Processing, microstructure and ageing behavior of in-situ submicron TiB2 particles reinforced AZ91 Mg matrix composites

Accepted Manuscript Processing, microstructure and ageing behavior of in-situ submicron TiB2 particles reinforced AZ91 Mg matrix composites Peng Xiao, Yimin Gao, Xirong Yang, Feixing Xu, Cuicui Yang, Bo Li, Yefei Li, Zhiwei Liu, Qiaoling Zheng PII:

S0925-8388(18)32089-9

DOI:

10.1016/j.jallcom.2018.05.351

Reference:

JALCOM 46327

To appear in:

Journal of Alloys and Compounds

Received Date: 7 April 2018 Revised Date:

28 May 2018

Accepted Date: 29 May 2018

Please cite this article as: P. Xiao, Y. Gao, X. Yang, F. Xu, C. Yang, B. Li, Y. Li, Z. Liu, Q. Zheng, Processing, microstructure and ageing behavior of in-situ submicron TiB2 particles reinforced AZ91 Mg matrix composites, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.05.351. 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 Processing, microstructure and ageing behavior of in-situ submicron TiB2 particles reinforced AZ91 Mg matrix composites Peng Xiaoa*, Yimin Gaoa∗, Xirong Yangb, Feixing Xua, Cuicui Yanga, Bo Lia, Yefei Lia, Zhiwei Liua, Qiaoling Zhenga

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a State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University, No. 28, Xianning West Road, Xi’an, Shaanxi Province 710049, P.R. China

b School of Metallurgy Engineering, Xi’an University of Architecture and Technology,

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Xi’an, Shaanxi Province 710055, P.R. China Abstract

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In this paper, in-situ submicron TiB2 particles (2.5 wt.%) reinforced AZ91 matrix composites are processed by the master alloy method combined with an optimized SHS technique. The effect of TiB2 particles on the microstructure, ageing behavior and mechanical properties of Mg matrix composites is investigated. The results reveal that only TiB2 phase is in-situ formed in the Al-TiB2 master alloy and exhibits the size

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distribution of 100 nm to 700 nm. Compared with the unreinforced AZ91 alloy, the grain size and morphology of Mg17Al12 phase in as-cast composites are much refined. The ageing process of composites is accelerated and the peak-aged time of composites

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decreases from 42 h to 22 h significantly with the introduction of TiB2 particles. Large numbers of fine Mg17Al12 in composites preferentially precipitate at grain boundaries and near TiB2 particles regions. In addition, the hardness, yield strength, ultimate

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tensile strength and elongation of as-cast composites are improved by 16.1%, 53.0%, 26.3% and 25.6%, respectively. After T6 heat treatment, the strength of composites is increased further due to large amounts of submicron Mg17Al12 phase, and the age hardening efficiency of composites is higher than that of AZ91 alloy. Keywords: In-situ; Mg matrix composites; Microstructure; Ageing behavior; Mechanical properties 1. Introduction *

Corresponding author. E-mail address: [email protected] (Yimin Gao), [email protected] (Peng Xiao)

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ACCEPTED MANUSCRIPT Mg-Al series magnesium alloys, due to their low density, high specific strength and fuel efficiency, have received considerable attention as the important lightweight structural materials in the automotive and aerospace industries [1, 2]. To improve the mechanical properties of Mg alloys for widening their application, particulate

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reinforced metal matrix composites (PRMMCs) combined high hardness of ceramics particles with high ductility of metals have attracted great interests because of increased specific elastic modulus, strength and wear resistance [3, 4]. There are many processing methods to fabricate PRMMCs, which can be divided into two types:

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in-situ and ex-situ methods [4, 5], including powder metallurgy, disintegrated melt deposition and stir casting [6-13]. Compared to the conventional ex-situ method, the

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advanced in-situ technique developed to prepare particulate reinforced Mg matrix composites possesses several highlighted advantages: a) the in-situ formed reinforcements are much finer in size; b) the reinforced particles are thermodynamically stable; c) the interfacial bonding between the reinforcements and matrix is strong and the wetting is improved [5].

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In-situ micron TiC (5 µm) particulate reinforced AZ91 matrix composites were fabricated by introducing the Al-TiC master alloy, and the hardness and ultimate tensile strength were improved significantly [14]. Hao et al. [15] investigated the

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effect of in-situ Al2Y particles (6.4~14.2 µm) on the microstructure and mechanical properties of AZ31 alloy in the as-cast and as-rolled condition, and it was found that the addition of micron Al2Y particles resulted in the finer grains and simultaneous

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increase in tensile strength and elongation. Nevertheless, most open literatures in previous reports [16-19] have concentrated on the microstructure and mechanical properties of in-situ micron scale particulate reinforced Mg matrix composites. It is well documented that a few ceramics particles with smaller size would lead to a remarkable influence on the microstructure and mechanical properties of particulate reinforced Mg matrix composites [9, 20, 21]. In Yang et al.’s [22] research of nanosized AlN/Mg composites developed by in-situ casting, the fine grain size of the composite resulted in a significant increase in the ultimate tensile strength and fracture elongation. Furthermore, ageing treatment is a simple and effective technique 2

ACCEPTED MANUSCRIPT to improve the strength of Mg alloy and its composites. Several reports [23-25] have studied the effect of reinforcements on the ageing behavior of Al/Mg matrix composites and found that the addition of reinforcements could change the ageing behavior during the age treatment. Zheng et al. [25] found that although the hardening

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response of SiCw/AZ91 composites was accelerated compared with AZ91 alloy, the age-hardening efficiency of the composite was decreased. Researchers [26] reported that the hybrid reinforcements (B4Cp and SiCw) could accelerate the ageing process and the age-hardening efficiency remained unchanged. Overall, the influence of

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reinforced particles on the ageing behavior of composites is quite complex, which is also a controversial issue. However, little attempt has been made to investigate the

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influence of in-situ submicron particles on the microstructure and mechanical properties of Mg matrix composites. In particular, the study on the ageing behavior of in-situ particles reinforced Mg matrix composites is rather limited [24]. Therefore, it is essential to investigate the processing, microstructure and ageing behavior of in-situ particulate reinforced Mg matrix composites systematically

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What’s more, TiB2 particles with low density, high modulus, high melting point and excellent chemical stability [27, 28] can refine the grain size of Mg alloy due to good lattice matching with magnesium [29, 30]. In-situ TiB2 reinforced Mg/Al matrix

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composites were fabricated successfully by introducing the Al-TiB2 master alloy synthesized by the well-known mixed-salt reaction (i.e. K2TiF6 and KBF4) [27, 31]. However, this mixed-salt method has some drawbacks, such as lower utilization of Ti

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and B salts, fluorine emission and dross disposal, which may cause high processing cost and environmental issues [5]. Owing to low cost, simplicity, high efficiency in time and energy, self-propagating high temperature synthesis (SHS) has been widely used to process PRMMCs [16, 32]. In the present work, in-situ TiB2/AZ91 composites with 2.5 wt.% TiB2 particles was developed by an optimized SHS technique combined with the master alloy method, which can reduce the burning loss of alloying elements and control the content of reinforcements precisely. The aim of this work is to study the effect of in-situ submicron TiB2 particles on the microstructure, ageing behavior and 3

ACCEPTED MANUSCRIPT mechanical properties of AZ91 Mg matrix composites systematically. 2. Experimental procedure 2.1 Raw materials Pure aluminum powders (purity 99.85%, ~40 µm), titanium powders (purity

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99.7%, ~40 µm) and boron powders (purity 96%, ~1 µm) were selected as raw materials. Commercial pure Mg, Al, Zn and Mg-Mn alloy were used to synthesis AZ91 matrix alloy in this work. 2.2 Fabrication of composites

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The fabrication of TiB2/AZ91 composites was composed of two steps as follows. First, the Al, Ti and B powders were mixed well with the ZrO2 ball-to-powder weight

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ratio of 10:1 and milled at 60 r/min for 8 h. The mixed powders with a composition of 50 wt.% Al and a molar ratio of Ti/B=1:2.2 were compacted into a columniform block with a dimension of Φ44 mm × 20 mm at 50 MPa. Subsequently, the block was heated to 880 ℃ and hold for 15 min to complete the SHS process in high vacuum condition of 10-4 Pa. Second, pure Mg, Al, Zn and Mg-Mn alloy ingots were melt in

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the steel crucible and kept at 750 ℃ for 15 min under the protective argon atmosphere. Then, the prepared Al-50 wt.% TiB2 master alloy was introduced into molten Mg alloy and the melt was hold at this temperature until the Al-TiB2 master alloy was

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dissolved completely. Next, the temperature of the melt was decreased to the semi-solid condition (590 ℃) and mechanical stirring was performed at 300 rpm for 20 min to make TiB2 particles distribute more uniformly in Mg alloy matrix. After

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stirring, the composites melt was reheated to 720 ℃ rapidly and then poured into a steel mould preheated to 300 ℃. For comparison, AZ91 matrix alloy was prepared with the same procedures. And the main chemical composition of AZ91 matrix alloy and TiB2/AZ91 composites was detected with X-ray fluorescence spectrometer (XRF, S8 TIGER, Germany) and listed in Table 1. 2.3 Ageing treatment The as-cast samples were solution treated at 410 ℃ for 24 h and followed by quenching in warm water. Then the solution treated samples were immediately aged at 170 ℃ for different time, such as 2 h, 4 h, 6 h, 8 h, 10 h, … , 48 h, and 50 h, and 4

ACCEPTED MANUSCRIPT then air-cooled to room temperature. For simplicity, the peak ageing was called T6 condition. 2.4 Extraction experiment for obtaining TiB2 particles To analyze the size distribution of in-situ formed TiB2 particles quantitatively in

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the Al-TiB2 master alloy, pure TiB2 particles could be obtained via extraction experiment. A small master alloy block about 5 g was added into 10 vol.% hydrochloric acid solution for 12 h to dissolve completely. The mixed supernatant solution was filtered using the quantitative filter paper, washed and vacuum-dried.

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2.5 Mechanical properties test

The Brinell hardness test of samples was carried out using a Brinell hardness

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tester at the load of 250 kgf for a dwell time of 30 s. For the reproducibility, five measurements were calculated to obtain an average hardness. The tensile samples were cut into the dimension with a gauge length of 15 mm and a cross section of 4 ×2 mm2 according to the ASTM-E8 standard. And the tensile test was conducted at the room temperature in MTS Universal Tensile Testing Machine (MTS, CMT5305, USA)

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with a tensile speed of 0.5 mm/min. 2.6 Material characterization

The phase analysis was performed at X-Ray Diffraction (XRD, Bruker D8

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Advance) with Cu Kα radiation at a scan rate of 10°/min in the range of 20°-90°. The microstructures were observed by using optical microscopy (OM, Leica DMI 5000M, Germany), scanning electron microscopy (SEM, SU3500, Japan), field emission

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scanning electron microscopy (FESEM, Zeiss Gemini SEM 500, Japan), electron probe micro-analysis (EPMA, JXA-8100) equipped with wavelength dispersive spectroscopy (WDS) and high-resolution transmission electron microscopy (HR-TEM, JEM-2100 Plus, Japan). These specimens were polished with diamond (2.5 µm) and etched with an etchant composed of 5 vol.% HNO3 + Ethanol solution. The grain size of specimens was measured by the software: Image-Pro Plus. And the measurement of size distribution for pure TiB2 particles diluted with ethanol was conducted by using the Malvern Zetasizer Nano ZS. In addition, the samples for TEM analysis were mechanically grinded, polished to 50 µm thickness foil, punched into a circular disc 5

ACCEPTED MANUSCRIPT with 3 mm diameter and then ion-thinned by Ion Beam Milling (Leica EM RES102, Germany). 3. Results and discussion 3.1 XRD analysis for Al-TiB2 master alloy

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To confirm the phase composition in the Al-TiB2 master alloy, a few powders ground from the master alloy block were used for the XRD test. Fig. 1 shows the XRD pattern of the Al-TiB2 master alloy. It is clear that Al and TiB2 diffraction peaks can be detected, suggesting that only TiB2 phase is in-situ formed in the Al-TiB2

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master alloy after SHS reaction. In other words, the intermediate products such as AlTi, Al3Ti and AlB2 phases are not detected in final products in the Al-TiB2 master

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alloy, which is ascribed to the reason that the SHS reaction is finished completely. For Al-Ti-B system, several possible chemical reactions between the reactants would take place as follows:

Al + 2B → AlB (1) Al + Ti → AlTi (2)

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3Al + Ti → Al Ti (3) Ti + 2B → TiB (4)

To explain the formation of in-situ formed TiB2 particles in the Al-TiB2 master

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alloy further, the change of Gibbs free energy (∆G) of these reactions (1~4) is calculated according to the thermodynamic data [33]. Fig. 2 presents the change of Gibbs free energy (∆G) of possible reactions in Al-Ti-B system as a function of

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different temperatures. From Fig. 2, it can be observed that the changes of Gibbs free energy for possible reactions (1~4) are negative, indicating these reactions would take place possibly in the temperature range of our present study. It is worth noting that TiB2 phase is much more thermodynamically stable than AlTi, Al3Ti and AlB2 phases due to its lowest ∆G. This reveals that the formation of TiB2 phase is the most favorable in Al-Ti-B system based on the thermodynamics consideration, which is well consistent with the result of Fig. 1. 3.2 Microstructure of Al-TiB2 master alloy Fig. 3 shows the microstructure of the Al-50 wt.%TiB2 master alloy. As shown in 6

ACCEPTED MANUSCRIPT Fig. 3 (a), a large number of TiB2 particles with a size of sub-micro scale are embedded in aluminum, which protects the surface of particles and improves the wetting of TiB2 in Mg matrix composites. One important feature should be noted that flaked Al3Ti particles can not be found in the Al-TiB2 master alloy in our study, which

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is different from previous reports [16, 18, 23]. In their research, the unfavorable Al3Ti phase is formed in the master alloy, which is detrimental to the mechanical properties of Mg matrix composites. Fig. 3 (b) is the WDS point analysis of in-situ formed particles (region A marked in Fig. 3 (a)). It can be seen that point A mainly consists of

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Ti and B element and the molar ration of Ti/B is approximately 1/2, confirming that the particle is TiB2 phase further.

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Fig. 4 (a) presents the XRD pattern of the products extracted from the Al-50 wt.%TiB2 master alloy. Only TiB2 diffraction peaks are detected, which reveals that no other undesired compounds appear in the master alloy. Besides, the morphologies of extracted TiB2 particles are mainly composed of hexagonal and rectangular shapes as shown in Fig. 4 (b).The size distribution of TiB2 particles is shown in Fig. 5 and the

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related data is embedded in the right part. It is evident that the size of all TiB2 particles ranges from 100 nm to 700 nm, indicating that the in-situ formed TiB2 particles exhibit a submicron scale size.

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3.3 Microstructure of TiB2/AZ91 composites Fig. 6 shows the optical micrographs of as-cast AZ91 matrix alloy and 2.5 wt. % TiB2/AZ91 composites. The microstructure of as-cast AZ91 alloy contains primary

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α-Mg and lamellar eutectic β-Mg17Al12 phase with network distribution along grain boundaries as shown in Fig. 6 (a) and (b). From Fig. 6 (c) and (d), it can be seen that TiB2 particles distribute relative homogeneously in the matrix and most of which are located at the grains interior of composites. Compared with the unreinforced AZ91 alloy, the grain size of composites is refined significantly from 112.4 ± 5.8 µm to 58.6 ± 4.1 µm with the introduction of TiB2 particles. In other words, the formation of in-situ TiB2 particles in composites has a significant effect on the grain refinement. According to the Bramfitt’s theoretical model [29, 34], the disregistry (δ) of two-dimensional lattices could be used to estimate the capacity of heterogeneous 7

ACCEPTED MANUSCRIPT nuclei. The nucleation site is effective when the disregistry δ between small index planes of heterogeneous nuclei and nucleating phase is less than 6%, medium effective for 6% < δ < 12% and invalid for more than 12% [30]. Based on the detailed calculation, the smallest two-dimensional disregistry between α-Mg and TiB2 is 5.6%,

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suggesting that TiB2 particles (within the grains of composites) could act as the heterogeneous nuclei of α-Mg grains. What’s more, some TiB2 particles are pushed to the front of solid/liquid interface during the solidification stage, which are distributed along the grain boundaries in composites as shown in Fig. 6 (d) and Fig. 7 (b). Instead

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of acting as the effective heterogeneous nucleating agent, these TiB2 particles can inhibit the growth of α-Mg grains during the solidification. Therefore, above two

composites as shown in Fig. 6.

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factors are contributed to the refinement of α-Mg grains in as-cast TiB2/AZ91

Fig. 7 shows the SEM micrographs of as-cast AZ91 matrix alloy and 2.5 wt. % TiB2/AZ91 composites. The Mg17Al12 phase distributes along the grain boundaries in as-cast AZ91 matrix alloy as shown in Fig. 7 (a). This can be explained as follows.

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Grain boundaries are regarded as higher energy zone with large amount of defects, so Al solute atoms tend to move to grain boundaries to reduce energy. Subsequently, the Mg17Al12 phase is prone to forming in these regions during the solidification. For the

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observation of Fig. 7 (b), it is clear that the size of Mg17Al12 phase in TiB2/AZ91 composites is much smaller than that of AZ91 matrix alloy. As discussed above section, more nucleation sites of the Mg17Al12 phase are provided as a result of much

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more grain boundaries in the composite. Consequently, the precipitated Mg17Al12 phase is refined obviously with the addition of TiB2 particles. Similar phenomenon has been reported by Zhang’s work [35], in which the morphology and size of Al2Ca phase in SiC/Mg-Al-Ca composites are strongly affected by the content and size of SiC particles. The Mg matrix composites reinforced with TiB2 particles, synthesized by in-situ method, were explored by HR-TEM to evaluate its interface characteristic. Fig. 8 presents the TEM morphology of TiB2/AZ91 composites. The TiB2 particle, with the nearly spherical shape and the size of ~100 nm (Fig. 8 (a)), is verified by the typical 8

ACCEPTED MANUSCRIPT selected area electron diffraction (SAED) pattern observed in Fig. 8 (b). It should be pointed out that the morphology (spherical) of tiny TiB2 particles here is different from the typical shape (hexagonal and rectangular) in Fig. 3 (a) and Fig. 4 (b), which probably results from the uncompleted growth of TiB2 crystal [36, 37]. The HR-TEM

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observation performed at the interface of TiB2/AZ91 composites (higher magnification in the blue solid box in Fig. 8 (a)) demonstrates that the interfacial bonding is tight, strong and free of impurities, oxides and voids as shown in Fig. 8 (c), which is different from that of Mg matrix composites prepared by the ex-situ method

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[25]. Besides, the interplanar spacing of TiB2 and Mg matrix is 0.262 nm and 0.245 nm, which is well consistent with interplanar spacing of ideal (1210) plane for TiB2

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crystal and ideal (1213) plane for Mg, respectively. Fig. 8 (d) shows an inverse fast Fourier transform (FFT) micrograph of the region inside the red dashed box in Fig. 8 (c). On the basis of the inverse FFT image, high density dislocations caused by the mismatch in the coefficients of thermal expansion (CTE) between TiB2 (6.4×10 -6 K -1) -6

K

-1

) [17, 27, 28], can be observed, where the red T

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and Mg matrix (26×10

demonstrates the position and orientation of dislocation as shown in Fig. 8 (d). These dislocations near the TiB2 particles would play a significant role in the improvement of mechanical properties and ageing behavior of composites [23, 25].

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3.4 Influence of TiB2 particles on ageing behavior Fig. 9 presents the optical micrographs of AZ91 matrix alloy and 2.5 wt. %

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TiB2/AZ91 composites in as-solutionized condition. After solution treatment, almost all coarse Mg17Al12 phase was dissolved into α-Mg matrix, leading to the formation of supersaturated solid solution of the matrix for both AZ91 alloy and composites. Moreover, it is evident that a majority of TiB2 particles distribute within the grains presented in Fig. 9 (b). The hardness variation of AZ91 matrix alloy and composites corresponding to different ageing times was used to evaluate the influence of TiB2 particles on ageing behavior of composites, as shown in Fig. 10. It can be found that the hardness of composites increases with the ageing time, reaches the peak at 22 h and then 9

ACCEPTED MANUSCRIPT decreases gradually after surpassing the peak ageing time. Similar trend is shown in AZ91 matrix alloy. However, one interesting phenomenon should be noted that the time to reach peak-aged condition for composites (22 h) is significantly reduced compared with the AZ91 matrix alloy (42 h). This indicates that the addition of TiB2

accordance with previous research [23, 24].

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particles can accelerate the age hardening behavior of AZ91 matrix, which is in good

The XRD patterns of the matrix and composites aged for 0 h, 22 h and 42 h are presented in Fig. 11. From Fig. 11 (a), the diffraction peaks of Mg17Al12 phase in

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AZ91 alloy and composites disappear after the solution treatment, agreeing with the observation of Fig. 9. After aged for 22 h as seen in Fig. 11 (b), although the peak

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intensity of Mg17Al12 phase in both AZ91 alloy and composites exhibits an obvious increase, the intensity in composites is much higher than that in AZ91 alloy, indicating that the precipitation rate of Mg17Al12 phase in composites is enhanced. With the increasing ageing time, the intensity of Mg17Al12 phase in AZ91 alloy continues to increase and reaches peak as shown in Fig. 11 (c). However, the content

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of Mg17Al12 phase in composites remains nearly constant with the increase of ageing time. These results further conform that the presence of TiB2 particles greatly reduces the peak-aged time.

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Fig. 12 shows the evolution of precipitation process of AZ91 matrix alloy and composites at different ageing time. According to previous reports [23, 24, 38], the Mg17Al12 phase precipitates with two models during ageing process, including

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discontinuous precipitation with a lamellar-like at grain boundaries and continuous precipitation with a lath-shaped morphology within the supersaturated α-Mg grains. For AZ91 alloy, discontinuous precipitations appear gradually at the grain boundaries after 12 h. With increasing ageing time, large numbers of discontinuous Mg17Al12 precipitates are observed and some continuous precipitations come into being within the grains as shown in Fig. 12 (b) and (c). Compared with AZ91 matrix alloy, the precipitation process of Mg17Al12 in composites is enhanced obviously as seen in Fig. 12 (d), (e) and (f). For the same ageing time, the amounts of precipitated Mg17Al12 phase in composites are much higher than that of AZ91 alloy, especially in the earlier 10

ACCEPTED MANUSCRIPT ageing period. Fig. 13 shows the SEM micrographs of T6-AZ91 matrix alloy (aged for 42 h) and T6-TiB2/AZ91 composites (aged for 22 h). Numbers of Mg17Al12 precipitates are observed at the grain boundaries and a few within the grains as shown in Fig. 13 (a)

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and (b). However, the ageing behavior of composites seen in Fig. 13 (c) and (d) is different from that of AZ91 alloy. It can be found that Mg17Al12 phase seems to preferentially precipitate near the TiB2 particles region, which are located at the interior of α-Mg grains. At the same time, a few Mg17Al12 precipitates appear in the

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TiB2 particles free region (far away from particles region). This indicates it is much easier for Mg17Al12 phase to nucleate and precipitate near TiB2 particles than that at

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the grain boundaries as shown in Fig. 13 (c) and (d). In other words, the addition of TiB2 particles in composites offers more nucleation sites for the precipitation and thus accelerates the ageing process.

The effect of in-situ TiB2 particles on the ageing behavior of composites can be explained with the following several possible reasons. It is well reported that the

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precipitate behavior of Mg17Al12 phase in Mg-Al alloy, in nature, is mainly controlled by the diffusion rate of Al atoms and the nucleation of Mg17Al12 phase [24, 25, 39]. Firstly, the grain size of composites is reduced compared with the unreinforced matrix

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alloy. Therefore, the number of grain boundaries is increased accordingly, which improves the nucleate rate of precipitated phase owing to the reduction of activation energy for heterogeneous nucleation [38]. Secondly, plenty of interfaces between TiB2

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reinforcements and Mg matrix are generated in composites, especially in the TiB2 particles clusters region. Thereby, these interfaces with higher energy can offer more diffusion paths for Al atoms and more favorable nucleation sites for the precipitate of Mg17Al12 phase [24], which is contributed to the formation of precipitates distributing in the vicinity of TiB2 reinforcements. Thirdly, high density dislocations are generated near the TiB2 reinforcements as presented in Fig. 8 (d). On the one hand, these dislocations can enhance the diffusion rate of solute Al atoms and reduce the activation energy for heterogeneous nucleation of precipitates at the age process [23-25], which leads to forming fine Mg17Al12 precipitates near TiB2 particles as 11

ACCEPTED MANUSCRIPT shown in Fig. 13 (c) and (d). On the other hand, high density dislocations formed near the TiB2 particles have a significant effect on improving the strength of composites according to related research [9, 17, 27] and the strengthening mechanism would be analyzed in next section.

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Fig. 14 (a) (b) and (c) (d) show the bright field TEM micrographs and corresponded SAED pattern of precipitates in T6-AZ91 matrix alloy and TiB2/AZ91 composites, respectively. A large number of submicron precipitates are observed in both T6-AZ91 matrix alloy and composites, which is identified to be Mg17Al12 phase

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according to the SAED as presented in Fig. 14 (b) and (d). In addition, the morphology of precipitates mainly exhibits rod-like with a length of 100-400 nm and

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a width of 100 nm, which is the typical feature of the precipitation in the aged AZ91 alloy [25, 39]. However, the size of Mg17Al12 precipitates in composites is smaller than that of AZ91 alloy. As discussed above, the number of nucleation sites for the precipitate of Mg17Al12 phase in composites is increased with the addition of TiB2 particles, leading to the refinement of the Mg17Al12 precipitates.

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3.5 Mechanical properties of TiB2/AZ91 composites

Fig. 15 shows the tensile engineering stress-strain curves of AZ91 matrix alloy and TiB2/AZ91 composites in as-cast and peak ageing (T6) conditions. And the

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mechanical properties including Brinell hardness, yield strength (YS), ultimate tensile strength (UTS) and elongation are summarized in Table 2. It can be seen clearly that the hardness, YS, UTS and elongation of as-cast composites are 72 HB (increased by

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16.1%), 127 MPa (increased by 53.0%), 211 MPa (increased by 26.3%) and 5.4% (increased by 25.6%) compared with AZ91 alloy, respectively. Deng et al. [40] fabricated the submicron size SiC/AZ91 composites using the stir casting, and found the addition of submicron particles could have a great effect on the tensile strength of AZ91 matrix composites. This is in good agreement with our results that submicron TiB2 particles lead to a significant improvement in hardness and strength of composites. It has been reported that the improvement of strength in PRMMCs is generally related

to

the

Hall-Petch

strengthening, 12

CTE

mismatch

strengthening,

ACCEPTED MANUSCRIPT load-transferring strengthening and Orowan strengthening [22, 27, 30, 41-43]. Firstly, the grain size of metal matrix composites is much refined with the addition of ceramic particles [22, 27, 30]. As discussed above, grain refinement of composites due to the presence of TiB2 particles generates more grains and grain boundaries, which would

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impede the motion of dislocations. This contributes to the improvement of the hardness and tensile strength in our work according to the Hall-Petch relationship [27, 30]. In addition, plenty of residual plastic strains are generated near the TiB2 particles in composites due to the difference of CTE between Mg matrix and TiB2

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reinforcements when the composites melt is cooled from the processing temperature to the room temperature. Meanwhile, high-density dislocations are formed around

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TiB2 particles to accommodate the mismatch of CTE as observed in Fig. 8, which would make it more difficult for further plastic deformation and then lead to the enhanced strength of composites [27, 41]. What’s more, the load applied in the metal matrix can be effectively transferred to the hard ceramic particles during the plastic deformation, which also contributes to the improved strength in composites [22, 42].

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At the same time, the strong interfacial bond between the Mg matrix and TiB2 reinforcement is beneficial to the capacity of the loading bearing and has a positive effect on the load transfer strengthening. At last, the fine second-phase particles would

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interact with the motion of dislocations in the matrix when the size of particles is smaller than 1 µm according to the Orowan strengthening [41, 43]. In our present research, all of TiB2 particles show a small size less than 1 µm, which would inhibit

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the motion of the dislocations during the tensile deformation and thus improve the strength of composites.

After T6 heat treatment, the hardness and strength of both AZ91 matrix and

composites are improved further due to the large amounts of Mg17Al12 precipitates, which are mainly responsible for the age hardening [24,39]. Besides, it should be especially emphasized that the increase of UTS in composites is 74 MPa (improved by 35%), which is higher than that 51 MPa (improved by 31%) in AZ91 matrix alloy in the peak ageing condition, as shown in Fig. 15 and Table 2. This reveals that the age hardening efficiency in composites is increased with the addition of TiB2 particles 13

ACCEPTED MANUSCRIPT in this study. What’s more, the elongation to fracture of composites is slightly higher than that of AZ91 matrix alloy in both as-cast and T6 conditions. That is to say, the strength of composites is improved significantly without the sacrifice of ductility compared with

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the unreinforced AZ91 matrix alloy. The grain refinement, uniform distribution of fine TiB2 particles and strong interfacial bond are beneficial to the ductility of composites [7, 19, 22]. 4. Conclusions

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In the present work, a developed master alloy method is proposed to process in-situ submicron TiB2/AZ91 composites by adding the Al-TiB2 master alloy into

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molten Mg alloy. The microstructure of the Al-TiB2 master alloy is analyzed, and the microstructure, ageing behavior and mechanical properties of TiB2/AZ91 composites are investigated in detail. Several significant conclusions are summarized as follows: 1.

Fine TiB2 particles are in-situ formed during the SHS process of Al-Ti-B system.

The morphologies of TiB2 particles in the Al-TiB2 master alloy are mainly composed

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of hexagonal and rectangular shape and TiB2 particles have the size distribution ranging from 100 nm to 700 nm. 2.

In the as-cast state, the grain size and the morphology of Mg17Al12 phase in

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composites are refined evidently compared with the unreinforced AZ91 alloy, which is due to the presence of TiB2 particles. 3.

HR-TEM analysis indicates that large amounts of dislocations are generated near

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TiB2 particles and the interface between reinforcements and Mg matrix is bonded well. 4.

The peak-aged time of composites is significantly reduced from 42 h to 22 h

compared with AZ91 alloy, indicating that the precipitation of Mg17Al12 phase is accelerated with the introduction of TiB2 particles. Numbers of fine Mg17Al12 in composites preferentially precipitate at grain boundaries and near TiB2 particles regions due to the grain refinement and high density dislocations. 5.

The mechanical properties of as-cast composites are improved significantly,

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Acknowledgments This work was supported by the Science and Technology Project of Guangdong Province in China (grant number 2015B010122003, 2015B090926009) and the Science and Technology Project of Guangzhou City in China (grant number

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201604046009, 7060024148074). We thank Mr Ren Zijun and Miss Li Jiao at Instrument Analysis Center of Xi'an Jiaotong University for their assistance with SEM

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and HR-TEM analysis. References

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ACCEPTED MANUSCRIPT Figure captions: Fig. 1. XRD pattern of the Al-TiB2 master alloy Fig. 2. The change of Gibbs free energy (∆G) of possible reactions in Al-Ti-B system as a function of different temperatures

of region A marked in (a).

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Fig. 3. (a) Microstructure of the Al-50 wt.%TiB2 master alloy; (b) EDS point analysis

Fig. 4. (a) XRD pattern of the products extracted from the Al-50 wt.%TiB2 master

Fig. 5. Size distribution of extracted TiB2 particles.

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alloy; (b) the morphologies of extracted TiB2 particles.

TiB2/AZ91 composites.

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Fig. 6. Optical micrographs of (a), (b) as-cast AZ91 matrix alloy and (c), (d)

Fig. 7. SEM micrographs of (a) as-cast AZ91 matrix alloy and (b) TiB2/AZ91 composites.

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Fig. 8. (a) Bright-field TEM morphology of TiB2/AZ91 composites; (b) Typical SAED pattern of TiB2 particle in (a); (c) High-resolution TEM morphology of the interface in the blue solid box in (a); (d) Inverse fast Fourier transform (FFT)

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micrograph of the region inside the red dashed box in (c). (The red T demonstrates the position and orientation of dislocation).

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Fig. 9. Optical micrographs of solution treated samples: (a) AZ91 matrix alloy; (b) TiB2/AZ91 composites. Fig. 10. The hardness variation of AZ91 matrix alloy and composites corresponding to different ageing times. Fig. 11. XRD pattern of the matrix and composites aged for (a) 0 h, (b) 22 h and (c) 42h. Fig. 12. Optical micrographs of AZ91 matrix alloy and composites at different ageing

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Fig. 14. Bright field TEM micrographs and corresponded selected area electron diffraction (SAED) pattern of precipitates in (a)(b) T6-AZ91 matrix alloy, (c)(d) T6-TiB2/AZ91 composites.

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Fig. 15. Tensile engineering stress-strain curves of AZ91 matrix alloy and TiB2/AZ91

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composites at as-cast and T6 conditions.

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Tables Table 1 Chemical composition of the AZ91 alloy and TiB2/AZ91 composites (wt. %). Al 9.45 9.19

Zn 0.76 0.73

Mn 0.27 0.24

Table 2

Ti 1.86

B 0.84

Mg Bal. Bal.

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Materials AZ91 alloy Composites

Materials

Brinell hardness

YS (MPa)

62±1.95

83±5

As-cast TiB2/AZ91 T6 AZ91 T6 TiB2/AZ91

72±1.81 86±1.47 100±1.54

127±4 129±3 161±5

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UTS (MPa)

Elongation, %

167±12

4.3±0.76

211±10 218±6 285±8

5.4±0.40 3.1±0.45 3.8±0.68

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As-cast AZ91

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ACCEPTED MANUSCRIPT Highlights TiB2/AZ91 composites are processed by an optimized SHS technique.

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Both α-Mg grains and Mg17Al12 phase are refined by TiB2.

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The ageing process of in-situ composites is accelerated.

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The mechanical properties of composites are improved significantly.

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