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Materials Chemistry and Physics 143 (2014) 1178e1190

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Materials Chemistry and Physics

journal homepage: www.elsevier .com/locate/matchemphys

Microstructural evolution and mechanical properties of Mgcomposites containing nano-B4C hybridized micro-Ti particulates

S. Sankaranarayanan a, R.K. Sabat b, S. Jayalakshmi a, S. Suwas b, M. Gupta a,*

aDepartment of Mechanical Engineering, National University of Singapore (NUS), 9 Engineering Drive 1, Singapore 117 576, SingaporebDepartment of Materials Engineering, Indian Institute of Science, Bangalore 560012, India

h i g h l i g h t s

� Micro-Ti particulates are hybridized with varying weight fractions of nano-B4C.� The hybrid mixture was used as hybrid reinforcements in magnesium.� Microstructure and mechanical properties of Mg-(5.6Ti þ x-B4C)BM are compared with Mg-5.6Ti.� Electron back scattered diffraction (EBSD) analysis conducted to study the microtexture evolution.

a r t i c l e i n f o

Article history:Received 18 July 2013Received in revised form30 October 2013Accepted 8 November 2013

Keywords:Composite materialsElectron microscopy (SEM)Electron diffraction (electron back scattereddiffraction)Mechanical properties

* Corresponding author. Tel.: þ65 6516 6358.E-mail address: mpegm@nus.edu.sg (M. Gupta).

0254-0584/$ e see front matter � 2013 Elsevier B.V.http://dx.doi.org/10.1016/j.matchemphys.2013.11.019

a b s t r a c t

In this work, the microstructural evolution and mechanical properties of extruded Mg compositescontaining micro-Ti particulates hybridized with varying contents of nano-B4C are investigated, andcompared with Mg-5.6Ti. Microstructural characterization showed the presence of uniformly distributedmicro-Ti particles embedded with nano-B4C particulates that resulted in significant grain refinement.Electron back scattered diffraction (EBSD) analyses of Mg-(5.6Ti þ x-B4C)BM hybrid composites showedthat the addition of hybridized particle resulted in relatively more recrystallized grains, realignment ofbasal planes and extension of weak basal fibre texture when compared to Mg-5.6Ti. The evaluation ofmechanical properties indicated improved strength with ductility retention in Mg-(5.6Ti þ x-B4C)BMhybrid composites. When compared to Mg-5.6Ti, the superior strength properties of the Mg-(5.6Ti þ x-B4C)BM hybrid composites are attributed to the presence of nano-reinforcements, the uniform distri-bution of the hybridized particles, better interfacial bonding between the matrix and the reinforcementparticles and the matrix grain refinement achieved by nano-B4C addition. The ductility enhancementobtained in hybrid composites can be attributed to the fibre texture spread and favourable basal planeorientation achieved due to nano B4C addition.

� 2013 Elsevier B.V. All rights reserved.

1. Introduction

Magnesium and its alloys have low density (w1.74 g cc�1), highspecific mechanical properties and excellent damping characteris-tics. Due to these reasons, they exhibit tremendous applicationpotential in automobile and aerospace industries, where weightreduction is critical. However, the inherent poor mechanical char-acteristics such as low elastic modulus, strength and ductility andpoor high temperature stability of magnesium and its alloys restrictits extensive utilization in critical engineering applications [1e4]. Inthis regard, magnesium metal matrix composites (Mg-MMCs)reinforced with ceramic elements in the form of fibres/particulates

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exhibit superior elevated temperature strength alongside improvedelastic modulus, hardness and wear resistance. However, theincorporation of such ceramic reinforcements in pure Mg and itsalloys often results in brittleness [3e8].

Recently, high strength, high modulus metallic elements like Ti,Ni, and Cuwere added to improve the mechanical properties of pureMg and its alloys [9e11]. When 5.6 wt.% Ti which is insoluble in Mgwas added to pure Mg, an overall improvement in yield strength by60% and ductility by 45% was reported [9]. When soluble metallicelements likeNi andCu (with limited solubility inMg)were added topure Mg, significant strength improvement was reported, but withpoor ductility [10,11]. On the other hand, an overall enhancement instrength and ductility was reported when nano-sized ceramic re-inforcements (such as Al2O3, ZrO2 or ZnO) were added to Mg and itsalloys [12e15]. Recent works have shown the positive influence ofhybrid reinforcement (prepared by mechanical alloying) on the

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mechanical response of Mgmaterials [16e19]. In our recent work onthe influence of hybrid particulate reinforcement addition (micron-sized metallic Ti þ nano-sized ceramic Al2O3) on the mechanicalresponse of Mg [19], we identified that the method of hybrid rein-forcement addition also played a major role in determining the me-chanical responseof theMg-composites (in addition to thepropertiesof the individual reinforcements and the strengthening mecha-nisms). When the hybrid reinforcement addition to Mg was carriedout after pre-processing of the reinforcements by ball milling (ratherthan direct addition of the reinforcements to Mg), better ductilityalongwith strength retentionwasachieved [19]. The improvement inmechanical properties in all the above mentioned studies wasattributed to possible change in texture and activation of non-basalslip systems [12e19]. However, no detailed investigations pertain-ing to texture evolution has been done. Such studies are relativelymeagre in the open literature, both for individual nanoscale rein-forcement addition as well as for hybrid reinforcement addition.

In the present work, micron-sized Ti particulates are hybridizedwith varying weight fractions of nano-sized B4C particulates, and thehybrid (5.6 wt.% micro-Ti þ x % nano-B4C) mixture was used as re-inforcements in pure Mg. The microtexture evolution of as-extrudedMg-composites containing (5.6Ti þ x-B4C)BM mixture has been stud-ied in detail in comparison to Mg-5.6Ti using electron back scattereddiffraction (EBSD) analysis. The results of mechanical propertiescharacterization are correlatedwith themicrostructural observationsto understand the mechanical behaviour of developed Mgmaterials.

Fig. 1. X-ray diffractograms of (a) ball milled powder and (b) bulk Mg-materials.

2. Materials and methods

2.1. Materials

Mg turnings of >99.9% purity supplied by ACROS Organics, NewJersey,USAwereusedas thematrixmaterial. Elemental titanium(Ti)particulates of size <140 mm (98% purity) from Merck and nano-sized boron carbide (B4C) particulates of particle size w50 nmsupplied byNaBond, Chinawere used as particulate reinforcements.

2.2. Preparation of hybrid reinforcements

A Retsch PM-400 mechanical alloying machine was used to ball-mill the hybrid (metal-ceramic) particulate mixture (Ti and B4C)with different B4C fraction, herein referred to as (5.6Ti þ x-B4C)BM.Prior to ball milling, the elemental particulates were blended for 1 h(with0.3wt.% stearic acidasprocess control agent) soas to ensure theuniform mixing of powder particulates. After blending, hardenedstainless steel balls of diameter 15 mmwere added and the blendedmixture was ball-milled for 2 h at 200 rpm, with a ball-to-powderratio of 20:1. Ball-milling was carried out to reduce the size of pow-der particulates and to allow it to hybridize with each other [19,20].

2.3. Melting and casting

Mg materials used in this study were synthesized through theliquid metallurgy route based disintegrated melt deposition tech-nique [12].Mg turnings togetherwitheither individualTiparticulatesor the hybrid (5.6Ti þ x-B4C)BM mixture were heated in a graphitecrucible to 800 �C inanelectrical resistance furnaceunder inert argongas protective atmosphere. In order to facilitate the uniform distri-bution of reinforcement particulates in Mg, the superheated moltenslurrywas stirred at 465 rpmusing a twinblade (pitch 45�)mild steelimpeller (coated with Zirtex 25) for 5 min. The composite melt wasthen bottom poured into the steel mould after disintegration by twojets of argon gas oriented normal to the melt stream. Followingdeposition, an ingot of 40 mm diameter was obtained. The Mg ma-terialsobtained fromtheDMDprocesswere thenmachined to36mmdiameter and soaked at 400 �C for 60 min. Hot extrusion was thencarried out using a 150 T hydraulic press at 350 �C with an extrusionratio of 20.25:1 to obtain rods of 8 mm in diameter.

2.4. Materials characterization

The presence and distribution of second phases in Mg matrix ofdevelopedMg-materialswere studied using an automated ShimadzuLAB-X XRD-6000 X-ray diffractometer (CuKa radiation,l ¼ 2.54056 �A), a Hitachi S-4300 field emission scanning electronmicroscope (FESEM) and a Jeol JXA-8530F Electron Probe Microanalyser (EPMA). An ESEM Quanta 200, FEI Field Emission GunScanning Electron Microscope (FEG-SEM) equipped with electronback scattereddiffraction (EBSD)detector,wasused to study thegraincharacteristics and crystallographic texture evolution of Mg-matrix.The grain characteristics of Mg-matrix were studied in terms ofgrain sizeand itsdistribution. The resultsofmicrostructuralevolutionstudies based on EBSD scanwere obtained in the formof inverse polefigure (IPF) maps, misorientation distribution (MD), grain boundarycharacter distribution (GBCD), kernel average misorientation distri-bution (KAM) andgrain orientation spread (GOS)map. For themicro-texture analysis, the pole figures (PF) and orientation distributionfunction (ODF)were calculated fromtheEBSDmaps. The test samplesfor EBSD scanningwerepreparedbyelectro polishing ina 3:5 volumefraction of ortho-phosphoric acid and ethanol in an electro-polishingunitwitha stainless steel cathodeat3V for30s followedimmediatelyby 2.5 V for 2 min. The EBSD scans were recorded on the surface

Fig. 2. SEM micrographs showing (a) as-received Ti particulates with sharp edges, (b) flattened (5.6Ti þ 1.5B4C)BM particulates with blunted edges after ball milling and (c) nano-B4C particulates adhered to micron-sized Ti.

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parallel to the extrusion direction (ED). TSL softwarewas used for theOIM data acquisition and analyses.

The mechanical properties of the developed Mg-materials weremeasured under indentation, tensile and compressive loading. AMatsuzawa MXT 50 automatic digital micro-hardness tester wasused to measure the microhardness of developed Mg-composites.The microhardness test was conducted on the as-polished speci-mens of extruded Mg materials in accordance with the ASTMstandard E384-99 with a Vickers indenter under a test load of25 gf and a dwell time of 15 s [21]. The tests were conducted onthree samples for each composition for 10 to 15 repeatable read-ings. To determine the tensile and compressive properties of theas-extruded samples, a fully automated servo-hydraulic mechan-ical testing machine, Model-MTS 810 was used in accordance withASTM test methods E8/E8M-08 and E9-09 respectively [22,23].The crosshead speed was set at 0.254 mm min�1 and0.04 mm min�1. For each composition, a minimum of 6 tests wereconducted to obtain repeatable values. The fractured samplesunder tensile and compressive loading of Mg-materials wereanalysed using a Hitachi S-4300 FESEM.

3. Results and discussion

3.1. X-ray diffraction studies

The results of X-ray characterization studies conducted on theball milled powder and bulk Mg-materials are shown in Fig. 1.

The X-ray diffractogram of ball milled (5.6Ti þ x-B4C)BM powder(Fig. 1a) shows prominent peaks corresponding to Ti and B4C.However, after the incorporation of ball milled powder into theMg-matrix, the crystalline peaks corresponding to Ti and B4C arenot prominently observed in the developed Mg-(5.6Ti þ x-B4C)BMbulk material (Fig. 1b). This could either be attributed to theabsence or relatively low volume fraction of intermetallic phasesin the Mg matrix [24].

3.2. Microstructure

3.2.1. Distribution of reinforcements/second phasesThe results of microstructural characterization studies con-

ducted on the powder reinforcements and the bulk Mg samples areshown in Figs. 2 and 3. Themicrostructure of as-received Ti and ballmilled (Ti þ B4C)BM powder particulates are shown in Fig. 2(aec). Itshows that the large sized (as-received) Ti particulates were flat-tened; their sharp edges were rounded off and broken down duringthe ball milling process. This can be attributed to the repeatedloading, flattening and breakdown of powder particulates duringball milling process [20]. From Fig. 2c, the nano-B4C particles areseen adhered to the Ti-particles. The distribution of Ti particles inMg-5.6Ti is shown in Fig. 3a which indicates fairly uniform distri-bution of Ti particles with less agglomeration in Mg matrix. Fig. 3bshows a representative micrograph of the Mg-(5.6Ti þ x-B4C)BMcomposite with 0.5 wt.% nano-B4C addition, in which the Ti-particles are seen to be of varying size and shape. The interface

Fig. 3. SEM micrographs showing the distribution of (a) Ti particles in Mg-5.6Ti, (b) Ti-particles of varying size and shape in Mg-(5.6Ti þ 0.5B4C)BM composite, good interfacialbonding between (Ti þ B4C)BM particles and the Mg-matrix in Mg-(5.6Ti þ 2.5B4C)BM composite.

Fig. 4. EPMA results showing the compositional distribution in Mg-(5.6Ti þ 2.5B4C)BM hybrid composite around Mg/Ti interface.

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Fig. 5. Results of EBSD measurements showing (a) grain size distribution in the developed Mg materials and (b) representative inverse pole figure micrographs of (i) Mg-5.6Ti and(ii) Mg-(5.6Ti þ 2.5B4C)BM composite.

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between the (Ti þ B4C)BM particle and the Mg-matrix is shown inthe representative image (Fig. 3c, Mg-(5.6Ti þ 2.5B4C)BM compos-ite), which indicates a good interfacial bonding between(Ti þ B4C)BM particle and the Mg-matrix.

Further, the interfacial characteristics of Mg-(5.6Ti þ 2.5B4C)BMcomposite studied using electron probe microscopic analysis areshown in Fig. 4. The results of EPMA show increased concentrationof B and C around the Ti particles which indicates the combinedpresence of Ti(B,C) phases on the Ti particle and at the Mg/Ti

Table 1Results of grain size measurements.

S. no. Material Grain size (mm)

1 Mg-5.6Ti 18 � 72 Mg-(5.6Ti þ 0.5B4C)BM 17 � 93 Mg-(5.6Ti þ 1.5B4C)BM 9 � 64 Mg-(5.6Ti þ 2.5B4C)BM 6 � 4

interface. This not only improves the reinforcement-matrix inter-facial bonding but would also efficiently contribute towards inter-face strengthening [3,25]. The chemical reaction between Ti andB4C was earlier reported by various researchers based on thefollowing favourable chemical reactions as the standard Gibbs freeenergy for these reactions is negative [25e28].

3 Ti þ B4C / 2 TiB2 þ TiC

5 Ti þ B4C / 4 TiB þ TiC

Further, it was also reported that when (Ti þ B4C) particles(containing TiB2 and TiC) were incorporated into AZ91 matrix, thephases remained unchanged and non-reactivewithMg (i.e they didnot form any new additional products with Mg) [26]. In the presentcase too, no additional Mg-based interfacial products wereobserved in the composites (Fig. 4). However, the absence of TiB2and TiC peaks in the ball-milled powder (Fig. 1, XRD) may be due to

Fig. 6. Graphs showing the (a) misorientation distribution, (b) Kernel average misorientation distribution and (c) grain boundary character distribution (GBCD) in the developed Mgmaterials.

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the low volume fraction of these phases. In contrast, in Mg-5.6Ti,there is no chemical bonding at the interface due to the negli-gible solubility of Ti in Mg [29] and the interfacial bonding presentin Mg-5.6Ti is only due to the wettability between Mg and Ti(physical bonding) [30,31].

3.2.2. Grain characteristicsThe results of EBSD measurements exhibiting the evolution of

grain structure in Mg materials are shown in Fig. 5. The averagegrain size of developed Mg-materials based on the results of EBSDmeasurements (Fig. 5) are listed in Table 1. It can be seen from Fig. 5that the microstructures of extruded Mg-materials are charac-terised by unimodal grain size distribution, i.e. one maxima for thegrain size distribution. Further, the microstructural characteristicssuch as misorientation distribution, grain boundary character dis-tribution and grain orientation spread of developed Mg-(5.6Ti þ x-B4C)BM composites are studied in detail vis-a-vis Mg-5.6Ti toidentify the influence of hybridizing micro-Ti with nano-B4C on themicrostructural evolution.

The distribution of misorientation angles which represents thechange in orientation of the specific crystal axis (reference) to theorientation of neighbour crystal axis due to deformation is shownin Fig. 6a, for all the developed extruded Mg-materials [32]. Adiffused population of (15�-180�) high angle grain boundaries(HAGBs) with small peak at w30� and large peak at w90� isobserved in Mg-5.6Ti. The small peak at w30� attributes to either

recrystallization or f1011g � f1012g double twinning (38�5�

h1120i) and the large peak at w90� is due to f1012g tensile twinboundaries (86�5� h1120i) [32,33]. In Mg-(5.6Ti þ x-B4C)BM com-posites, the intensity of the peak at 30� increases with increasingcontent of nano-B4C, thus indicating enhanced recrystallization.However, at 90�, the change in peak intensity does not vary withthe nano-B4C content. Considering that the peak at 90� representstensile twin boundaries, the higher intensity peak for the com-posite containing 0.5% B4C indicates increased formation of tensiletwins. On the other hand, the relatively lower values obtained in 1.5and 2.5B4C imply a lower tensile twin fraction. This could bepossibly due to the relatively fine grains observed in these com-positions, as it is known that the grain refinement retards twinformation and growth [34,35]. From the data analysis, nearly 3%fraction of tensile twins from the extrusion process is observed tobe present in Mg-5.6Ti while the tensile twins are not significantlyobserved in the case of Mg-(5.6Ti þ 2.5B4C)BM.

The grain boundary character distribution (GBCD) profile of Mg-5.6Ti and Mg-(5.6Ti þ x-B4C)BM composites calculated in terms ofthe individual number fraction of low angle grain boundaries(LAGBs) and high angle grain boundaries (HAGBs) are shown inFig. 6b. In Mg-5.6Ti, the number fraction of HAGBs is found to bew48%. While in case of Mg-(5.6Ti þ x-B4C)BM composites, it isfound to increase with the increase in nano-B4C content withmaximum value (w68%) observed in Mg-(5.6Ti þ 2.5B4C)BM com-posite. In general, large HAGB (15e180�) fraction corresponds to

Fig. 7. EBSD generated (a) GOS maps showing combined presence of recrystallized grains and deformed grains in (i) Mg-5.6Ti and (ii) Mg-(5.6Ti þ 2.5B4C)BM composite, (b) GOSmaps after the separation of DRX and deformed grains in (i) Mg-5.6Ti and (ii) Mg-(5.6Ti þ 2.5B4C)BM composite and (c) GBCD profile of Mg-5.6Ti and Mg-(5.6Ti þ 2.5B4C)BMcomposites after the separation of (i) DRX and (ii) deformed grains.

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Fig. 7. (continued).

Fig. 8. High magnification inverse pole figure maps superimposed with grain bound-aries and HCP unit cells of (a) Mg-5.6Ti and (b) Mg-(5.6Ti þ 2.5B4C)BM composite.

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recrystallized grains [32]. Hence, an increase in HAGB fraction(dominant recrystallized grains) observed due to nano-B4C addi-tion indicates the presence of relatively more recrystallized grains(less deformed grains) in hybrid Mg-(5.6Tiþ x-B4C)BM composite incomparison to Mg-5.6Ti. This can further be verified from theKernel Average Misorientation (KAM) criterion which indicates thedistribution of local misorientation within a grain due to the strainimparted in the material during extrusion [36]. This involved thecalculation of average orientation of each pixel with respect to itsfirst nearest neighbour (with a provision that excludes themisorientation exceeding 5�) [36]. In the present case, it is clearlyobserved that the KAM number fraction increases and its peakvalue shifts towards left with the increase in B4C content (Fig. 6c).This confirms the presence of more recrystallized grains Mg-(5.6Ti þ x-B4C)BM composite in comparison to Mg-5.6Ti, thuscorroborating with the GBCD distribution (Fig. 6b).

In order to separate the dynamic recrystallized (DRX) grains anddeformed grains, the EBSD micrographs of Mg-5.6Ti and selectedMg-(5.6Ti þ 2.5B4C)BM composite were partitioned based on thegrain orientation spread (GOS) criterion as shown in Fig. 7a. Thisinvolved the separation of deformed grains and recrystallizedgrains based on the average misorientation between all pixels in agrain Refs. [36e38]. In general, a high average GOS value corre-spond to higher geometrically necessary dislocation (GND) contentand more deformation in the sample (deformed microstructure);while low GOS value indicates lower dislocation content due torecrystallization (recrystallized microstructure). In the GOS map,the grains with orientation spread less than the average orientationspread are shown in blue colour. It represents the recrystallizedgrains without stress concentration and lattice distortion. Similarly,the deformed grains with orientation spread greater than theaverage orientation spread of the total microstructure are shown inred colour indicating the area of larger lattice distortion accom-panied by the high dislocation density [36]. Thus, the combinedpresence of near equi-axed recrystallized grains and elongatedgrains (deformed grains) can be seen in the (extrusion) micro-structures of developed Mg-materials after GOS partition (Fig. 7a).For better understanding, the partitioned GOS maps and individualboundary fractions of DRX and deformed grains (based on averageGOS value) of Mg-5.6Ti and selected Mg-(5.6Ti þ 2.5B4C)BM com-posite are shown separately in Fig. 7b. It indicates the presence ofrelatively more localized recrystallized grains in Mg-(5.6Ti þ 2.5B4C)BM composite in comparison to Mg-5.6Ti. The val-idity of the partitioning based on average GOS criterion is alsochecked by the following conditions: (a) the dynamically recrys-tallized grains should be equiaxed and having aspect ratio closer to

Fig. 9. (0002) and (10-10) pole figures of (a) Mg-5.6Ti, (b) Mg-(5.6Ti þ 0.5B4C)BM, (c)Mg-(5.6Ti þ 1.5B4C)BM and (d) Mg-(5.6Ti þ 2.5B4C)BM composite.

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1 and (b) The DRX grains contain high fraction of HAGBs anddeformed grains have low fraction of HAGBs as seen in Fig. 7c[39,40].

Further, to study the role of hybrid reinforcements on the dynamicrecrystallization, themagnifiedviewof imagequality (IQ)mapsofMg-5.6Ti and Mg-(5.6Ti þ 2.5B4C)BM composites are superimposed with

Fig. 10. 4 ¼ 90� section of orientation distribution functions after partition of DRX

grain boundaries and HCP unit cells as shown in Fig. 8(a & b)[40,41].The boundary fractions of individual grain boundaries of Mg-5.6Ti and Mg-(5.6Ti þ 2.5B4C)BM are shown as inset in Fig. 8. In theIQ maps, the LAGBs (2e5�) and HAGBs (15e180�) are respectivelyrepresented in red (inweb version) and blue colour (inweb version).Similarly, the intermediate (5e15�) medium angle grain boundariesare shown in green colour (inwebversion). Further, the second phaseparticles/reinforcements and its nearby regions are mis-indexed dueto the high dislocation density (GOS value) surrounding the particlesand are represented as black regions [36,37,40,41]. From Fig. 8, it canbe seen that the very LAGBs (2e5�) and HAGBs (15e180�) are prom-inent in comparison to theboundarieswith (5e15�). Further, incaseofMg-5.6Ti, only a few LAGB are present adjacent to the medium anglegrain boundaries (5e15�) [41]. Hence, it can be anticipated that only afew LAGB are converted to HAGB during the high temperatureextrusionprocess as the fraction ofmediumangle grain boundaries isless. This could result in an increased formation of twins near Ti par-ticles to accommodate the strain during deformation (marked as ‘1’and ‘2’ in Fig. 8a). In case of Mg-(5.6Tiþ 2.5B4C)BM, themagnified IQFmaps indicate the presence of recrystallized grains (near the secondphases/reinforcements) which are rotated by w30� around c-axisfrom the parent grains (marked as ‘1’ and ‘2’ in Fig. 8b) [41]. This couldbe attributed to the fact that the nano-B4C particles are founddistributed around the micron-sized Ti (i.e. surrounding Ti) as seenfrom EPMA results (Fig. 4), which could probably assist in dislocationpileups at these sites. In such cases, the dynamic recrystallization isexpectedtobedominant [37]. Further,manyLAGBsare foundadjacentto the medium angle grain boundaries (MAGB) which could essen-tially indicate the conversion of LAGB to MAGB and then to HAGB asthe deformation progress during extrusion. Hence, in Mg-(5.6Tiþ 2.5B4C)BM, newMg-matrix grains are formedby bothparticlestimulatednucleation (PSN)andby theconversionof LAGBstoHAGBs.As a result, an increased fraction of recrystallization is observed in thecase of Mg-(5.6Tiþ 2.5B4C)BM in comparison to Mg-5.6Ti.

3.2.3. Crystallographic textureThe results of crystallographic texture measurements based on

the EBSD scan conducted on the plane parallel to the extrusiondirection of Mg-5.6Ti and Mg-(5.6Ti þ x-B4C)BM composites areshown in Fig. 9 in the form of (0002) and ð1010Þ pole figures. It canbe seen from Fig. 9a that most of the basal planes are alignedparallel to the extrusion direction in case of Mg-5.6Ti. However, thenano-B4C hybridized Ti addition is observed to promote thespreading of weak basal texture in Mg-(5.6Ti þ x-B4C)BM compos-ites as the basal planes of Mg-(5.6Ti þ 0.5B4C)BM, Mg-(5.6Tiþ 2.5B4C)BM and Mg-(5.6Tiþ 2.5B4C)BM composites are tiltedrespectively by w(20e45)�, w(10e35)� and w(20e30)� withrespect to the extrusion direction (Fig. 9(bed)) [2,36,41].

For the quantitative texture analysis of dynamic recrystallizationbehaviour, the 4 ¼ 90� orientation distribution functions of Mg-

and deformed grains of (a) Mg-5.6Ti and (b) Mg-(5.6Ti þ 2.5B4C)BM composite.

Table 2Room temperature mechanical properties of developed Mg materials.

S. no. Material Micro-hardness

Tensileproperties

Compressiveproperties

0.2% tensileyield strength[MPa]

Ultimatetensile strength[MPa]

Tensilefracturestrain [%]

0.2% compressionyield strength[MPa]

Ultimatecompressionstrength [MPa]

Compressionfracture strain[%]

[Hv]

1 Pure Mg 48 � 1 120 � 9 169 � 11 6.2 � 0.7 65 � 4 248 � 8 19.2 � 1.12 Mg-5.6Ti 71 � 2 158 � 6 226 � 6 8.0 � 1.5 77 � 3 347 � 5 13.5 � 0.83 Mg-(5.6Ti þ 0.5B4C)BM 70 � 4 156 � 9 228 � 12 11.7 � 0.4 77 � 7 372 � 3 15.8 � 0.54 Mg-(5.6Ti þ 1.5B4C)BM 87 � 5 180 � 5 238 � 6 9.8 � 0.7 96 � 7 395 � 5 15.6 � 1.35 Mg-(5.6Ti þ 2.5B4C)BM 92 � 7 215 � 9 260 � 8 8.1 � 0.3 118 � 5 419 � 7 13.0 � 0.5

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5.6Ti and Mg-(5.6Ti þ 2.5B4C)BM composite are plotted as shown inFig. 10(a and b). It shows a maximum texture intensity at 41 ¼ 0�,4 ¼ 90� and 42 ¼ 0e60� in case of Mg-5.6Ti which indicates atypical basal fibre texture in Mg-5.6Ti [32,36]. However, in case ofMg-(5.6Ti þ 2.5B4C)BM, the texture intensity is maximum at42 ¼ 30� compared to basal fibre texture in Mg-5.6Ti. Themaximum intensity observed along 42 ¼ 0� and 42 ¼ 30� indicatesthe combined presence of DRX grains and deformed grains in bothMg-5.6Ti and Mg-(5.6Ti þ 2.5B4C)BM composite. Further, the4 ¼ 90� ODF sections are partitioned as shown in Fig. 10(c and d) inorder to separate the deformed grains and recrystallized grains[36e41]. In case of Mg-5.6Ti, the maximum texture intensity is

Fig. 11. (a) Tensile and (b) compressive response of developed Mg materials.

observed at 42 ¼ 0� and 42 ¼ 60� for the partitioned deformedgrains and no peak intensity is observed at 42 ¼ 30� for recrystal-lized grains (Fig. 10c). This indicates that during extrusion, most ofthe matrix grains underwent deformationwithout recrystallizationand hence a higher fraction of tensile twins (3%) are formed.However, in Mg-(5.6Ti þ 2.5B4C)BM, the peak intensity is observedat an angle near to 42 ¼ 30� for both the deformed grains andrecrystallized grains (Fig. 10d). This conclusively indicates thepresence of relatively more recrystallized grain in Mg-(5.6Ti þ 2.5B4C)BM and the influence of nano-B4C hybridized Tiparticulates and other related phases on the Mg-matrix grainrefinement by recrystallization.

3.3. Mechanical properties

The mechanical properties of developed Mg-materials underindentation, tensile and compressive loading conditions are listedin Table 2. It can be seen from the table that while the addition ofeither individual or hybridized Ti particulates increased thehardness values, superior hardness values are obtained when theTi particulates are added after hybridizing with nano-B4C. Amongthe hybrid addition, Mg-(5.6Ti þ 2.5B4C)BM shows the highestmean microhardness value, i.e. w90% and 30% increase whencompared to pure Mg and Mg-5.6Ti respectively. Such an increasein microhardness values can be attributed to the higher constraintoffered by the hard reinforcements and other intermetallic phasestowards localized matrix deformation during indentation [1,42e44]. The flow curves of developed Mg-materials under tensileand compressive loading are shown in Fig. 11. It clearly indicatesthe superior strength properties of developed Mg materials (incomparison with pure Mg) and the strengthening effects fromindividual Ti and hybrid (5.6Ti þ x-B4C)BM reinforcements. FromTable 2, the yield strength and ultimate strength of Mg-(5.6Ti þ x-B4C)BM composites are found to increase with the increase in B4Caddition and amongst the composites, Mg-(5.6Ti þ 2.5B4C)BMcomposite exhibit the highest strengths (Table 2). The strengthproperties improvement in Mg-(5.6Ti þ x-B4C)BM composites asmentioned above can primarily be ascribed to the presence andrelatively uniform distribution of hybrid (5.6Ti þ x-B4C)BM re-inforcements/second phases and the combined presence of Ti(B, C)intermetallic phases at the Mg/Ti interface (Fig. 4) [1,3]. In the Mg-(5.6Ti þ x-B4C)BM composites, the dominant strengtheningmechanisms are the coupled effects of (i) the presence of finehybrid reinforcement/intermetallic phases and the effective loadtransfer from matrix to second phases and (ii) the grain refine-ment due to localized dynamic recrystallization (Figs. 7 and 8). Inaddition, the mismatch in elastic and thermal expansion co-efficients between matrix and reinforcements/second phases(aMg �28.4 � 10�6 �C�1, aTi �9.17 � 10�6 �C�1, aB4C �5� 10�6 �C�1) would also contribute towards the strengtheningeffects in these composites [1,3,42]. The presence of additional

Fig. 12. Representative tensile fracture surfaces of (a) Mg-5.6Ti and (b) Mg-(5.6Ti þ x-B4C)BM (c) higher magnification image showing Mg-matrix cracking extending into Ti-particlein Mg-(5.6Ti þ 2.5B4C)BM composite.

S. Sankaranarayanan et al. / Materials Chemistry and Physics 143 (2014) 1178e11901188

(B4C) reinforcements in Mg-(5.6Ti þ x-B4C)BM composites incomparison to Mg-5.6Ti is expected to improve the strengths atthe expense of ductility. However, the strength enhancement inMg-(5.6Ti þ x-B4C)BM is accompanied by no significant adverseeffect on the ductility (under both tension and compression)which could be due to the absence/relatively lower volume frac-tion of Mg based brittle secondary phases [1,3,12].

Further, the results of EBSD based texture evolution studies aresignificant to comprehend the role of hybrid (5.6Ti þ x-B4C)BM re-inforcements on microstructural features and hence on thestrength/ductility improvement. In general,Mgmaterials deformbybasal slip which is the dominant deformation mechanism based onthe lower critical resolved shear stress (CRSS) value at room tem-perature [2,41,45]. In the present case, the results of texture analysisindicate that most of the basal planes inMg-5.6Ti are aligned nearlyparallel to the extrusion direction. For such crystallographic texture,the basal slip is difficult to be activated at room temperature and themechanical twinning is an important deformation mode whichcould result in poor ductility [45,46]. In contrast, the nano-B4C hy-bridized Ti addition is observed to promote the spreading of weakbasal texture in Mg-(5.6Ti þ x-B4C)BM composites. This couldpossibly result in a slip transition (non-basal slip) which couldimprove the deformation behaviour of the Mg materials [2,38,45e47]. Further, the higher ductility value of Mg-(5.6Ti þ 0.5B4C)BMamongst the hybrid composites could be attributed to the alignmentof basal planes which are more favourable for slip. The weak basal

fibre texture extending between w(10e90)� in case of Mg-(5.6Ti þ 1.5B4C)BM composite could result in higher Schmid factorwhich in turn improves/retains the ductilitywhen compared toMg-5.6Ti [2,38,45e47]. Similar results were obtained by Goh et al. [47]wherein the presence of carbon nanotubes in Mg-matrix resultedin the rotation of basal planes byw20� from the extrusion directionand favoured the activation of non-basal slip.

3.4. Fractography

The results of fractographic analysis conducted on the fracturedsamples of Mg-materials under tensile loading are shown in Fig. 12.It shows mixed mode fracture with prominent features showingevidence of plastic deformation in contrast to the typical cleavagetype brittle fracture as reported in case of pure Mg [30]. Alongsideductile features, Ti particle debonding is also prominently seen inMg-5.6Ti (Fig. 12a) which can be attributed to the lack of chemicalbonding at the Mg/Ti interface [30]. However, in the case of Mg-(5.6Ti þ x-B4C)BM composites, the higher magnification image(Fig. 12c) near the Ti particle show particle cracking extending fromthe Mg-matrix into the Ti particle, without any cracking at thematrix-reinforcement interface. This indicates good interfacialbonding and that the particle acting as a load-bearing member intheseMg-(5.6Tiþ x-B4C)BM composites. In such cases, the completeload transfer from the matrix to the particles would lead to largestresses on the particles. However, as particle deformation would

Fig. 13. Representative compressive fracture surfaces showing shear bands in (a) Mg-5.6Ti and (b) Mg-(5.6Ti þ 1.5B4C)BM composite.

S. Sankaranarayanan et al. / Materials Chemistry and Physics 143 (2014) 1178e1190 1189

be restricted (considering the higher dislocation density presentaround these particles, as explained before in Section 3.2), theselarge local internal stresses, upon reaching a critical value wouldlead to particle fracture. This behaviour is unlike those seen inceramic reinforcements, wherein matrix cracking, interfacecracking and debonding are prominent fracture features [48].

Under compressive loading, fracture in all the samples wasinitiated near the specimen ends and occurred at w45� anglewith respect to the compression test axis at the maximum stress.This indicates that the overall failure mechanism in all the Mgmaterials under compression loading is unchanged. Further, therepresentative fractographic evidences of Mg-5.6Ti and its hybridcomposites samples under compressive loading conditions revealthe presence of shear bands as shown in Fig. 13. It attributes tothe heterogeneous deformation and the work hardening behav-iour as the work hardening rate is faster in the case of samplesfailed by shear bands [49,50].

4. Conclusions

The microstructural evolution and mechanical properties of Mghybrid nanocomposites containing micron-sized Ti, hybridizedwith different weight fractions of nano-sized-B4C were studied.Based on the processestructureeproperty correlation, thefollowing conclusions are drawn.

1. When compared to Mg-5.6Ti, the addition of nano-B4C hybrid-ized micron-sized Ti particles promoted localized dynamicrecrystallization and resulted in refinement in grain size andincrease in microhardness values.

2. Based on microstructural evolution studies by EBSD analyses(misorientation angle distribution, grain boundary characterdistribution, grain orientation spread), theMg-(5.6Tiþ x-B4C)BMhybrid composites, show more recrystallized grains and lesstensile twins compared to Mg-5.6Ti.

3. Texture studies confirmed that the addition of nano-B4C hy-bridizedmicroTi addition toMg resulted in realignment of basalplanes and extension of weak basal fibre texture whencompared to Mg-5.6Ti, where the basal planes are nearly par-allel to the extrusion direction.

4. Under tensile loads and compressive loads, the Mg-(5.6Ti þ x-B4C)BM hybrid composites exhibit the best combination ofproperties with enhanced strength and improved/retainedductility.

5. The enhancement in strength properties under both tension andcompression of Mg-(5.6Ti þ x-B4C)BM hybrid composites can beattributed to the presence of nano-reinforcements, the uniformdistribution of the hybridized particles, better interfacialbonding between the matrix and the reinforcement particlesand texture modification achieved by nano-B4C addition.

6. The higher/retained ductility values of Mg-(5.6Ti þ x-B4C)BMcomposites in comparison to Mg-5.6Ti corresponds to thefavourable basal planes alignment and the spreading of weakfibre texture resulted from the nano-B4C hybridized Ti addition.

Acknowledgements

One of the authors, Mr. S. Sankaranarayanan, sincerely thanksthe NUS Research scholarship for supporting this research for hisgraduate program.

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