On the notch ductility of a magnesium-rare earth alloy

B. Kondori a,n, A.A. Benzerga a,b

a Department of Materials Science & Engineering, Texas A&M University, College Station, TX 77843, USAb Department of Aerospace Engineering, Texas A&M University, College Station, TX 77843, USA

a r t i c l e i n f o

Article history:Received 18 May 2015Received in revised form21 August 2015Accepted 23 August 2015Available online 25 August 2015

Keywords:MagnesiumWE43FractureAnisotropyTriaxialityDuctility

a b s t r a c t

The room-temperature notch ductility of magnesium-rare earth alloy WE43 is investigated for twoloading orientations. This material is endowed with quasi-isotropic plastic flow properties, higherstrength and similar uniaxial ductility in comparison with other commercially available Mg alloys. Theauthors have recently shown that the notch ductility of a Mg–Al–Zn alloy is greater than its uniaxialductility over a wide range of notch geometries. This paper investigates whether the same trends hold forWE43, discusses the orientation dependence of ductility and the propensity for intergranular fracture athigh levels of hydrostatic tension. The latter mode of fracture is analyzed by means of detailed fracto-graphy in order to elucidate the role of grain-boundary particles and precipitates in the fracture process.

& 2015 Elsevier B.V. All rights reserved.

1. Introduction

Magnesium alloys exhibit high specific strength, good ma-chinability and damping properties. They suffer, however, frompoor formability at low temperatures. This sets severe limits totheir insertion as structural materials in weight-critical applica-tions. The lack of ductility is attributed to the limited number ofindependent deformation systems owing to the hexagonal closepacked crystal structure of Mg [1]. In strongly textured poly-crystals, stress enhancement at grain boundaries has been re-ported in activating non-basal slip [2], grain boundary sliding [3]and extension twinning, even under unfavorable loading orienta-tions. Details aside, however, it is the net plastic anisotropy thatresults either from crystallography or texture which is held re-sponsible for the limited ductility of wrought Mg alloys [4,5].

Following this rationale, major research efforts have aimed attexture weakening using various techniques. One generic methodconsists of strengthening the basal system either via solid solutionwith various alloying elements [6,7] or directed precipitation onselected habit crystallographic planes [8,9]. For instance, plate-likeprecipitates on basal planes have been shown to be effective inprohibiting twin growth and thus reducing tension–compressionasymmetry and plastic anisotropy [8]. Another method for textureweakening that has received wider interest in recent years is al-loying with rare earth (RE) elements [10–13]. Addition of RE

elements not only weakens the texture, but also changes its qua-litative character [14]. To date, the development of Mg alloyscontaining RE elements stands out as the most promising route forproducing ductile and formable Mg alloys for structural applica-tions. At dilute limits, RE-containing alloys exhibit remarkablestrain to failure in compression parallel to the rolling direction(strain to maximum load higher than 0.50) and good tensileductility (between 0.13 and 0.25 depending on the level of REelements in the alloy) [13,15]. However, these alloys suffer fromlow yield strength (∼ 60 MPa). To achieve higher strength, yttriumand higher concentrations of other RE elements are typically used.With such an increase in strength, the uniaxial ductility of these Y-and RE-concentrated alloys is at best equal to that of AZ31 withsimilar heat treatment condition (compare the data in [16–18]with those in [19,20]). Combination of very high specific strength(ultimate strength of ∼400 MPa) and good tensile ductility (∼0.15)a priori makes these alloys ideal candidates for aerospace anddefense applications where manufacturing cost is of limitedimportance.

Current understanding of damage and fracture in Mg alloys islimited to uniaxial loading conditions [13,15,20–23]. Efforts aimedat understanding fracture during processing [24] or at the crack tip[17,25–27] are limited. In most, emphasis is laid on deformationmechanisms. For instance, the role of texture design or precipitateengineering in fracture is viewed under such perspective [13,24].On the other hand, the coupling of deformation mechanisms to theinherent processes of damage initiation and accumulation is lar-gely unexplored. There are areas in which further research is

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Materials Science & Engineering A

http://dx.doi.org/10.1016/j.msea.2015.08.0770921-5093/& 2015 Elsevier B.V. All rights reserved.

n Corresponding author.E-mail address: bkondori.13034@tamu.edu (B. Kondori).

Materials Science & Engineering A 647 (2015) 74–83

necessary. First, the effects of multiaxial stress state on damageaccumulation to fracture need to be elucidated. Efforts in this di-rection include experimental [17,25,26] as well as fundamental[27] investigations of crack-tip processes and their relation tofracture toughness. Similar studies for RE-containing Mg alloysremain scarce [19,28]. One way to investigate damage mechanismsunder controlled stress triaxiality conditions is to utilize notchedspecimens. These are commonly employed in ductile fracturecharacterization to generate various states of stress differingthrough the amount of superimposed hydrostatic tension [29,30].The authors have recently used round notched bars of AZ31 alloy[18,31]. They found a striking effect of hydrostatic tension on thefracture behavior at ambient temperature, in that the notch duc-tility was consistently larger than uniaxial ductility. This behavioris in contrast with metallic alloys such as steels and aluminumalloys. The increase in notch ductility was associated with theactivation of void growth to coalescence mechanisms, the absenceof macroscopic shear failure, and a hypothesized activation ofmultiple deformation systems, including extension twinning. Onthe other hand, the effect of stress state on the ductility of RE-containing alloys remains to be elucidated.

Another important issue is whether mere texture weakeningand reduced anisotropy, which are key characteristics of the RE-containing family of alloys, can enhance ductility in a broadersense. Fundamentally, the effects of load triaxiality and texture arerelated since the active deformation systems favored by a giventexture inevitably depend on triaxiality. Also, it should be men-tioned that alloying and processing affect the anisotropic flowproperties of polycrystalline Mg. Primary processing, such as ex-trusion or rolling, generally leads to a strong basal texture.Alloying affects the strength, texture and possibly the propensityfor twinning, but also leads to the formation of second-phaseparticles, some of which can play a role in the ductile fractureprocess [17,21,32]. Knowledge of how texture, twinning and sec-ond-phase particles affect the damage process across a wide rangeof stress states is still lacking. This work is set out to addressthese issues by means of experiments designed to investigatethe effect of stress triaxiality in a commercially available RE-containing alloy.

2. Experimental procedure

2.1. Material and microstructural analyses

Test specimens were taken from a hot rolled plate of Elektron43,1 a modified WE43 alloy, in T5 condition (strain hardened andartificially aged) with a thickness of 1.5 in (38 mm) and nominalcomposition given in Table 1. The principal directions of the plateare labeled as L (longitudinal/rolling), T (transverse) and S (short-transverse/normal). Metallographic sections of as-received mate-rial were prepared. They were mechanically ground using SiCpaper and fine polished using 1, 0.3 and 0.05 mμ alumina sus-pensions. The use of water was restricted to grinding only. Forrinsing the samples, isopropyl alcohol was used and ultrasoniccleansing was done employing acetone. To reveal the micro-structure, acetic picral solution (4.2 g picric acid, 10 ml acetic acid,70 ml ethanol and 10 ml water) was used as etchant for 5 s. Mi-crostructural observations were carried out utilizing optical andScanning Electron Microscopy (SEM). Energy Dispersive X-RaySpectroscopy (EDS) was utilized to study the chemical composi-tion of different phases in the initial microstructure. The grain sizewas measured using the line intercept method and corroborated

by area measurements using ImageJ, an open source software forimage analysis. Crystallographic texture measurements were car-ried out using a Bruker-AXS D8 X-ray diffractometer (XRD) withCu Kα radiation on a sample from the plate's mid-section to get(0002) and (1010¯ ) pole figures using a 5° grid size and an 85°sample tilt.

2.2. Mechanical behavior and anisotropy

In order to reduce the propensity for shear localization, cy-lindrical specimens were employed, Fig. 1. Initial shapes were cutout using wire electric discharge machining (EDM). All specimenswere deformed to fracture. One principal direction was system-atically marked on both ends of each specimen to track the evo-lution of deformation anisotropy.

Round compression pins and tensile specimens were cut alongtwo principal directions, L and T. Compression pins were also ta-ken through the plate thickness (S direction). The compressiontests were carried out on a servo-hydraulic MTS machine (Model318.25) with a load cell capacity of 250 kN at a nominal strain rateof 10!3 s!1. A pure nickel anti-seize lubricant was used to preventearly barreling. The true axial strain is calculated as

⎛⎝⎜

⎞⎠⎟

HH

ln1

axial0

ε =( )

where H and H0 are the current and the initial height, respectively.The accuracy of strain measurement is 0.001. In all tests, a distinctload drop was observed before the specimen failed in shear.

The uniaxial tension tests were carried out at an initial strainrate of 10!3 s!1 on a servo-hydraulic MTS machine (Model380.50) equipped with a 250 kN load cell. True axial strain wasmeasured using a laser extensometer over a gauge length ofL 30 mm0 = . Strain to complete fracture, fε , was defined on thebasis of cross-sectional area variation:

⎛⎝⎜

⎞⎠⎟

AA

ln2

f0

fε =

( )

The area of the fractured specimen, Af , was measured post-mor-tem using top-view photographs of the fractured specimen as-suming an elliptical shape.

In order to study the effect of stress triaxiality on deformationand fracture, round notched (RN) specimens with three differentnotch geometries were used to provide a range of triaxialities[30,31]. Stress triaxiality is quantified by the ratio of hydrostatictension to some deviatoric stress measure. Both the longitudinal(L) and transverse (T) directions were characterized. All test con-ditions along L were doubled to check for scatter, which was foundto be small. Inside the notch, the stress state is triaxial: in additionto the major axial stress, there are two equal minor (principal)stresses. Each notched bar can be characterized based on the notchseverity parameter, ζ, equal to ten times the notch radius to spe-cimen diameter at the notch. Three values of ζ were explored andthe corresponding specimens were denoted by RNζ (Fig. 1). Thereis a direct relation between notch severity and stress triaxiality.The lower the value of ζ the higher the levels of stress triaxiality.Taking the notch height as gauge length, a nominal strain rate of3 10 s4 1× − − was imposed in all tests. In the notched bars, the useof an axial extensometer would render limited data about the

Table 1Nominal chemical composition of WE43 alloy used in this study.

Element Yttrium Rare earth Zirconium Magnesium

WE43 3.7–4.3% 2.3–3.5% 0.2% min. Bal.

1 Obtained from Magnesium Elektron company.

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–83 75

deformation inside the notch unless the gauge is restricted to theheight of the notch itself, which would be difficult given the size ofthe specimens used here. Instead, the instantaneous diameteralong S direction was continuously measured thanks to a custom-made radial extensometer. Refer to Ref. [31] for more details onthe radial extensometer. All tests were carried out to completefracture. Unlike in initially smooth tensile bars or compressionpins, the strains are spatially nonuniform in the gauge section of anotched bar. Hence, the following definitions are typically adop-ted; see Ref. [30]:

⎛⎝⎜

⎞⎠⎟ln ,

3X

0

Xε Φ

Φ¯ =( )

, 4f X1 f X2 fε ε ε¯ = ¯ | + ¯ | ( )

where the over-bar stands for spatial averaging over the mini-mum-diameter section, Φ0 and ΦX are the initial and the currentdiameter along X, respectively, X1 and X2 are two perpendiculardirections transverse to the loading axis, and subscript ‘f’ indicatesvalues at complete fracture. These definitions are the counterpartof Eq. (2) in uniaxial bars.

2.3. Fractography

After each fracture experiment and in order to prevent oxida-tion, the fracture surfaces of broken specimens were sprayed im-mediately with a silicone mold release spray then placed and heldin a manually vacuumed desiccator prior to being examined inSEM. It is worth noting that, even with extreme care, oxidation issuch a major problem in magnesium and its alloys that fracturesurfaces can only be observed once. For this reason, the testingcampaign has been paced to accommodate SEM observations ofoxide-free fracture surfaces. Occasionally, EDS analysis of thesecond phases on the surface were recorded.

3. Results

3.1. Microstructure

The microstructure of the material consists of a uniform dis-tribution of equiaxed grains, Fig. 2a. Pancake-like grains are oc-casionally observed in through-thickness planes. Grain-size mea-surements over a sample of more than 350 grains suggest asomewhat dual grain size distribution with 10 m∼ μ small grainsand 25 m∼ μ larger grains, Fig. 2c. The 1010( ¯ ) and 0002( ) polefigures depicted in Fig. 2b reveal a weak texture in comparisonwith pure Mg or other wrought alloys such as AZ31. The intensity

of the basal pole is significantly reduced and shifted towards therolling (L) direction. Texture weakening due to Y and Nd additionsis expected based on previous studies [10,11,33].

Fig. 1. Geometry of compression pins, round smooth and round notched (RN) tensile bars.

0

20

40

60

80

0 5 10 15 20 25 30 35 40

Freq

uenc

y

Grain diameter (µm)

Fig. 2. (a) Microstructure of as-received WE43 hot-rolled plate (rolling direction Lis horizontal, thickness direction S is vertical). (b) Pole figures corresponding to1010( ¯ ) and 0002( ) planes. (c) Grain size distribution for a sample of over 350 grains.

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–8376

Second-phase particles in the matrix or at grain boundaries(GB) can be observed in Fig. 2a as tiny black or bright dots de-pending on their composition and microscope settings. A betterrendering of the particles is obtained in high-resolution SEM,Fig. 3. Some particles are found inside the grains but most arelocated near or at GBs. Closer examination reveals that the grainboundaries are decorated with finer second phase particles. Insome locations, the preferred orientation of particles along the Ldirection is evident. Using EDS, the particles are identified as fa-ceted particles that contain Mg–Y, Mg–Nd and Mg–Y–Nd. In ad-dition, there are irregularly shaped, fine Mg–Zr particles whichoften appear as clusters. Elemental composition, phase diagrams(Mg–Nd and Mg–Y) and information in the literature [34–37] areused to qualitatively identify these particles as Mg Nd41 5, Mg Y2 ,Mg Y24 5, β-Mg Nd Y14 2 and Mg–Zr. These particles are distributed inan α-Mg matrix and its grain boundaries. High number density ofparticles at GBs is in accord with other experimental observations[20] and is rationalized by segregation of alloying elements atthese interfaces [12]. Other nano-sized precipitates, such asMg NdY2 and Mg Nd3 , may be present in the matrix but are notvisible in SEM [20,38].

White contrast bands are also observed in Fig. 3. EDS mapping(Fig. 4) indicates that these regions are Zr rich solid solution(Fig. 4b). As shown in part (a) of this figure, the boundaries ofthese contrast bands are also decorated with second phase parti-cles, which according to Fig. 4c are rich with Nd element. It isworth noting that Nd is also present in the matrix as solid solutionelement although with lower concentration. As shown in part(d) of the figure, yttrium is uniformly distributed in the matrix andsecond phase particles of the analyzed section.

3.2. Mechanical behavior

The mechanical response of WE43 is reported in Fig. 5. Nominalstress–strain curves in compression along the three principal di-rections of the plate are shown in Fig. 5a. In contrast to the re-sponse of pure Mg and common magnesium alloys, such as AZ andZK series [16,18], there is no significant anisotropy in the flowstress and hardening of WE43. The S-shaped curve typical of in-plane compression (here L or T directions) is also absent in theseplots. The nominal response typically exhibits a maximum fol-lowed by a gradual, albeit short, decrease in the load before thespecimen fails by splitting in two parts (shear failure).

The uniaxial tension and compression responses along therolling direction are compared in Fig. 5b. As depicted, there is notension–compression asymmetry to speak of in this alloy. In thisfigure, the true (Cauchy) stress versus true (logarithm) straincurves are shown. Post-load-drop, the stress is corrected by ex-trapolating the hardening rate just prior to localization. In uniaxialtension, the specimen fractures catastrophically with neither loaddrop prior to fracture nor a well-developed neck. Absence ofnecking is consistent with Ref. [20]. The lateral strains in WE43during tension are independent of direction and the initial circularcross-section retains its shape, which is indicative of isotropicbehavior.

The orientation and loading mode dependence of the yield andultimate flow strengths is summarized in Table 2. Accordingly,alloy WE43 is isotropic strength-wise with the transverse direc-tion being somewhat harder. Also, the tension–compressionasymmetry is marginal with the flow stress being higher incompression.

Fig. 3. Scanning electron micrograph of as-received WE43 alloy revealing secondphase particles and grain-boundary precipitates (rolling direction L is horizontal.)

Fig. 4. (a) SEM micrograph of WE43 alloy showing a typical contrast band in L–Splane. The boundaries of these contrast bands are decorated with second phaseparticles. EDS map of (b) zirconium (Zr); (c) neodymium (Nd); and (d) yttrium (Y).

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–83 77

3.3. Fracture behavior

To investigate the effect of stress state triaxiality on the de-formation and fracture of WE43, the uniaxial tension and com-pression results are now supplemented with the data of notchedbar experiments. Curves of applied load, divided by initial mini-mum cross-section area, versus reduction of diameter are pre-sented in Fig. 6. Two sets of curves are shown. One set correspondsto loading along the rolling direction L (Fig. 6a), the other toloading along the transverse direction T (Fig. 6b). Diameter re-duction was consistently recorded along the S direction for alltensile specimens, including some smooth tensile bars. In situmeasurement of lateral diameter during compression of pins wasnot possible; thus, the figure only includes experiments in tension.Subsequent to yielding, each specimen deforms until it fracturescatastrophically. The uniaxial bar breaks after shallow neckingwhile the RN10 and RN4 specimens break at the limit load orthereabout. On the other hand, the RN2 specimen (sharpest notch)breaks before the limit load is attained. This behavior contrastswith that of alloy AZ31 previously investigated by the authors [31].As depicted in Fig. 6, the specimens that experience highertriaxiality require higher (axial) loads for similar reduction in

diameter. For instance, at 0.01 diameter reduction, the normalizedload increases from 300 MPa to 350, 400 and 450 MPa startingfrom the uniaxial bar to RN10, RN4 and RN2, respectively. In somestrain intervals, a nearly constant load was measured whereas inothers, there were load fluctuations to sustain the applied nominalstrain rate. Also evident in Fig. 6 is the fact that notched specimenswith lower triaxialities accommodate higher strains to fracture.Incidentally, the diameter reduction to fracture does not changesignificantly between the RN10 and uniaxial bars.

Postmortem measurements confirm that the strain to fracturefε̄ is strongly dependent on the stress state triaxiality. The resultsare reported for both loading orientations in Fig. 7a. Note that thetotal strain to fracture is the sum of two contributions (see Eq. (4)):one due to diameter reduction along S (shown in Fig. 6) and an-other due to diameter reduction along T or L for loading along Land T. Although stress triaxiality is a field, it is useful to associate anominal value of triaxiality to each RN specimen based on pre-valent values over the deformation history, as inferred from finiteelement calculations. Hence, approximate figures are reported inFig. 7a (top abscissa). Clearly, the lowest fracture strains ( 0.04∼ )are realized at the highest triaxiality (i.e., in RN2 specimens). Byreducing stress triaxiality (i.e., going from RN2 to RN10 specimens)

fε̄ increases to over 0.15 (L loading). The trend of increasing duc-tility with decreasing triaxiality does not continue upon furtherreduction of triaxiality, as the fε̄ of uniaxial bars is lower than thatof RN10 bars. This is in contrast with the behavior of most metallicalloys where ductility is usually higher under uniaxial loading. Theabove mentioned trends persist for all in-plane specimens. Inter-estingly, the T orientation is less ductile than the L orientation atlow triaxialities. However, this trend is reversed at higher triaxi-alities although the fracture anisotropy is noticeably small. For

0

100

200

300

400

500

600

0 2 4 6 8 10 12 14 16 18

F/A 0

(MPa

)

e (%)

compression-Lcompression-Tcompression-S

0

100

200

300

400

500

600

0 2 4 6 8 10 12 14 16

σ (M

Pa)

ε (%)

compression-Ltension-L

Fig. 5. Nominal stress–strain compression response of WE43 along three principaldirections (L, T and S). (b) True stress–strain response of WE43 in tension andcompression along the rolling direction.

Table 2Yield and ultimate tensile or compressive strength values for three principalloading orientations (see text).

Property Direction L T S

Yield strength (MPa) Tension 231.5 255.0 –

Compression 264.0 281.0 216.5

Ultimate strength (MPa) Tension 418.4 430.6 –

Compression 438.2 450.3 454.0

0

100

200

300

400

500

600

0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

F/A 0

∆ΦS/Φ0

RN2-LRN4-L

RN10-Ltension-L

0

100

200

300

400

500

600

0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

F/A 0

∆ΦS/Φ0

RN2-TRN4-T

RN10-Ttension-T

Fig. 6. Force divided by initial cross-sectional area versus normalized reduction indiameter along S direction for uniaxial and notched specimens loaded along (a) Land (b) T directions.

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–8378

completeness, the fracture strains of compressed pins are alsoreported in Fig. 7a. The strain to failure in compression is con-sistently higher than in tension; its value is close to that obtainedin RN10 specimens.

In order to compare the fracture locus of WE43 with that ofcommon magnesium alloys, results of the same experiments on

hot-rolled AZ31 from [31] are shown in Fig. 7b. Although the twoalloys exhibit comparable uniaxial ductility, the strain to failure ofWE43 under triaxial loading is significantly lower than that ofAZ31. For instance, AZ31-RN2 specimens have an average strain tofailure four times higher than their WE43 counterparts. This figureshows the significantly different sensitivity of WE43 and AZ31 tostress triaxiality. Note that AZ31 used in this comparison is an alloywith strong anisotropy, tension–compression asymmetry andstrong basal texture whereas WE43 exhibits no tension–com-pression asymmetry and has weakened texture and isotropic be-havior in tension.

3.4. Fractography

The fracture surfaces of various WE43 specimens were ana-lyzed in SEM to infer the microscopic mechanisms controlling theoverall fracture response. Unlike in AZ31 [31], it was not possibleto interrupt the experiments just prior to fracture so that fracto-graphy is limited to post-mortem examinations. Under uniaxialloading, fracture was slightly slanted (much less than in AZ31) andthe fracture surface exhibited mixed characteristics of trans- andintergranular fracture, TGF and IGF, respectively, as shown in Fig. 8.The observation of IGF is in accord with the literature [26]. Occa-sionally, some large dimples are seen. In addition, the presence ofshallow dimples and ductile ridges covering a significant portionof the studied area points to the extent of plasticity prior to failureand indicates that void growth in this alloy is truncated by earlycoalescence of voids/microcracks.

An increase in triaxiality enhances the IGF features on thefracture surface, as shown in Figs. 9 and 10. Smoother facets areobserved more frequently, rendering a brittle-like appearance tothe fracture surface. The presence of grain pull-out provides clearevidence for IGF. Traces of particles along grain boundaries (notshown for brevity) show their involvement in the IGF process. Inaddition to grain boundaries (GBs), twin boundaries (TBs) may alsobe preferred locations for damage initiation [22,27,31] and idealpath for macroscopic crack advancement [23].

When brittle-like microcracks encounter such boundaries asGBs or TBs, the crossing of the latter may lead to crack front seg-mentation as has been reported in the literature. This is illustratedin Fig. 11, which presents a high resolution micrograph of thecentral region in Fig. 10c. Cleavage facets are observed, althoughsporadically on the fracture surface. It is clear on the same mi-crograph of an RN2 specimen that ductile ridges and a series ofparallel cracks in their vicinity coexist. This clearly demonstratesthe mixed nature of the fracture process in this alloy. It is worthnoting that second phase particles are frequently observed on thefracture surface, located on flat and faceted features and, some-times, at the center of shallow dimples. In many occasions, out-lines of grains on the fracture surface are decorated with secondphase particles. Presence of second phase particles on the fracturesurface suggests that these particles are actively involved in da-mage initiation, notably favoring IGF (also refer to Fig. 3).

4. Discussion

In this work, the mechanics and mechanisms of fracture in hot-rolled WE43-T5 alloy were studied. No attempt has been made toevaluate fracture toughness or investigate crack propagation. In-stead, focus was laid on crack initiation in nominally crack-freespecimens. To this end, the stress-state dependence of the fracturestrain, which characterizes ductility, was determined. An im-portant parameter that influences the fracture strain of materialsis stress state triaxiality [30,31]. The effect is generally associatedwith the strong dependence of void growth rates on triaxiality

Fig. 8. Fracture surface of a uniaxial tension specimen.

0

0.04

0.08

0.12

0.16

0.2-0.66 -0.33 0 0.33 0.66 0.99 1.32 1.65

– ε f

Specimen Type

Nominal Triaxiality

compression

tension RN10 RN4 RN2

L-directionT-direction

0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

0.4-0.66 -0.33 0 0.33 0.66 0.99 1.32 1.65

– ε f

Specimen Type

Nominal Triaxiality

compression

tension RN10 RN4 RN2

WE43-LAZ31-L

Fig. 7. (a) Strain to complete fracture for two different in-plane directions (L and T).(b) Fracture loci of WE43 and AZ31. Values of stress triaxiality are indicative only.The data for AZ31 is taken from [31].

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–83 79

Fig. 9. Fracture surface of a RN10 specimen showing a mix of intergranular failure,ductile tearing and other brittle features (see text for details).

Fig. 10. Fracture surface of a RN2 specimen showing intergranular failure, brittlefeatures and grain-boundary precipitates.

B. Kondori, A.A. Benzerga / Materials Science & Engineering A 647 (2015) 74–8380

[39]. In a notched bar, a triaxial stress state develops in the not-ched region, the intensity of which depends on the notch geo-metry. One advantage of using round notched bars in this regard isthat triaxiality variations upon straining are minimized, thusleading to nearly proportional stressing paths at failure locationseven if the material is anisotropic [40]. Another advantage of usingthese specimens in the context of Mg alloys is that the axisym-metric notch reduces the propensity for shear localization, thusenabling a thorough analysis of damage progression mechanisms.

From a macroscopic viewpoint, the main finding of this work iscaptured by the qualitative and quantitative differences betweenthe fracture loci of alloys WE43 and AZ31, Fig. 7. On one hand,WE43 has higher strength, quasi-isotropic plastic properties,marginal tension–compression asymmetry (Table 2) and a tensileductility that is close to that of AZ31. On the other hand, the notchductility of WE43 is much lower than that of AZ31, irrespective ofnotch acuity, Fig. 7b. In AZ31, the notch ductility is consistentlylarger than uniaxial ductility. Kondori and Benzerga [31] attributedthe enhanced notch ductility of AZ31 to the activation of voidgrowth to coalescence mechanisms (macroscopically normalfracture). By way of contrast, the round tensile unnotched speci-mens of that material exhibited macroscopic shear failure. Diffuseplastic flow, which is needed for void growth, was rationalized inthe notched bars by the activation of multiple deformation sys-tems, including extension twinning, favored by the multiaxialstress state prevailing therein. In WE43, however, the void growthand coalescence mechanisms are not operative, at least not to thesame extent as in AZ31. Fractography reveals a mix of ductile in-tergranular failure, ductile tearing, cleavage, and possibly twin-sized cracks, Figs. 9 and 10. Since twin-sized cracks and cleavagewere also observed in AZ31, it is hypothesized that it is the in-creased propensity for intergranular failure (IGF) which adverselyaffects ductility in WE43. The observation of IGF is consistent witha large concentration of GB particles and precipitates, as observedin this alloy (Fig. 3) as well as some Al alloys [41] and possiblyothers. Note that the decrease in WE43 fracture strain with in-creasing notch severity is consistent with the effect of stresstriaxiality in ductile IGF [41]. Similarly, the decrease in AZ31fracture strain is in keeping with triaxiality effects in ductilefracture by cavitation [39].

The relatively low notch ductility of WE43 is an obvious con-cern. In order to mitigate it, either in WE43-like compositions or

other RE alloys, it is important to understand its origins. The ex-periments reported herein show that the plastic and fracture be-havior of Mg alloys is ultimately determined by intricate factors. Itis of utmost importance to combine concepts from the mechanicsof ductile and brittle fracture with fundamental materials sciencenotions of alloying effects on plastic deformation in HCP metals inorder to put forth a consistent theory of fracture in Mg alloys.Thus, it would be of interest to develop a broader database on thenotch sensitivity of other RE alloys. Unfortunately, it is not easy toremove second phase particles from GBs of Mg-RE alloys. Aging at210 °C for 48 h of hot-rolled WE43 sheets leads to a ductility drop,presumably due to an increase in the density of GB precipitates[20]. However, the as-rolled sheets had a ductility of 0.04 (along L),much less than in the present plate. In addition, whether alter-native heat treatment could dissolve all GB precipitates remains tobe seen. In fact, Hadorn et al. [12] showed that annealing at 673 Kdoes not “desegregate” the RE elements from grain boundaries;also see [42]. Note that Kumar et al. [20] have shown evidence of asignificant increase in ductility when the sheets were processed byfriction stirring, probably due to the dissolution/redistribution ofGB precipitates. These authors, however, have not examinedtriaxiality effects.

It is worth noting that while the notch ductility of WE43 ismuch lower than in AZ31, their tensile ductilities are comparable.There are two aspects to this issue: (i) why is the WE43 uniaxialductility not higher than in AZ31, given that in the dilute limit Mg–Y alloys exhibit a noticeable ductility enhancement [15] and (ii)why does IGF not affect the uniaxial ductility in WE43 as much asit affects its notch ductility? We argue that the answers are es-sentially rooted in mechanics with details associated with thecomplex deformation mechanisms entering indirectly throughtheir net effect on strain-hardening, plastic anisotropy and/or non-Schmid behavior.

Consider first the increased ductility of magnesium alloyscontaining RE elements in dilute concentrations. It is believed tooriginate from a weakened texture. Concurrent activation ofmultiple deformation systems (i.e., contraction twinning, pris-matic and pyramidal c a⟨ + ⟩ slip) is proposed to contribute to thisimproved ductility. In addition, homogeneously distributed shear/slip bands are observed in RE-containing Mg alloys with morefrequency compared to other Mg alloys [15]. Presumably, thisdistributed shearing delays failure by macroscopic shear localiza-tion. Higher activity of contraction and secondary twins is ob-served in RE-containing alloy which can be related to the changein CRSS of these deformation mechanisms by alloying elements.Ease of activation of prismatic slip and its effect on easier ac-commodation of strain caused by deformation twinning has alsobeen invoked as the reason for increased contraction twinningactivity [12,43]. Clearly, alloy WE43 has the same weakened tex-ture as dilute alloys; yet its tensile ductility is not as large. Notethat the ductility of dilute alloys continues to increase even aftertexture weakening by RE elements has saturated [11,44]. Thissuggests that other factors contribute to the improved ductility, inaddition to texture weakening. Therefore, a simpler explanation onpure mechanistic grounds is as follows. As a measure of tensileductility, the fracture strain in uniaxial tension is the sum of twoterms:

5f u pnε ε ε= + ( )

where uε is the true strain at the onset of necking and pnε is thepost-necking strain. Typical values for uε are 0.12 (this work) and0.17 (Ref. [15]) for WE43 and Mg–3 wt%Y, respectively. Typicalvalues for pnε are on the other hand 0.005< (this work) and 0.07>[15]. The relative values of uniform elongation scale with the strainhardening capacity: the Mg–Y alloy has a low strength (UTS <

Fig. 11. Mixed cleavage–ductile fracture surface in a RN2 specimen showing crackfront segmentation for a main cleavage crack traversing either a grain- or twin-boundary.

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250 MPa) and large hardening exponent whereas WE43 has a highstrength (UTS¼420 MPa) and relatively low hardening. Part of thedecrease in ductility with increasing RE element concentration istherefore due to a decrease in hardening capacity. It is unclear,however, why the activation of multiple deformation systems inWE43 does not reflect on post-necking ductility.

With the above as basis, it becomes clear that (i) the tensileductilities of WE43 and AZ31 are close because they are both de-termined by uniform elongation. The ductility of AZ31 is largerbecause the post-necking strain is somewhat higher 0.03∼ [31];(ii) IGF does not affect the uniaxial ductility of WE43 as much as itaffects its notch ductility because its uniaxial ductility is basicallylimited by uniform elongation, which is controlled to first order byhardening. Another fact that hints at mechanistic aspects of frac-ture is that uniaxial axisymmetric bars of both materials exhibitshear failure. In AZ31, the fracture surface (whether planar orconical) is slanted at about 45° to the loading direction. In WE43the slanting is much less (about 20° to the plane normal to theaxial load; also see [20] who used rectangular-prismatic speci-mens). The reduction in slant angle is presumably due to the muchreduced plastic anisotropy of WE43. Therefore the fundamentalquestion is why are neither AZ31 nor WE43 capable of large post-necking deformation? In AZ31 there is a tendency for enhancedductility upon increasing triaxiality (which is retrieved in notchedbars). It is unclear why such tendency does not reflect on higherpost-necking deformation in neither alloy. Possible reasons in-clude the isolated or concurrent effects of plastic anisotropy(mostly for AZ31), propensity for IGF (only for WE43), exhaustionin hardening capacity, and strong deviations from Schmid's lawand associated dilatation. Regarding the latter, significant volumechange was measured in AZ31 in the absence of significant da-mage [31]. At present, a complete theory of fracture in Mg alloys isstill lacking.

In addition to the proposed rationale for fracture in WE43, thepresence of flat and facet-like surfaces on the fracture surfacesuggest the possibility of cleavage fracture in this alloy. Featuressimilar to what was observed on the fracture surfaces of notchedbars and especially those depicted in Fig. 11 have been referred toas cleavage planes in the literature [45,46]. This might requiremore in-depth investigation similar to that in Ref. [27], with at-tention focused on thermal and alloying effects. Wu and Curtin[27] showed that at 0 K, cleavage in pure magnesium is dominantfor most crack orientations whereas crack blunting via dislocationemission from crack tip is very limited. This finding shows theimportance of accounting for brittle fracture in Mg alloys. Theabove-mentioned study, however, does not consider the effects oftemperature as well as changes in unstable stacking fault energy,which would be expected upon adding RE elements.

5. Conclusions

The notch ductility of magnesium alloy WE43 containing REelements was determined for two loading orientations at ambienttemperature under quasi-static loading conditions. The material isendowed with quasi-isotropic plastic flow properties, due to aweak non-basal texture, higher strength and similar uniaxialductility in comparison with commercially available Mg alloys,such as AZ31. The main findings are as follows:

# The authors have recently shown that the notch ductility ofAZ31 Mg alloy was greater than its uniaxial ductility over a widerange of notch geometries. The same is not true for WE43. Forsmall amounts of superimposed hydrostatic tension, the samepositive trend is found for both loading orientations. However,upon increasing further the hydrostatic tension a sharp

decrease in fracture strain is measured, irrespective of loadingorientation.

# The decrease in ductility with increasing triaxiality is associatedwith an increased propensity for intergranular fracture at highlevels of hydrostatic tension. This mode of fracture is favored bythe formation of particles and precipitates at grain-boundaries,most likely due to strong propensity for segregation of Y andother rare earth elements to these boundaries. This propensityis believed to originate from a combination of high atomic misfitbetween these elements and Mg atoms and their high bulksolubility in the matrix.

# Although the damage mechanisms in AZ31 and WE43 are quitedifferent at high stress triaxiality, the present findings illustratethat mere texture weakening does not lead to enhancedductility.

# The results provide the groundwork for understanding the ef-fects of microstructural and loading variables on damage andfracture in magnesium alloys, in particular their RE containingfamily.

Acknowledgments

This research was supported by NPRP Grant no 4-1411-2-555from the Qatar National Research Fund (a member of QatarFoundation). The statements made herein are solely the respon-sibility of the authors. B.K. also gratefully acknowledges a TexasA&M University Dissertation Fellowship. The authors thank Mag-nesium Elektron for supplying the material and Dr. E. Dogan for hisassistance with texture measurements.

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