PT

of TiCp and TiBw reinforcements.

Fig. 21 Cyclic oxidation kinetics plots of the Ti6Al4V alloy and three different composites at (a)

EP

600 oC; (b) 700 oC; (c) 800 oC [93].

Fig. 22 [93] showed the calculated results of thermal stress within alloy and

AC C

composites during cyclic oxidation process. The thermal stresses within the Ti6Al4V alloy and three different composites increased linearly with increasing temperature. At 873 K (corresponding to the temperature difference ∆T=575 K), the thermal stresses in all the four materials were lower than the failure critical value (σc), illustrating that protective oxide scales formed on the substrates. This result was agreement with the kinetics results in Fig. 21a. At 973 K, the value of thermal stress within Ti6Al4V alloy increased beyond the σc value, inferring that cracks or spallation tend to form in oxide scales. However, the thermal stresses in the composites were close to the failure stress. 31

ACCEPTED MANUSCRIPT The oxide scales on composites still played a role in inhibiting oxidation process to some extent. The oxidation resistance of composites were higher than that of alloy. At 1073 K, the thermal stresses in all four materials exceed their threshold values,

M AN U

SC

mass-gain kinetics as depicted in Fig. 21c.

RI PT

obvious macro-scale spallation occurred on the oxide scales, leading to the linear

Fig. 22 The relationship between thermal stress versus temperature difference during cyclic oxidation [93].

Furthermore, metal carbide, such as Mo2C and silicide are beneficial for improving the oxidation resistance of Ti alloys. Liu et al. [94] investigated the

TE D

oxidation behaviors of Mo2C/Ti composites at 600 oC and 700 oC. The addition of Mo2C could improve the oxidation resistance of composites, due to the formation of a

EP

dense and protective oxidation layer. Silicide coatings were usually deposited on the surfaces of Ti alloys, owing to their excellent oxidation resistance at high temperature

AC C

[95].

High oxidation resistance of DRTMCs is one of the crucial properties in the

aspects of high-temperature application. Some conclusions were deduced from the past investigations, such as oxidation products, oxide scales and anti-oxidation mechanisms. In order to improve the oxidation resistance of TMCs, two points are essential, on the one hand, the formation of continuous, dense and thermodynamically stable oxide scales; on the other hand, the strong bonding between oxide scales and Ti alloys. 32

ACCEPTED MANUSCRIPT 4.4 Creep resistance Creep behaviors and creep mechanisms of TMCs have been reported in literatures, and the creep resistance of composites was improved by reinforcements. Ma et al. [96] investigated the creep behaviors of TiCp/Ti6Al4V composites with a homogeneous

RI PT

microstructure at 550-650 oC and reported that the presence of three stress regions and threshold stress, and different creep mechanisms for different stress regions. A previous literature showed that the addition of B enhanced both elevated temperature strength

SC

and creep properties of Ti6Al4V alloy [97]. Boehlert [98] also evaluated the creep deformation of the extruded Ti6Al4V1B(wt.%) alloy and found that the addition of B

M AN U

element decreased the steady-state creep rate and improved the creep resistance of Ti6Al4V alloy. However, the extruded alloy showed similar creep rates with that of the as-cast Ti6Al4V1B alloy. Furthermore, Imayev et al. enhanced the creep resistance of VT18U alloy by introducing the TiB and TiC reinforcements. The creep

TE D

rupture time of the (TiB+TiC)/VT18U composites exceeded 50 h under the conditions of 600 oC and 300 MPa. The refined and randomly oriented TiB whiskers was beneficial for the improvement in creep resistance of composites [77].

EP

Wang et al. [33] significantly enhanced the creep resistance of in-situ

AC C

TiBw/Ti6Al4V composite by network architecture design. Compared with the traditional TMCs with homogeneous microstructure [96], the improvement in creep resistance of the TiBw/Ti6Al4V composite with one-level network structure is much more evident. Fig. 23 presented the typical creep curves of the Ti6Al4V alloy and TiBw/Ti6Al4V composites with different TiBw fractions. The steady-state creep rates were decreased drastically with increasing TiBw fractions and network sizes (Ti6Al4V powder size), which caused by the increased local fractions of the reinforcement in the TiBw-rich network boundary. Although the two TiBw/Ti6Al4V 33

ACCEPTED MANUSCRIPT composites (65 µm and 150 µm) had equal overall fractions, the local fraction of

RI PT

TiBw was different. Larger network size could lead to higher TiBw fraction.

SC

Fig. 23 Creep curves of the Ti6Al4V alloy and 3vol.%, 5vol.%, 8vol.%TiBw/Ti6Al4V composites at 600 oC and 200 MPa: (a) network size of 150 µm; (b) network size of 65 µm [33].

M AN U

Fig. 24 showed the stress exponent, activation energy and creep data compensation results of the casted (TiB+TiC+La2O3)/Ti composites after hot forging and heat treatments, in the low stress region and high stress region, respectively [99]. The stress exponents were about 2.3 in the low stress region and 5.1 in the high stress

TE D

region of the composites (Fig. 24a). In the high stress region, the activation energy of the composite were 434-448 kJ/mol. However, the activation energy of the composite (247 kJ/mol) is lower than that of the matrix alloy in the low stress region (Fig. 24b).

EP

The values of activation energy and stress exponent of the composites in the high stress region indicated that the creep mechanism is dislocation climb. The threshold stresses

AC C

for the composites could explain the higher stress exponents by introducing the stress exponent of the pure metal matrix, which was commonly used in the discontinuously reinforced titanium matrix composites. The threshold stresses in the high stress region were much higher than that in the low stress region, as shown in Fig. 24c and 24d.

34

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 24 The stress exponent results (a) and activation energy results (b) of the matrix alloys and the (TiB+TiC+La2O3)/Ti composites; the creep data compensation of the composites with threshold stresses: low stress region (c) and high stress region (d) [99].

In order to further improve the creep resistance of TiBw/Ti6Al4V composite,

TE D

nano-Ti5Si3 particles and two-level hierarchical structure were introduced to the composite [21]. The rupture time of the as-sintered (Ti5Si3+TiBw)/Ti6Al4V composite (73.5 h) was more than 6 times longer than that of the as-sintered Ti6Al4V

EP

alloy at 600 oC/200 MPa. The steady-state creep rate also underwent a significant decrease from 7.28×10-7 s-1 to 2.59×10-7 s-1, once compared with that of

AC C

8vol.%TiBw/Ti6Al4V composite with solely first-scale network [33]. Fig. 25 showed the microstructure of the (Ti5Si3+TiBw)/Ti6Al4V composite after creep tests at 600 o

C and 200 MPa and 250 MPa [21]. When the TiBw reinforcement was added, the

cracks formed around the TiBw reinforcement and a part of TiB whiskers fractured and deboned (Fig. 25a), which was due to the deformation incompatibility between matrix and reinforcement. Compared with the tensile process, the creep is a long and slow process, therefore, the phenomenon of deformation incompatibility is obvious. However, no cracks were observed for the nano-Ti5Si3 particles distributed within the 35

ACCEPTED MANUSCRIPT β phases and at the α/β interfaces, due to the fine characteristics (Fig. 25b). The reason of the improved creep resistance for the (Ti5Si3+TiBw)/Ti6Al4V composite was as follows: the TiBw reinforcement distributed around the Ti6Al4V matrix could block grain boundary sliding; and the Ti5Si3 particles located near the α/β interfaces

TE D

M AN U

SC

impeded the dislocations motion (Fig. 25c and 25d).

RI PT

could hinder phase boundary sliding. Moreover, both TiBw and Ti5Si3 reinforcements

Fig. 25 SEM micrographs of the (Ti5Si3+TiBw)/Ti6Al4V composites after creep tests at 600 o

C/200 MPa (a and b); TEM images of the composites creeping at 600 oC/250 MPa (c and d) [21].

The understanding and tailoring of the creep performance are preconditions for

EP

an effective use of TMCs in structural applications at high temperatures. The creep

AC C

behaviors and creep mechanisms of traditional TMCs with homogeneous reinforcement distributions have been so far extensively investigated [96-98]. However, limited systemic creep researches are available for DRTMCs [21, 99]. Moreover, more analysis of microstructural evolution during creep process should be studied, in order to understand the creep mechanisms of composites deeply.

5. Modeling and Calculation Modeling of TMCs is a challenging work, mainly due to the complexity of the multi-component system. Hence, few theoretical studies on TMCs were reported in 36

ACCEPTED MANUSCRIPT the available literatures. However, there are increasing interests in the computational modeling of TMCs to investigate the structure, thermodynamics, kinetics, and mechanical properties. 5.1 Analytical model

RI PT

Wang et al. [2] evaluated the strengthening effects of TiB whiskers and TiC particles on the casted TMCs by a comprehensive model. The tensile yield strength enhancement was mainly attributed to the grain refinement and solution strengthening.

SC

However, the load-transfer of reinforcements made few contributions to the strength

M AN U

improvement. Hu et al. [10] reported that the yield strength of TiB-Ti composite with a 3D quasi-continuous network microstructure can be modelled by accumulating the following mechanisms: grain refinement, load-transfer of reinforcements and dislocation strengthening.

The TiB reinforcement accumulating at the Ti grain boundaries resulted in the

TE D

grain refinement of matrix, which increased the yield strength, as expressed in equation (1), basing on the Hall-Petch relationship. In the equation, k is the Hall-Petch constant for Ti; and d2 and d1 are the average grain sizes of TiB-Ti composite and Ti

EP

part, respectively.

AC C

∆σ HP = k(

1 1 ) d2 d1

(1)

The increment in yield strength caused by load-transfer of TiB was expressed as

equation (2), σy-Ti is the yield strength of Ti matrix, VTiB is the volume fraction of TiB and l/d is the aspect ratio of TiB. Co is the orientation factor and its value is 0.125, according to the 3D random array model [100].

1 l ∆σ TiB = σ y-TiVTiB C0 2 d

(2)

The dislocation strengthening brought in the increment in yield strength, which 37

ACCEPTED MANUSCRIPT could be described as: ∆σ dis =  ( ∆σ oro ) 2 + ( ∆σ the ) 2 + ( ∆σ geo ) 2 

1/ 2

(3)

In the equation, the ∆σoro is the Orowan stress, which needed for the dislocations to

RI PT

cross the TiB whiskers. The ∆σthe is the stress increment caused by the thermal expansion mismatch between reinforcement and matrix, which could be neglected since the coefficient of thermal expansion of TiB and Ti were similar. The ∆σgeo represents the stress increment due to strain gradient effects resulted from geometrical

SC

distributions of dislocations.

∆σ oro =

M AN U

The Orowan stress was expressed by the Orowan-Ashby equation:

0.13GTibTi ln(d TiB / 2bTi )

(4)

λ

where, GTi is the shear modulus of the Ti matrix; bTi is the Burgers vector of the Ti matrix; and dTiB is the equivalent diameter of TiB reinforcement; λ is the interparticle

TE D

spacing.

The stress increment due to strain gradient effects could be described as equation (5), and ε is the compressive true strain of Ti matrix.

∆σ geo = 0.4GTi VTiBε bTi / dTiB

EP

(5)

AC C

To sum up, the yield strength (σyc) of TiB-Ti composite could be depicted as:

σ yc = σ y-Ti (1 +

∆σ HP

σ y-Ti

)(1 +

∆σ dis

σ y-Ti

)(1 +

∆σ TiB

σ y-Ti

)

(6)

According to the equations (1)-(6), the calculated yield strengths of TiB-Ti

composite fabricated at different levels of laser power were presented in Fig. 26 [10]. Besides, the experimental values of TiB-Ti composite and Ti part were also shown in Fig. 26. The calculated yield strength values were in good agreement with the experimental values of TiB-Ti composite. With increasing the laser power, both the 38

ACCEPTED MANUSCRIPT

RI PT

calculated and experimental yield strength values of TiB-Ti composite were increased.

Fig. 26 The calculated and experimental yield strength values of Ti parts and TiB-Ti composites

SC

fabricated at different laser power [10].

5.2 First-principles calculation

M AN U

First-principles method, which based on the quantum theory, does not rely on any empirical parameters to predict the physical and chemical properties of materials. The method has been widely used in designing materials, owing to the low cost and high efficiency. The first-principles calculation with density functional theory (DFT)

TE D

is an effective method to provide fundamental information of bulks, surfaces and interfaces at atom or even electron levels [101, 102]. Li et al. investigated the interfacial properties of α-Ti(0001)/TiC(111) interface

EP

using first-principles calculation [102]. Three possible stacking sites (center-, hollowand top-sited) and C-terminated TiC(111) were investigated, as shown in Fig. 27a-c

AC C

[102]. The possible negative interface energy indicated the interfacial diffusion, and even new phase formation, was likely to occur across the interface. The largest interfacial fracture toughness was estimated about 4.8 MPa m1/2. The valence electron density and partial density of states (PDOS) (Fig. 27d) demonstrated that its interfacial bonding is mainly contributed from C-Ti covalent bonds and Ti-Ti metallic interaction.

39

SC

RI PT

ACCEPTED MANUSCRIPT

Fig. 27 Three stacking sites of Ti(0001) on C-terminated TiC(111): (a) center, (b) hollow and (c) top sites, the cyan and dark gray balls denote Ti and C atoms, respectively; (d) The partial density

M AN U

of states (PDOS) of C-terminated interface with hollow site stacking. The vertical solid line represents Fermi level [102].

In addition, the first-principles method was used in the aspects of hydrogen embrittlement, oxidation resistance and creep resistance of Ti alloys. Han et al. [103] reported that the presence of different mechanisms and anisotropy for hydrogen

TE D

diffusion along the C axis and in the basal plane of α-Ti. The calculated activation energy was consistent with the experimental data. Moreover, the mechanical

EP

properties of α-Ti were drastically affected by the presence of oxygen impurities. Bhatia et al. [104] found that a 1/6 monolayer of oxygen addition increases the Peierls

AC C

stress by 4 times and reduces the width of dislocation core by 18%. The calculated hardening effect resulted from oxygen and oxygen diffusion barriers were consistent with experiment results. Li et al. [105] provided fundamental understandings on the interactions between Si and other elements (X, such as Mo) in both α-Ti and β-Ti phases. The strong attraction of the alloying atoms to Si increased the solid solubility of Si in Ti alloys, which led to an increase in creep resistance of Ti alloys. Although some previous literatures investigated the mechanical properties of metal matrix composites using modeling and calculation, few studies are about TMCs, 40

ACCEPTED MANUSCRIPT especially the simulation of DRTMCs with inhomogeneous structure using finite element methods. Therefore, the inhomogeneous structure should be gradually and systematically analyzed using numerical methods.

RI PT

6. Summary and Future outlook (1) For the DRTMCs without secondary processing, designing hierarchical and multi-scale architecture may be an effective way to overcome the strength-ductility trade-off. The addition of proper reinforcements is beneficial for improving the

SC

overall properties of DRTMCs.

M AN U

(2) Novel processing strategies, such as additive manufacturing and laser deposition can be developed to prepare the structures with tailored unique distributions of reinforcement, in order to obtain the DRTMCs with optimized mechanical characteristics.

(3) The numerical methods, such as analytical models, first-principles and finite

TE D

element methods, should be used for fundamental research, in order to predicate deformation, explain mechanisms and instruct experiments effectively.

EP

(4) The DRTMCs with excellent properties should be widely developed to replace conventional materials, which can in turn lead to further progress in fundamental

AC C

theories and fabrication techniques of composites.

Acknowledgement

This work is financially supported by the National Key R&D Program of China

(No. 2017YFB0703100), the National Natural Science Foundation of China (NSFC) under Grant Nos. 51671068, 51731009 and 51471063.

References [1] L.J. Huang, L. Geng, H.X. Peng, Microstructurally inhomogeneous composites: Is a 41

ACCEPTED MANUSCRIPT homogeneous reinforcement distribution optimal?, Prog Mater Sci, 71 (2015) 93-168. [2] J. Wang, X. Guo, J. Qin, D. Zhang, W. Lu, Microstructure and mechanical properties of investment casted titanium matrix composites with B4C additions, Mat Sci Eng A, 628 (2015) 366-373. [3] S.C. Tjong, Y.-W. Mai, Processing-structure-property aspects of particulate- and whisker-reinforced titanium matrix composites, Compos Sci Technol, 68 (2008) 583-601.

composites, J Mater Sci, 35 (2000) 241-248.

RI PT

[4] L. Lu, M.O. Lai, H.Y. Wang, Synthesis of titanium diboride TiB2 and Ti-Al-B metal matrix [5] S. Gorsse, D.B. Miracle, Mechanical properties of Ti-6Al-4V/TiB composites with randomly oriented and aligned TiB reinforcements, Acta Materialia, 51 (2003) 2427-2442.

TiB/Ti composite, Mat Sci Eng A, 239–240 (1997) 647-651.

SC

[6] T. Yamamoto, A. Otsuki, K. Ishihara, P.H. Shingu, Synthesis of near net shape high density

[7] C.J. Zhang, F.T. Kong, S.L. Xiao, E.T. Zhao, L.J. Xu, Y.Y. Chen, Evolution of microstructure and tensile properties of in situ titanium matrix composites with volume fraction of (TiB+TiC)

M AN U

reinforcements, Mat Sci Eng A, 548 (2012) 152-160.

[8] F. Ma, S. Lu, P. Liu, W. Li, X. Liu, X. Chen, K. Zhang, D. Pan, W. Lu, D. Zhang, Microstructure and mechanical properties variation of TiB/Ti matrix composite by thermo-mechanical processing in beta phase field, J Alloy Compd, 695 (2017) 1515-1522.

[9] A. Azarniya, A. Azarniya, S. Sovizi, H.R.M. Hosseini, T. Varol, A. Kawasaki, S. Ramakrishna, Physicomechanical properties of spark plasma sintered carbon nanotube-reinforced metal matrix

TE D

nanocomposites, Prog Mater Sci, 90 (2017) 276-324.

[10] Y. Hu, W. Cong, X. Wang, Y. Li, F. Ning, H. Wang, Laser deposition-additive manufacturing of TiB-Ti composites with novel three-dimensional quasi-continuous network microstructure: Effects on strengthening and toughening, Composites Part B: Engineering, 133 (2018) 91-100.

EP

[11] S. Liu, Y.C. Shin, The influences of melting degree of TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites, Mater Design, 136 (2017) 185-195.

AC C

[12] S. Wang, L. Huang, Q. An, L. Geng, B. Liu, Dramatically enhanced impact toughness of two-scale laminate-network structured composites, Mater Design, 140 (2018) 163-171. [13] L.J. Huang, L. Geng, A.B. Li, F.Y. Yang, H.X. Peng, In situ TiBw/Ti–6Al–4V composites with novel reinforcement architecture fabricated by reaction hot pressing, Scripta Mater, 60 (2009) 996-999.

[14] S. Gorsse, Y.L. Petitcorps, S. Matar, F. Rebillat, Investigation of the Young's modulus of TiB needles in situ produced in titanium matrix composite, Mat Sci Eng A, 340 (2003) 80-87. [15] J. Zhu, A. Kamiya, T. Yamada, W. Shi, K. Naganuma, Influence of boron addition on microstructure and mechanical properties of dental cast titanium alloys, Mat Sci Eng A, 339 (2003) 53-62. [16] L.J. Huang, L. Geng, Y. Fu, B. Kaveendran, H.X. Peng, Oxidation behavior of in situ 42

ACCEPTED MANUSCRIPT TiCp/Ti6Al4V composite with self-assembled network microstructure fabricated by reaction hot pressing, Corros Sci, 69 (2013) 175-180. [17] N. Dalili, A. Edrisy, K. Farokhzadeh, J. Li, J. Lo, A.R. Riahi, Improving the wear resistance of Ti–6Al–4V/TiC composites through thermal oxidation (TO), Wear, 269 (2010) 590-601. [18] S.C. Tjong, Recent progress in the development and properties of novel metal matrix nanocomposites reinforced with carbon nanotubes and graphene nanosheets, Materials Science and

RI PT

Engineering: R: Reports, 74 (2013) 281-350. [19] M.E. Maja, O.E. Falodun, B.A. Obadele, S.R. Oke, P.A. Olubambi, Nanoindentation studies on TiN nanoceramic reinforced Ti–6Al–4V matrix composite, Ceram Int, 44 (2018) 4419-4425.

[20] M. Moradi, M. Moazeni, H.R. Salimijazi, Microstructural characterization and failure mechanism of vacuum plasma sprayed Ti-6Al-4V/B4C composite, Vacuum, 107 (2014) 34-40.

SC

[21] Y. Jiao, L.J. Huang, S. Wang, X.T. Li, Q. An, X.P. Cui, L. Geng, Effects of first-scale TiBw on secondary-scale Ti5Si3 characteristics and mechanical properties of in-situ (Ti5Si3+TiBw)/Ti6Al4V composites, J Alloy Compd, 704 (2017) 269-281.

M AN U

[22] S.C. Tjong, Z.Y. Ma, Microstructural and mechanical characteristics of in situ metal matrix composites, Materials Science and Engineering: R: Reports, 29 (2000) 49-113. [23] B.X. Liu, L.J. Huang, L. Geng, B. Wang, C. Liu, W.C. Zhang, Fabrication and superior ductility of laminated Ti–TiBw/Ti composites by diffusion welding, J Alloy Compd, 602 (2014) 187-192.

[24] L. Jiang, H. Yang, J.K. Yee, X. Mo, T. Topping, E.J. Lavernia, J.M. Schoenung, Toughening of

TE D

aluminum matrix nanocomposites via spatial arrays of boron carbide spherical nanoparticles, Acta Mater, 103 (2016) 128-140.

[25] J.C. Wong, M. Paramsothy, M. Gupta, Using Mg and Mg–nanoAl2O3 concentric alternating macro-ring material design to enhance the properties of magnesium, Compos Sci Technol, 69 (2009)

EP

438-444.

[26] H.D. Espinosa, J.E. Rim, F. Barthelat, M.J. Buehler, Merger of structure and material in nacre and bone – Perspectives on de novo biomimetic materials, Prog Mater Sci, 54 (2009) 1059-1100.

AC C

[27] P. Fratzl, R. Weinkamer, Nature’s hierarchical materials, Progress in Materials Science, 52 (2007) 1263-1334.

[28] L. Jiang, Z. Li, G. Fan, D. Zhang, A flake powder metallurgy approach to Al2O3/Al biomimetic nanolaminated composites with enhanced ductility, Scripta Mater, 65 (2011) 412-415. [29] Y.M. Wang, T. Voisin, J.T. McKeown, J. Ye, N.P. Calta, Z. Li, Z. Zeng, Y. Zhang, W. Chen, T.T. Roehling, R.T. Ott, M.K. Santala, P.J. Depond, M.J. Matthews, A.V. Hamza, T. Zhu, Additively manufactured hierarchical stainless steels with high strength and ductility, Nat Mater, 17 (2018) 63-71. [30] G. Liu, G.J. Zhang, F. Jiang, X.D. Ding, Y.J. Sun, J. Sun, E. Ma, Nanostructured high-strength molybdenum alloys with unprecedented tensile ductility, Nat Mater, 12 (2013) 344-350. [31] S. Abkowitz, S.M. Abkowitz, H. Fisher, P.J. Schwartz, CermeTi® discontinuously reinforced 43

ACCEPTED MANUSCRIPT Ti-matrix composites: Manufacturing, properties, and applications, JOM, 56 (2004) 37-41. [32] H. Hu, L. Huang, L. Geng, B. Liu, B. Wang, Oxidation behavior of TiB-whisker-reinforced Ti60 alloy composites with three-dimensional network architecture, Corros Sci, 85 (2014) 7-14. [33] S. Wang, L.J. Huang, L. Geng, F. Scarpa, Y. Jiao, H.X. Peng, Significantly enhanced creep resistance of low volume fraction in-situ TiBw/Ti6Al4V composites by architectured network reinforcements, Sci Rep, 7 (2017) 40823.

RI PT

[34] L. Huang, L. Wang, M. Qian, J. Zou, High tensile-strength and ductile titanium matrix composites strengthened by TiB nanowires, Scripta Mater, 141 (2017) 133-137.

[35] Z. Yang, W. Lu, L. Zhao, J. Qin, D. Zhang, Microstructure and mechanical property of in situ synthesized multiple-reinforced (TiB+TiC+La2O3)/Ti composites, J Alloy Compd, 455 (2008) 210-214.

SC

[36] A. Dorri Moghadam, E. Omrani, P.L. Menezes, P.K. Rohatgi, Mechanical and tribological properties of self-lubricating metal matrix nanocomposites reinforced by carbon nanotubes (CNTs) and graphene – A review, Composites Part B: Engineering, 77 (2015) 402-420.

M AN U

[37] K.S. Munir, Y. Zheng, D. Zhang, J. Lin, Y. Li, C. Wen, Improving the strengthening efficiency of carbon nanotubes in titanium metal matrix composites, Mat Sci Eng A, 696 (2017) 10-25. [38] S. Li, B. Sun, H. Imai, T. Mimoto, K. Kondoh, Powder metallurgy titanium metal matrix composites reinforced with carbon nanotubes and graphite, Composites Part A: Applied Science and Manufacturing, 48 (2013) 57-66.

[39] X.N. Mu, H.M. Zhang, H.N. Cai, Q.B. Fan, Z.H. Zhang, Y. Wu, Z.J. Fu, D.H. Yu,

TE D

Microstructure evolution and superior tensile properties of low content graphene nanoplatelets reinforced pure Ti matrix composites, Mat Sci Eng A, 687 (2017) 164-174. [40] K.S. Munir, D.T. Oldfield, C. Wen, Role of Process Control Agent in the Synthesis of Multi-Walled Carbon Nanotubes Reinforced Titanium Metal Matrix Powder Mixtures, Adv Eng

EP

Mater, 18 (2016) 294-303.

[41] H. Kwon, M. Estili, K. Takagi, T. Miyazaki, A. Kawasaki, Combination of hot extrusion and spark plasma sintering for producing carbon nanotube reinforced aluminum matrix composites,

AC C

Carbon, 47 (2009) 570-577.

[42] L. Zhang, J. Wu, Ti5Si3 and Ti5Si3-based alloys: Alloyingbehavior, microstructure and mechanical property evaluation, Acta Mater, 46 (1998) 3535-3546. [43] M. Sumida, K. Kondoh, In-Situ Synthesis of Ti Matrix Composite Reinforced with Dispersed Ti5Si3 Particles via Spark Plasma Sintering, Materials Transactions, 46 (2005) 2135-2141. [44] L. Liu, J. Xu, P. Munroe, Z.-H. Xie, Microstructure, mechanical and electrochemical properties of in situ synthesized TiC reinforced Ti5Si3 nanocomposite coatings on Ti–6Al–4V substrates, Electrochim Acta, 115 (2014) 86-95. [45] S. Shu, B. Xing, F. Qiu, S. Jin, Q. Jiang, Comparative study of the compression properties of TiAl matrix composites reinforced with nano-TiB2 and nano-Ti5Si3 particles, Mat Sci Eng A, 560 (2013) 596-600. 44

ACCEPTED MANUSCRIPT [46] K.P. Rao, Y.J. Du, In situ formation of titanium silicides-reinforced TiAl-based composites, Mat Sci Eng A, 277 (2000) 46-56. [47] L.J. Huang, S. Wang, L. Geng, B. Kaveendran, H.X. Peng, Low volume fraction in situ (Ti5Si3+Ti2C)/Ti hybrid composites with network microstructure fabricated by reaction hot pressing of Ti–SiC system, Compos Sci Technol, 82 (2013) 23-28. [48] Y. Jiao, L.J. Huang, Q. An, S. Jiang, Y.N. Gao, X.P. Cui, L. Geng, Effects of Ti5Si3 adjustment

on

microstructure

and

tensile

properties

of

in-situ

RI PT

characteristics

(Ti5Si3+TiBw)/Ti6Al4V composites with two-scale network architecture, Mat Sci Eng A, 673 (2016) 595-605.

[49] Z. Zhang, X. Shen, C. Zhang, S. Wei, S. Lee, F. Wang, A new rapid route to in-situ synthesize TiB–Ti system functionally graded materials using spark plasma sintering method, Mat Sci Eng A,

SC

565 (2013) 326-332.

[50] L. Li, J. Wang, P. Lin, H. Liu, Microstructure and mechanical properties of functionally graded TiCp/Ti6Al4V composite fabricated by laser melting deposition, Ceram Int, 43 (2017)

M AN U

16638-16651.

[51] F.-C. Wang, Z.-H. Zhang, Y.-J. Sun, Y. Liu, Z.-Y. Hu, H. Wang, A.V. Korznikov, E. Korznikova, Z.-F. Liu, S. Osamu, Rapid and low temperature spark plasma sintering synthesis of novel carbon nanotube reinforced titanium matrix composites, Carbon, 95 (2015) 396-407. [52] B.X. Liu, L.J. Huang, L. Geng, B. Wang, X.P. Cui, C. Liu, G.S. Wang, Microstructure and tensile behavior of novel laminated Ti–TiBw/Ti composites by reaction hot pressing, Mat Sci Eng A,

TE D

583 (2013) 182-187.

[53] D. Zhang, W. Zhang, J. Gu, T. Fan, Q. Liu, H. Su, S. Zhu, Inspiration from butterfly and moth wing scales: Characterization, modeling, and fabrication, Prog Mater Sci, 68 (2015) 67-96. [54] Y. Jiao, L.J. Huang, L. Geng, X.T. Li, Y.N. Gao, M.F. Qian, R. Zhang, Nano-scaled Ti 5 Si 3

EP

evolution and Strength Enhancement of titanium matrix composites with two-scale architecture via heat treatment, Mat Sci Eng A, 701 (2017) 359-369. [55] L.J. Huang, L. Geng, H.X. Peng, In situ (TiBw+TiCp)/Ti6Al4V composites with a network

AC C

reinforcement distribution, Mat Sci Eng A, 527 (2010) 6723-6727. [56] R.O. Ritchie, The conflicts between strength and toughness, Nat Mater, 10 (2011) 817-822. [57] K. Lu, The Future of Metals, Science, 328 (2010) 319. [58] Z.F. Yang, W.J. Lu, J.N. Qin, D. Zhang, Oxidation behavior of in situ synthesized (TiC+TiB+Nd2O3)/Ti composites, Mat Sci Eng A, 472 (2008) 187-192. [59] X. Tan, Y. Kok, Y.J. Tan, M. Descoins, D. Mangelinck, S.B. Tor, K.F. Leong, C.K. Chua, Graded microstructure and mechanical properties of additive manufactured Ti–6Al–4V via electron beam melting, Acta Mater, 97 (2015) 1-16. [60] B. Zhang, H. Liao, C. Coddet, Selective laser melting commercially pure Ti under vacuum, Vacuum, 95 (2013) 25-29. [61] Y. Zhang, Z. Wei, L. Shi, M. Xi, Characterization of laser powder deposited Ti–TiC composites 45

ACCEPTED MANUSCRIPT and functional gradient materials, J Mater Process Tech, 206 (2008) 438-444. [62] V. Imayev, R. Gaisin, E. Gaisina, R. Imayev, H.J. Fecht, F. Pyczak, Effect of hot forging on microstructure and tensile properties of Ti–TiB based composites produced by casting, Mat Sci Eng A, 609 (2014) 34-41. [63] S. Li, K. Kondoh, H. Imai, B. Chen, L. Jia, J. Umeda, Microstructure and mechanical properties of P/M titanium matrix composites reinforced by in-situ synthesized TiC–TiB, Mat Sci Eng A, 628

RI PT

(2015) 75-83. [64] A. Sabahi Namini, M. Azadbeh, M. Shahedi Asl, Effect of TiB2 content on the characteristics of spark plasma sintered Ti–TiB w composites, Adv Powder Technol, 28 (2017) 1564-1572.

[65] K. Kondoh, T. Threrujirapapong, H. Imai, J. Umeda, B. Fugetsu, Characteristics of powder metallurgy pure titanium matrix composite reinforced with multi-wall carbon nanotubes, Compos

SC

Sci Technol, 69 (2009) 1077-1081.

[66] M. Frary, S. Abkowitz, S.M. Abkowitz, D.C. Dunand, Microstructure and mechanical properties of Ti/W and Ti–6Al–4V/W composites fabricated by powder-metallurgy, Mat Sci Eng A,

M AN U

344 (2003) 103-112.

[67] L.J. Huang, S. Wang, Y.S. Dong, Y.Z. Zhang, F. Pan, L. Geng, H.X. Peng, Tailoring a novel network reinforcement architecture exploiting superior tensile properties of in situ TiBw/Ti composites, Mat Sci Eng A, 545 (2012) 187-193.

[68] L. Huang, M. Qian, Z. Liu, V.T. Nguyen, L. Yang, L. Wang, J. Zou, In situ preparation of TiB nanowires for high-performance Ti metal matrix nanocomposites, J Alloy Compd, 735 (2018)

TE D

2640-2645.

[69] B. Ya, B. Zhou, H. Yang, B. Huang, F. Jia, X. Zhang, Microstructure and mechanical properties of in situ casting TiC/Ti6Al4V composites through adding multi-walled carbon nanotubes, J Alloy Compd, 637 (2015) 456-460.

EP

[70] M.A. Lagos, I. Agote, G. Atxaga, O. Adarraga, L. Pambaguian, Fabrication and characterisation of Titanium Matrix Composites obtained using a combination of Self propagating High temperature Synthesis and Spark Plasma Sintering, Mat Sci Eng A, 655 (2016) 44-49.

AC C

[71] P. Zhou, J.N. Qin, W.J. Lu, D. Zhang, Microstructure and mechanical properties ofin situsynthesised (TiC+TiB)/Ti–6Al–4V composites prepared by powder metallurgy, Mater Sci Tech-Lond, 27 (2013) 1788-1792. [72] P. Qiu, H. Li, X. Sun, Y. Han, G. Huang, W. Lu, D. Zhang, Reinforcements stimulated dynamic recrystallization behavior and tensile properties of extruded (TiB+TiC+La2O3)/Ti6Al4V composites, J Alloy Compd, 699 (2017) 874-881. [73] S. Decker, J. Lindemann, L. Krüger, Synthesis and mechanical properties of TiAl particle reinforced Ti-6Al-4V, Mat Sci Eng A, 674 (2016) 361-365. [74] Y.-J. Kim, H. Chung, S.-J.L. Kang, Processing and mechanical properties of Ti–6Al–4V/TiC in situ composite fabricated by gas–solid reaction, Mat Sci Eng A, 333 (2002) 343-350. [75] Z.-Y. Hu, X.-W. Cheng, S.-L. Li, H.-M. Zhang, H. Wang, Z.-H. Zhang, F.-C. Wang, 46

ACCEPTED MANUSCRIPT Investigation on the microstructure, room and high temperature mechanical behaviors and strengthening mechanisms of the (TiB+TiC)/TC4 composites, J Alloy Compd, 726 (2017) 240-253. [76]

C.

Liu,

L.J.

Huang,

L.

Geng,

Y.

Jiao,

A.

Tang,

In

Situ

Synthesis

of

(TiC+Ti3SiC2+Ti5Si3)/Ti6Al4V Composites with Tailored Two-scale Architecture, Adv Eng Mater, 17 (2015) 933-941. [77] V.M. Imayev, R.A. Gaisin, R.M. Imayev, Microstructure and mechanical properties of near α

RI PT

titanium alloy based composites prepared in situ by casting and subjected to multiple hot forging, J Alloy Compd, 762 (2018) 555-564.

[78] R.A. Gaisin, V.M. Imayev, R.M. Imayev, Effect of hot forging on microstructure and mechanical properties of near α titanium alloy/TiB composites produced by casting, J Alloy Compd, 723 (2017) 385-394.

SC

[79] B. Wang, L.J. Huang, H.T. Hu, B.X. Liu, L. Geng, Superior tensile strength and microstructure evolution of TiB whisker reinforced Ti60 composites with network architecture after β extrusion, Mater Charact, 103 (2015) 140-149.

M AN U

[80] R. Zhang, D.J. Wang, L.J. Huang, S.J. Yuan, Effects of heat treatment on microstructure and high temperature tensile properties of TiBw/TA15 composite billet with network architecture, Mat Sci Eng A, 679 (2017) 314-322.

[81] C. Zhang, F. Kong, S. Xiao, H. Niu, L. Xu, Y. Chen, Evolution of microstructural characteristic and tensile properties during preparation of TiB/Ti composite sheet, Materials & Design (1980-2015), 36 (2012) 505-510.

TE D

[82] B.-J. Choi, I.L.Y. Kim, Y.-Z. Lee, Y.-J. Kim, Microstructure and friction/wear behavior of (TiB+TiC) particulate-reinforced titanium matrix composites, Wear, 318 (2014) 68-77. [83] H. Attar, S. Ehtemam-Haghighi, D. Kent, I.V. Okulov, H. Wendrock, M. Bӧnisch, A.S. Volegov, M. Calin, J. Eckert, M.S. Dargusch, Nanoindentation and wear properties of Ti and Ti-TiB

EP

composite materials produced by selective laser melting, Mat Sci Eng A, 688 (2017) 20-26. [84] Q. An, L. Huang, S. Jiang, X. Li, Y. Gao, Y. Liu, L. Geng, Microstructure evolution and mechanical properties of TIG cladded TiB reinforced composite coating on Ti-6Al-4V alloy,

AC C

Vacuum, 145 (2017) 312-319.

[85] J.P. Quast, C.J. Boehlert, R. Gardner, E. Tuegel, T. Wyen, A microstructure and sonic fatigue investigation of Ti–TiB functionally graded materials, Mat Sci Eng A, 497 (2008) 1-9. [86] W. Chen, C.J. Boehlert, The elevated-temperature fatigue behavior of boron-modified Ti–6Al– 4V(wt.%) castings, Materials Science and Engineering: A, 494 (2008) 132-138. [87] P. Majumdar, S.B. Singh, M. Chakraborty, Fatigue behaviour of in situ TiB reinforced β-titanium alloy composite, Mater Lett, 64 (2010) 2748-2751. [88] L. Xie, Q. Zhou, X. Jin, Z. Wang, C. Jiang, W. Lu, J. Wang, Q. Jane Wang, Effect of reinforcements on rolling contact fatigue behaviors of titanium matrix composite (TiB+TiC)/Ti– 6Al–4V, Int J Fatigue, 66 (2014) 127-137. [89] E. Zhang, G. Zeng, S. Zeng, Effect of in situ TiB short fibre on oxidation behavior of Ti–6Al– 47

ACCEPTED MANUSCRIPT 1.2B alloy, Scripta Materialia, 46 (2002) 811-816. [90] X.N. Zhang, C. Li, X.C. Li, L.J. He, Oxidation behavior of in situ synthesized TiC/Ti–6Al composite, Mater Lett, 57 (2003) 3234-3238. [91] Y. Qin, W. Lu, D. Zhang, J. Qin, B. Ji, Oxidation of in situ synthesized TiC particle-reinforced titanium matrix composites, Materials Science and Engineering: A, 404 (2005) 42-48. [92] X.T. Li, L.J. Huang, S.L. Wei, Q. An, X.P. Cui, L. Geng, Cycle oxidation behavior and

network microstructure, Sci Rep, 8 (2018) 5790.

RI PT

anti-oxidation mechanism of hot-dipped aluminum coating on TiBw/Ti6Al4V composites with

[93] S.L. Wei, L.J. Huang, X.T. Li, Q. An, L. Geng, Interactive effects of cyclic oxidation and structural evolution for Ti-6Al-4V/(TiC+TiB) alloy composites at elevated temperatures, J Alloy Compd, 752 (2018) 164-178.

SC

[94] L. Yanbin, L. Yong, Z. Zhongwei, C. Yanhui, T. Huiping, Effect of addition of metal carbide on the oxidation behaviors of titanium matrix composites, J Alloy Compd, 599 (2014) 188-194. [95] M. Mitoraj-Królikowska, E. Godlewska, Silicide coatings on Ti-6Al-1Mn (at.%) alloy and

M AN U

their oxidation resistance, Surface and Coatings Technology, 334 (2018) 491-499. [96] Z.Y. Ma, R.S. Mishra, S.C. Tjong, High-temperature creep behavior of TiC particulate reinforced Ti–6Al–4V alloy composite, Acta Mater, 50 (2002) 4293-4302. [97] G. Singh, D.V.V. Satyanarayana, R. Pederson, R. Datta, U. Ramamurty, Enhancement in creep resistance of Ti–6Al–4V alloy due to boron addition, Mat Sci Eng A, 597 (2014) 194-203. [98] C.J. Boehlert, The creep behavior of powder-metallurgy processed Ti–6Al–4V–1B(wt.%),

TE D

Materials Science and Engineering: A, 510 (2009) 434-439.

[99] L. Xiao, W. Lu, J. Qin, Y. Chen, D. Zhang, M. Wang, F. Zhu, B. Ji, Creep behaviors and stress regions of hybrid reinforced high temperature titanium matrix composite, Compos Sci Technol, 69 (2009) 1925-1931.

EP

[100] X. Guo, L. Wang, M. Wang, J. Qin, D. Zhang, W. Lu, Effects of degree of deformation on the microstructure, mechanical properties and texture of hybrid-reinforced titanium matrix composites, Acta Mater, 60 (2012) 2656-2667.

AC C

[101] L. Liu, S. Wang, H. Ye, First‐principles study of metal/nitride polar interfaces: Ti/TiN, Surf Interface Anal, 35 (2003) 835-841. [102] J. Li, Y. Yang, G. Feng, X. Luo, Q. Sun, N. Jin, Adhesion and fracture toughness at α-Ti(0001)/TiC(111): A first-principles investigation, Appl Surf Sci, 286 (2013) 240-248. [103] X.L. Han, Q. Wang, D.L. Sun, T. Sun, Q. Guo, First-principles study of hydrogen diffusion in alpha Ti, Int J Hydrogen Energ, 34 (2009) 3983-3987. [104] M.A. Bhatia, X. Zhang, M. Azarnoush, G. Lu, K.N. Solanki, Effects of oxygen on prismatic faults in α-Ti: a combined quantum mechanics/molecular mechanics study, Scripta Mater, 98 (2015) 32-35. [105] Y. Li, Y. Chen, J.R. Liu, Q.M. Hu, R. Yang, Cooperative effect of silicon and other alloying elements on creep resistance of titanium alloys: insight from first-principles calculations, Sci Rep, 6 48

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

(2016) 30611.

49

ACCEPTED MANUSCRIPT The manuscript entitled “Progress on discontinuously reinforced titanium matrix

AC C

EP

TE D

M AN U

SC

RI PT

composites” is a review article. The data is unavailable.

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT