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Journal of Alloys and Compounds 460 (2008) 386–391

An investigation of the quasi-static fracture behaviorof a rapidly solidified magnesium alloy

T.S. Srivatsan a,∗, Satish Vasudevan a, M. Petrorali b

a Division of Materials Science and Engineering, Department of Mechanical Engineering,The University of Akron, Akron, OH 44325-3903, USA

b Division of Research, The TIMKEN Company, Canton, OH 44706-0930, USA

Received 5 December 2006; received in revised form 19 June 2007; accepted 20 June 2007Available online 29 June 2007

bstract

In this paper, is presented and discussed the quasi-static fracture behavior of a rapid solidification processed magnesium alloy. Test specimens of

he magnesium alloy were deformed under quasi-static loading. The resultant tensile properties and fracture behavior are presented and discussedn light of the competing and synergistic influences of nature of loading, intrinsic microstructural effects, matrix deformation characteristics, and

acroscopic fracture.2007 Elsevier B.V. All rights reserved.

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eywords: Rapid solidification; Magnesium alloy; Microstructure; Deformatio

. Introduction

The use of magnesium alloys in a variety of technology-elated applications has seen a progressive growth during theast two decades, and both magnesium and its alloy counterpartsontinue to make their impact in a spectrum of automotive prod-cts. The preferential disposition towards the selection and usef magnesium is attributed primarily to its lightweight, i.e. 30%ighter than aluminum, 75% lighter than zinc, and 70% lighterhan steel [1,2]. Besides, magnesium has the highest strength-o-weight ratio [σ/ρ] of any of the commonly used non-ferrousnd ferrous metallic materials [1]. Other noteworthy advantagesn favor of choosing magnesium include its good cast abil-ty, high die casting rates, electromagnetic inference, shieldingroperties, part consolidation, dimensional accuracy, and over-ll excellent machinability, all of which favor its selection andtilization in automobile products [3–5].

The magnesium alloys produced by conventional ingot metal-

urgy (IM) technique exhibit the drawbacks of less than desirabletrength, inferior formability, low thermal stability, inadequatereep resistance, poor oxidation resistance, and inferior corro-

∗ Corresponding author. Tel.: +1 330 972 6196.E-mail address: TSRIVATSAN@uakron.edu (T.S. Srivatsan).

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925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.jallcom.2007.06.079

cture; Mechanisms

ion resistance. Furthermore, a limited number of slip systemslaces limitations on achieving enhanced strengthening coupledith a degradation in ductility or formability of the alloy. Conse-uently, extensive use of the conventional ingot metallurgy (IM)rocessed magnesium-base alloys has been restricted [6,7]. Theddition of alloying elements, such as: molybdenum, titaniumnd chromium, having high melting points that far exceed theoiling point of magnesium, was found to be beneficial. Con-equently, alloying by traditional ingot metallurgy methods wasifficult and had its limitation [7]. Further, the most commonlysed alloying elements have limited solid solubility in magne-ium and tend to form intermetallic compounds as a result ofhe electropositive nature of magnesium [8]. These limitationsere overcome by use of the technique of rapid solidificationrocessing [9–11].

Rapid solidification processing of magnesium-base alloysacilitates a departure from thermodynamic equilibrium and aidsn the preparation of high purity alloys having compositionalexibility [9–11]. The capability for extended solid solubilitiesnd improved chemical homogeneity that is achievable by rapidolidification, enabled in the preparation of alloy compositions

hat cannot be easily made using the traditional ingot metallurgyechnique. The improved compositional flexibility facilitated inmproving the corrosion behavior by minimizing the galvanicoupling between the microscopic in-homogeneities. Also, the

T.S. Srivatsan et al. / Journal of Alloys an

Table 1Nominal chemical composition of the magnesium alloy (in wt%)

Al 5.72Zn 2.96NM

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oof the Mg–5.72 Al–2.96 Zn–6.05 Nd alloy reveals the powderparticles to be deformed and elongated in the direction of defor-mation, i.e. extrusion direction (Fig. 2). The average size of thegrain is about 3 �m. High magnification observation revealed the

d 6.05agnesium Balance

resence of new non-equilibrium phases facilitates in improvinghe corrosion behavior. A conjoint and mutually interactive influ-nce of these effects results in improved mechanical and physicalroperties coupled with an elimination of redundant metal work-ng and finishing operations. The development and emergence ofigh strength magnesium-base alloys having lightweight serveds an attractive and potentially viable alternative to aluminumlloys with a concomitant savings in weight.

During the last three decades, i.e. since the early 1980s,here has been a preponderance of research activity onapidly solidified magnesium alloys [9–16]. This paper presentsicrostructural influences on the quasi-static deformation and

racture behavior of a rapid solidification processed (RSP)agnesium alloy. The mechanisms governing the quasi-static

racture behavior of the alloy are presented and discussed inight of the conjoint influences of nature of loading, intrinsic

icrostructural effects, matrix deformation characteristics, andacroscopic aspects of fracture.

. Materials and processing

The magnesium-base alloy used in this study was provided by Allied Signalorporation (Morristown, New Jersey, USA). The chemical composition of thelloy (in wt%) is given in Table 1. The material was manufactured using theowder metallurgy (PM)/rapid solidification (RS) technique. Planar flow cast-ng was used to produce rapidly solidified ribbons of the magnesium alloy. Therocessing of ribbons was conducted in an inert vacuum environment with thebjective of: (i) preventing oxidation of the liquid metal surface and (ii) pre-enting the entrapment of air under the liquid film. The ribbons, about 50 mm inidth, were then reduced to powder [sieve size 500–250 �m] using a series ofigh speed mechanical communition processes [10,12,13,17]. The final prod-ct consisted of irregularly shaped flat platelets with a thickness equal to theriginal ribbon thickness. The powder particles obtained from the process haveuniform microstructure irrespective of the particulate size [14]. To prepare a

onsolidated body the powders were either out-gassed in a can and then sealednder vacuum, or subjected to hot pressing at 473–573 K for different lengths ofime, ranging from 1–24 h, depending upon size of the billet. The cans were thenxtruded at temperatures between 473 and 573 K to round bars at an extrusionatio of 18:1. A marginal rise in temperature of the billet occurred due to exten-ive deformation induced by the mechanical working operation, i.e. extrusion.recise details of the processing technique and the precautionary methods usedre described in detail elsewhere [14].

. Experimental techniques

The initial microstructure of the as-received material, in the extruded con-ition, was characterized by optical microscopy after standard metallographicreparation techniques. The etched specimens were observed in an optical micro-cope and photographed using standard bright field technique.

Tensile specimens were precision machined such that the longitudinal direc-

ion, or major stress axis, of each specimen was parallel to the extrusion direction.hus, in each case, the gross fracture plane was essentially perpendicular to

he extrusion direction. The cylindrical test specimens, with threaded endsnd had a gage section, which measured 25.4 mm in length and 6.25 mm iniameter, conformed well with the standards specified in ASTM E-8 [18]. To

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d Compounds 460 (2008) 386–391 387

inimize the effects of surface irregularities and finish the gage section of allest specimens were mechanically ground using 600 grit silicon carbide impreg-ated emery paper in order to remove any and all circumferential scratches andurface machining marks. Uniaxial tensile tests were performed on a fully auto-ated computer controlled servo-hydraulic test machine in the room temperature

27 ◦C) laboratory air (relative humidity of 55%) environment. The specimensf the magnesium alloy were deformed at a constant strain rate of 10−4 s−1.

Fracture surfaces of the deformed and failed test specimens were comprehen-ively examined in a scanning electron microscope (SEM) to (a) determine theacroscopic final fracture mode and (b) characterize the fine-scale topography

nd microscopic mechanisms governing quasi-static fracture. The distinctionetween macroscopic and microscopic fracture mechanism(s) is based entirelyn the magnification level at which the observations are made. The samplesor observation in the SEM were obtained from the failed tensile specimens byectioning parallel to the fracture surface.

. Results and discussion

.1. Initial microstructure

A triplanar optical micrograph illustrating the grain structuref the magnesium alloy is shown in Fig. 1. The microstructure

ig. 1. Triplanar optical micrograph illustrating the grain structure of the mag-esium alloy along the three orthogonal directions.

388 T.S. Srivatsan et al. / Journal of Alloys and Compounds 460 (2008) 386–391

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ig. 2. Triplanar optical micrograph illustrating the morphology of the grains asdirect consequence of the mechanical deformation resulting from extrusion.

owder particles to comprise of well-defined grains (Fig. 3). Theverage grain size was between 2 and 3 �m. At the higher allow-ble magnifications of the optical microscope fine particles wereound to be distributed in the alloy matrix (Fig. 4). These parti-

ig. 3. Optical micrograph showing the presence of well-defined grains withinach powder particle of the Mg–5.72 Al–2.96 Zn–6.05 Nd alloy.

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ig. 4. Optical micrograph showing the distribution of fine second-phase parti-les in the microstructure of the Mg–5.72 Al–2.96 Zn–6.05 Nd alloy.

les have been identified and reported as being the dispersoids14]. This earlier study used micro-diffraction to identify the dis-ersoid particles and found them to contain a significant portionf aluminum (Al) and neodymium (Nd), i.e. the A12Nd [14].he formation and presence of the A12Nd dispersoids (meltingoint temperature of 1733 K) instead of the Mg–RE (rare-earth)ispersoids, which have a lower melting point temperature, inhis rapidly solidified Mg–Zn–Al–RE alloy is interesting. Thel2Nd dispersoids are thermally stable and help to pin the grainoundaries and prevent the coarsening of grains during high tem-erature consolidation and hot extrusion. The microstructure ofhe Mg–5.72 Al–2.96 Zn–6.05 Nd alloy revealed a non-uniformrain size along the three orthogonal directions of the extruded

late. The grains were flattened and elongated in the directionf mechanical deformation, i.e. extrusion (Fig. 2).

able 2oom temperature tensile properties of the rapidly solidified magnesium alloy

(GPa) 45

.2% yield strengthksi 62MPa 430

ltimate tensile strengthksi 67MPa 461

ngineering fracture stress (σf)ksi 66MPa 454

longation (G.L. = 12.5 mm)% 5.4

eduction in area ln(A0/Af)% 12

esults are the mean values based on duplicate tests.

T.S. Srivatsan et al. / Journal of Alloys and Compounds 460 (2008) 386–391 389

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ig. 5. Engineering stress vs. engineering strain curve for the magnesium alloy.Fig. 6. Monotonic stress vs. strain curve for the rapidly solidified magnesiumalloy.

ig. 7. Scanning electron micrographs of the tensile fracture surface of the magnesium alloy showing: (a) overall morphology; (b) array of macroscopic cracksistributed through the fracture surface; (c) high magnification of (b) showing non-linear nature of macroscopic cracks; (d) shallow dimples, macroscopic crack andne microscopic cracks.

3 ys and Compounds 460 (2008) 386–391

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Fig. 8. Scanning electron micrographs of the fracture surface of the magnesiumabo

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90 T.S. Srivatsan et al. / Journal of Allo

.2. Tensile response

The ambient temperature tensile properties of the rapidlyolidified magnesium alloys, in the as-extruded condition, areummarized in Table 2. The results reported are the mean val-es based on duplicate tests. The yield strength of the alloy is2 ksi. The high yield strength of this rapid solidification pro-essed alloy is ascribed to the conjoint and mutually interactivenfluences of (a) fine grain size, (b) solid solution strengthen-ng of the magnesium matrix, and (c) dispersion strengtheningrising from the presence of Al2Nd particles. The ultimate ten-ile strength of the alloy is 67 ksi and only marginally higherhan the yield strength, indicating that the work hardening rateast yielding is low. The ductility, measured by elongation over2.7 mm gage length of the specimen, is low and only 5.4%. Theeduction-in-area, another measure of tensile ductility, is 12%nd more noticeable than tensile elongation. The engineeringtress versus engineering strain curve is shown in Fig. 5. Thetrain hardening characteristics of the alloy was evaluated fromxamining the variation of stress with plastic strain and is shownn Fig. 6. The variation of monotonic stress with plastic strainbeyed the relationship σ = K(εp)n, where K is the monotonictrength coefficient and n is the strain hardening exponent. Aow degree of strain hardening can be inferred from this figure.

.3. Mechanisms governing tensile fracture behavior

The tensile fracture surfaces are helpful in elucidating usefulnformation on the role and/or influence of intrinsic microstruc-ural features on strength, ductility and fracture properties ofhe rapid solidified magnesium alloy. For the duplicate samplesested fracture occurred at the gage section and was essentiallyormal to the far-field stress axis. Representative features arehown in Figs. 7 and 8.

On a macroscopic scale fracture of the test specimen wasrittle in appearance (Fig. 7a) with the presence of an arrayf fine microscopic cracks distributed randomly through theracture surface (Fig. 7b). Examination of the crack path mor-hology at higher magnifications revealed the macroscopicrack to be essentially non-linear as it propagated through thelloy microstructure (Fig. 7c). The macroscopic cracks wereurrounded by a random distribution of very fine microscopicracks. Examination of the fracture surface at higher magnifica-ions revealed a population of dimples of varying size and shapemmediately adjacent to the macroscopic and fine microscopicracks, features reminiscent of locally ductile and brittle failureechanisms (Fig. 7d). The higher magnifications also revealedrandom distribution of fine microscopic voids (Fig. 8a). Sincerack extension under quasi-static loading occurs at high stressntensities, comparable to the fracture toughness of the mate-ial, the presence of a population of fine microscopic voidsegrades the actual strain-to-failure associated with ductile frac-ure. Although the exact nucleation of the fine microscopic

oids is difficult to pin-point, their near-equiaxed shape sug-ests that they may have nucleated around the Al2Nd dispersoidarticles during the later stages of tensile deformation, with-ut undergoing appreciable growth that would result in their

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lloy deformed in tension, showing: (a) view of a microscopic crack surroundedy shallow dimples of varying size; (b) dimples and cracking in the region ofverload.

valization. The nucleation of a microscopic void at a disper-oid particle and other second-phase particles present in theicrostructure occurs when the elastic energy in the particle

xceeds the surface energy of the newly formed void surfaces.hile this is a necessary condition, it must also be aided bystress at the matrix-second phase particle interface that is in

xcess of the interfacial strength. When a critical value of thenterface stress is reached void nucleation is favored to occur.he coalescence of the fine microscopic voids is the last stage

n the ductile fracture process. The halves of these voids arehe shallow dimples observed on the tensile fracture surfaceFig. 8b).

During tensile deformation the presence of dislocation pile

ps and grain boundary dislocations aids in nucleating voids athe second-phase particles distributed within the alloy matrixnd also along the grain boundary regions with the initiation oficro-cracking along the grain boundaries (Fig. 7b). The fine

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[16] T.S. Srivatsan, W. Li, C.F. Chang, J. Mater. Sci. 30 (1995) 1832–1838.

T.S. Srivatsan et al. / Journal of Allo

icroscopic and macroscopic cracks were observed traversinghe grain boundaries extending in the direction of the tensiletress axis, suggesting the importance of normal stress in enhanc-ng tensile deformation. Microscopically, the magnesium alloypecimens revealed features reminiscent of

(i) Locally ductile mechanisms, namely fine microscopic voidsand shallow dimples, and

ii) Brittle mechanisms, i.e. an array of fine microscopic andmacroscopic cracks both through the matrix and along thegrain boundary regions.

The very fine microscopic voids coalesce and it is the halvesf these voids that are the isolated pockets of shallow dimplesbserved on the tensile fracture surface. The growth of the micro-copic voids is dictated by localized plastic deformation. Theimited growth of the fine microscopic voids coupled with lackf their coalescence, as a dominant fracture mode, suggests therittle nature of the microstructure that governs the deformationroperties.

. Conclusions

A study aimed at understanding the quasi-static fractureehavior of a rapidly solidified magnesium alloy provides theollowing useful highlights:

1) The grains in the alloy were small in size and elongatedin the direction of deformation, i.e. extrusion. Overall, themicrostructure consisted of well-defined powder particles.The dispersoids were found to be distributed through thealloy microstructure.

2) The yield strength of the alloy was high and the tensilestrength was only marginally higher than the yield strength

indicating the tendency for strain hardening beyond yield tobe low.

3) Tensile fracture surface morphology revealed an over-all macroscopically brittle appearance and microscopi-

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d Compounds 460 (2008) 386–391 391

cally features reminiscent of ductile and brittle failuremechanisms.

eferences

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317–409.[3] T.S. Srivatsan, T.S. Sudarshan, Rapid Solidification Technology: An Engi-

neers Guide, Technomic Publishing Company, Lancaster, PA, 1993, pp.603–720.

[4] J.J. Lewandowski, in: T.W. Clyne (Ed.), Metal Matrix Composites, vol. 3,Elsevier Publishers, 2000, pp. 151–187.

[5] S. Schumann, F. Friedrich, The Use of Magnesium in Cars—Todayand in Future, Magnesium Alloys and their Applications, Werkstoff-Informatinsgesellschaft mbH, 1998, pp. 3–14.

[6] F.H. Froes, Y.W. Kim, S. Krishnamurthy, Mater. Sci. Eng. A 117 (1989)19–32.

[7] S.K. Das, L.A. Davis, Mater. Sci. Eng. A 98 (1988) 1–12.[8] L.A. Carapelia, Met. Prog. 48 (1947) 297–307.[9] A. Joshi, R.E. Lewis, in: H.G. Paris, W.H. Hunt Jr. (Eds.), Advances

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10] M.C. Flemings, A. Mortensen, Rapid Solidification Processing of Magne-sium Alloys, AMMRC Technical Report, TR-84-37, September 1984.

11] F. Hehmann, H. Jones, in: T.S. Srivatsan, T.S. Sudarshan (Eds.), RapidSolidification Technology: An Engineering Guide, Technomic PublishingInc., 1993, pp. 441–471.

12] S.K. Das, High Strength Rapidly Solidified Magnesium-Base Metal Alloys,U.S. Patent 4,675,157, June 1986.

13] S.K. Das, C.F. Chang, in: S.K. Das, B.H. Lear, C.M. Adams (Eds.), RapidlySolidified Crystalline Alloys, AIME, Warrendale, PA, USA, 1985, pp.137–156.

14] C.F. Chang, S.K. Das, D. Raybould, R.L. Bye, E.V. Limoncelli, Advancesin Powder Metallurgy, vol. 1, Metal Powder Industries Federation, NewJersey, 1989, pp. 331–346.

15] W. Li, The Cyclic Fatigue and Fracture Behavior of Magnesium Alloys,Master of Science Thesis, The University of Akron, Akron, OH, USA,1994.

17] S.K. Das, Apparatus for Casting High Strength Magnesium-Base MetalAlloys, U.S. Patent No. 4,718,475, January 1988.

18] ASTM, Standard E-8-93: Tension Testing of Metallic Materials, AmericanSociety for Testing and Materials, Philadelphia, PA, 1993.