Effect of Friction Stir Processing on the Microstructural Evolutionand Tensile Behaviors of an ¡/¢ Dual-Phase Mg­Li­Al­Zn Alloy

Chung-Wei Yang+

Department of Materials Science and Engineering, National Formosa University,No. 64, Wunhua Road, Huwei, Yunlin 63201, Taiwan, ROC

The effect of friction stir processing followed by an aging heat treatment on an extruded ¡/¢ dual-phase Mg­8.5Li­2.8Al­1.1Zn (LAZ931)alloy is investigated. The aim of present study is to explore its microstructural characteristic and tensile mechanical properties. The friction stirprocess reduces the extruded texture and causes a significant grain size refining effect on the coarse Mg-rich ¡-phase and Li-rich ¢-phase grains.An apparent decrease in the volume fraction of ¡-phase is confirmed, and it is resulted from the dissolution and solid solution of the ¡-phasewithin the ¢-phase matrix. The ¡/¢-phase grain size refining and solid solution strengthening effects by the friction stir process can not onlyimprove the microhardness within the stir zone, but also enhance the tensile strength of extruded dual-phase LAZ931-F alloy. With performingan aging heat treatment to friction stirred LAZ931-FSP alloy at 150°C, the microhardness is further increased with an age hardening effect by theprecipitation of ¡-phase and metastable ª-MgLi2Al precipitate in the stir zone. However, the coarsening of precipitated ¡-phase at ¢-phase grainboundaries and the phase decomposition of metastable ª-MgLi2Al into a AlLi compound reduces the precipitation strengthening effect to thetensile strength of aged LAZ931-FSP/A alloy. Regardless of being aged state or not, applying a tensile force perpendicular to the stir processingdirection causes a decrease in the tensile strength and elongation to friction stirred LAZ931-FSP alloys. [doi:10.2320/matertrans.M2013344]

(Received September 9, 2013; Accepted November 11, 2013; Published December 20, 2013)

Keywords: magnesium­lithium alloy, dual-phase, friction stir process, microstructure, tensile mechanical properties, texture

1. Introduction

Magnesium (Mg) alloys are known as lightweight metallicconstructional materials in industrial applications.1,2) Mgalloys have high specific strength, high specific stiffness, wellrecyclability and radiation absorption of electromagneticwaves. It also provides high damping capacities, high thermaland electrical conductivities.2) Based on these advantages, Mgalloys are now widely used in computer, consumer electronicindustries, sports and biomedical applications.3,4) There arenumerous applications in the automobiles and aerospace ofnowadays for the purpose of reducing vehicle weight andfuel consumption.1,5,6) But the commercial applications ofMg alloys are limited due to the poor formability resultedfrom its hexagonal close-packed (hcp) structure.

The Mg­Li alloys have attracted attention as a basis forultra-lightweight metals, which can be good candidates formaking components for aerospace vehicles, such as the skinof fuselage, wings and landing frame.7,8) Adding lithium (Li)to Mg can transform the hcp structure to a body-centeredcubic (bcc) structure, substantially improving the ductilityand further reducing the density of Mg­Li alloys.7) Accord-ing to the equilibrium Mg­Li phase diagram, Li has a highsolubility in Mg. When Li content is less than 5mass%, onlyMg-rich ¡-phase (hcp) exists. Mg alloying with the additionof about 5­11mass% Li content exhibits a dual-phaseeutectic structure, which consists of a Mg-rich ¡-phase anda Li-rich ¢-phase (bcc). The Mg­Li alloy with a single Li-rich ¢-phase is obtained if Li content greater than 11mass%is added. Mg­Li alloys have better formability and vibrationresistance than commonly used Mg­Al­Zn (AZ-series)alloys.7,9) However, Mg­Li alloys exhibit low mechanicalstrength, which limits their engineering and structuralapplications. To overcome this drawback, some studiesinvestigated the effects of cold working, addition of alloying

Al, Zn and rare earth (RE) elements to improve themechanical strength of Mg­Li alloys.8,10­15)

In the present study, we use the friction stir processing tomodify the dual-phase microstructure and achieve bettertensile mechanical properties of a Mg­Li alloy. Friction stirprocessing (FSP) has been developed as a thermo-mechanicalmicrostructural modification technique of metallic materialsbased on the basic principles of friction stir welding(FSW),16) which was invented by The Welding Institute(TWI) of UK in 1991.17) During the FSP/FSW, thecontribution of intense plastic deformation and elevatedprocessing temperature result in generation of a dynamicallyrecrystallized equiaxed grains, development of texture andprecipitate dissolution within the stir zone (SZ).16) Somestudies indicated that the grain refinement and control ofprecipitation by a thermo-mechanical process, such as theequal channel angular extrusion (ECAE), can enhance thestrength of Mg­Li alloys.12,18) The effects of FSP onmechanical properties, which related to the microstructuralevolutions, grain size refining effect, precipitation anddynamic recrystallization (DRX), of Mg alloys have beenextensively reported by many studies.19­24) B. Mansoor andA. K. Ghosh reported that the tensile strength of FSPmodified ZK60 extruded plate is improved by the ultra-finegrains and precipitates.20) In addition to the grain refiningeffect, the enhanced tensile strength is also affected by thestrong basal texture within the stir zone20) because thenumber of independent slip systems in the Mg hcp structureis limited. Recently, the texture of FSP/FSWAZ-series alloysis examined by the X-ray diffraction (XRD), neutron dif-fraction and electron backscattered diffraction (EBSD).23­27)

W. Woo et al.23) and J. Yang et al.24,25) reported that boththe tensile strength and elongation of friction stir processedand welded AZ31 rolled-plate are significantly reduced. Thebasal texture variation during FSP/FSW significantly affectstensile mechanical properties, and a failure location isoccurred at the incompatible boundary between the SZ and+Corresponding author, E-mail: cwyang@nfu.edu.tw

Materials Transactions, Vol. 55, No. 2 (2014) pp. 371 to 377©2013 The Japan Institute of Metals and Materials EXPRESS REGULAR ARTICLE

the thermo-mechanically affected zone (TMAZ). S. H. C.Park et al.26,27) indicated the severe plastic deformationduring FSW results in a texture evolution which the basalplane normal is surrounding the welding tool surface in theSZ of FSW AZ61 joint. The tensile yield strength isdecreased because the failure location in the SZ near thetransition region of FSW AZ61 joint is associated with thebasal plane texture lies at 45° to the tensile direction.

Since the microstructure can be effectively refined, re-distributed and a specific texture is established with perform-ing the FSP, it is important to investigate the effect of FSP onimproving mechanical properties of the dual-phase Mg­Lialloys. However, the strengthening and failure behaviors ofFSP-modified Mg­Li alloys are still not widely clarified andstudied. Based on the above-mentioned concepts, the aimof present study is to investigate the tensile mechanicalproperties and failures of the ¡/¢ dual-phase Mg­Li­Al­Zn(LAZ) extruded alloy correlated with the re-distribution andgrain refining of ¡/¢ phases via the FSP modification andstabilization heat treatment. In addition, the effects ofdifferent tensile orientations on FSP-modified ¡/¢ dual-phase LAZ alloy are also presented.

2. Experimental Materials and Methods

The base metal (BM) used in this study was 3mm-thick as-extruded Mg­Li­Al­Zn sheets with a chemical compositionof 8.5 Li, 2.8 Al and 1.1 Zn (mass%), which was determinedby inductively coupled plasma-atomic emission spectrometry(ICP-AES) and these sheets were labeled as LAZ931-F.

The as-extruded LAZ931-F sheets were machined intoFSP specimens with dimensions of 100mm (l) © 60mm (w).During the FSP, a cylindrical rotating tool made of AISI H-13tool steel with a protruding probe plunged into therectangular LAZ931-F specimens. The shoulder and thestirring probe of rotating tool were 17-mm diameter and6-mm diameter, respectively. The stirring probe was 2-mmdepth. The tool rotational speed was set at 2500 rpm, andthe downward push force was controlled at about 2.6 kN.The downward push force and the rotational speed weremaintained for an appropriate time to generate sufficientfrictional heat. The generated frictional heat softened theLAZ931-F base metal, and the stirring probe caused materialplastic flow in both circumferential and axial directions. Therotating tool was tilted by 1.5° from the workpiece normal,and the stirring probe moved along the extruded direction(ED) of the LAZ931-F specimens at a traverse speed of about150mm·min¹1. Then the as-processed specimens, whichwere denoted by LAZ931-FSP, were 5°C-water quenchedimmediately. The plane normal of the processed direction(PD), normal direction (ND) and transverse direction (TD)of these prepared LAZ931-FSP specimens were defined asshown in Fig. 1.

Considering the Mg­Li alloys system, it is noted that asignificant natural aging of ¡-phase, metastable ª-phase orAlLi compounds from ¢-phase matrix with the addition of Li,Al and Zn elements.28­30) Therefore, the friction stirredspecimens used for microstructural observation and subse-quent tensile tests were then given a stabilization heattreatment to eliminate additional precipitation effect of

various degrees of natural aging on the microstructure andmechanical properties. The friction stirred specimens were allstabilized by artificial aging at 423K, held for 3 h and waterquenched. These specimens were designated as LAZ931-FSP/A.

The microstructures of LAZ931-F, LAZ-FSP andLAZ931-FSP/A specimens were examined with opticalmicroscope (OM). The phase composition and crystallo-graphic texture measurements were identified by X-raydiffractometry (XRD), using CuK¡ radiation at 30 kV,20mA with a scan speed of 1° (2ª)min¹1. Micro-Vickershardness test (Hv) was applied to evaluate the variations ofhardness after the FSP. The micro-Vickers hardness testacross the cross-sections (on the PD plane) of LAZ931-FSPand LAZ931-FSP/A specimens was applied using a Vickersindenter with a 980mN load for 10 s dwell time. Eachmeasured hardness datum was the average of three tests.

The tensile mechanical properties of extruded LAZ931-Fand friction stir processed specimens were measuredaccording to the standard tension testing of ASTM E8M-04.31) Figure 1 shows the dimensions of the tensile speci-mens. Uniaxial tensile tests of the LAZ931-F were conductedparallel to the ED. For the friction stir processed specimens,the tensile tests were carried out in the directions of parallel(0°) and perpendicular (90°) to the PD, and the arrangementof tensile specimens is illustrated in Fig. 1. The specimenswere tested at room temperature with an initial strain rate of1.67 © 10¹3 s¹1. Each tensile testing datum was the averageof five tests. The samples that tensile failed were examinedusing a scanning electron microscope (SEM) to observe thefracture surfaces.

3. Results and Discussion

3.1 Microstructural evolution by the FSP and agingtreatment

Figure 2 shows the microstructural features and thedifference between the extruded LAZ931-F and the FSPmodified specimens. The observed direction of the microsco-

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Fig. 1 Schematic illustrations of the arrangement of friction stir processedtensile specimens and the tensile specimen dimension.

C.-W. Yang372

py is on the TD plane and on the PD plane for the extrudedLAZ931-F base metal and the friction stirred specimens,respectively. Figure 2(a) shows the optical microstructureof the as-extruded LAZ931-F base metal, which displays atypical extruded texture with an elongated microstructurealong the extrusion direction. We can see the LAZ931-F iscomposed of an ¡/¢ dual-phase structure. The white anddark gray regions correspond to the hcp Mg-rich ¡-phaseand the bcc Li-rich ¢-phase, respectively. It can be seen thatthe ¡-phase is elongated and surrounded by the ¢-phase.Figure 2(b) shows a high magnification micrograph of theLAZ931-F. The size of single ¡-phase is observed in a rangeof about 50 to 100 µm long. It is recognized that dynamicrecrystallization of the ¢-phase should be occurred duringextrusion,30,32) and the average grain size of ¢-phase is about30.7 « 3.5 µm. In addition, the volume fraction of ¡-phase,which is quantitatively calculated by an image analyzer(OPTIMAS 6.0), is measured to be about 21.8 « 2.3 vol% forthe LAZ931-F alloy.

After the FSP, Fig. 2(c) shows the microstructural featureof LAZ931-FSP specimens within the stir SZ region in theas-processed condition for illustration. It can be seen thatthe ¢-phase is the major phase represented within the SZ.Comparing Fig. 2(c) with Fig. 2(b), it is apparent that thefriction stir process modified the coarse extruded texture ofLAZ931-F to a refined equal-axial ¢-phase microstructure,and the average grain size of equal-axial ¢-phase fine grainsis significantly reduced to about 7.9 « 0.6 µm. Moreover,only a few fine ¡-phase grains are shown in the SZ as thoseencircles in Fig. 2(c), and most of the ¡-phase grains arelocated in the ¢-phase grain boundaries. The average grainsize and the volume fraction of these ¡-phase grains are about1.2 « 0.6 µm and 1.2 « 0.4 vol%, respectively. As a result, itis worth noted that the coarse and elongated ¡-phase grainsare significantly refined and redistributed in an equal-axialsize by the FSP. During the FSP, it is reported that thepeak working temperatures on both the advancing side(AS) and retreating side (RS) can reach about 160 to

200 μm

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Fig. 2 Optical microstructures of (a) the extruded LAZ931-F base metal and (b) high magnification micrograph of ¡/¢-phases. (c) Theas-received LAZ931-FSP and (d) the LAZ931-FSP/A specimens within SZ. (e) High magnification microstructural feature of the ¡-phase precipitation within the SZ of LAZ931-FSP/A.

Effect of Friction Stir Processing on the Microstructural Evolution and Tensile Behaviors of an ¡/¢ Dual-Phase Mg­Li­Al­Zn Alloy 373

180°C.33) However, the peak temperature can reach 550°Cwithin the SZ region,34) which is generally higher than thesolid solution temperature of Mg­Li alloy.29,30) Therefore,the significant reduction in ¡-phase volume fraction can beresulted from the dissolution of ¡-phase in the ¢-phase duringthe FSP.

Figures 2(d) and 2(e) display the friction stirred speci-mens, which are given stabilized artificial aging treatment(LAZ931-FSP/A) for illustrations. Comparing Fig. 2(d) withFig. 2(c), we can see the ¡-phase is precipitated from the ¢-phase matrix after the aging heat treatment. The difference indetail of these precipitated ¡-phase grains in morphologies isexamined with a high magnification optical micrograph asdisplayed in Fig. 2(e). Precipitation of the ¡-phase grains(indicated by the triangular marks) is observed at the ¢-phasegrain boundaries. In addition, a large number of needle-like¡-phase fine grains (as those encircled in Fig. 2(e)) are alsoprecipitated within the ¢-phase grains. The precipitationeffect significantly increases the volume fraction of ¡-phaseto about 9.6 « 1.0 vol% compared with the microstructure offriction stirred LAZ931-FSP specimen displayed in Fig. 2(c).According to the above-mentioned observations, it can berecognized that the ¡-phase is dispersed and dissolved duringFSP, and then it will be precipitated after performing anartificial aging treatment.

Figure 3 shows XRD patterns of the extruded ¡/¢ dual-phase LAZ931-F alloy on the ND, ED and TD planes.Compared with the standard powder diffraction of Mg(JCPDS 35-0821), the prismatic planes of ð11�20Þ and ð10�10Þare the preferred orientations of hcp ¡-phase on the ND andED planes, respectively. The (0002) basal plane is thepreferred orientation of ¡-phase on the TD plane. As for thebcc ¢-phase, the (200) and (110) planes (compared with theJCPDS 01-1131)35) are identified as the preferred orientationsof the ND plane and of both the ED and TD planes,respectively. In addition to the strong diffraction peaks ofmajor ¡-phase and ¢-phase, relatively weak peaks are also

identified as the AlLi equilibrium phase (JCPDS 71-0362)35)

in the extruded LAZ931-F base metal.The XRD patterns obtained from the SZ region of friction

stirred specimens in LAZ931-FSP and LAZ931-FSP/Aconditions are given in Figs. 4(a) and 4(b), respectively.The diffraction pattern of LAZ931-FSP only shows sharppeaks of the bcc ¢-phase, whatever on the ND, PD or TDplanes, as depicted in Fig. 4(a). Comparing Fig. 3 withFig. 4(a), we can see the extruded texture is changed,meanwhile, the diffraction peaks of ¡-phase and AlLiphase are significantly reduced after the FSP modification.Referring to the Fig. 2(c), the XRD analysis result entirelycorresponds to the microstructural feature, which almostdisplays refined uniaxial ¢-phase grains. The disappearanceof ¡-phase and AlLi phase peaks can be resulted from thesevere dispersion, refining and dissolving effect with SZthrough the FSP. Figure 4(b) shows XRD patterns of the agedLAZ931-FSP/A specimen on the ND, PD and TD planes. Itis noted that diffraction peaks of the hcp ¡-phase are clearlyidentified and the pyramidal plane of ð10�11Þ is the preferredorientation of precipitated ¡-phase on the ND, PD and EDplanes. Moreover, another new phase is detected, which isrecognized as a precipitation of the metastable ª-phase (aMgLi2Al compound) reported in the literatures.11,28) Theappearance of the AlLi peak (see Fig. 4(a) cf. Fig. 4(b)) canbe resulted from the phase decomposition of the ª-phase aftera heat treatment.11)

The precipitation effect of LAZ931-FSP through an agingheat treatment can also be demonstrated from the peakposition of ¢-phase (200), (211) and (220) planes, as listed inTable 1. Song et al.11) indicated that the addition 3mass% Alelement causes the shift of the Li-rich ¢-phase peaks to lowerdiffraction angle because of the difference in atomic radiussize between Al (rAl = 0.1387 nm) and Li (rLi = 0.1520 nm)elements. According to the present results listed in Table 1,

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Fig. 3 X-ray diffraction patterns of the extruded ¡/¢ dual-phase LAZ931-Fbase metal on the ND, ED and TD planes.

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Fig. 4 XRD patterns of (a) the as-processed LAZ931-FSP and (b) theLAZ931-FSP/A conditions on the ND, PD and TD planes.

C.-W. Yang374

therefore, it is recognized that the diffraction peaks shift tohigher angles with the solid solution of ¡-phase after the FSP,and the peaks reduce to lower angles with the precipitation of¡-phase, ª-MgLi2Al and the presence of AlLi compound afterthe aging heat treatment. In addition to the texture variation,the above microstructural observation (see Fig. 2(e)) is alsoverified by the XRD analysis, as depicted in Fig. 4(b), inwhich a significant precipitation of the ¡-phase and theª-MgLi2Al is confirmed after performing an aging heattreatment on the friction stirred dual-phase LAZ931 alloy.However, the AlLi and metastable ª-MgLi2Al compounds arenot observed under optical micrographs probably due to theirminute quantity and small particle size.

3.2 Effect of FSP and aging treatment on mechanicalproperties

Figure 5 displays an overall cross-sectional microstructure(i.e., on the plane with a plane normal parallel to the PD) andvariation of microhardness (Hv) of the friction stir processedand stabilization aged LAZ931 alloy. The microhardnessprofiles recorded along the PD plane of friction stirredspecimens, and the dash line in the cross-section representsthe Vickers indenter testing area, which is located at 1mmdepth from the surface. The micro-Vickers hardness testshows the average microhardness of ¡-phase to be aboutHv67.5 « 2.2, while that of ¢-phase is Hv55.3 « 2.3. Theresult of LAZ931-FSP condition represents that the micro-hardness within SZ and thermo-mechanical affected zone(TMAZ) is increased and significantly higher than theaverage level of the non-stir zone (the LAZ931-F base metalregion). The increased microhardness within the SZ can beresulted from the significant grain refinement of ¢-phaseand the solid solution effect of ¡-phase during the FSPmodification. Moreover, the microhardness of LAZ931-FSP/A is higher than that of friction stirred condition (LAZ931-FSP) because the aging heat treatment results in a fine ¢-phase matrix uniformly dispersed with fine ¡-phase grainsand metastable ª-MgLi2Al precipitates (referring to the XRDpatterns in Fig. 4(b)). It can be recognized that the abilityfor age hardening within SZ of the friction stir processeddual-phase LAZ931 alloy is attributed to the precipitation of¡-phase and a few metastable ª-MgLi2Al.

A comparison on the yield strength (YS), ultimate tensilestrength (UTS) and elongation of the extruded LAZ931-F,LAZ931-FSP and LAZ931-FSP/A specimens is shown inFig. 6. It can be seen that the YS and UTS, which values areobtained from the tensile force along the PD, of the extrudeddual-phase LAZ931-F alloy is significantly enhanced by the

FSP. Referring to the above-mentioned measuring results ofaverage grain sizes and volume fraction of ¡-phase and ¢-phase from Fig. 2, the grain size of ¢-phase is significantlyreduced from 30.7 µm to about 7.9 µm. In addition, the ¡-phase is also significantly refined in an ultrafine equal-axialgrain size (1.2 µm), and a reduced volume fraction resultedfrom the dissolution of ¡-phase within the ¢-phase matrixduring the FSP. Therefore, the increase in tensile strength tothe LAZ931-FSP specimen can be related to the combinationof the solid solution strengthening of ¡-phase and the grainsize refining effect of ¢-phase (see Fig. 2(c)) according to theHall­Petch relation.

As mentioned earlier, the precipitation of ¡-phase and ª-MgLi2Al from the ¢-phase matrix within the SZ is confirmedafter performing the aging heat treatment, however, noimpressive precipitation strengthening effect on the averagetensile stress of LAZ931-FSP/A specimens. This phenome-non can be resulted from the phase decomposition ofcoherent metastable ª-MgLi2Al. The reduction in strengthen-ing effect for the LAZ931-FSP/A should be related to the

Table 1 A comparison of Bragg’s angles for the ¢-phase (200), (211) and(220) reflections in the XRD patterns of the extruded LAZ931-F alloy,friction stirred specimens in LAZ931-FSP and LAZ931-FSP/A con-ditions.

Peak position 2ª (°)

(200) (211) (220)

LAZ931-F 51.92 64.84 76.43

LAZ931-FSP 52.08 64.97 76.60

LAZ931-FSP/A 51.96 64.88 76.52

1.5 mm

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Fig. 5 The overall cross-sectional microstructure (on the PD plane) andvariation of microhardness of the friction stir processed and stabilizationaged LAZ931 alloy. The indentations are made with a spacing of 0.5mmalong the parallel dash line (AS: advancing side; RS: retreating side).

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Fig. 6 Room temperature tensile mechanical properties of the ¡/¢ dual-phase LAZ931 extruded alloy and friction stirred specimens underdifferent tensile directions, (the initial strain rate = 1.67 © 10¹3 s¹1).

Effect of Friction Stir Processing on the Microstructural Evolution and Tensile Behaviors of an ¡/¢ Dual-Phase Mg­Li­Al­Zn Alloy 375

phase transformation from the metastable ª-MgLi2Al into theequilibrium AlLi compound.28,36) In addition, the coarseningof precipitated ¡-phase at ¢-phase grain boundaries is alsoharmful to the tensile strength because the decohesion-induced cracks propagation is occurred at the weak ¡/¢-phase interface boundaries.37) The interfacial cracking willfurther reduces the elongation of friction stirred specimens.

While the applying tensile force is perpendicular to the PDfor those LAZ931-FSP and LAZ931-FSP/A aged specimens,we can see the tensile strength is lower than those FSP-modified specimens with applying tensile force parallel to thePD. In addition, the elongation of these conditions is evenlower than the LAZ931-F base metal, as illustrated in Fig. 6.This phenomenon is related to the change in tensile directionwith respect to the friction stir processing texture. Figure 7shows a typical tensile failure specimen of the friction stirredLAZ931-FSP alloy, and the fracture position is generallylocated at the TMAZ (referring to Fig. 5 cf. Fig. 7) close tothe advancing side (AS; i.e., the tool rotating direction is thesame as the PD). Since the friction stir process causes a

processing texture which results in an apparent micro-structural change on the overall cross-section, the failurebehavior of friction stirred specimens along the cross-sectional direction should be closely correlated with thetexture evolution. W. Woo et al.23) and S. H. C. Parket al.26,27) reported that the basal plane normal is roughlysurrounding the rotation pin surface. When the applied tensileforce is perpendicular to the PD, the failure generallyoccurred at the weak zone of SZ/TMAZ boundary due tothe texture change in the TMAZ.24,25) Although the micro-hardness is significantly increased with applying the FSP andaging heat treatment, the fracture occurred at the TMAZ(Fig. 7) can be resulted from the weak discontinuousinterface (referring to the cross-sectional microstructure inFig. 5) between the base metal and the SZ region with asevere plastic flow on the AS. It is recognized as thedetrimental factor to the tensile properties of those frictionstirred specimens.

Figure 8 compares the fracture morphologies of extrudedLAZ931-F alloy, LAZ931-FSP and LAZ931-FSP/A tensilefailure specimens. Figure 8(a) shows a mixed fracture surfaceof the ¡/¢ dual-phase LAZ931-F alloy, which displays abrittle cleavage at the ¡-phase (as those indicated by thetriangular marks)37) and a dimpled ruptures at the ¢-phase(as the encircled region). Figure 8(b) shows a fracture surfaceof the LAZ931-FSP specimen. It can be see that the brittlefracture region is disappeared with the reduction of ¡-phasevolume fraction, which is resulted from the dissolution of¡-phase in the ¢-phase within the SZ during the FSP (seeFig. 2(b)). As seen from the fracture morphology of agedLAZ931-FSP/A specimens shown in Fig. 8(c), it is recog-

3 mm

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Fig. 7 An illustration of the failure friction stirred LAZ931 specimens witha tensile testing direction perpendicular to the PD.

50 μμm

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(a)

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Fig. 8 SEM tensile fracture surface of the (a) LAZ931-F, friction stirred specimens in (b) LAZ931-FSP and (c) LAZ931-FSP/Aconditions, which tensile direction is parallel to the PD. (d) A typical tensile fracture surface of the friction stirred specimens tensiletested perpendicular to the PD.

C.-W. Yang376

nized that the intergranular fracture is occurred from the ¡/¢-interface decohesion and the cracks propagation along thegrain boundaries of ¢-phase. Different from those frictionstirred specimens tensile failed along the PD, Fig. 8(d)displays a typical fracture surface, which shows a flat and lessductile feature, of the specimens tensile tested perpendicularto the PD. The fracture morphology is fairly related to a weakdiscontinuous interface during FSP, and it results in a failurebehavior as shown in Fig. 7.

4. Conclusions

The effect of friction stir process and aging heat treatmenton the microstructural feature and tensile mechanical proper-ties of the extruded ¡/¢ dual-phase Mg­8.5Li­2.8­Al­1.1Zn(LAZ931) alloy is studied. The following conclusions aredrawn based on the above results and discussion:(1) The extruded texture of dual-phase LAZ931 is reduced

and the coarse Mg-rich ¡-phase and Li-rich ¢-phasegrains can be effectively refined in an equal-axial sizeby friction stir process.

(2) Friction stir process significantly results in dissolutionand solid solution effects of the ¡-phase within the¢-phase matrix. The microhardness of friction stirredLAZ931 alloy is increased with the solid solutionstrengthening of ¡-phase and the grain refining of ¢-phase.

(3) A significant precipitation of the ¡-phase and themetastable ª-MgLi2Al is demonstrated with performingan aging heat treatment on the friction stirred dual-phase LAZ931 alloy. The increasing microhardnesswithin the stir zone is resulted from the precipitationhardening effect.

(4) The tensile strength of LAZ931 alloy can be enhancedby the friction stir process. However, the tensilestrength of friction stirred LAZ931 is decreased whileapplying a tensile direction perpendicular to the stirprocessing direction.

(5) The phase decomposition of metastable ª-phase intoequilibrium AlLi compound is a cause of reducing theprecipitation strengthening on the LAZ931 alloy afterperforming an aging heat treatment.

Acknowledgement

This study was financially supported by the NationalScience Council of Taiwan for which we are grateful(Contract No. NSC 101-2221-E-150-028).

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