Journal of Alloys and Compounds 563 (2013) 165–170

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Journal of Alloys and Compounds

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Synthesis of Mg–Al2O3 nanocomposites by mechanical alloying

Jinling Liu a,b, C. Suryanarayana c, Dipankar Ghosh d, Ghatu Subhash e, Linan An a,b,⇑a Advanced Materials Processing and Analysis Center, University of Central Florida, Orlando, FL 32816-2455, USAb Department of Materials Science and Engineering, University of Central Florida, Orlando, FL 32816-2455, USAc Department of Mechanical and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USAd Department of Materials Science and Engineering, University of Florida, Gainesville, FL 32611-6400, USAe Department of Mechanical and Aerospace Engineering, University of Florida, Gainesville, FL 32611-6250, USA

a r t i c l e i n f o

Article history:Received 21 December 2012Received in revised form 26 January 2013Accepted 28 January 2013Available online 9 February 2013

Keywords:MagnesiumAluminaMechanical alloyingPhase evolutionX-ray diffractionScanning electron microscopy

0925-8388/$ - see front matter � 2013 Elsevier B.V. Ahttp://dx.doi.org/10.1016/j.jallcom.2013.01.113

⇑ Corresponding author at: Advanced Materials ProUniversity of Central Florida, Orlando, FL 32816, USA.407 882 1462.

E-mail address: Linan.An@ucf.edu (L. An).

a b s t r a c t

Mg–Al2O3 nanocomposite powders, with Al2O3 particles of 50 nm size, were synthesized by mechanicalalloying starting from a mixture of 70 vol.% pure Mg and 30 vol.% Al2O3 powders. A steady-state conditionwas obtained on milling the powder mix for about 20 h, when the crystallite size of the Mg powder wasabout 10 nm. The structural evolution during milling was monitored using scanning electron microscopy,energy dispersive spectrometry, and X-ray diffraction methods. The results showed that a mixture of Mg,Al2O3, and MgO phases were obtained on mechanical alloying. On annealing the milled powders at 600 �Cfor 30 min, a displacement reaction occurred between the Mg and Al2O3 phases, when the formation of amixture of pure Al and MgO phases was observed. Also, a reaction occurred between the initial Mgpowder and Al formed as a result of the displacement reaction, leading to the formation of Mg17Al12,Al0.58Mg0.42, and Al3Mg2 phases. Thus, the powder annealed after milling the Mg + Al2O3 powder mixfor 25 h consisted of Al, MgO and Al3Mg2 phases.

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1. Introduction

Metal-matrix composites (MMCs) combine the high-perfor-mance of the metal matrix and the superior strength and stiffnessof the ceramic reinforcement, and these find important engineeringapplications including automotive and aerospace [1]. Althoughdifferent forms of reinforcements can be incorporated into the me-tal matrix, MMCs reinforced with particles are of great interestsince they exhibit isotropic properties, are easier to manufactureand often cheaper in comparison to the continuous fiber-reinforcedMMCs. For most of the particulate MMCs, however, the sizes of theceramic particles are typically in tens to hundreds of microns,resulting in considerable reduction in ductility and toughness andineffective utilization of the strength and stiffness of the reinforce-ment. This is mainly due to the easy initiation and propagation ofcracks in the ceramic particles or at the interface.

If enhanced mechanical properties are desired, very smallparticles must be used for reinforcement. The particle type, size,morphology, volume fraction, and distribution of reinforcing parti-cles in the metal matrix play an important and critical role inenhancing or limiting the overall properties of the composite

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material. Reinforcing the matrix with much smaller particles, inthe submicron or nanometer-sized range, is one of the key factorsin producing high-performance composites, which yield improvedmechanical properties. Further, the ductility and toughness of suchMMCs can be significantly improved with simultaneous increase instrength by reducing the particle size to the nanometer range, inthe so-called nanocomposites [2,3].

Solidification processing route has been the most popular andinexpensive method for the fabrication of composites. However,it is not easy to uniformly disperse nanoscale ceramic particles ina metal matrix and clusters are formed, thus defeating the verypurpose of using them. In particular, incorporating nanoparticlesvia the liquid metallurgy route is very difficult since wetting is sopoor that even the most vigorous stirring is unable to break theagglomerates. A more effective method, especially for high volumefractions of the reinforcement, is to disperse individual nanoparti-cles through solid-state processing methods, e.g., mechanicalalloying of a mixture of metal particles and ceramic nanoparticles.

Mechanical alloying (MA) is a solid-state powder processingtechnique that involves repeated cold welding, fracturing, andrewelding of powder particles in a high-energy ball mill. In thisprocess, a small quantity of the blended elemental powder mixtureis loaded into a container along with the grinding media, and thewhole mass is agitated at a high speed for a predetermined lengthof time. During the milling process, the metallic particles gothrough the repeated sequence of cold welding, fracturing, and

166 J. Liu et al. / Journal of Alloys and Compounds 563 (2013) 165–170

rewelding. The ceramic particles are then mixed with the metallicparticles during the process. The size of the composite particlesalso gets refined during the milling process. A balance can beachieved between the welding and fracturing after a certain time,leading to a steady-state particle size distribution. The continuousmilling beyond this stage still helps to refine the grain structureand improve the uniformity of dispersion. A high degree of uniformdispersion of the reinforcement is likely to be achieved by the re-peated collisions between the grinding medium and the powders.One of the disadvantages of MA is that usually it takes a long time.Because of this it is possible that the milled powder could get con-taminated. However, when dealing with materials containing oxi-des, this should not be a serious problem.

Compared to other processing techniques, MA has been shownto be a promising technique for producing high volume fractionmetal matrix nanocomposites [4]. MA has also been shown to re-sult in displacement/exchange reactions and phase transforma-tions due to the application of mechanical energy [5,6]. Anexcellent and exhaustive review [7] and a recent book [8] summa-rize the recent information on MA.

Magnesium-based metal matrix composites (MMCs) arepromising materials for structural applications in the aerospaceand automotive industries due to their low density, reasonablyhigh strength, superior creep resistance, high damping capacity,and good dimensional stability [9,10]. Nanocomposites in whicheither or both of the components are at nanometer scales canfurther enhance the strength and creep resistance [2].Consequently, recent efforts have been increasingly devoted todevelop nanoparticles-reinforced magnesium matrix composites.These efforts have led to a variety of magnesium matrixnanocomposites reinforced with either oxide or carbide ceramicparticles, such as Al2O3 [11], Y2O3 [12], ZrO2 [13], and SiC[14,15]. But, there has been no effort on the synthesis of Mg–Al2O3 nanocomposites by MA methods. It was also felt that itwould be worthwhile to explore the synthesis of a compositecontaining a large volume fraction of the reinforcement, sinceit was shown earlier that the mechanical properties of nanocom-posites are significantly improved with increasing volume frac-tion of the reinforcement [16].

Fig. 1. X-ray diffraction patterns of the pure Mg powder after milling for 10 and20 h. It may be noted that the diffraction peak width increases with increasingmilling time. A small amount of the MgO phase has also formed in the milledpowder, possibly due to slight oxidation during milling.

2. Experimental procedure

Commercially available magnesium powder of 99.8% phase-purity and a meshsize of �325 from Alfa Aesar Corporation (Ward Hill, MA, USA) and 50 nm sizeAl2O3 powder from Buehler Ltd. (Lake Bluff, IL, USA) were used in the present study.70 vol.% of Mg powder and 30 vol.% of Al2O3 (the mole ratio is about 5:1) weremixed under argon atmosphere inside an argon-filled glove box to minimize anycontamination resulting from handling of powders in the atmospheric air. The pow-der mixture was then milled in a SPEX 8000M mill at room temperature using tung-sten carbide grinding vial and zirconia balls of 10 mm size. The weight ratio of thezirconia balls to the total powder was approximately 10:1. The powders and theballs were loaded into the vials inside an argon-filled glove box. To avoid anyunwarranted and excessive cold welding of powder particles amongst themselves,onto the internal surfaces of the vial, and/or onto the surface of the grinding med-ium during milling, about 2 wt.% of stearic acid of 98% purity from Alfa Aesar Cor-poration (Ward Hill, MA, USA) was added to the powder mix as the process controlagent. The vials were cooled by fan during milling to minimize the temperature riseof the vial to less than 50 K. Milling of the powder was carried out for times rangingfrom 5 to 25 h.

The milled powders were characterized for their crystal structure and micro-structure. X-ray diffraction (XRD, Rigaku X-ray diffractometer, Tokyo, Japan) pat-terns were recorded with Cu Ka radiation to obtain information about thenumber and nature of the phases. The peak width of the diffraction peaks was uti-lized to obtain the crystallite size of the matrix phase. A scanning electron micro-scope (SEM, Hitachi-3500N), equipped with an energy dispersive spectrometer(EDS), was used for elemental analysis. X-ray mapping was also carried out onthe milled powders to evaluate the elemental distribution.

The thermal stability of the phases in the milled powders was evaluated byannealing the milled powders for 30 min at 600 �C under a vacuum of 10 Pa or low-er in a spark plasma sintering furnace (DR SINTER�, Model SPS-1030, SPS SyntexInc., Kanagawa, Japan). The annealed powders were also analyzed using XRD.

3. Results and discussion

3.1. Crystal structure

3.1.1. Milled powderFig. 1 shows the XRD patterns of pure Mg powder milled for 0,

10 and 20 h. All the diffraction peaks in the pattern obtained fromthe unmilled Mg powder (0 h) can be identified as belonging to Mgwith the HCP crystal structure and with the lattice parametersa = 0.32089 nm and c = 0.52101 nm, which are the same as listedin standard books. With increasing milling time, the peaks werenoted to broaden and no other change was observed. From ananalysis of the peak widths in the milled powder using the Scherrerformula, the crystallite size was calculated to be 21.5 nm aftermilling for 10 h, and 20.1 nm after milling for 20 h. Thus, it couldbe concluded that the milled powder becomes nanocrystalline innature. It was also noted that a small amount of the MgO phasehad formed on milling; possibly due to the slight oxidation of Mgduring high-energy ball milling.

Fig. 2 shows the XRD patterns of the Mg–30 vol.% Al2O3 powdermixture after milling for different times. The XRD pattern of theblended Mg–30 vol.% Al2O3 powder mixture in the un-milledcondition indicated the presence of only the Mg and a-Al2O3

phases. Like in the case of pure Mg powder, here also the peakwidth increased with increasing milling time and a small amountof the MgO phase had also formed. Using the Scherrer formula,the crystallite size of the Mg phase was evaluated and the resultsare shown in Table 1. It is noted that, for equivalent milling times,the crystallite size in the composite was smaller than in the puremetal. This is understandable since the powder blend containsAl2O3, which is much harder than the Mg powder. Consequently,this also acts like a grinding medium leading to faster comminu-tion process and consequently smaller crystallite sizes. A similarsituation of a second hard phase acting as a grinding mediumand reducing the grain size down to an amorphous phase in Sihas been reported earlier [17].

From the positions and intensities of the diffraction peaks inFig. 2, it is clear that the milled powder contains Mg, Al2O3 andMgO phases. It is possible that some of the MgO phase had formed

J. Liu et al. / Journal of Alloys and Compounds 563 (2013) 165–170 167

as a result of oxidation, like in the case of pure Mg. But, it is alsopossible that, as a result of milling, a reaction has occurred be-tween Mg and Al2O3 according to the exchange/displacementreaction:

3Mgþ Al2O3 ! 2Alþ 3MgO

Thus, it is possible that the MgO formed in the milled Mg + Al2O3

powder blend is a result of two possible mechanisms. Firstly, someMgO would have formed as a result of oxidation of Mg. Secondly,some MgO would have also formed as a result of the displacementreaction. However, it appears that during milling, the intensity ofthe MgO phase diffraction peaks has increased with increasing mill-ing time and therefore it is safe to assume that the above exchangereaction has certainly occurred during milling. Hence, majority ofthe MgO formed in the powder blend would have formed as a resultof the exchange reaction.

Fig. 2. X-ray diffraction patterns of the Mg-30 vol.% Al2O3 powder mixture milledfor 5, 10, 15, 20, and 25 h. While only Mg and a-Al2O3 phases were present in theunmilled powder blend, the milled powder shows the presence of Mg, a-Al2O3 anda small amount of the MgO phase.

Table 1Crystallite size of the Mg phase in pure Mg and Mg–30 vol.% Al2O3 powder blendmilled for different times.

Milling time (h) 5 10 20Pure Mg powder – 21.5 20.1Mg–30 vol.% Al2O3 powder blend 16.1 14.7 10.0

The above exchange reaction has a large negative free energychange (DG = �125 kJ/mol at 298 K) and therefore the reaction isthermodynamically feasible at room temperature. But, it does notnormally occur at room temperature due to kinetic constraints.Solid-state reactions involve the formation of a product phase atthe interfaces of the reactants; further growth of the product phaseinvolves diffusion of atoms of the reactant phases through theproduct phase, which acts as a barrier layer preventing furtherdiffusion. Hence, the reactions usually do not occur at roomtemperature under normal conditions; and high temperatures areoften required for the reaction to happen at a reasonable rate.However, such reactions could occur during high-energy ball mill-ing at room temperature.

During milling, the particle size is reduced and clean and freshsurfaces are produced as a result of fracturing of powder particles,and the defect density is increased due to the heavy deformationinvolved. As a result of the combined effect of all these processes,diffusion is enhanced and consequently formation of the productphases occurs easily, i.e., the kinetics of the reaction are signifi-cantly faster.

3.1.2. Annealed powderFig. 3 shows the XRD patterns of the powder milled for different

times after annealing for 30 min at 600 �C. As a result of the pres-ence of very fine (nanometric-size) particles of Mg and Al2O3 in themilled product, annealing induced a further chemical reaction be-tween these two phases and one could observe the presence ofboth Al and MgO phases. It also appears that the exchange reactionreferred to above has taken place to a larger extent during anneal-ing. This is inferred from the fact that pure Al diffraction peaks

Fig. 3. X-ray diffraction patterns of the Mg–30 vol.% Al2O3 powder mixture milledfor 5, 10, 15, 20, and 25 h and then annealed for 30 min at 600 �C. As a result of theexchange reaction between Mg and Al2O3 phases, the phases present in theannealed powder are Al and MgO. Some chemical reaction also seems to have takenplace during the high-temperature annealing resulting in the formation of Mg3Al2,and Mg17Al12 intermetallic phases.

Table 2Crystal structure data and lattice parameters of the different equilibrium phases in the binary Mg–Al system [21].

Phase Approximate composition range (at.% Al) Crystal structure Pearson symbol Space group Lattice parameters

a (nm) c (nm) c/a or a

Mg(Al) 0–12.9 HCP hP2 P63/mmc 0.32094a 0.52112a 1.624a

c (Mg17Al12) 39.5–55.0 Cubic cI58 I4�3m 1.056 – –R (e) 58.0 Rhombohedral hR53 R�3m 1.03625 – 76�27.70

b (Mg2Al3) 59.7–61.5 Cubic cF1168 Fd�3m 2.8239 – –Al(Mg) 81.4–100.0 FCC cF4 Fm�3m 0.40494a – –

a Lattice parameters listed are for the pure metals.

Fig. 4. Scanning electron microscope images of the powder mixture milled for (a)15, (b) 20, and (c) 25 h.

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could be observed in the annealed powder, whereas they were notpresent at all in the as-milled powder. The intensities of the MgO

and Al phases increased with increasing milling time, due to thefurther occurrence of the exchange reaction. Additionally, it is alsonoted that due to a reaction between the Al that has formed andthe remaining Mg in the milled powder, some Mg–Al intermetallicphases, including Mg17Al12, Al0.58Mg0.42, and Al3Mg2, phases haveformed. Since the diffraction peak widths of all these phases arereasonably large, it is safe to conclude that the phases continueto be nanometric in nature. This is especially remarkable consider-ing that both Mg and Al are relatively low-melting point metalsand that annealing of the milled powders was done at 600 �C.

It is also useful to notice from Fig. 3 that the powder containsmultiple phases – some intermetallics and some oxide reinforce-ments (MgO and Al2O3). Materials with such multi-phase constitu-tion have been referred to as hybrid nanocomposites in theliterature [18,19]. Although, it was not the intention of this study,it is important to note that such a hybrid nanocomposite could beeasily obtained by MA methods. Further, the presence of suchmicrostructures is expected to make the material stronger andharder. Evaluation of the mechanical properties of these materialsis under progress and these results will be published later.

The Mg–Al binary system contains two solid solutions and sev-eral intermediate phases [20]. The crystal structure details of thesephases are listed in Table 2 [21]. While the crystal structures andlattice parameters of the b and c intermediate phases are wellestablished, there is significant confusion in the literature regard-ing the R phase. While some investigators refer to this as the Rphase, others refer to this as the e phase or the Al0.58Mg0.42 phase[22,23]. There have also been some studies [24] on mechanicallyalloyed Mg–Al powders where it was shown that the phase bound-aries between different phases get altered under non-equilibriumconditions of processing. Most recently, there have also been someinvestigations to determine the temperature range in which the Rphase is stable [25].

It is well known that non-equilibrium phases, including super-saturated solid solutions, metastable intermediate phases, andamorphous phases could form in mechanically alloyed powders.Thus, it is possible that the constitution of mechanically alloyedpowders could differ from that of equilibrium alloys. That is, thecrystal structures and lattice parameters of the phases present inmechanically alloyed powders could be different from those ofthe expected phases. Thus, in the annealed powder, we have anumber of phases, some of which are not expected under equilib-rium conditions. The fact that the alloy, even in the annealed con-dition, contains all the possible intermediate phases in the Mg–Alsystem is an indication that complete equilibration is perhapsnot achieved. This is because, for any given composition, the binaryalloy is expected to contain only two phases under equilibriumconditions.

3.2. Microstructure

Fig. 4 shows the SEM micrographs of the powder mixture milledfor 15, 20 and 25 h. It is seen that the particle size of the powdermilled for longer times (20 h – Fig. 4b and 25 h – Fig. 4c) is smaller

J. Liu et al. / Journal of Alloys and Compounds 563 (2013) 165–170 169

and more uniform than that milled for shorter time (15 h – Fig. 4a).The inset images of Fig. 4a–c show the morphology of the largeparticles in the powder mixture. The large particles shown inFig. 4b and c are due to agglomeration and/or clustering of smallerparticles, while most of the large particles in Fig. 4a are monoliths.Close examination of Fig. 4b and c reveals that the size and the sizedistribution of the powders milled for 20 h and 25 h are very sim-ilar, suggesting that the steady state, i.e., balance between fractur-ing and rewelding processes during milling, was achieved at amilling time of about 20 h. The tiny bright spots are Al2O3/MgOnanoparticles. It is seen that such particles are much more uni-formly distributed in Fig. 4b and c than in Fig. 4a. This suggests thatthe uniform distribution of nanoparticles was achieved by millingthe powder for 20 h.

Even though one assumes that the heavy deformation involvedin the MA process ensures that the constituent phases are uni-formly distributed in the microstructure, it can be easily confirmedwhether it is so by conducting elemental mapping in the SEM.Fig. 5 shows the distribution of Mg and Al in the different phasesin the powders milled for 15, 20, and 25 h. The bright spots inthe ‘‘microstructure’’ indicate the presence of the element. It is ofinterest to note that in the powder milled for 15 h, some areasseem to lack the presence of both Al and Mg elements. It is possiblethat these areas represent either porosity or heavy concentration ofoxygen as oxide particles. But, on milling for a longer time, e.g., 20or 25 h, both Al and Mg are more uniformly distributed, suggestingthat the distribution of Mg and Al2O3 is continuously improved.This also is an indirect confirmation that at the steady state condi-tion, when the particle size is stabilized, the distribution of theconstituent elements is very uniform.

Fig. 5. Elemental maps of Mg (left) and Al (right) for the powder mixture milled for(a) 15, (b) 20, and (c) 25 h.

Fig. 6 shows the SEM micrographs of the powder milled for dif-ferent times and annealed for 30 min at 600 �C. It is seen that themorphology and size of the powder after annealing were almostthe same as those in the milled powder. The size of the powdermilled for longer times (20 h – Fig. 6b and 25 h – Fig. 6c) is smallerand more uniform than that milled for shorter time (15 h – Fig. 6a)after annealing. The large particles, which were shown in the insetimages of Fig. 6a–c, arising from agglomeration and/or clusteringof smaller particles, coexist with the small particles in the nano-composite powders. The tiny spots uniformly distributed in thepowder clearly suggest that a uniform distribution of nanoparticleswas retained after annealing also.

Fig. 6. Scanning electron microscope images of the powder milled for differenttimes and subsequently annealed for 30 min at 600 �C: (a) 15, (b) 20, and (c) 25 h.

170 J. Liu et al. / Journal of Alloys and Compounds 563 (2013) 165–170

4. Concluding remarks

A homogenous distribution of the Al2O3 nanoparticles in the Mgmatrix was obtained by mechanically alloying a mixture of 70 vol.%of Mg and 30 vol.% of Al2O3 powders. The phase evolution and theirdistribution were evaluated as a function of the milling time. It wasnoted that the Mg, Al2O3, and MgO phases were present in the as-milled powder with the exchange reaction partially occurring be-tween Mg and Al2O3. All the phases had nanometric dimensionsas a result of milling. On annealing the milled powder for 30 minat 600 �C, the exchange reaction between Mg and Al2O3 had oc-curred to a large extent resulting in the formation of Al and MgOphases. Additionally, the reaction between Al and un-reacted Mgled to the formation of Mg–Al intermetallics. Formation of nano-structured phases was observed by scanning electron microscopyand the uniform distribution of the phases was confirmed by X-ray elemental mapping method. The SEM and EDS results indicatethat that a uniform distribution of the hybrid (Al2O3 + MgO) rein-forcement can be achieved after milling the powder blend for20 h. The thermal stability of the formed nanocomposite was eval-uated by annealing it at a high temperature.

It has been shown that MA can lead to the formation of nano-composites with ultrafine sizes and uniform dispersion of the rein-forcement particles. Such composites are expected to be strongerand perhaps even more ductile. Such a combination is particularlyuseful in the Mg–Al system, to improve the properties and perfor-mance of these composites.

Acknowledgement

Jinling Liu would like to thank the China Scholarship Council(Grant 2008629106) for partial support of this work.

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