Effect of Heat Treatment on Microstructure and Mechanical Properties of the AZ31/WE43 Bimetal Composites

metals

Article Effect of Heat Treatment on Microstructure and Mechanical Properties of the AZ31/WE43 Bimetal Composites

Dexing Xu 1, Kangning Zhao 1, Changlin Yang 2 , Hongxiang Li 1,* and Jishan Zhang 1,*

1 State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, 100083 Beijing, China; [email protected] (D.X.); [email protected] (K.Z.) 2 State Key Laboratory of Solidiﬁcation Processing, Northwestern Polytechnical University, Xi’an 710072, China; [email protected] * Correspondence: [email protected] (H.L.); [email protected] (J.Z.); Tel.: +86-10-62332350 (H.L.)

Received: 30 October 2018; Accepted: 15 November 2018; Published: 20 November 2018

Abstract: Effect of heat treatment on the microstructure and mechanical properties of the AZ31/WE43 bimetal composites was investigated, and the optimized solution treatment and ageing treatment parameters were achieved. After heat treatment, it was found that that the high melting-point Y-rich phases and Mg-RE (rare earth) phases at the interface were distributed more uniformly, so the interfacial strength and plasticity were improved. Meanwhile, the shear strength and plasticity of AZ31 were increased, and the shear strength of WE43 was also improved by heat treatment. The evolution of interface morphologies and enhancement of the mechanical properties at the interface will be discussed in detail.

Keywords: Mg/Mg bimetal composites; heat treatment; interfacial microstructure; mechanical properties; fracture behaviors

1. Introduction Magnesium alloys are becoming increasingly attractive for potential applications in aerospace, automobile, and electronics, due to the advantages of light weight, high specific strength, and high specific stiffness [1–3]. However, the conventional magnesium alloys, such as AZ31, AZ61, and ZK60, have many undesirable properties, such as poor wear resistance and relatively low mechanical strength, which has restricted their wide application in many industrial fields [4,5]. On the other hand, Mg-based alloys containing rare earth (RE) elements, such as WE43, WE54, and GW103, are achieving ultra-high strength, due to the combination of solid solution strengthening and precipitation hardening [6–8]. However, the high cost and poor toughness of Mg-RE alloys greatly limit their range of applications. In order to save the manufacturing cost and improve the material properties, a new structural design for creating laminated composites with excellent ductility and high strength, and decreasing the amount of RE elements, has been intended. Thus, an Mg/Mg hybrid metal is a desirable lightweight composite that benefits from the advantages of both heavy rare earth Mg alloys and commercial high-strength Mg alloys. The Mg/Mg bimetallic materials can be achieved by ECAE (equal channel angular extrusion) [9], ARB (accumulative roll bonding) [10], compound casting [11,12]. Recently, a new AZ31/WE43 bimetal composite with a good interfacial metallurgical bonding was fabricated using an insert molding method by the authors [11]. Moreover, the as-cast microstructure and bonding mechanism at the interface have also been investigated. However, although good metallurgical bonding was achieved for the Mg/Mg bimetal composites, they have poor overall properties because of uneven microstructures and casting defects in the interface and/or basal metals.

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In contrast to the as-cast studies, recently investigating the interfacial microstructure and mechanical properties of the bimetal composites after heat treatment is becoming a focus, because it can improve the interfacial bonding strength and the overall properties of bimetal composites. Dezellus et al. [13] investigated the effect of T6 heat treatment on the mechanical properties of Ti/Al–7Si–0.3Mg bimetals, and found that the size and number of Si particles decreased significantly at the interface, and the interfacial strength was improved. Liu et al. [14] reported the microstructure evolution and mechanical properties of 6101/A356 bimetal composites after T6 heat treatment, and identified that the morphology of silicon particles was changed from long, coarse plate-like, to fine spherical shape, and Zn, Mg elements were diffused more homogeneously across the interface. Thus, the interfacial shear strength can be improved by about 30%. Yan et al. [15] investigated the microstructure evolution and mechanical properties of 7050/6009 bimetal slabs during homogenization annealing, and developed an optimal homogenizing annealing process for subsequent processing. Wang et al. [16] characterized the microstructure evolution at the bonding interface of Al/Al–Mg–Si alloy after the heat treatment, similarly leading to the improved interface bonding strength. Li et al. [17] investigated the diffusion behavior of the alloying elements of 2024/3003 gradient composite at 768 K for 45 min, and then aged at 458 K for 12 h, and discovered that it included two kinds of diffusion behaviors, i.e., long-distance diffusion and short-distance diffusion at the interface with a heat treatment. From the above studies, it is seen that heat treatment is an effective method to improve the overall properties and interfacial strength of bimetal composites. However, the heat treatment process of the Mg/Mg bimetal composites is difficult to control, even if it is so significant. During the heat treatment process, the effect of the elemental diffusion, the solution, and precipitation of interface phases on the interfacial properties, is rather important [13,18]. Moreover, it is necessary for optimizing parameters to achieve a desirable overall property, by considering the consistency between basal metals and interface layers [18]. Thus, as for the microstructure evolutions and mechanical properties at the interface of Mg/Mg bimetal composites after heat treatment, no studies have been reported thus far. For this purpose, the AZ31/WE43 bimetal composites are selected as an example, and the heat treatment process of this kind of bimetal composite is optimized. Meanwhile, the microstructure and mechanical properties after heat treatment were analyzed in detail. According to studies, the microstructure and mechanical properties of WE43 would be strengthened during heat treatment, and T6 treatment is an effective way [6,19]. However, AZ31 has no obvious solid solution and ageing strengthening effect, due to the low content of alloying elements [20]. Therefore, the heat treatment process of WE43 is used as a reference for optimizing that of the AZ31/WE43 bimetal composite, and the solution time is fixed at 12 h, in order to shorten the process time. Eventually, an optimized heat treatment parameter is selected to improve the interfacial uniformity and the overall properties of this novel bimetal composite. The study can offer useful instruction for the heat treatment process of Mg/Mg bimetal composites, with good industrial application prospects.

2. Experimental Procedures

2.1. Specimen Treatment The principle of Mg/Mg bimetal composite material produced by insert molding was introduced in our previous works [11] in detail. Commercially available AZ31(Mg–3Al–Zn) magnesium alloy and WE43 (Mg–4Y–3RE–0.5Zr) alloy were used as casting material and solid insert material. Before the inset molding procedure, the inserts were cut into cylindrical bars with 40 mm diameter and 90 mm height. The surface of the WE43 bars was mechanically polished with SiC papers before use and then treated by alkaline cleaning, acid pickling, and ultrasonically degreasing with acetone. Similarly, the AZ31 ingots were cut into pieces and pre-treated by the same procedure, except mechanical polishing. Afterwards, the pre-treated AZ31 pieces were melted in the crucible located in a resistance furnace (SG-10-10, Jinan Changda Thermal Industry Co., Ltd., Jinan, China). During the melting process, a protective gas mixture containing SF6 and CO2 were used for preventing oxidation. When the melt reached Metals 2018, 8, 971 3 of 15 the preset temperature of 933 K, WE43 bars were immersed into the AZ31 melt, and kept for 110 s. After that, the entire assembly containing mold was taken out of the furnace and cooled to room temperature. A detailed description of the insert molding procedure and the experimental device is given in reference [11]. In this study, T6 treatment refers to a solution treatment at 673–793 K for several hours, followed by quenching and artificial ageing treatment. The samples were firstly treated under an argon atmosphere and the procedure of solution treatment was designed as 698, 723, 748, 773 and 788 K, with a fixed time of 12 h, followed by water-immersion quenching with a maximum delay of 5 s. Shortly thereafter, the specimens were subjected to an artificial ageing treatment at 483 K for different times, and quenched by cooling water.

2.2. Materials Characterization To analyze the interfacial formation, the specimens were cut from the middle part of the samples parallel to the cylindrical insert with a size of 15 mm × 15 mm by setting the interface in the middle. The microstructures of the specimens were observed using a ZEISS EVO-18 (Jena, Germany) scanning electron microscope (SEM) equipped with an energy-dispersive X-ray spectrometer (EDS, Thermo scientiﬁc, Madison, WI, USA). The elemental distributions were analyzed by JEOL JXA-8230 (JEOL, Tokyo, Japan) electron probe microanalyzer (EPMA). The precipitated phases on bonding interface were also characterized by Rigaku D/max 2400 X-ray diffractometer with Cu Kα target and a Tecnai G2 F30 (FEI, Boston, MA, USA) transmission electron microscope (TEM) operated at 300 kV, respectively. The interfacial shear test was performed using a self-deﬁned method in a universal testing machine, which was described in reference [11]. Interfacial shear strength (τ) can be calculated using Equation (1) introduced in reference [11]. The Vickers hardness was measured across the interface by the microhardness tester at a load of 50 g and a dwell time of 15 s.

3. Results

3.1. Solution-Treated Microstructure As expected, the elemental mixing and diffusion at the interface of AZ31/WE43 bimetal composites can accelerate the formation of sound metallurgical bonding. There is a continuous transition region between the AZ31 matrix and the WE43 insert. The interface is taintless and free from oxide inclusions, pores, and other undesirable defects. As shown in Figure1a, a relatively uniform transition interface layer with a thickness of 150–180 µm consists of α-Mg grains and other intermetallic phases formed due to the contact between the WE43 alloy insert and AZ31 melt. Figure1b–f shows the interfacial microstructure evolution of AZ31/WE43 bimetal composite at different solution temperatures of 698 K, 723 K, 748 K, 773 K, and 788 K for 12 h. It can be seen from Figure1b–f that the interfacial thickness after solution treatment increases from 316 µm to 782 µm, because of the long-distance diffusion of Mg, RE, Al, Zn element with increasing annealing temperature, from 698 K to 788 K. It is found that RE-rich eutectic phases at WE43 side were dissolved gradually, and ﬁnally disappeared with the increase of solution-treatment temperature. When the temperature reached 773 K, almost all of the RE-rich eutectic phase and Mg17Al12 phase were dissolved into the α-Mg matrix, and only the high melting-point white particles were evenly distributed in the interface layer. Figures2–4 present the concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn across the interface under the temperatures of 723 K, 773 K, and 788 K for 12 h. When the solution temperature was 723 K, the amount of eutectic phase at the interface was reduced and partially dissolved, as shown in Figure2. When the solution temperature was above 773 K, eutectic compounds in the interface were almost dissolved, as shown in Figures3 and4, and the elements were also evenly distributed. When solution heat treatment is carried out at 773 K or 788 K for 12 h, the coarse RE-rich phases in the interface adjacent to WE43 undergo fragmentation, necking, spheroidization, and solution. The lamellar-like (Al, Mg)2RE phase in the interface adjacent to the AZ31 side becomes Metals 2018, 8, 971 4 of 15 more uniform and ﬁner. Meanwhile, the transition region of AZ31/WE43 bimetal composite become smooth and no evident gradient exists. Metals 2018, 8, x FOR PEER REVIEW 4 of 15 Metals 2018, 8, x FOR PEER REVIEW 4 of 15

Figure 1. Microstructure in the interfacial region of AZ31/WE43 bimetal after different solution FigureFigure 1. 1. MicrostructureMicrostructure inin thethe interfacialinterfacial regionregion of AZ31/WE43 bimetal bimetal after after different different solution solution treatment and corresponding thickness of the transition zone :(a) As-cast, 185 μm; (b) 698 K, 12 h, 316 treatmenttreatment and and correspondingcorresponding thickn thicknessess of of the the transition transition zone zone: :(a) (As-cast,a) As-cast, 185 185μm; µ(bm;) 698 (b) K, 698 12 K,h, 316 12 h, μm; (c) 723 K, 12 h, 443 μm; (d) 748 K, 12 h, 575 μm; (e) 773 K, 12 h, 681 μm, (f) 788 K, 12 h, 782 μm. 316μm;µm; (c) (723c) 723 K, K,12 12h, 443 h, 443 μm;µ m;(d)( d748) 748 K, K,12 12h, 575 h, 575 μm;µ m;(e) (773e) 773 K, 12 K, h, 12 681 h, 681 μm,µ m,(f) 788 (f) 788 K, 12 K, h, 12 782 h, 782 μm.µ m.

Figure 2. Concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution Figure 2. Concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution Figuretreatment 2. Concentration conditions ofmappings 723 K for 12 of h. Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution treatment treatment conditions of 723 K for 12 h. conditions of 723 K for 12 h.

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Figure 3. Concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution Figure 3.3. ConcentrationConcentration mappingsmappings ofof Mg,Mg, RE RE (Y, (Y, Gd, Gd, Nd), Nd), Al, Al, Zn, Zn, and and Mn underMn under the solution the solution treatment treatment conditions of 773 K for 12 h. conditionstreatment conditions of 773 Kfor of 773 12 h.K for 12 h.

Figure 4. Concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution Figure 4. Concentration mappings of Mg, RE (Y, Gd, Nd), Al, Zn, and Mn under the solution Figuretreatment 4. Concentration conditions of mappings 788 K for 12 of Mg,h. RE (Y, Gd, Nd), Al, Zn, and Mn under the solution treatment treatment conditions of 788 K for 12 h. conditions of 788 K for 12 h. It is known that the other factor that determines the solution treatment process of the bimetal It is known that that the the other other factor factor that that determines determines the the solution solution treatment treatment process process of the of bimetal the bimetal compositecomposite is isthe the dissolving dissolving degree degree of ofphases phases in inAZ AZ3131 and and WE43 WE43 matrix. matrix. It canIt can be be seen seen from from Figure Figure composite1b–f that is the the Mg dissolving17Al12 phases degree in the of phasesAZ31 side in AZ31 are eliminated and WE43 when matrix. the It temperature can be seen was from above Figure 6981b–f 1b–f that the Mg17Al12 phases in the AZ31 side are eliminated when the temperature was above 698 that the Mg17Al12 phases in the AZ31 side are eliminated when the temperature was above 698 K, K, K,and and no no significant significant changes changes were were found found for for th eth microstructuree microstructure of ofAZ31 AZ31 side, side, even even when when the the and no signiﬁcant changes were found for the microstructure of AZ31 side, even when the temperature temperaturetemperature was was increased increased to to788 788 K, K,particul particularlyarly without without overheating. overheating. Meanwhile, Meanwhile, the the

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wasMetals increased 2018, 8, x FOR to 788 PEER K, REVIEW particularly without overheating. Meanwhile, the microstructure at6 of WE43 15 side shows no significant evolution until the temperature reaches 773 K. The RE-rich eutectic phase almostmicrostructureMetals disappears 2018, 8, x atFOR under WE43 PEERthe sideREVIEW solution shows no conditions significant of evolution 788 K for until 12 h. the From temperature Figure4, eutecticreaches 773 compounds6 K.of 15The wereRE-rich all dissolved, eutectic phase no matter almost thedisappears interface under or matrix the solution and only conditions some highof 788 melting-point K for 12 h. From phases Figure are evenly4, eutecticmicrostructure distributed compounds inat theWE43 interfacewere side all shows dissolved, and no WE43 significant no matrix. matter evolution By the XRD interface until analysis, the temperatureor thematrix high and reaches melting-point only 773 some K. The high phases atmelting-point theRE-rich interface eutectic afterphases phase solution are almost evenly treatment disappears distributed are under Y-richin the the interface phasessolution and conditionsand otherWE43 of Mg-REmatrix. 788 K for cuboid-shapedBy 12 XRD h. From analysis, Figure phases, the 4, eutectic compounds were all dissolved, no matter the interface or matrix and only some high ashigh shown melting-point in Figure5 .phases The high at the melting-point interface after phases solution at thetreatment interface are areY-rich further phases analyzed and other by Mg- TEM. RE melting-pointcuboid-shaped phases phases, are asevenly shown distributed in Figure in 5. the Th interfacee high melting-point and WE43 matrix. phases By XRDat the analysis, interface the are The TEM images with the corresponding selected area electron diffraction (SAED) patterns are shown furtherhigh analyzedmelting-point by TEM. phases The at TEM the interface images withafter sothelution corresponding treatment are selected Y-rich area phases electron and other diffraction Mg- in Figure6. From Figure6a,c, it is found that the distribution of high melting-point phases mainly (SAED)RE cuboid-shaped patterns are shownphases, inas Figureshown in6. FigureFrom Figure5. The high6a,c, melting-pointit is found that phases the distributionat the interface of arehigh includes two different types, i.e., one is uniformly dispersed in the grain interior, and another one melting-pointfurther analyzed phases by mainly TEM. The includes TEM images two different with the types, corresponding i.e., one is selected uniformly area dispersed electron diffraction in the grain is clustered(SAED) onpatterns the grain are shown boundary. in Figure From 6. FigureFrom Figure6b,d, these6a,c, it cuboid-shaped is found that the phases distribution are identified of high to interior, and another one is clustered on the grain boundary. From Figure 6b,d, these cuboid-shaped melting-point phases mainly includes two different types, i.e., one is uniformly dispersed in the grain bephases a face-centered are identified cubic to (fcc)be a structure,face-centered and cubic their (fcc) chemical structure, composition and their chemical is identified composition as the Y-rich is interior, and another one is clustered on the grain boundary. From Figure 6b,d, these cuboid-shaped phase.identified From as the the above Y-rich results, phase. theFrom optimized the above solution results, treatmentthe optimized process solution of the treatment AZ31/WE43 process bimetal of phases are identified to be a face-centered cubic (fcc) structure, and their chemical composition is materialthe AZ31/WE43 can be determined bimetal material to be 788 can K forbe 12determined h, at which to thebe bimetal788 K for composite 12 h, at showswhich morethe bimetal uniform identified as the Y-rich phase. From the above results, the optimized solution treatment process of microstructure, no matter at the interface or in the WE43 matrix. compositethe AZ31/WE43 shows more bimetal uniform material microstructure, can be determined no matter to beat the788 interface K for 12 orh, inat the which WE43 the matrix. bimetal composite shows more uniform microstructure, no matter at the interface or in the WE43 matrix.

Figure 5. XRD patterns of solution-treated (788 K for 12 h) and ageing-treated (483 K for 48 h) FigureFigure 5. XRD 5. patternsXRD patterns of solution-treated of solution-treated (788 K(788 for K 12 for h) and12 h) ageing-treated and ageing-treated (483 K(483 for 48K for h) specimens48 h) specimens at the transition interface. at thespecimens transition at interface. the transition interface.

FigureFigure 6. (a ,6.c), (a TEM,c), TEM bright bright ﬁeld field images images of of the the highhigh melting-pointmelting-point phases phases at atthe the transition transition interface; interface; Figure(b,d ),6. the (a, ccorresponding), TEM bright SAED field imagespatterns. of the high melting-point phases at the transition interface; (b,(db),,d the), the corresponding corresponding SAED SAED patterns. patterns.

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3.2.Metals Ageing2018, 8 ,Hardening 971 Response and Microstructure Evolution 7 of 15 Figure 7 shows the age hardening curves of the interface zone, AZ31, and WE43 matrix after solution3.2. Ageing treatment Hardening at 788 Response K for and 12 h, Microstructure and aged at Evolution483 K for different hours. The shape of the ageing curves shows different characterizations for the matrix materials and interface layer of the bimetal composite.Figure 7The shows hardness the age of hardeningAZ31 is almost curves cons of thetant interface in the whole zone, ageing AZ31, range, and WE43 which matrix cannot after be strengthened.solution treatment The athardness 788 K for of 12WE43 h, and increases aged at 483before K for 36 differenth, and then hours. shows The shapea wide of thepeak-aged ageing platformcurves shows after different36 h, indicating characterizations an excellent for ther themal matrix stability. materials Compared and interface with WE43, layer the of the hardening bimetal ratecomposite. of interface The zone hardness is faster, of AZ31and the is peak almost value constant can be in reached the whole at 24 ageing h and, range,then, after which peak-ageing, cannot be thestrengthened. value decreases The hardness again. The of WE43maximum increases hardening before value 36 h, andat peak-ageing then shows is a about wide peak-agedHV100. To platform analyze precipitatesafter 36 h, indicating after ageing an excellenttreatment, thermal XRD and stability. TEM analysis Compared of the with samples WE43, theat the hardening interface rate were of checked.interface From zone Figure is faster, 5, some and the small peak peaks value marked can be “△ reached” and “▽ at” 24 appear h and, in then,the XRD after pattern, peak-ageing, which suggeststhe value that decreases some again.new precipitates The maximum are hardeningformed in valuethe aged at peak-ageing specimens. is Figure about HV100.8 shows To the analyze TEM micrographsprecipitates afterwith ageingthe corresponding treatment, SAED XRD and pattern, TEM and analysis the HTEM of the micrographs samples at of the precipitates interface wereafter ageingchecked. at From483 K Figurefor 48 h.5, someAs seen small in Figure peaks 8a, marked the lens-shaped “ 4” and “ precipitates5” appear inwith the a XRD diameter pattern, about which 20– suggests that some new precipitates are formed in the aged specimens. Figure8 shows the TEM 80 nm are dispersed in the interface. From the SAED taken from the [2110] zone axis, and the micrographs with the corresponding SAED pattern, and the HTEM micrographs of precipitates after HTEM image shown in Figure 8c, it can be confirmed that the precipitates are β’ phases and α matrix, ageing at 483 K for 48 h. As seen in Figure8a, the lens-shaped precipitates with a diameter about i.e., (100)' //(1210) , [001] //[0001] . It was identified as an orthorhombic structure and 20–80 nm areβ dispersedα in theβ interface.' α From the SAED taken from the [2 110] zone axis, and the HTEMa = 2a image = 0.640 shown nm, in b Figure= 8b8c, it can= 2.223 be conﬁrmed nm, c = c that =0.525 the precipitates nm [19,21]. are Figureβ’ phases 8b shows and α Mg {1010}Mg Mg matrix, i.e., (100)β0 // 1210 , [001]β0 //[0001]α. It was identiﬁed as an orthorhombic structure and the plate shaped precipitates αwith a length of 100–500 nm, and the corresponding SAED pattern with a = 2aMg = 0.640 nm, b = 8b{10 10}Mg = 2.223 nm, c =cMg = 0.525 nm [19,21]. Figure8b shows the theplate [0001] shaped zone precipitates axis, indicating with a that length the of precipitates 100–500 nm, are and β1 phases, the corresponding with a fcc structure SAED pattern (a = 0.74 with nm), the and the orientation relationship between β1 and α matrix of (112)β / /(1010)α and [0001] zone axis, indicating that the precipitates are β1 phases, with a fcc structure (a1 = 0.74 nm), and the orientation relationship between β and α matrix of (112) //(1010) and [110] //[0001] [19,21]. [110] //[0001] [19,21]. 1 β1 α β1 α βα1

Figure 7. TheThe ageing hardening curves of WE43, AZ31, and the interface.

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Figure 8. TEM analysis of the precipitates at the interface after ageing treatment (483 K for 48 h) (a) Figure 8. TEM analysis of the precipitates at the interface after ageing treatment (483 K for 48 h) TEM image of the lens-shaped precipitates with corresponding SAED patterns; (b) TEM image of the (a) TEM image of the lens-shaped precipitates with corresponding SAED patterns; (b) TEM image of plate-shaped precipitates with corresponding SAED patterns; (c,d) corresponding HTEM image. the plate-shaped precipitates with corresponding SAED patterns; (c,d) corresponding HTEM image.

3.3.3.3. Mechanical Properties afterafter AgeingAgeing TheThe effecteffect ofof artificial artificial ageing ageing time time on on mechanical mechanical properties properties was was investigated investigated with with reference reference to the to hardnessthe hardness analysis. analysis. As shown As shown in Figure in 9Figurea, the shear 9a, the mechanical shear mechanical properties showproperties a good show consistency a good withconsistency hardness with variation hardness as the variation ageing as time the prolongs. ageing time The shearprolongs. strength The valueshear ofstrength the WE43 value increases of the rapidlyWE43 increases before 24 rapidly h and, then,before slowly 24 h and, increases then, after slow 24ly h.increases When the after ageing 24 h. time When reaches the ageing 48 h, it time has areaches maximum 48 h,shear it has strength a maximum value shear of 221 strength MPa for value the of WE43 221 MPa matrix. for Thethe WE43 shear strengthmatrix. The value shear of thestrength interface value reaches of the ainterface maximum reaches value a ofmaximum 132 MPa value at the of peak-ageing 132 MPa at the condition, peak-ageing and there condition, is no significantand there is change no significant for the change shear strength for the shear after ageingstrength treatment. after ageing When treatment. the ageing When time the increasesageing time to 48increases h, the shear to 48 strength h, the shear value ofstrength the interface value decreasesof the interface a little todecreases about 120 a MPa.little However,to about 120 the shearMPa. mechanicalHowever, the properties shear mechanical of AZ31 matrix properties is substantially of AZ31 unchanged,matrix is substantially which is consistent unchanged, with thewhich trend is ofconsistent hardness with variation. the trend of hardness variation. Aiming at the comparison of the mechanical behaviors of AZ31/WE43 bimetal composites for different heat treatment states, three typical displacement-loading force curves, which covered solution state (0 h), peak-aged state of interface zone (24 h), peak-aged state of WE43 value, but over-ageing state of interface zone (48 h), were selected. As shown in Figure9b–d, the displacements of the interface layers and WE43 matrix decrease with the increase of ageing time, while the AZ31 matrix is reversed. The change of displacement also reflects the evolution of plasticity to a certain extent. The plasticity of WE43 decreases as the time prolongs but the shear strength increases. The shear strength of interface zone is slightly improved without the dramatical decrease of plasticity. The shear strength of AZ31 matrix remains a constant while the plasticity is enhanced. By the above analysis, ageing treatment can improve the mechanical properties of AZ31/WE43 bimetal composites. In order to achieve the overall properties of the AZ31/WE43 bimetal composite, i.e., combining the high strength of WE43 and the high plasticity of AZ31, the optimized artificial ageing parameter can be defined to be 483 K for 48 h.

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Figure 9. 9. (a) The(a) Theshear shear properties properties of the interface, of the interface, WE43 and WE43 AZ31; and(b–d) AZ31; three typical (b–d) displacement- three typical loadingdisplacement-loading force curves covering force curves solution covering state (0 solution h), peak-aged state (0state h), of peak-aged interface statezone of(24 interface h), and peak- zone aged(24 h), state and of peak-aged WE43 value, state but of WE43 over-ageing value, butstate over-ageing of interface state zone of (48 interface h): (b) zonethe typical (48 h): displacement- (b) the typical loadingdisplacement-loading force curve of force WE43; curve (c) ofthe WE43; typical (c )displacement-loading the typical displacement-loading force curve force of interface curve of zone; interface (d) thezone; typical (d) the displacement-loading typical displacement-loading force curve force of AZ31. curve of AZ31. 3.4. Fracture Behaviors at the Interface Layer Aiming at the comparison of the mechanical behaviors of AZ31/WE43 bimetal composites for differentFigure heat 10 showstreatment the states, SEM micrographs three typical of displacement-loading interfacial fracture surfaces force forcurves, the solution-treatedwhich covered solutionsamples. state From (0 h), Figure peak-aged 10a,b, state it shows of interface that thezone cracks (24 h), are peak-aged generated state in of the WE43 zone value, with but a stressover- ageingconcentration state of parallelinterface to zone the (48 load h), direction, were selected. and propagated As shown in with Figure the 9b–d, increase the ofdisplacements shear strength. of theFrom interface Figure layers10c, the and shear WE43 fracture matrix surface decrease is composed with the increase of cleavage of ageing planes. time, Moreover, while the there AZ31 are matrix some whiteis reversed. phases The distributed change of surrounding displacement the also cleavage reflec planes.ts the evolution From Figure of plasticity10d, many to small a certain dimples extent. can Thebe observed plasticity near of WE43 the white decreases phases. as the time prolongs but the shear strength increases. The shear strength of interface zone is slightly improved without the dramatical decrease of plasticity. The shear strength of AZ31 matrix remains a constant while the plasticity is enhanced. By the above analysis, ageing treatment can improve the mechanical properties of AZ31/WE43 bimetal composites. In order to achieve the overall properties of the AZ31/WE43 bimetal composite, i.e., combining the high strength of WE43 and the high plasticity of AZ31, the optimized artificial ageing parameter can be defined to be 483 K for 48 h.

3.4. Fracture Behaviors at the Interface Layer Figure 10 shows the SEM micrographs of interfacial fracture surfaces for the solution-treated samples. From Figure 10a,b, it shows that the cracks are generated in the zone with a stress concentration parallel to the load direction, and propagated with the increase of shear strength. From Figure 10c, the shear fracture surface is composed of cleavage planes. Moreover, there are some white

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Metalsphases2018 distributed, 8, 971 surrounding the cleavage planes. From Figure 10d, many small dimples can10 of be 15 observed near the white phases.

Figure 10.10. FailureFailure sequencesequence occurringoccurring onon thethe solution-treatmentsolution-treatment AZ31/WE43AZ31/WE43 interface of the shearshear specimen.specimen. ( (aa)) A A side side face face view view of the failed interfaceinterface between AZ31 and WE43 fracturefracture surfacesurface indicatingindicating crackcrack arrestarrest lineslines andand thethe initiationinitiation ofof crackingcracking alongalong thethe lamellar-likelamellar-like zone.zone. (b)) HighHigh magniﬁcationmagnification viewsviews of of the the location location of cracks. of cracks. (c) The (c SE) The and SE BSE and images BSE of images the macroscopic of the macroscopic appearance ofappearance the fracture of surfacethe fracture and (surfaced) high and magniﬁcation (d) high magnification views. views.

Figure 1111 shows shows the the SEM SEM micrographs micrographs of interfacial of interf fractureacial fracture surfaces surfaces for the aged-treatedfor the aged-treated samples. Figuresamples. 11 a,bFigure represent 11a,b represent the cross-section the cross-section image after image shear after fracture shear offracture T6-treated of T6-treated AZ31/WE43 AZ31/WE43 bimetal composites.bimetal composites. A side A face side view face view of the of failedthe failed interface interface between between AZ31 AZ31 and and WE43 WE43 fracture fracture surfacesurface indicatesindicates thatthat aa mainmain crackcrack cancan initializeinitialize inin thethe WE43WE43 matrix, and thenthen propagatepropagate along the interfaceinterface zone. Figure Figure 11c–f 11c–f show show the the secondary secondary electron electron (SE) (SE)and andback backscattering scattering electron electron (BSE) images (BSE) images for the forinterfacial the interfacial fracture. fracture. Figure 11c Figure indicates 11c indicates that the frac thatture the mode fracture is mainly mode iscleavage mainly fracture cleavage with fracture river withpatterns river and patterns white andphases white that phases are evenly that distribu are evenlyted distributedsurrounding surrounding the cleavage the planes. cleavage Figure planes. 11d, Figureshows 11anotherd, shows type another of fracture type mode, of fracture i.e., quasi-cleavage mode, i.e., quasi-cleavage fracture with fracture tear ridge. with The tear tear ridge. ridge The along tear ridgegrain boundary along grain included boundary some included high melting-point some high melting-pointphases, shown phases, in the BSE shown images. in the The BSE magnified images. Theview magniﬁed of the grain view interior of the is shown grain interiorin Figure is 11e, shown in which in Figure some 11 smalle, in cleavage which some planes small and cleavagesteps are planesobserved. and Similarly, steps are from observed. the magnified Similarly, view from of the the magniﬁed tripe junctions view of at the grain tripe boundary junctions shown at grain in boundaryFigure 11f, shown some inkink Figure belts 11 andf, some dimples kink are belts detected. and dimples That areis to detected. say, the Thatinterfacial is to say, fracture the interfacial mode of fractureT6-treated mode AZ31/WE43 of T6-treated bimetal AZ31/WE43 composites bimetal is the composites mixed mode is theof the mixed cleavage mode ofand the quasi-cleavage cleavage and quasi-cleavagefractures. fractures.

Metals 2018, 8, 971 11 of 15 Metals 2018, 8, x FOR PEER REVIEW 11 of 15

Figure 11. Failure sequence occurring on theT6-treated AZ31/WE43 interface of the shear specimen. Figure 11. Failure sequence occurring on theT6-treated AZ31/WE43 interface of the shear specimen. (a) A side face view of the failed interface between AZ31 and WE43 fracture surface indicating crack (a) A side face view of the failed interface between AZ31 and WE43 fracture surface indicating crack arrest lines and the initiation of cracking. (b) High magnification views of the location of cracks. (c,d) arrest lines and the initiation of cracking. (b) High magniﬁcation views of the location of cracks. The SE and BSE images of the fracture surface macroscopic appearance. (e,f) High magnification (c,d) The SE and BSE images of the fracture surface macroscopic appearance. (e,f) High magniﬁcation views of the Figure 10d. views of the Figure 10d.

4.4. Discussion Discussion InIn our our previous previous study study [11], [11], the the solidification solidification of of the the interface interface zones zones and and bonding bonding mechanism mechanism have have beenbeen clarified.clarified. DueDue to to the the elemental elemental mixing mixing in the in fusedthe fused WE43 WE43 surface surface layer and layer the and elemental the elemental diffusion diffusionduring the during solidification, the solidification, the interface the zone interface tends zone to form tends Mg–Al-RE to form systems. Mg–Al-RE The systems. study reported The study that reportedthe transition that the interface transition is mainly interface composed is mainly of lamellar-like composed of and lamellar-like particulate Aland2RE particulate phases, Mg Al242REY5 phases,eutectic Mg phases,24Y5 eutectic coarse phases, Y-rich phases, coarse Y-rich and Mg phases,17Al12 andeutectic Mg17Al phases.12 eutectic A good phases. transition A good interface transition is interfaceformed, andis formed, the average and interfacethe averag sheare interface strength shear is 108 strength MPa. However, is 108 MPa. the distributionHowever, the of thesedistribution phases ofat these the transition phases at interface the transition is not interface uniform, is and not theuniform, properties and the are properties not stable. are Thus, not stable. it is necessary Thus, it tois necessaryoptimize theto optimize interfacial the microstructure interfacial microstr and propertiesucture and by properties heat treatment. by heat treatment.

Metals 2018, 8, 971 12 of 15

4.1. Strengthening Mechanism after Heat Treatment During the solution treatment, dissolution of intermetallic compounds, homogenization of dendrite, and diffusion of alloying elements will occur. The optimal solution parameter of 788 K for 12 h is chosen to gain a more uniform microstructure for the bimetal composites. After solution treatment, the long-distance diffusion of RE (Y, Gd, Nd), Al and Zn across the interface would broaden the transition interfacial zone. The Mg17Al12 and Mg24Y5 eutectic phases at the interface are dissolved due to the low eutectic point of 709 K and 839 K, while coarse RE-rich phases at the interface undergo fragmentation, necking, spheroidization, and solution, and then distribute uniformly at the interface. Moreover, some new cuboid-shaped Mg-RE phases are generated during solution treatment. By solution treatment, the plasticity is increased due to the dissolution of the eutectic phases, and the homogenous distribution of the high melting-point Y-rich phases and Mg-RE phases, and shear strength is also enhanced, due to the solution strengthening of Al and Y. Meanwhile, it is noted that the shear strength of AZ31 is improved about 40 MPa after solution treatment at 788 K for 12 h, due to the disappearance of the eutectic Mg17Al12 phase and the reduction of casting defects. During the ageing treatment at 483 K for different times, these high melting-point RE-rich phases and Mg-RE phases are almost unchanged, indicating that these phases have relatively good thermal stability. Generally speaking, the alloy containing fine, shear-resistant, evenly oriented (the planes parallel to the prismatic or basal planes of Mg matrix) and uniformly distributed particles, with adequate volume fraction, can exhibit a very high strength. The precipitation sequences in similar alloys, including Mg-Y-RE and WE43 during ageing, have been investigated in detail 0 by other researchers [19,21–23], which involved Mg(SSSS)→β”→β →β1→β (SSSS: supersaturated solid solution). The β” metastable phase is coherent with the α-Mg matrix in a DO19 hexagonal structure (a = 0.642 nm, c = 0.521 nm). The metastable β0 precipitate, which is formed during ageing at 473–523 K, is semi-coherent, with the α-Mg in a base-centered orthorhombic structure. The β1 phase is face-centered cubic structure (a = 0.74 nm). The sequence ends with the equilibrium β phase (fcc, a = 2.223 nm). In order to obtain high strength and high plasticity of the AZ31/WE43 bimetal composite, the optimized artificial ageing parameter, i.e., 483K for 48h, is obtained. In the case of 0 the alloy aged at 483 K for 48 h, the fine, dense, and uniformly dispersed β and β1 phases at the interface are the main precipitates, as shown in Figure8. In previous studies [ 19,22], the β0 phase formed on the prismatic planes vertical to the basal plane provides the most effective obstacle to basal dislocation movement, which leads to the high strength of the T6-treated WE43 alloy. When the ageing 0 time increases, from 24 h to 48 h, which reached the over-aged time at the transition interface, β →β1 transformation is generated, and β0 precipitates are larger and reduced. Thus, the shear strength of the interface decreased a little, due to the over-aging treatment. However, the more β0 precipitates generated after peak-ageing at 483 K for 48 h, leads to the high strength for BM WE43. However, the precipitates containing Al were unable to be found at interface layer after ageing treatment by the authors. It is possible that the amount of Al at interface is small, or the precipitates were too small 0 to be found. After ageing, β and β1 precipitates were generated in the interfacial zone and pinned dislocation slip and grain boundary movement. Thus, some kink belts and dimples were detected on the tripe junctions at grain boundary. Moreover, the strength of the interface was improved.

4.2. Fracture Behaviors The shear strength of the interface zone, WE43, and AZ31, are summarized in Table1. It is found that AZ31 can be strengthened by solution, but not by ageing treatment, while the interface zone and WE43 can be strengthened by solution and ageing treatment. Moreover, the WE43 is more signiﬁcantly strengthened than the interface zone. In our previous study [11], the shear strength of the as-casting interface zone, WE43, and AZ31, was 108, 127, and 68 MPa, respectively. Cracks can initiate at the side of AZ31, due to it having the lowest shear strength. After solution treatment, the shear strength of the interface zone, WE43, and AZ31 was 118 MPa, 160 MPa, and 110 MPa, respectively, which are much higher than those of the as-cast samples. By solution treatment, the shear strength of the interface Metals 2018, 8, 971 13 of 15 zone, WE43, and AZ31 can be increased about 10, 33, and 42 MPa, respectively. The shear strength of the interface zone is similar as that of the basal material AZ31, which leads to the initial cracks of the interface generated in the zone of stress concentration, shown in Figure 10a,b.

Table 1. The shear strength of the interface zone, WE43, and AZ31 of the bimetal at as-cast, solution treatment, and aging treatment, for 24 h and 48 h.

Treatment Interface AZ31 WE43 Reference As-cast 108 MPa 68 MPa 127 MPa reference [11] Solution treatment 118 MPa 110 MPa 160 MPa This study Ageing treatment 120 MPa 101 MPa 221 MPa This study

After ageing treatment, dispersed second phases are re-precipitated along the grain boundaries, and dense β’ phases are precipitated from α-Mg matrix, which results in the increase of shear strength at the interface zone and that of WE43. The AZ31 cannot be strengthened by ageing treatment. The plasticity of the WE43 is signiﬁcantly decreased, and that of the AZ31 is increased. Due to the lowest plasticity of WE43, the interfacial cracks of bimetal composites can be initialized near the WE43, and then propagated along the interface zone, which is in accordance with the observation of fracture morphologies shown in Figure 11. Thus, it seems that the interfacial cracks of bimetal composites are generated and propagated, depending on the cooperation of the strength and plasticity of the basal materials and interface.

5. Conclusions In this study, the WE43/AZ31 bimetal composites have been prepared by inserting WE43 alloy rod into AZ31magnesium alloy melt, and subsequent T6 heat treatment. Interfacial microstructure evolution and bonding strength were evaluated and analyzed. Based on the observation and analyses, the conclusions can be summarized as follows:

(1) After solution treatment at 788 K for 12 h, most eutectic phases of matrix materials and interface are dissolved into α-Mg matrix. Some high melting-point Mg-RE cuboid-shaped phases are distributed uniformly at the interface and WE43 matrix. The strength and plasticity are increased for the bimetal composites after solution treatment. Meanwhile, it is noted that the shear strength of AZ31 could improve about 40 MPa after solution treatment at 788 K for 12 h, due to the disappearance of the eutectic Mg17Al12 phase, and the reduction of casting defects. (2) In order to achieve high strength and high plasticity of the bimetal composites, the optimized artiﬁcial ageing parameter of about 483 K for 48 h is obtained. After ageing treatment at 483 K for 0 48 h, β and β1 precipitates are generated at the interface. The interfacial strength decreased a little 0 for β →β1 transformation, due to the over-aging treatment. However, the WE43 matrix achieves the highest strength, since more β0 phases can be generated at the peak-aged treatment. The AZ31 matrix achieves better plasticity with the increase of the ageing time. Therefore, the WE43/AZ31 bimetal composite will combine the high strength of WE43, and the high plasticity of AZ31 after heat treatment, exhibiting good application prospects for industry.

Author Contributions: D.X. and K.Z. conceived and designed the experiments; D.X. and K.Z. performed the experiments; D.X., K.Z. and H.L. analyzed the data; C.Y., H.L. and J.Z. contributed reagents/materials/analysis tools; D.X. wrote the paper. Funding: This research was funded by The National Natural Science Foundation of China grant number 51671017, the Fundamental Research Funds for the Central Universities grant number FRF-GF-17-B3 and the fund of the State Key Laboratory of Solidification Processing in NWPU grant number SKLSP201835, and was fund by Beijing Laboratory of Metallic Materials and Processing for Modern Transportation. Conflicts of Interest: The authors declare no conflict of interest. Metals 2018, 8, 971 14 of 15

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