Interfacial characteristics and normal mechanical strength of a NiTi shape memory alloy fiber reinforced Mg3AlZn (SMAFR-AZ31) composite sheet

Materials Science & Engineering A 765 (2019) 138283

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Interfacial characteristics and normal mechanical strength of a NiTi shape memory alloy fiber reinforced Mg3AlZn (SMAFR-AZ31) composite sheet

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Xiang Li, Fengchun Jiang, Zhenqiang Wang , Chunhuan Guo, Jiandong Wang, Zhongyi Niu, Yunpeng Chang Key Laboratory of Superlight Materials & Surface Technology, Ministry of Education, College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, 150001, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Shape memory alloy Magnesium alloy Composite Microstructure characterization Interface properties

In this paper, a novel NiTi shape memory alloy fiber reinforced Mg3AlZn (SMAFR-AZ31) composite sheet was successfully fabricated by laminate structure design combined with hot pressing method using Mg3AlZn (AZ31) foils and continuous NiTi shape memory alloy fibers. The interfacial microstructure of the SMAFR-AZ31 composite sheet was systematically characterized by a high resolution transmission electron microscopy (FEI-TEM Talos F100) combined with an FEI Helios 600i focused ion beam/scanning electron microscopy (FIB/SEM) system. The normal tensile strength of the SMAFR-AZ31 composite sheet was tested using a novel direct tensile sample configuration with the tensile direction perpendicular to prior AZ31 foils. Results showed that a welldensified AZ31 matrix and uniformly distributed NiTi fiber reinforcements were obtained in the SMAFR-AZ31 composite sheet after hot pressing. A continuous interfacial reaction layer consisting of nanocrystalline-amorphous mixture and a few intermetallic phases were formed around the NiTif/AZ31 interface. The nanocrystalline was identified as dual phase coexistence grains of MgO and TiO2, and the amorphous phase was a mixture of Mg–Ti–O and Mg–O. A newly Ti2Ni intermetallic phase that obeys a specific orientation relationship with NiTi fiber ([421] NiTi// [112] Ti2Ni, (132) NiTi// (111) Ti2Ni with an angle difference of 6.27°) precipitated adjacent to the interface reaction layer. A Ni-rich intermetallic phase identified as AlNi formed at the NiTif/AZ31 interface. Furthermore, it is found that large plastic deformation occurred near the interface reaction region in the AZ31 matrix, which is responsible for the successful embedding of NiTi fibers. Compared with the AZ31 laminate sheet without NiTi fibers, the SMAFR-AZ31 composite sheet possessed a superior normal tensile strength, which is attributed to the formation of the amorphous phase at the NiTif/AZ31 interface.

1. Introduction In recent years, due to the increasingly serious energy crisis and environmental pollution, the demand and development of high performance lightweight structural materials have become particularly urgent. As the lightest metallic structural material, Mg and its alloys have shown a growing interest for weight critical applications, including aerospace, automotive, communication electronics and other structural engineering parts [1–4]. Furthermore, Mg and its alloys are commonly used for vibration and noise control as well as energy absorption devices because of their high damping capacity [5,6]. However, some obvious disadvantages, such as relatively low absolute strength, limited ductility, insufficient stiffness, poor corrosion and creep resistance restrict their extensive applications [1,3,7]. In the last few decades, compositing method has been successfully

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employed to enhance the properties of Mg and its alloys via introducing some reinforcements into the matrix, such as ceramic particles and whiskers [8–12], and metallic reinforcements [13–22]. Shen et al. [8] fabricated a series of Mg matrix composites reinforced with three volume fractions (3, 5 and 10 vol%) of submicron-SiC particles by semisolid stirring assisted ultrasonic vibration method. They found that the ultimate tensile strength and yield strength of the 10 vol% SiCp/AZ31B composites were simultaneously improved. Zheng et al. [10] synthesised an Al18B4O33 w/Mg composite using squeeze casting technique, and found that the tensile yield strength increased by 147% compared to the Mg matrix without Al18B4O33 whisker reinforcement. Although the strength and elastic modulus of Mg and its alloys are effectively increased by the ceramic reinforcements, the ductility decreases sharply due to the brittle phases generated at the interface area or pores appeared within the interfaces [21–23]. For metallic reinforcements, Ni

Corresponding author. E-mail address: [email protected] (Z. Wang).

https://doi.org/10.1016/j.msea.2019.138283 Received 27 June 2019; Received in revised form 12 August 2019; Accepted 12 August 2019 Available online 13 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

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30 HP furnace (see Fig. 1b) and sintered under an appropriate processing condition, shown in Fig. 1c. To avoid the AZ31 foils reacting with the stainless steel punch during the sintering process, the upper and lower surfaces of the stack were covered with graphite papers. Before heating, the vacuum degree of the HP furnace should be controlled below 3 × 10−2 Pa, and the stepped heating mode should be employed to ensure the uniform heat transfer. During sintering, the temperature was firstly raised to 500 °C at a heating rate of 5 °C/min, and held for 40 min under a pressure of 6 MPa. Then, in order to ensure the adequate reaction of the foils with fibers, the stacked materials were heated to 560 °C at a heating rate of 2 °C/min, and held for 480 min. In this step, as the AZ31 foils became soft, the pressure was adjusted down to 2 MPa to prevent large deformation of AZ31. Finally, the stacked foils were cooled down to ambient temperature with the furnace, and the pressure was raised to 6 MPa after the temperature was below 520 °C. To guard against the volatilization of Mg alloy sintered at a high vacuum degree, the vacuum pumps were closed and high-purity argon was filled with the pressure of 300 Pa when the temperature was higher than 200 °C. In order to study the effects of the NiTi fibers on the properties of the SMAFR-AZ31 composite, the AZ31 laminate sheet without NiTi fibers was prepared with the same processing parameters.

[13,14] and Cu [15] were confirmed to be effective in strengthening Mg and its alloys. But, the ductility of composites are also adversely affected and the specific weights of the composite parts are increased due to the high density of these metallic reinforcing agents [21]. Whereas, the addition of pure Ti [16,17] and NiTi [18–22] alloy reinforcements can lead to a minimum reduction of ductility, and even when the contents of reinforcements are low, the ductility is increased to a certain extent. As an extensively used metallic functional material, NiTi alloy has been proven to be an effective reinforcement of Mg and its alloys due to its unique superelasticity and shape memory effect [24]. Li et al. [18] fabricated a NiTip/Mg composite by combining pore-forming technique, powder sintering and pressureless infiltration methods. They found that the NiTip/Mg composite still possessed superelasticity and shape memory effect similar to NiTi alloy. In addition, the composite had superior compressive strength and excellent linear superelasticity over 2%. Aydogmus et al. [22] investigated the room and high temperature compressive properties of interpenetrating NiTip/Mg composites prepared by spark plasma sintering. Their results indicated that by adding 30 vol% NiTi powders, the room temperature compressive strength and yield strength of pure Mg could be increased by 88% and 117%, respectively. Moreover, for case 10 vol% NiTi, the composite exhibited the optimal ductility with the elongation up to 20%. It is well known that the interface between reinforcement and matrix generally plays a vital role in the comprehensive mechanical properties of the composites due to its load transfer effect [25]. To our knowledge, however, the systemic studies on interfacial microstructures characterization and bonding properties of the NiTi/Mg composite have not been reported as far. Consequently, our works mainly focus on the microstructure characterization, phase identification and bonding strength in a NiTi/Mg composite. A continuous NiTi alloy fiber was selected to synthesize the “shape memory alloy fiber reinforced AZ31 matrix (SMAFR-AZ31) composite sheet”. A hot pressing sintering method was applied to fabricate the composite. The microstructure observation and interfacial characteristics analysis were performed by scanning electron microscopy (SEM) and high resolution transmission electron microscope (HRTEM). In order to study the bonding strength of AZ31/AZ31, the normal tensile strength of the SMAFR-AZ31 composite sheet was tested using a novel direct tensile sample configuration [26]. Moreover, the fracture behaviors of the composite sheet was discussed in details.

2.2. Microstructure characterization Microstructure observations of the SMAFR-AZ31 composite sheet was performed on a Hitachi SU-70 field emission scanning electron microscopy (FE-SEM) with an energy dispersive spectroscopy (EDS). The interface microstructures were carefully characterized by FEI Talos F100 × field emission transmission electron microscopy (FE-TEM) coupled with energy dispersive spectroscopy (EDS) at an accelerating voltage of 200 kV. TEM sample was prepared by FEI Helios 600i focused ion beam/scanning electron microscopy (FIB/SEM) system and the liftout technique. 2.3. Normal tensile tests The normal tensile strengths of the SMAFR-AZ31 composite sheet and AZ31 laminate sheet without NiTi fibers were tested using a novel direct tensile sample configuration proposed by Zhou et al. [26]. As detailed illustrated in Fig. 2a, the normal tensile sample has a gauge section with the length of 4 mm including three AZ31/NiTif/AZ31 (or AZ31/AZ31) interfaces, two transition steps and two mutually perpendicular cube-shaped gripping heads with the same dimensions of 15 mm (length) × 9 mm (width) × 4 mm (thickness). Fig. 2b shows the special fixture for fixing normal tensile specimen. The fixture has two mutually perpendicular supports, and the distance between the supports is 9.5 mm. During tensile test, one end of fixture was against the indenter of the universal testing machine and the other end of the fixture connected with the loading surfaces of the normal tensile sample. The two support faces of the fixture transferred the compressive stress to the loading faces, while the gauge portion of the sample was in a tensile stress state.

2. Experimental procedures 2.1. Materials fabrication The shape memory alloy NiTi fiber reinforced AZ31 (SMAFR-AZ31) composite sheet was synthesised via hot pressing (HP) method using Mg3AlZn (AZ31) foils (foil thickness 1 mm) and continuous NiTi fibers (fiber diameter 0.3 mm) as the raw materials. The chemical compositions of the AZ31 foil and NiTi fiber are given in Table 1. Prior to sintering, the AZ31 foils and NiTi fibers were ground with silicon carbide paper to remove surface oxide layers, and then cleaned in alcohol bath using ultrasonic cleaning machine for 20 min to get rid of surface organics. After that, all the foils and fibers were dried rapidly for sintering, and then stacked in an order of “AZ31-NiTi-AZ31 …...AZ31”. The top and bottom layers of the laminate were always AZ31 foils (see Fig. 1a). Afterwards, the stacked materials were placed into the ZRY-60-

3. Results 3.1. Microstructure observation Fig. 3 displays the representative SEM images of the SMAFR-AZ31

Table 1 Chemical compositions of AZ31 and NiTi fiber in the composite. Materials

Compositions (wt%)

Mg3AlZn (AZ31) NiTi fiber

Mg: balance, Al: 2.5–3.5, Zn: 0.6–1.4, Mn: 0.2–1.0, Si:0.08, Ca: 0.04, Cu: 0.01, Fe 0.003, Ni: 0.001. Ti: balance, Ni: 55.83, C: 0.04, N: 0.004–0.005, H: 0.001, other elements < 0.4.

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Fig. 1. (a) Schematic illustration of the stacking of the starting materials; (b) Schematic illustration of fabrication apparatus, and (c) processing parameters.

AZ31 matrix are calculated by the following equations:

fNiTi =

XNiTi ⋅π⋅(rNiTi ) 2 WW ⋅WL

fAZ31 = 1 − fNiTi

(1) (2)

where XNiTi is the numbers of NiTi fibers, rNiTi represents the radius of NiTi fiber, WL and Ww indicate the length and width of the AZ31 matrix, respectively. The volume fractions of NiTi fibers and AZ31 matrix calculated using Eqs. (1) and (2) are 8% and 92%, respectively. 3.2. Interfacial characteristics In order to deeply analyze the interfacial characteristics of NiTif/Mg in the SMAFR-AZ31 composite sheet, TEM observation was carefully conducted. Fig. 4a and b shows the overall TEM image and the schematic diagram of the interface zone between NiTi fiber and AZ31 matrix, respectively. According to the morphological difference, the interface zone can be divided into four parts: (Ⅰ) NiTi fiber region, (Ⅱ) continuous interface reaction layer region, (Ⅲ) interfacial granular phases region, and (Ⅳ) AZ31 matrix region. The high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) and EDS mapping analysis displayed in Fig. 5a–f shows that region I is mainly rich in Ni and Ti, indicating that it is NiTi fiber. The continuous reaction layer i.e. region II is composed of Mg, Ti, O and a trace of Al. The interfacial granular phases in region Ⅲ mainly contain Al and Ni elements (see Fig. 5c and d), which is consistent with the result in our previous study [27]. Next, detailed characterizations of the interface regionsⅠ, Ⅱ, Ⅲ and Ⅳ will be carried out separately. Fig. 6 shows the microstructure characterization of the interface region I. Fig. 6a and d depict the TEM image of the NiTi fiber region near the interface and its corresponding selected area electron diffraction (SAED) pattern, respectively. The left side of this region is

Fig. 2. (a) The dimension of the normal tensile test sample and (b) structure of the special fixture [26].

composite sheet and the AZ31 laminate sheet. As shown in Fig. 3a, the SMAFR-AZ31 composite sheet consists of a well-densified AZ31 matrix and uniformly distributed NiTi fiber reinforcements. It is noteworthy that the interfacial bond between NiTi fibers and AZ31 matrix is excellent without visible voids, and the boundaries of the original AZ31 foils disappear completely after sintering. Moreover, NiTi fibers keep its original circular state, indicating that no excessive reaction occurred between the matrix and fibers during sintering, which is beneficial to maintaining the excellent properties of the fiber. As shown in the magnified illustration in Fig. 3a, inhomogeneous dispersed granular phases form along the interface. For the AZ31 laminate sheet, the original boundaries of AZ31 laminate sheet also disappear completely and the microstructure is compact without defects (see Fig. 3b). Here, the volume fractions of the NiTi fibers reinforcements and

Fig. 3. SEM images of (a) SMAFR-AZ31 composite sheet and (b) AZ31 laminate sheet. The yellow short lines indicate the original boundary positions of AZ31 foils. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) 3

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Fig. 4. (a) Overall TEM image of the interface region between NiTi fiber and AZ31 matrix and (b) the schematic diagram of this region.

composed of MgO and TiO2 particles. It indicates that the interfacial microstructure is consistent top to down along this layer. Fig. 7d shows the TEM magnified image of the interface zone, which contains Mg and O elements only (marked as area “D” in Fig. 4b). From the SAED pattern (see Fig. 7h) and the HRTEM image (see Fig. 7j), it can be concluded that this area also consists of a coexisting microstructure of nanocrystalline-amorphous phases. Nevertheless, the difference is that the MgO nanocrystalline particles distributed within Mg–O amorphous matrix in this area rather than the nanocrystalline (MgO–TiO2)/amorphous Mg–Ti–O structure. In order to accurately identify the phase components of nanocrystallines, the FFT of nanocrystallines are captured in Fig. 7k, l and m. Based on the FFT images, the patterns show the characteristic (200), (020) and (220) diffraction spots of MgO at “L” and “N” areas, and (221), (213) and (012) diffraction spots of TiO2 at “M” area. Therefore, the interfacial nanocrystalline phases are unambiguously confirmed to be MgO and TiO2. According to the noisefiltered IFFT image (see Fig. 7n, o and p), the lattice fringes are regular, indicating no dislocations within the nanocrystalline particles. Fig. 8 shows the TEM images and SAED analysis of Al–Ni phase. The Al–Ni phases in “O” and “P” areas (Fig. 4b) are identified as AlNi phase with a cubic structure. According to the Al–Ni binary phase diagram [28], AlNi intermetallics with a wide range of components of 42–69 at. % Ni is more likely to form congruently from the liquid phase. Sun et al. [29] found that the Al atoms in AZ31 alloy dissolved into the Mg liquid and reacted with the Ni atoms to form AlNi particles between the Ni interlayer and AZ31 matrix. Moreover, AlNi phase was also confirmed

identified as monoclinic structured NiTi phase. Fig. 6b presents the TEM magnified image of the red arrow position in Fig. 6a. It is seen that several irregular precipitated particles are dispersed along the right boundary of NiTi fiber. To accurately determine the crystal structure of the precipitated phases, HRTEM analysis (see Fig. 6c) was performed on the particle marked by the white arrow in Fig. 6b. The Fast Fourier transform (FFT) pattern shown in Fig. 6e indicates that the precipitated particle is Ti2Ni phase with a cubic structure. The corresponding inverse Fast Fourier transform (IFFT) image in Fig. 6f demonstrates that the atomic arrangement of lattice fringes of the Ti2Ni particle is tortuous and irregular, and some screw dislocations can be observed. The SAED pattern (Fig. 6g) collected at the red dashed circular area in Fig. 6b shows that NiTi phase has a specific orientation relationship with the precipitated phase Ti2Ni: [421]NiTi//[112 ]Ti2Ni, (132)NiTi// (111)Ti2Ni with an angle difference of 6.27°. Fig. 7a, b and c display the TEM magnified images of the “A”, “B” and “C” positions in Fig. 4b, respectively. It is worth noting that the interface region Ⅱ is not uniform. The thickness is ~2 μm for top area, and only ~0.1 μm for the bottom zone. But it is continuous from top to bottom along the interface, meaning that the NiTi fiber and AZ31 matrix are bonded well through this reaction layer. Fig. 7i shows the HRTEM image of the “E” area in Fig. 7a. It is seen that the interface region Ⅱ has a coexisting microstructure of nanocrystalline and amorphous phases. The nanocrystalline particles are dispersed uniformly in the amorphous matrix. The multi-crystalline SAED patterns in Fig. 7e, f and g show that the nanocrystalline phases in the three positions are all

Fig. 5. (a)–(f) HAADF-STEM-EDS mapping images of the interface region between NiTi fiber and AZ31 matrix. 4

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Fig. 6. (a) TEM image of the NiTi fiber region near the continuous interface reaction layer region; (b) TEM magnified image of the precipitated particles corresponding to the position marked by red arrow in (a); (c) HRTEM lattice image taken from the white arrow area in panel b; (d) SAED pattern of the NiTi phase; (e) FFT patterns of the precipitated particles taken from red dashed square area in panel c; IFFT image of the precipitated particle; (g) SAED pattern taken from the red dashed circle in panel b. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Fig. 7. (a)–(c) TEM magnified images from top to bottom along the NiTif/AZ31 interface; (d) TEM magnified image of the interface region containing Mg and O elements only; (e)–(h) multi-crystalline SAED patterns obtained from the “E”, “F”, “G” and “H” areas in (a), (b), (c) and (d), respectively; (i) and (j) HRTEM lattice image taken from “E” and “H” areas in (a) and (d), respectively; (k) and (l) FFT patterns of the nanocrystallines obtained from the “L″and“M” areas in (i), respectively; (m) FFT patterns of the nanocrystalline obtained from the “N” area in (j); (n) and (o) IFFT image of the nanocrystallines corresponding to (k) and (l), respectively and (p) IFFT image of the nanocrystalline corresponding to (m). 5

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Fig. 10. The typical normal tensile stress-strain curves of the SMAFR-AZ31 composite sheet and AZ31 laminate sheet.

process. This is mainly responsible for the successful embedding of NiTi fiber into Mg matrix during sintering. Some studies demonstrated that the highly dense dislocation tangles within grain or at grain boundaries tended to occur in Mg alloy matrix after undergoing large plastic deformation [32,33]. In this work, the sintering pressure was reduced to 2 MPa while isothermal holding at 560 °C to prevent the excessive thermal deformation of AZ31 matrix. Thus, the whole composite did not undergo large plastic deformation. The massive amounts of deformed structures of AZ31 matrix around the interface are most likely caused by the hard NiTi fiber embedding into soft AZ31 matrix during hot pressing process. The SAED pattern of AZ31 matrix depicted in Fig. 9b indicates that the matrix is typical closely packed hexagonal structure of Mg. Besides the large numbers of dislocation entanglements distributing at the interface, the AZ31 matrix contains a certain number of stacking faults. Fig. 9c and d shows the HRTEM image of the stacking faults and the IFFT image of the yellow dashed square area of Fig. 9c, respectively. The images demonstrate that the atomic arrangement of lattice fringes of the AZ31 matrix are crooked and show visible distortion and mismatch (see Fig. 9c and d), which proved that the AZ31 matrix adjacent to interface experiences obvious large plastic deformation.

Fig. 8. (a) and (b) TEM images of Al–Ni phase in the upper (“O” area in Fig. 4b) and middle positions (“P” area in Fig. 4b) of the interface, respectively; (c) and (d) SAED patterns of the AlNi phases corresponding to the “O” and “P″areas in Fig. 4b, respectively.

3.3. Mechanical properties Fig. 10 shows the typical normal tensile stress-strain curves of the SMAFR-AZ31 composite sheet and AZ31 laminate sheet, and the tensile Table 2 Normal tensile mechanical properties of the SMAFR-AZ31 composite sheet and AZ31 laminate sheet.

Fig. 9. (a) TEM micrograph of AZ31 matrix with high magnification adjacent to the NiTif-AZ31 interface; (b) SAED pattern of the Mg matrix; (c) HRTEM lattice image taken from the yellow arrow area in (a) and (d) IFFT image of the yellow dashed square area in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Materials

AZ31 laminate sheet

to exist at the Mg/Ni interface in other previous studies [30,31]. Fig. 9a shows the TEM image of AZ31 matrix adjacent to the NiTif/ AZ31 interface (see red dashed circle area in Fig. 4a). It can be clearly observed that the AZ31 matrix has a high density of heterogeneously distributed dislocations. It indicates that the matrix near NiTif/AZ31 interface underwent severe plastic deformation during the hot pressing

SMAFR-AZ31 composite sheet

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Specimen

1 2 3 Average 1 2 3 Average

Normal tensile mechanical properties Ultimate tensile strength, σ UTS (MPa)

Fracture strain, ε (%)

21.9 20.3 19.9 20.7 67.7 45.6 59.5 57.6

0.88 1.11 1.18 1.06 5.1 3.51 4.86 4.49

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original NiTi fibers and AZ31 foils start to break due to extrusion and friction at the interface (see Fig. 11a). When the temperature exceeds the eutectic reaction point (over 506 °C), Ni atoms from NiTi fiber can gradually diffuse into the Mg alloy to form MgNielp. The oxides from the surface of as-received materials start to blend within MgNielp. During the isothermal holding process at 560 °C, the MgNielp containing Mg, Ni, Ti, Al and O atoms gradually surrounds the NiTi fiber and forms a continuous liquid layer along the interface under the sintering pressure (see Fig. 11b and c). As the holding stage progresses, the alloy element Al in AZ31 matrix tends to react with solute Ni element from MgNielp to form the Al–Ni intermetallics in the liquid layer [29]. Due to the consumption of Ni via Al–Ni reaction, liquid phases around the interface solidify gradually and form the solid-phase regions containing Mg, Ti and O elements. The increase of the solid-phase regions hinder the further diffusion of Ni elements into AZ31 matrix, so the MgNielp around the interface gradually disappears (see Fig. 11d). When Al atoms react with Ni atoms completely from the remaining liquid phases, the liquid phases disappear totally and a continuous reaction layer rich in Mg, Ti and O elements is formed (see Fig. 11e). The formation of Al–Ni intermetallics near the reaction layer confirms that the diffusion of Ni atoms at the interface are hindered and do not react with the AZ31 matrix far from the interface. It is well known that Ti prefers to nucleate and precipitate directly from the liquid and separate from Mg during the liquid phase solidification due to the limited mutual solubility and no compounds are formed at any temperature between Ti and Mg elements. However, in this work, the notable phenomenon is that O elements combine with Mg and Ti elements to form a continuous amorphous-nanocrystalline reaction layer. In previous studies [38–40], it was found that a critical level of oxygen is effective in enhancing the glass-forming ability of the materials by increasing the crystallization resistance via suppressing the precipitation of the competitive primary phase and stabilizing the liquid via depressing the liquidus temperature [40]. In the initial stage of MgNielp solidification, the O elements around the interface are mainly from the as-received materials, thus the contents are relatively small. The moderate O contents can resist the crystallization of Ti and Mg, and promote the formation of the single amorphous phase with Ti and Mg elements. With the prolonging of the holding time, the O atoms from the sintering atmosphere diffuse into the interface area, and the excessive O contents will destabilize the liquid and induce formation of the more stable O-containing crystalline phase [40], thus resulting in the formation of the TiO2 and MgO

properties are given in Table 2. The ultimate tensile strength (σUTS) of SMAFR-AZ31 composite sheet is 57.6 MPa, which is much greater than that (20.7 MPa) of AZ31 laminate sheet. Meanwhile, the fracture strain (ε) of SMAFR-AZ31 composite sheet also increases from 1.06% to 4.49%. In addition, it can be seen that there is no obvious plastic deformation for the two composite sheets before fracture.

4. Discussion 4.1. Interface formation mechanism As shown in Fig. 5a–f, the diffusion distance of Ni is larger than that of Ti across the interface. As we known, for the sake of the solid-state diffusion of elements, a certain solid solubility is required for the diffusing atoms in the matrix metal [34]. Due to the limited solid solubility of Ti in Mg (the solubility of Ti in Mg is about 0.08% at 560 °C) and no compounds are generated at the sintering temperature (560 °C) [35], the diffusion rate and distance of Ti in AZ31 matrix are greatly restricted. However, for Ni, as the sintering temperature (560 °C) exceeds the eutectic point (506 °C) of Mg and Ni elements [36], Mg–Ni eutectic liquid phase (MgNielp) is formed around the NiTif/AZ31 interface. The formation of MgNielp significantly accelerates the diffusion of Ni. Therefore, Ni diffusion is farther away from the NiTi fiber than Ti. The different diffusion rates of Ni and Ti elements can lead to the obvious depletion of Ni atoms along the edge of the NiTi fiber boundary as well as the excessive non-equilibrium of element distribution between Ti and Ni atoms in this region. Finally, the decrease in the relative content of Ni elements promotes the nucleation and growth of Ti2Ni phase at the certain areas along NiTi fiber boundary (see Fig. 6b) [25,37]. It is noteworthy that the two-phase eutectic structures (αMg + Mg2Ni) predicted by Mg–Ni binary phase diagram are not observed at the interface, and the granular precipitated particles at the interface are mainly composed of Al and Ni elements, rather than Mg and Ni elements. The absence of two-phase eutectic structures and Mg2Ni phase around the interface imply that the MgNielp has transformed into other phases completely during the isothermal holding process before cooling down below the eutectic temperature. The schematic diagram of the interface formation is shown in Fig. 11. Initially, with the increase of heating temperature, the AZ31 matrix gradually softens, and NiTi fibers are embedded into the AZ31 foils under pressure. Ti–O and Mg–O oxides distributed on the surface of the

Fig. 11. Schematics diagram of the interface formation: (a) NiTi fibers are closely connected with AZ31 matrix under pressure; (b) and (c) Ni, Mg, Ti and Al atoms diffuse to the interface area and form MgNielp containing Mg, Ni, Al, Ti and O elements over 506 °C; (d) Al atoms react with Ni atoms to form AlNi phase, and partially MgNielp solidifies to form Mg–Ti–O amorphous; (e) a continuous Mg–Ti–O amorphous layer is formed at the interface with moderate O contents after Ni atoms are completely consumed in MgNielp; (f) excessive O contents lead to the formation of nanocrystalline TiO2 and MgO.

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Fig. 12. (a) SEM microstructure on upper portion of the fracture surface of the interface tensile sample; (b) SEM fracture surface micromorphology of the upper diffusion bonding region of the AZ31 foil; (c) AlNi precipitated particles in pits formed by debonding of fiber and AZ31 matrix, and the corresponding EDS mapping images are shown in (d); (e) SEM microstructure of the precipitate-free areas in pit corresponding to the yellow rectangle area in panel c; (f) SEM microstructure of the lower part of the interface tensile sample; (g) surface micromorphology of the debonded NiTi fiber, and (h) shows the corresponding EDS mapping image; (i) nanoparticles and NiTi matrix regions corresponding to the yellow rectangle area in panel g; (j) SEM fracture surface micromorphology of the lower diffusion bonding region of the AZ31 foil. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

SMAFR-AZ31 composite sheet are significantly superior to those of AZ31 laminate sheet without NiTi fibers. In order to further study the fracture mode of the composite, the fracture surfaces of the normal tensile fractured sample were observed via scanning electron microscope (SEM) in SE mode. Fig. 12 represents the fracture surfaces analysis of the normal tensile samples. From Fig. 12a and f, it can be

nanocrystalline particles in the Mg–Ti–O amorphous layer (see Fig. 11f).

4.2. Fracture mode analyses and reinforcing mechanism As shown in Fig. 10 and Table 2, the normal tensile properties of 8

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respectively. (2) TEM observation showed that a continuous nanocrystalline-amorphous reaction layer with a width of 0.1–2 μm is formed between NiTi fibers and AZ31 matrix. The cubic Ti-rich Ti2Ni precipitated particles with some screw dislocations are formed at the edge of NiTi fibers due to the excessive non-equilibrium of element distribution between Ti and Ni atoms in this region. Additionally, NiTi phase has a specific orientation relationship with the precipitated phase Ti2Ni: [421]NiTi//[112 ]Ti2Ni, (132)NiTi//(111)Ti2Ni with an angle difference of 6.27°. The AlNi intermetallics particles are formed congruently from the MgNielp adjacent to the continuous nanocrystalline-amorphous reaction layer at the side of AZ31 matrix. (3) Compared with the AZ31 laminate sheet, the σUTS of SMAFR-AZ31 composite sheet is increased from 20.7 MPa to 57.6 MPa. The fracture morphology of NiTif/AZ31 interface is rougher and more tortuous than that of AZ31 foils diffusion bonding interface, which means that the cracks propagation require more extra energy. The Mg matrix adhering to the surface of NiTi fibers indicates that the interfacial amorphous layer bonding between NiTi fibers and AZ31 matrix is stronger than partial neighboring Mg matrix. Consequently it can be concluded that the NiTi fibers can increase the bonding strength of the adjacent AZ31 diffusion interface by the formation of amorphous phase at NiTif/Mg interface.

observed that the failure of samples occurs at the diffusion bonding interface of two adjacent AZ31 foils. As illustrated in Fig. 12a, the upper portions of NiTi fibers are separated from the AZ31 matrix under tensile stress, while the lower portions are well bonded with the matrix (see Fig. 12f). Fig. 12b and j shows that the step-like patterns are distributed at the fracture surface, which indicates the typical cleavage fracture at the AZ31 foils diffusion bonding regions under tensile stress. The pits formed by debonding of fibers and AZ31 matrix were studied in details. It is seen that granular precipitated particles are distributed uniformly on the side of the AZ31 matrix (see Fig. 12c). The SEM-EDS mapping analysis (Fig. 12d) shows that the granular precipitated particles mainly consist of Al and Ni elements, which corresponds to the interfacial AlNi phase. To further investigate the microstructures of pits, the precipitatefree areas in pits were observed at higher magnification (see Fig. 12e). The fracture morphology of NiTif-AZ31 interface is rougher and more tortuous than that of AZ31 foils diffusion bonding interface. In addition, it should be noted that the AZ31 matrix interface is covered with a certain number of nanocrystalline particles, whose micro fracture planes caused by the grain breakage can be clearly observed. Fig. 12g shows the surface micromorphology of the debonded NiTi fiber. The attachments on NiTi fiber surface are confirmed as Mg from the matrix by SEM-EDS mapping analysis (Fig. 12h). It means that the interfacial bonding between NiTi fibers and AZ31 matrix is excellent, even stronger than the neighboring Mg matrix. Corresponding to the micro fracture planes on AZ31 matrix in pits, the nanocrystalline particles which are broken and extracted from the fracture planes are distributed on the surface of NiTi fiber (see Fig. 12i). The SEM-EDS point analysis of the yellow circular area in Fig. 12i demonstrates that the nanoparticles and matrix regions are rich in Mg, Ti and O, which corresponds to the interfacial nanocrystalline and amorphous layer. Previous studies indicated that the existence of amorphous phases at interface significantly increased the bonding strength between the coating and the matrix, which demonstrated that the amorphous phase has stronger bonding ability to metal than the crystalline phase [41–43]. Unlike the brittle intermetallics-metal interface [44–47], amorphous phases existed at the amorphous-metal interface have been proved to play a toughening role in the microstructure [48–50]. During deformation process, the amorphous layer can absorb dislocations effectively, thus changing the crack nucleation and crack growth rates. In addition, due to the limited driving force of crack growth in amorphous, crack propagation in amorphous layer is slower than in a clean grain boundary [48]. It is noteworthy that in the study of Yang at el. [50], the glass alumina distributed between grain boundaries can even match the deformation of Al matrix like a liquid without any cracks and spallation when stretching the pure aluminum nanotips under O2 gas environments at a moderate strain rate. This phenomenon fully confirms that the amorphous phase has a superior bonding effect with the metal matrix. Therefore, it can be concluded that the NiTi fibers can effectively enhance the bonding strength of the adjacent AZ31 diffusion interface by forming amorphous phase at NiTif/Mg interfaces.

Acknowledgement The authors appreciate the financial support by the Natural Sciences Foundation of China (Grant No. 51671065), the Heilongjiang Postdoctoral Fund of China (No. LBH-Z16046), the Heilongjiang Natural Science Foundation (No. QC2018051), the China Postdoctoral Science Foundation (2017T100227) and the Central University Foundation of Harbin Engineering University (No.3072019CF1012). References [1] B.L. Mordike, T. Ebert, Magnesium properties-applications-potential, Mater. Sci. Eng. A 302 (2001) 37–45. [2] H. Friedrich, S. Schumann, Research for a new age of magnesium in the automotive industry, J. Mater. Process. Technol. 117 (2001) 276-28. [3] Alan A. Luo, Magnesium casting technology for structural applications, J. Magnes. Alloy. 1 (2013) 2–22. [4] Fusheng Pan, Mingbo Yang, Xianhua Chen, A review on casting magnesium alloys: modification of commercial alloys and development of new alloys, J. Mater. Sci. Technol. 32 (2016) 1211–1221. [5] Yujie Cui, Yunping Li, Shihai Sun, Huakang Bian, Hua Huang, Zhongchang Wang, Yuichiro Koizumi, Akihiko Chiba, Enhanced damping capacity of magnesium alloys by tensile twin boundaries, Scr. Mater. 101 (2015) 8–11. [6] Chi Xu, Jinghuai Zhang, Shujuan Li, Yongbin Jing, Yufeng Jiao, Longjiang Xu, Li Zhang, Fengchun Jiang, Milin Zhang, Ruizhi Wu, Microstructure, mechanical and damping properties of Mg–Er–Gd–Zn alloy reinforced with stacking faults, Mater. Des. 79 (2015) 53–59. [7] X.J. Wang, D.K. Xu, R.Z. Wu, X.B. Chen, Q.M. Peng, L. Jin, Y.C. Xin, Z.Q. Zhang, Y. Liu, X.H. Cheng, G. Chen, K.K. Deng, H.Y. Wang, What is going on in magnesium alloys? J. Mater. Sci. Technol. 34 (2018) 245–247. [8] M.J. Shen, X.J. Wang, C.D. Li, M.F. Zhang, X.S. Hu, M.Y. Zheng, K. Wu, Effect of submicron size SiC particles on microstructure and mechanical properties of AZ31B magnesium matrix composites, Mater. Des. 54 (2014) 436–442. [9] M.J. Shen, X.J. Wang, M.F. Zhang, X.S. Hu, M.Y. Zheng, K. Wu, Fabrication of bimodal size SiCp reinforced AZ31B magnesium matrix composites, Mater. Sci. Eng. A 601 (2014) 58–64. [10] Mingyi Zheng, Kun Wu, Congkai Yao, S. Kamado, Y. Kojima, Squeeze cast Al18B4O33 whisker-reinforced magnesium matrix composite, J. Mater. Sci. Lett. 21 (2002) 533–535. [11] M. Habibnejad-Korayem, R. Mahmudi, W.J. Poole, Enhanced properties of Mgbased nano-composites reinforced with Al2O3 nano-particles, Mater. Sci. Eng. A 519 (2009) 198–203. [12] Q.C. Jiang, X.L. Li, H.Y. Wang, Fabrication of TiC particulate reinforced magnesium matrix composites, Scr. Mater. 48 (2003) 713–717. [13] S.F. Hassan, M. Gupta, Development of a novel magnesium/nickel composite with improved mechanical properties, J. Alloy. Comp. 335 (2002) L10–L15. [14] S.F. Hassan, O.O. Nasirudeen, N. Al-Aqeeli, N. Saheb, F. Patel, M.M.A. Baig, Magnesiume nickel composite: preparation, microstructure and mechanical properties, J. Alloy. Comp. 646 (2015) 333–338. [15] S.F. Hassan, M. Gupta, Development of a novel magnesium-copper based composite

5. Conclusions An innovative NiTi shape memory alloy fiber reinforced AZ31 (SMAFR-AZ31) composite sheet was successfully fabricated using vacuum hot pressing method. The microstructure characterizations, phase identifications, normal tensile property and fracture characteristics of the composite were investigated systematically. The main conclusions are summarized as follows: (1) The interface between NiTi fiber and AZ31 matrix is well bonded without obvious macro defects, and the fibers are evenly arranged in the matrix with their original circular state. Furthermore, the original AZ31 foils boundaries of SMAFR-AZ31 composite sheet and AZ31 laminate sheet disappear completely after sintering. The volume fractions of NiTi fibers and AZ31 matrix are 8% and 92%, 9

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under large strain deformation, Mater. Char. 58 (2007) 162–167. [34] Z.X. Cui, Metallography and Heat Treatments, Mechanical Industry Press of China, 2000. [35] J.L. Murray, The Mg-Ti (Magnesium-Titanium) system, J. Phase Equilibria 7 (1986) 245–247. [36] A.A. Nayeb-Hashemi, J.B. Clark, The Mg-Ni (Magnesium-Nickel) system, J. Phase Equilibria 6 (1985) 238–244. [37] J. Hu, G. Wu, Q. Zhang, P. Kang, Y. Liu, Microstructure of multilayer interface in an Al matrix composite reinforced by TiNi fiber, Micron 64 (2014) 57–65. [38] D.J. Sordelet, X.Y. Yang, E.A. Rozhkova, M.F. Besser, M.J. Kramer, Oxygen-stabilized glass formation in Zr80Pt20 melt-spun ribbons, Appl. Phys. Lett. 83 (2003) 69. [39] C.T. Liu, Z.P. Lu, Effect of minor alloying additions on glass formation in bulk metallic glasses, Intermetallics 13 (2005) 415–418. [40] H.X. Li, J.E. Gao, Z.B. Jiao, Y. Wu, Z.P. Lu, Glass-forming ability enhanced by proper additions of oxygen in a Fe-based bulk metallic glass, Appl. Phys. Lett. 95 (2009) 161905. [41] K.A. Khor, Z.L. Dong, C.H. Quek, P. Cheang, Microstructure investigation of plasma sprayed HA/Ti6Al4V composites by TEM, Mater. Sci. Eng. A 281 (2000) 221–228. [42] Pavitra Bansal, Nitin P. Padture, Alexandre Vasiliev, Improved interfacial mechanical properties of Al2O3-13wt%TiO2 plasma-sprayed coatings derived from nanocrystalline powders, Acta Mater. 51 (2003) 2959–2970. [43] Y. Wang, W. Tian, Y. Yang, Thermal shock behavior of nanostructured and conventional Al2O3/13 wt% TiO2 coatings fabricated by plasma spraying, Surf. Coat. Technol. 201 (2007) 7746–7754. [44] Yuqiang Han, Fengchun Jiang, Chunfa Lin, Yuan Ding, Hao Huang, Enhao Wang, Zhenqiang Wang, Chunhuan Guo, Microstructure and mechanical properties of continuous ceramic SiC and shape memory alloy NiTi hybrid fibers reinforced Ti-Al metal-intermetallic laminated composite, J. Alloy. Comp. 729 (2017) 1145–1155. [45] Yao Tian, Enhao Wang, Wei Li, Zhongyi Niu, Yunpeng Chang, Chunhuan Guo, Zhenqiang Wang, Zhe Leng, Fengchun Jiang, Improved fracture toughness of NiTi shape memory alloy fiber reinforced Ti-Al metal-intermetallic-laminate composite, J. Alloy. Comp. 739 (2018) 669–677. [46] Chunfa Lin, Yuqiang Han, Chunhuan Guo, Yunpeng Chang, Xiaoxiao Han, Lin Lan, Fengchun Jiang, Synthesis and mechanical properties of novel Ti-(SiCf/Al3Ti) ceramic-fiber-reinforced metal-intermetallic-laminated (CFRMIL) composites, J. Alloy. Comp. 722 (2017) 427–437. [47] Enhao Wang, Chunhuan Guo, Peijun Zhou, Chunfa Lin, Xiaoxiao Han, Fengchun Jiang, Fabrication, mechanical properties and damping capacity of shape memory alloy NiTi fiber-reinforced metal-intermetallic-laminate (SMAFR-MIL) composite, Mater. Des. 95 (2016) 446–454. [48] Zhiliang Pan, Timothy J. Rupert, Amorphous intergranular films as toughening structural features, Acta Mater. 89 (2015) 205–214. [49] Anandh Subramaniama, Christoph T. Koch, Rowland M. Cannon, Manfred Ruhle, Intergranular glassy films: an overview, Mater. Sci. Eng. A 422 (2006) 3–18. [50] Yang Yang, Akihiro Kushima, Weizhong Han, Huolin Xin, Ju Li, Liquid-like, selfhealing aluminum oxide during deformation at room temperature, Nano Lett. 18 (2018) 2492–2497.

with improved mechanical properties, Mater. Res. Bull. 37 (2002) 377–389. [16] S.F. Hassan, M. Gupta, Development of ductile magnesium composite materials using titanium as reinforcement, J. Alloy. Comp. 345 (2002) 246–251. [17] P. Perez, G. Garces, P. Adeva, Mechanical properties of a Mg-10 (vol.%)Ti composite, Compos. Sci. Technol. 64 (2004) 145–151. [18] D.S. Li, X.P. Zhang, Z.P. Xiong, Y.-W. Mai, Lightweight NiTi shape memory alloy based composites with high damping capacity and high strength, J. Alloy. Comp. 490 (2010) L15–L19. [19] Biao Yan, Gang Li, Mg alloy matrix composite reinforced with TiNi continuous fiber prepared by ball-milling/hot-pressing, Composites Part A Appl. Sci. Manuf. 36 (2005) 1590–1594. [20] K. Mizuuchi, K. Inoue, K. Hamada, M. Sugioka, M. Itami, M. Fukusumi, M. Kawahara, Processing of TiNi SMA fiber reinforced AZ31 Mg alloy matrix composite by pulsed current hot pressing, Mater. Sci. Eng. A 367 (2004) 343–349. [21] Ziya Esen, The effect of processing routes on the structure and properties of magnesium-TiNi composites, Mater. Sci. Eng. A 558 (2012) 632–640. [22] Tarik Aydogmus, Processing of interpenetrating Mg-TiNi composites by spark plasma sintering, Mater. Sci. Eng. A 624 (2015) 261–270. [23] N. Hort, H. Dieringa, S.T. Kumar, K.U. Kainer, Magnesium Technology, Springer Verlag Berlin Heidelberg, 2006, pp. 315–334. [24] Brian T. Lester, Theocharis Baxevanis, Yves Chemisky, Dimitris C. Lagoudas, Review and perspectives: shape memory alloy composite systems, Acta Mech. 226 (2015) 3907–3960. [25] Zichuan Lu, Fengchun Jiang, Yunpeng Chang, Zhongyi Niu, Zhenqiang Wang, Chunhuan Guo, Interfacial microstructure characterization and mechanical behavior of NiTi fiber reinforced Al3Ti composite, Mater. Des. 143 (2018) 274–284. [26] Peijun Zhou, Chunhuan Guo, Enhao Wang, Zeming Wang, Ye Chen, Fengchun Jiang, Interface tensile and fracture behavior of the Ti/Al3Ti MetalIntermetallic Laminate (MIL) composite under quasi-static and high strain rates, Mater. Sci. Eng. A 665 (2016) 66–75. [27] Li Xiang, Yuhang Sun, Zhenqiang Wang, Fengchun Jiang, Fabrication, microstructure characterization and mechanical property of the Ti6Al4V-(NiTif/ Mg3AlZn) laminate composite, J. Alloy. Comp. 774 (2019) 656–667. [28] H. Okamoto, Al-Ni (Aluminum-Nickel), J. Phase Equilibria 14 (1993) 257–259. [29] N. Sun, S.T. Niknejad, G. Zhang, M.K. Lee, L. Wu, Y. Zhou, Microstructure and mechanical properties of resistance spot welded AZ31/AA5754 using a nickel interlayer, Mater. Des. 87 (2015) 905–913. [30] A.M. Nasiri, P. Chartrand, D.C. Weckman, N.Y. Zhou, Thermochemical analysis of phases formed at the interface of a Mg alloy-Ni-plated steel joint during laser brazing, Metall. Mater. Trans. A 44 (2012) 1937–1946. [31] A.M. Nasiri, D.C. Weckman, Y. Zhou, Interfacial microstructure of diode laser brazed AZ31B magnesium to steel sheet using a nickel interlayer, Weld. J. 92 (2013) 1s–10s. [32] K. Janecek, M. Popov, M.G. Krieger, R.J. Hellmig, Y. Estrin, Mechanical properties and microstructure of a Mg alloy AZ31 prepared by equal-channel angular pressing, Mater. Sci. Eng. A 462 (2007) 116–120. [33] Q. Guo, H.G. Yan, Z.H. Chen, H. Zhang, Grain refinement in as-cast AZ80 Mg alloy

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