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Journal of Alloys and Compounds 696 (2017) 875e883

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Journal of Alloys and Compounds

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Microstructural evolution and mechanical properties of MgeZneYeZralloy during friction stir processing

Yaobin Wang, Yongxian Huang*, Xiangchen Meng, Long Wan, Jicai FengState Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, PR China

a r t i c l e i n f o

Article history:Received 22 September 2016Received in revised form2 December 2016Accepted 4 December 2016Available online 7 December 2016

Keywords:Friction stir processingMg alloysMicrostructurePhase transformationMechanical properties

* Corresponding author.E-mail address: yxhuang@hit.edu.cn (Y. Huang).

http://dx.doi.org/10.1016/j.jallcom.2016.12.0680925-8388/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

Here, a Mge6Zne1Ye0.5Zr casting was subjected to friction stir processing (FSP). The effect of thermaland mechanical effects on the microstructural evolution and mechanical properties of the casting wasstudied. FSP resulted in remarkable grain refinement, dissolution and dispersion of intergranular eutecticI-phase (Mg3Zn6Y) networks and strong basal texture. Based on mechanically activated effect of FSP, I-phase transformed to W-phase (Mg3Zn3Y2) and dispersed particles with a core-shell structure formed.The increase of travel speed caused greater grain refinement and higher fraction of dispersed particles,which greatly improved yield strength, ultimate tensile strength, and elongation, 93.1%, 53.0%, and151.4% higher than that of the cast materials, respectively.

© 2016 Elsevier B.V. All rights reserved.

1. Introduction

As the lightest structural metallic materials, magnesium alloysare potential in transportation industries because of their highspecific strength [1]. Rare-earth (RE) elements are added intomagnesium alloys in order to achieve excellent high-temperaturestrength and creep resistance, and thus to broaden their elevatedtemperature application [2]. It should be noticed that currently,magnesium alloys containing RE elements are fabricated by castingprocessing. The obtained cast components often have coarse grainsand intergranular eutectic networks, resulting in their poor me-chanical performance. Therefore, it is necessary to modify the mi-crostructures of cast Mg alloys to enhance their strength andductility.

Friction stir processing (FSP) is a novel severe plastic deforma-tion technology derived from friction stir welding (FSW) [3]. DuringFSP, coupling of thermal and mechanical effects generated by therotating tool can efficiently homogenize and refine microstructure.Not only grain refinement, intergranular eutectic phase can also bebroken up into tiny particles and some can dissolve into Mgmatrix.As a result, FSP can refine coarse casting microstructure andstrengthen Mg alloys with a combination of fine-grain

strengthening, second phase strengthening (including dispersionstrengthening or/and precipitation strengthening) and solutionstrengthening [4e7]. Combined with heat treatment, FSP is aneffective method to control microstructure evolution such as grainsize, morphology and amount of the precipitates and solution de-gree [8,9]. In addition, FSP develops a special texture characteristicof Mg alloys that basal planes alignedwith the pin surface at the stirzone (SZ), which exerts a great role on mechanical properties[10,11].

MgeZn-RE alloys are attractive because of different types ofsecond phase related to various Zn/RE ratios [12,13]. FSP of MgeZn-RE alloys has been widely studied to investigate the influence ofsevere plastic strain on second phasemodification and achieve highmechanical properties. Yang et al. [14] intensively studied the in-fluence of temperature, plastic strain and cooling rate on the dis-tribution of long-period stacking ordered (LPSO) phase in FSPedMgeGdeYeZneZr alloy, and successfully developed a specialmicrostructure of fine LPSO lamellae only existedwithin fine grains,which greatly improved strength and ductility [15]. Xie et al. [16,17]investigated FSW and FSP of MgeZneYeZr alloy and found thatbulky I-phase were broken up and dispersed with some of themtransformed to W-phase. Fine equiaxed recrystallized grains anddispersed W-phase particles were beneficial to the enhanced ten-sile properties and the excellent superplasticity of MgeZneYeZralloy. However, transformation mechanism from I-phase to W-phase still remains unanswered. The intrinsic relationship between

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883876

processing parameters, microstructure modification and propertiesimprovement needs to be further explored as well. In this paper,different processing parameters of FSP were applied to MgeZ-neYeZr castings to study the influence of thermal and mechanicaleffects on the microstructure modification and tensile propertiesimprovement.

2. Experimental procedures

Mge6Zne1Ye0.5Zr (wt.%) cast billet was used as base metal(BM). A tool with a shoulder 16 mm in diameter and a threadedcylinder pin 7 mm in diameter and 2.5 mm in length was used.4 mm thick plates were cut from the as-received BM and thensubjected to one-pass FSP at a constant rotation rate of 800 rpm andvaried travel speeds of 20, 80 and 200 mm/min (hereafter denotedas 800/20, 800/80 and 800/200, respectively). The principle di-rections of FSP samples were marked as PD (process direction,along the FSP direction), TD (transverse direction, transverse to theFSP direction), and ND (normal direction, normal to the platesurface).

Microstructural evolutions were conducted on the cross-sectionof the samples, perpendicular to the PD, by optical microscopy(OM; Olympus-MPG3), scanning electron microscopy (SEM; FEIQuanta 200FEG) equipped with energy dispersive spectroscopy(EDS) and electron backscatter diffraction (EBSD), X-ray diffraction(XRD; D/max-RB), differential scanning calorimetry (DSC;STA449F3), and transmission electron microscopy (TEM; TalosF200x). The specimens for OM were prepared by mechanical pol-ishing and etching using a solution of 6 g picric acidþ10 mL aceticacidþ70 mL ethanolþ10 mL water. The SEM specimens were pre-pared by mechanical polishing. The grain sizes and particle volumefractions were counted by Image-pro plus 6.0 software. The spec-imens for EBSDwere prepared bymechanical polishing followed byelectropolishing (40 vol% H3PO4 and 60 vol% ethanol at 20 �C) usinga voltage of 2e3 V for 10e20 s. Thin foils for TEM were ion-milledby a PIPS691 miller.

Tensile specimens with a gauge length of 2.5 mm, a width of1.4mm, and a thickness of 0.8mmweremachined from BM and FSPsamples. The FSP specimens weremachined along TDwith the gagebeing completely at SZ. The room temperature tensile tests wereperformed using an INSTRON 5965 mini tester at an initial strainrate of 1 � 10�3 s�1. The fracture surfaces were examined by SEM.

3. Results

3.1. Microstructure characterization

3.1.1. Grain refinementFig. 1 shows the microstructures of BM and center of SZ. The BM

is characterized as typical as-cast microstructure with coarse grainsand intergranular eutectic networks (Fig.1a). The SZs are composedof equiaxed and fine grains. The average grain size decreases withthe increase of travel speed. Fig. 2 presents the grain size distri-butions of FSPed samples at various parameters. The average grainsizes of 800/20, 800/80 and 800/200 samples are determined to be3.20 ± 1.42 mm, 2.37 ± 1.48 mm, 1.65 ± 1.12 mm, respectively.

3.1.2. Transformation of the second phasesFig. 3 shows the morphology transformation of second phase

after FSP. Coarse eutectic networks have been broken up into smallparticles and dispersed in SZ. Fig. 4 shows the variation tendencythat particle volume fraction increases with the increase of travelspeed. Table 1 lists the EDS analyses of second phases of BM and SZof 800/200 sample in Fig. 3. In BM, EDS results of several locationsshow that the Zn/Y ratio of eutectic phase is close to 6, indicating

that the eutectic networks mainly consist of I-phase. However, Zn/Yratio of large particles in SZ was close to 2.5, which implied that I-phase might transform to other phases. Fig. 5(a) shows the XRDresults of BM and SZ of three FSPed samples. In BM, a-Mg and I-phase are detected, whereas a-Mg andW-phase are detected in SZ.It indicates that I-phase transforms to W-phase during FSP. InFig. 5(b), DSC analyses show that there are two endothermic peakslocating at 454 �C and 520 �C in BM and one endothermic peaklocating at 513 �C in SZ of 800/200 sample. The endothermic peak of454 �C and 520 �C corresponds to the dissolution of I-phase and theeutectic temperature of W-phase, respectively [16,18]. The DSCresults reveal that all I-phase transforms to W-phase after FSP. Inaddition, W-phase detected by DSC in BM is few, and cannot befigured out by EDS and XRD.

Fig. 6(a) is a high angle annular dark field (HAADF) image of SZof 800/200 sample, in which large and tiny particles distributediscretely. Fig. 6(bef) shows the element distribution of HAADFimage. Despite of Mg matrix, it is obvious that all particles containZn element. There is a distinct difference between distributions of Yand Zr elements that nearly all Y element exists in large particleswhile Zr element mainly exists in tiny particles. Fig. 7(a) and (b) arethe typical EDX spectrums of the representative particles marked inFig. 6(a). The corresponding EDX result of point 1 shows moderateamount of Mg, Zn and Y elements, and a small quantity of Zr ele-ments. The Zn/Y ratio is 2.6, corresponding to EDS results in Table 1.Combined with the XRD and DSC results, large particles are W-phase. Point 2 contains moderate amount of Mg, Zn and Zr ele-ments, and a small quantity of Y elements. It could be speculatedthat these tiny particles contain Zr elements.

Fig. 8 is the high-resolution TEM (HRTEM) image of a typicallarge particle. Obviously, the large particle show a core-shellstructure, where core part is pointed by large arrow. The FastFourier Transformation (FFT) pattern from the periphery of theparticle shows that the ‘shell’ of the particle is typical face-centeredcubic W-phases.

3.1.3. Texture evolutionFig. 9 shows the {0002}, {10-10} and {11-20} pole figures in the

SZ center of 800/200 sample. Because of limited slip systems,texture can be easily formed in magnesium during hot processing[19]. During FSP, the texture distribution is affected by the toolshoulder and pin. The shoulder provides the compressive stress butonly affects the top part of the processing zone [20]. The rotatingpin generates a shear deformation and results in the typical texturein the SZ with the (0002) planes surrounding the pin surface [21].The present EBSD result displays strong basal texture with the c-axis nearly parallel to PD, inferring the formation of typical FSP-Mgtexture in SZ. In addition, Yelement canweaken and randomize thetexture of SZ and result in the deviation of c-axis of some grainsfrom PD (shown in {0002} pole figure).

3.2. Tensile test results

The room temperature uniaxial tensile properties of BM andFSPed samples are summarized in Fig. 10. The as-cast BM exhibits ayield strength (YS) of 88.6 MPa, an ultimate tensile strength (UTS)of 196 MPa and an elongation (El) of 10.9%. All FSPed samplesexhibit higher strength and elongation than BM and the total me-chanical properties improve with the increase of travel speed. Forthe parameter of 800/200 sample, the yield strength, ultimatetensile strength and elongation improves by 93.1%, 53.0% and151.4%, respectively, compared with BM.

Fig. 11(a) shows the engineering tensile curves of BM and FSPedsamples. Clearly FSP favors the enhancement of the mechanicalproperties, and the mechanical properties gradually increases as

Fig. 1. Microstructures of MgeZneYeZr samples: (a) BM, (b) 800/20, (c)800/80 and (d) 800/200.

Fig. 2. Grain size distributions of FSPed samples at various parameters.

Fig. 3. SEM images of second phase i

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883 877

the travel speed increases. All samples show a distinct strainhardening stage during tensile tests, which exerts a great role onmechanical properties. Fig. 11(b) is normalized work hardeningcurves of BM and FSPed samples. Normalized work hardening rateis defined by the following equation [22]:

Q ¼ 1s

�vs

vε

�ε

·(1)

where s, ε and ε

·, are the true stress, true strain and strain rate

(1 � 10�3 s�1 in this work), respectively. In Fig. 11(b), all curvesexhibit an initial rapid drop due to the short period of elastoplastictransition. After then, all curves enter a steady stage III hardeningrange [23]. The inset of Fig. 11(b) shows the evident comparison ofthe stage III of various samples. It can be seen that with the increaseof travel speed, the slope of normalized work hardening rate curvesdecreases, exhibiting more uniform strain.

3.3. Fractography

Fig. 12(a) and (b) show the top surface of fractured tensile

n (a) BM and (b) SZ of 800/200.

Fig. 4. vol fraction of particles with different FSP parameters.

Table 1Results of EDS analysis of locations marked in Fig. 3 (At. %).

A B C D E F G

Mg 55.87 56.74 63.60 65.27 67.41 80.71 38.32Zn 36.79 35.95 30.97 29.60 26.97 13.69 44.60Y 7.34 7.05 5.43 5.13 5.61 5.60 17.08Zn/Y 5.01 5.10 5.70 5.77 4.81 2.44 2.61

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883878

sample of BM. The crack propagation path is roughly perpendicularto the tensile axis (Fig. 12(a)) and is mainly propagated along grainboundaries (Fig. 12(b)). A detail observation of the fracture asshown in Fig. 12(c) reveals the intergranular fracture characteristicssuch as cleavage steps (pointed by black arrow) and pits whosesizes are parallel to coarse cast grains (pointed by white arrow).

Fig.13(a) shows the top surface of the fractured tensile sample ofFSPed 800/200 sample. The crack propagation path is 45� deviatedfrom tensile axis. Fig.13(b) shows that themorphology of fracture isnot uniform, and can be roughly divided into two parts. Region I isclose to the edge of SZ and it is full of dimples (Fig. 13(c) and (d)),inferring the ductile fracture behavior and crack initiation. Region IIis close to the center of SZ with relatively smooth fracture surface. It

Fig. 5. Phase analysis results of (a) X

can be speculated that the crack initiates near the edge of SZ andpropagates along the path of 45�deviated from tensile axis to thecenter of SZ.

4. Discussions

4.1. Microstructural evolution

4.1.1. Relationship between microstructure evolution and FSPparameters

Rotation rate and travel speed are crucial parameters of FSP,which directly influence the thermal cycles and strain rate in SZ andlead to the final microstructure. In the present study, only travelspeed increased from 20 to 200 mm/min with a constant rotationrate, and both heat input and strain rate decreased in this trend[24]. The final microstructure including grain size and fraction ofdispersion phase are closely affected by these thermal-mechanicalparameters. The final grain size is the competition result of grainrefinement and grain coarsening. Increasing strain rate promotesthe grain refinement while decreasing heat input helps to reducethe grain coarsening. The results in Fig. 2 show that the effect ofheat input is larger than strain rate in controlling grain size.

During FSP process, drastic stirring effect breaks up the eutecticnetworks into small particles and part of the second phase dis-solves into matrix. Decreasing travel speed will improve both heatinput and strain rate, which strengthens the solid solution of sec-ond phase. It corresponds to the particle volume fraction variationwith travel speed shown in Fig. 4. In addition, the dispersion par-ticles helps to hinder grain coarsening and larger fraction ofdispersed particles contributes to the smaller grain size at the travelspeed of 200 mm/min. As a result, the microstructure of highertravel speed is characterized by finer grain size and larger fractionof dispersion particles.

4.1.2. Transformation of the second phase in MgeZneYeZr systemSome researches have reported the transformation of I-phase to

W-phase during heat treatment process [25,26]. Liu et al. [26] tooka systematical investigation on the transform of I-phase and foundthat I-phase began to transform to W-phase at 400 �C in TEMsamples and 447 �C in bulk samples and the growth of W-phasewas controlled by diffusion. During the FSP of cast rare-earth Mg,the peak temperature of SZ can reach 520 �C [24] corresponding tothe eutectic temperature of W-phase, which guarantees the tem-perature requirement of phase transformation during FSP. Zhang

RD patterns and (b) DSC curves.

Fig. 6. (a) HAADF image of SZ of 800/200, (b) corresponding mixed EDX mapping, and (cef) EDX element mappings for Mg, Zn, Y and Zr elements.

Fig. 7. EDX spectrums of particles in Fig. 6 (a).

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883 879

Fig. 8. HRTEM image of large particle at SZ of 800/200. The inset is FFT patternrecorded from the framed region.

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883880

et al. [27] compared the reaction of Al and Ti during annealing andFSP, and found that mechanically activated effect of FSP substan-tially accelerated the reaction. They ascribed the reaction promo-tion to the pipe diffusion by dislocations, lower activation energycaused by more reaction interface of finer particles and activatedatoms of reactants during severe plastic deformation process. Thecore-shell structure of large particles in SZ (shown in Fig. 8)

Fig. 9. Pole figures of the cen

demonstrates the element diffusion during phase transformation.During FSP, eutectic I-phase is broken up into small particles andmore reaction interfaces are generated. High density of dislocationswould be introduced near small particles and diffusion would besubstantially accelerated through pipe diffusion. The activatedatoms in FSP decrease activation energy and promote the phasetransformation from I-phase toW-phase. Enough high temperatureand mechanically activated effect of FSP substantially promote thephase transformation and result in the transformation of all I-phaseto W-phase in SZ. The core part of the large particle is likely to bethe remnant of I-phasewith high Zn containment, which causes theZn/Y ratio of EDS and EDX results higher than 1.5. Although someresearchers have reported the existence of Zr-containing particlesin MgeZneYeZr alloy systems [28e30], the characterization andunderstanding of their effects still need further research.

4.2. Improvement of mechanical properties

4.2.1. Strengthening of MgeZneYeZr materialThe BM sample exhibits the lowest yield strength and ultimate

tensile strength. It is mainly caused by its coarse grains and inter-granular eutectic networks where cracks nucleation and debondingfrom matrix are preferred. FSP improves total mechanical proper-ties in both strength and ductility, which is strongly related to themicrostructural modification. During FSP, coarse grains are refinedto several micrometers through dynamic recrystallization, whichintroduces great fine-grain strengthening. Eutectic networks arebroken up into small particles and disperses. The dispersed parti-cles could pin the mobile dislocations and thus hinder their

ter part at SZ of 800/200.

Fig. 10. Yield strength, ultimate tensile strength and elongation of BM and FSPedsamples.

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883 881

movement, which introduces dispersion strengthening. In addition,the combination of thermal and mechanical effects dissolves aportion of the second phase into matrix, which introduces solutionstrengthening. The strong basal texture also contributes to thestrengthening of FSPed materials. The Schmid factor on basal slipsystem with respect to TD is relatively low because of strong basaltexture [31], which impedes the operation of basal slip and resultsin higher yield strength. All above factors strengthen the f FSPedmaterials compared with as cast material.

The higher travel speed possesses higher yield and ultimatetensile strength. With the increase of travel speed, grain size

Fig. 11. Tensile properties of BM and FSPed samples. (a) engineering stress-

Fig. 12. Fractography of BM. (a) macrography of top surface,

becomes smaller and the fraction of dispersed particles increases. Itimproves the fine-grain strengthening and dispersion strength-ening. However, the increase of travel speed decreases the disso-lution of the second phase and results in a lower degree of solutionstrengthening. According to the tensile results, the increase of fine-grain strengthening and dispersion strengthening are obviouslygreater than the decrease of solution strengthening. Based on thediscussion of strengthening related to microstructure, it could beconcluded that FSP parameter of high travel speed with low heatinput contributes to the fine-grain strengthening and dispersionstrengthening, which is beneficial to the total mechanicalproperties.

4.2.2. Toughening of MgeZneYeZr materialThe BM sample is characterized by coarse grains and eutectic

networks which are detrimental to ductility. However, the elon-gation of BM sample in the present study is pretty high comparedwith other cast rare earth magnesium such as GW103 [5] andMgeNdeY alloy [32]. The excellent ductility of BM is mainlyascribed to the existence of I-phase. I-phase is quasicrystal phaseand has a coherent interface structure with Mg matrix [33]. Theatomic bonding between I-phase and a-Mg is rigid which woulddelay the debonding between I-phase and matrix during tensiletest.

By FSP, dispersed particles pin the mobile dislocations andweaken the dynamic recovery of grains. With the accumulation andpropagation of dislocations, the work hardening ability improves. Itresults in the reduction of slop in the stage III, which could enhancethe ductility of materials [7]. Furthermore, Xin et al. [31] studied theinfluence of the strong basal texture around the pin surface on

strain curves and (b) work hardening rates as a function of true strain.

(b) micrography of top surface and (c) fracture surface.

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883882

ductility and found that such texture combined with fine grainscould help to activate non-basal slip systems during tensile tests,which could promote the deformation compatibility and enhanceductility.

As compared with curves in Fig. 11(b), the stage III of 800/20sample finishes before normalized hardening rate reaches 1 whichis regarded as necking demarcation while the stage III of 800/200sample lasts further longer. It can be concluded that with the in-crease of travel speed, the slope of the stage III decreases and leadsto the increase of uniform stain, which is mainly due to theincreased fraction of dispersed particles. The increase of dispersedparticles helps to reduce low strength and defect-rich regionswhich provides stress concentration and degraded ductility. Yuanet al. [34] investigated the influence of grain size on work hard-ening behavior and the results showed that work hardeningbehavior is almost identical for grain sizes ranging 2.8e5.4 mmwhen tested in TD. Hence the influence of grain size in the presentFSPed samples on the work hardening behavior doesn't differ a lot.Work hardening behavior is mainly improved by the dispersed

Fig. 13. Fractography of FSPed sample of 800/200. (a) macrography of t

particles and results in the enhancement of ductility.

4.3. Fracture behavior

Fracture behavior can reflect the mechanical performance. Theintergranular fracture behavior (Fig. 12) is caused by as castmicrostructure of coarse grains and eutectic networks which woulddamage themechanical properties especially ductility. This fracturebehavior suggests a brittle fracture even though the elongation canreach 12%. FSPed samples exhibit a heterogeneous fracture surfacewhich suggests the fracture initiates near the edge of SZ andpropagates to the center of SZ. This fracture behavior is associatedwith the texture distribution of SZ. The present study shows thestrong basal planes surrounding the pin surface in SZ. The basal slipis easy to activate at the edge of SZ but difficult at the center of SZ[31]. During tensile testing, basal slip is firstly operated at the edgeof SZ and dislocations interacting with particles results in thedimples on the fracture surface (Fig. 13(c) and (d)). The propagationpath of cracks is roughly along the onion ring layers at initiating

op surface, (b) fracture surface, (c)(d) Region I, and (e)(f) Region II.

Y. Wang et al. / Journal of Alloys and Compounds 696 (2017) 875e883 883

position, as shown in Fig. 13 (a). It might be ascribed to heteroge-neous dissolution extent and particle fraction in different positionof a layer [6]. After crack propagates, the propagation directionremains unchanged and the propagation path is not along the on-ion layer in the latter process. The basal slip operation is difficult inthe center of SZ and dimples are less in fracture surface. Addi-tionally, the rapid crack propagation is also a reason for a smoothfracture surface as shown in Fig. 13 (e) and (f).

5. Conclusions

Microstructural evolution and its effect on mechanical proper-ties of MgeZneYeZr alloy after FSP are investigated in the presentstudy. The following conclusions are drawn:

1. FSP causes great grain refinement, dissolution and dispersion ofeutectic networks and strong basal texture around pin surface.The dynamic recrystallized grain size decreases and volumefraction of dispersed particles increases with the increase oftravel speed, mainly resulting from the decrease of heat input.

2. The I-phase in BM transforms to W-phase during FSP and formslarge particles with a core-shell structure. Mechanically acti-vated effect induced by FSP promotes the transformation.

3. FSPed Mge6Zne1Y-0.5Zr alloy samples exhibit better me-chanical properties including yield strength (93.1%), ultimatetensile strength (53.0%) and elongation (151.4%) than the castone, due to the fine-grain refinement strengthening, solutionstrengthening, dispersion strengthening and texturestrengthening.

4. The greater grain refinement and higher fraction of dispersedparticles are beneficial for the pronounced ductility.

Acknowledgements

This work was jointly supported by the National Natural ScienceFoundation of China (No. 51575132) and the Fund of National En-gineering and Research Center for Commercial AircraftManufacturing (No. COMAC-SFGS-2016-33214).

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