On the notch ductility of a magnesium-rare earth alloy

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

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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

On the notch ductility of a magnesium-rare earth alloy B. Kondori a,n, A.A. Benzerga a,b a b

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

art ic l e i nf o

a b s t r a c t

Article history: Received 18 May 2015 Received in revised form 21 August 2015 Accepted 23 August 2015 Available online 25 August 2015

The room-temperature notch ductility of magnesium-rare earth alloy WE43 is investigated for two loading orientations. This material is endowed with quasi-isotropic plastic flow properties, higher strength and similar uniaxial ductility in comparison with other commercially available Mg alloys. The authors have recently shown that the notch ductility of a Mg–Al–Zn alloy is greater than its uniaxial ductility over a wide range of notch geometries. This paper investigates whether the same trends hold for WE43, discusses the orientation dependence of ductility and the propensity for intergranular fracture at high levels of hydrostatic tension. The latter mode of fracture is analyzed by means of detailed fractography in order to elucidate the role of grain-boundary particles and precipitates in the fracture process. & 2015 Elsevier B.V. All rights reserved.

Keywords: Magnesium WE43 Fracture Anisotropy Triaxiality Ductility

1. Introduction Magnesium alloys exhibit high specific strength, good machinability and damping properties. They suffer, however, from poor formability at low temperatures. This sets severe limits to their insertion as structural materials in weight-critical applications. The lack of ductility is attributed to the limited number of independent deformation systems owing to the hexagonal close packed crystal structure of Mg [1]. In strongly textured polycrystals, stress enhancement at grain boundaries has been reported in activating non-basal slip [2], grain boundary sliding [3] and extension twinning, even under unfavorable loading orientations. Details aside, however, it is the net plastic anisotropy that results either from crystallography or texture which is held responsible for the limited ductility of wrought Mg alloys [4,5]. Following this rationale, major research efforts have aimed at texture weakening using various techniques. One generic method consists of strengthening the basal system either via solid solution with various alloying elements [6,7] or directed precipitation on selected habit crystallographic planes [8,9]. For instance, plate-like precipitates on basal planes have been shown to be effective in prohibiting twin growth and thus reducing tension–compression asymmetry and plastic anisotropy [8]. Another method for texture weakening that has received wider interest in recent years is alloying with rare earth (RE) elements [10–13]. Addition of RE n

Corresponding author. E-mail address: [email protected] (B. Kondori).

http://dx.doi.org/10.1016/j.msea.2015.08.077 0921-5093/& 2015 Elsevier B.V. All rights reserved.

elements not only weakens the texture, but also changes its qualitative character [14]. To date, the development of Mg alloys containing RE elements stands out as the most promising route for producing ductile and formable Mg alloys for structural applications. At dilute limits, RE-containing alloys exhibit remarkable strain to failure in compression parallel to the rolling direction (strain to maximum load higher than 0.50) and good tensile ductility (between 0.13 and 0.25 depending on the level of RE elements in the alloy) [13,15]. However, these alloys suffer from low yield strength (∼ 60 MPa). To achieve higher strength, yttrium and higher concentrations of other RE elements are typically used. With such an increase in strength, the uniaxial ductility of these Yand RE-concentrated alloys is at best equal to that of AZ31 with similar 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 and defense applications where manufacturing cost is of limited importance. Current understanding of damage and fracture in Mg alloys is limited to uniaxial loading conditions [13,15,20–23]. Efforts aimed at understanding fracture during processing [24] or at the crack tip [17,25–27] are limited. In most, emphasis is laid on deformation mechanisms. For instance, the role of texture design or precipitate engineering in fracture is viewed under such perspective [13,24]. On the other hand, the coupling of deformation mechanisms to the inherent processes of damage initiation and accumulation is largely unexplored. There are areas in which further research is

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

necessary. First, the effects of multiaxial stress state on damage accumulation to fracture need to be elucidated. Efforts in this direction include experimental [17,25,26] as well as fundamental [27] investigations of crack-tip processes and their relation to fracture toughness. Similar studies for RE-containing Mg alloys remain scarce [19,28]. One way to investigate damage mechanisms under controlled stress triaxiality conditions is to utilize notched specimens. These are commonly employed in ductile fracture characterization to generate various states of stress differing through 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 the fracture behavior at ambient temperature, in that the notch ductility was consistently larger than uniaxial ductility. This behavior is in contrast with metallic alloys such as steels and aluminum alloys. The increase in notch ductility was associated with the activation of void growth to coalescence mechanisms, the absence of macroscopic shear failure, and a hypothesized activation of multiple deformation systems, including extension twinning. On the other hand, the effect of stress state on the ductility of REcontaining alloys remains to be elucidated. Another important issue is whether mere texture weakening and reduced anisotropy, which are key characteristics of the REcontaining family of alloys, can enhance ductility in a broader sense. Fundamentally, the effects of load triaxiality and texture are related since the active deformation systems favored by a given texture inevitably depend on triaxiality. Also, it should be mentioned that alloying and processing affect the anisotropic flow properties of polycrystalline Mg. Primary processing, such as extrusion or rolling, generally leads to a strong basal texture. Alloying affects the strength, texture and possibly the propensity for twinning, but also leads to the formation of second-phase particles, some of which can play a role in the ductile fracture process [17,21,32]. Knowledge of how texture, twinning and second-phase particles affect the damage process across a wide range of stress states is still lacking. This work is set out to address these issues by means of experiments designed to investigate the effect of stress triaxiality in a commercially available REcontaining alloy.

2. Experimental procedure 2.1. Material and microstructural analyses Test specimens were taken from a hot rolled plate of Elektron 43, a modified WE43 alloy, in T5 condition (strain hardened and artificially aged) with a thickness of 1.5 in (38 mm) and nominal composition given in Table 1. The principal directions of the plate are labeled as L (longitudinal/rolling), T (transverse) and S (shorttransverse/normal). Metallographic sections of as-received material were prepared. They were mechanically ground using SiC paper and fine polished using 1, 0.3 and 0.05 μm alumina suspensions. The use of water was restricted to grinding only. For rinsing the samples, isopropyl alcohol was used and ultrasonic cleansing was done employing acetone. To reveal the microstructure, 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. Microstructural observations were carried out utilizing optical and Scanning Electron Microscopy (SEM). Energy Dispersive X-Ray Spectroscopy (EDS) was utilized to study the chemical composition of different phases in the initial microstructure. The grain size was measured using the line intercept method and corroborated 1

1

Obtained from Magnesium Elektron company.

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Table 1 Nominal 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.

by area measurements using ImageJ, an open source software for image analysis. Crystallographic texture measurements were carried out using a Bruker-AXS D8 X-ray diffractometer (XRD) with Cu Kα radiation on a sample from the plate's mid-section to get ¯ ) pole figures using a 5° grid size and an 85° (0002) and ( 1010 sample tilt. 2.2. Mechanical behavior and anisotropy In order to reduce the propensity for shear localization, cylindrical specimens were employed, Fig. 1. Initial shapes were cut out using wire electric discharge machining (EDM). All specimens were deformed to fracture. One principal direction was systematically marked on both ends of each specimen to track the evolution of deformation anisotropy. Round compression pins and tensile specimens were cut along two principal directions, L and T. Compression pins were also taken through the plate thickness (S direction). The compression tests were carried out on a servo-hydraulic MTS machine (Model 318.25) with a load cell capacity of 250 kN at a nominal strain rate of 10
3 s
1. A pure nickel anti-seize lubricant was used to prevent early barreling. The true axial strain is calculated as

⎛ H⎞ εaxial = ln ⎜ ⎟ ⎝ H0 ⎠

(1)

where H and H0 are the current and the initial height, respectively. The accuracy of strain measurement is 0.001. In all tests, a distinct load drop was observed before the specimen failed in shear. The uniaxial tension tests were carried out at an initial strain rate of 10
3 s
1 on a servo-hydraulic MTS machine (Model 380.50) equipped with a 250 kN load cell. True axial strain was measured using a laser extensometer over a gauge length of L 0 = 30 mm . Strain to complete fracture, εf , was defined on the basis of cross-sectional area variation:

⎛A ⎞ εf = ln ⎜ 0 ⎟ ⎝ Af ⎠

(2)

The area of the fractured specimen, Af , was measured post-mortem using top-view photographs of the fractured specimen assuming an elliptical shape. In order to study the effect of stress triaxiality on deformation and fracture, round notched (RN) specimens with three different notch geometries were used to provide a range of triaxialities [30,31]. Stress triaxiality is quantified by the ratio of hydrostatic tension to some deviatoric stress measure. Both the longitudinal (L) and transverse (T) directions were characterized. All test conditions along L were doubled to check for scatter, which was found to be small. Inside the notch, the stress state is triaxial: in addition to the major axial stress, there are two equal minor (principal) stresses. Each notched bar can be characterized based on the notch severity parameter, ζ, equal to ten times the notch radius to specimen diameter at the notch. Three values of ζ were explored and the corresponding specimens were denoted by RNζ (Fig. 1). There is 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 of 3 × 10−4 s−1 was imposed in all tests. In the notched bars, the use of an axial extensometer would render limited data about the

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

deformation inside the notch unless the gauge is restricted to the height of the notch itself, which would be difficult given the size of the specimens used here. Instead, the instantaneous diameter along S direction was continuously measured thanks to a custommade radial extensometer. Refer to Ref. [31] for more details on the radial extensometer. All tests were carried out to complete fracture. Unlike in initially smooth tensile bars or compression pins, the strains are spatially nonuniform in the gauge section of a notched bar. Hence, the following definitions are typically adopted; see Ref. [30]:

⎛Φ ⎞ ε¯ X = ln ⎜ 0 ⎟, ⎝ ΦX ⎠

(3)

ε¯f = ε¯ X1|f + ε¯ X2 |f ,

(4)

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

where the over-bar stands for spatial averaging over the minimum-diameter section, Φ0 and ΦX are the initial and the current diameter along X, respectively, X1 and X2 are two perpendicular directions transverse to the loading axis, and subscript ‘f’ indicates values at complete fracture. These definitions are the counterpart of Eq. (2) in uniaxial bars. 2.3. Fractography After each fracture experiment and in order to prevent oxidation, the fracture surfaces of broken specimens were sprayed immediately with a silicone mold release spray then placed and held in a manually vacuumed desiccator prior to being examined in SEM. It is worth noting that, even with extreme care, oxidation is such a major problem in magnesium and its alloys that fracture surfaces can only be observed once. For this reason, the testing campaign has been paced to accommodate SEM observations of oxide-free fracture surfaces. Occasionally, EDS analysis of the second phases on the surface were recorded.

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3. Results 3.1. Microstructure The microstructure of the material consists of a uniform distribution of equiaxed grains, Fig. 2a. Pancake-like grains are occasionally observed in through-thickness planes. Grain-size measurements over a sample of more than 350 grains suggest a somewhat dual grain size distribution with ∼10 μm small grains ¯ ) and (0002) pole and ∼25 μm larger grains, Fig. 2c. The (1010 figures depicted in Fig. 2b reveal a weak texture in comparison with pure Mg or other wrought alloys such as AZ31. The intensity

Frequency

60

40

20

0

0

5

10

15 20 25 30 Grain diameter (μm)

35

40

Fig. 2. (a) Microstructure of as-received WE43 hot-rolled plate (rolling direction L is horizontal, thickness direction S is vertical). (b) Pole figures corresponding to ¯ ) and (0002) planes. (c) Grain size distribution for a sample of over 350 grains. (1010

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Second-phase particles in the matrix or at grain boundaries (GB) can be observed in Fig. 2a as tiny black or bright dots depending on their composition and microscope settings. A better rendering of the particles is obtained in high-resolution SEM, Fig. 3. Some particles are found inside the grains but most are located near or at GBs. Closer examination reveals that the grain boundaries are decorated with finer second phase particles. In some locations, the preferred orientation of particles along the L direction is evident. Using EDS, the particles are identified as faceted particles that contain Mg–Y, Mg–Nd and Mg–Y–Nd. In addition, there are irregularly shaped, fine Mg–Zr particles which often appear as clusters. Elemental composition, phase diagrams (Mg–Nd and Mg–Y) and information in the literature [34–37] are used to qualitatively identify these particles as Mg41Nd5, Mg2Y , Mg24 Y5, β- Mg14 Nd2 Y and Mg–Zr. These particles are distributed in an α-Mg matrix and its grain boundaries. High number density of particles at GBs is in accord with other experimental observations [20] and is rationalized by segregation of alloying elements at these interfaces [12]. Other nano-sized precipitates, such as Mg2NdY and Mg3Nd, may be present in the matrix but are not visible 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 of these contrast bands are also decorated with second phase particles, which according to Fig. 4c are rich with Nd element. It is worth noting that Nd is also present in the matrix as solid solution element although with lower concentration. As shown in part (d) of the figure, yttrium is uniformly distributed in the matrix and second phase particles of the analyzed section.

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The uniaxial tension and compression responses along the rolling direction are compared in Fig. 5b. As depicted, there is no tension–compression asymmetry to speak of in this alloy. In this figure, the true (Cauchy) stress versus true (logarithm) strain curves are shown. Post-load-drop, the stress is corrected by extrapolating the hardening rate just prior to localization. In uniaxial tension, the specimen fractures catastrophically with neither load drop prior to fracture nor a well-developed neck. Absence of necking is consistent with Ref. [20]. The lateral strains in WE43 during tension are independent of direction and the initial circular cross-section retains its shape, which is indicative of isotropic behavior. The orientation and loading mode dependence of the yield and ultimate flow strengths is summarized in Table 2. Accordingly, alloy WE43 is isotropic strength-wise with the transverse direction being somewhat harder. Also, the tension–compression asymmetry is marginal with the flow stress being higher in compression.

3.2. Mechanical behavior The mechanical response of WE43 is reported in Fig. 5. Nominal stress–strain curves in compression along the three principal directions of the plate are shown in Fig. 5a. In contrast to the response of pure Mg and common magnesium alloys, such as AZ and ZK series [16,18], there is no significant anisotropy in the flow stress and hardening of WE43. The S-shaped curve typical of inplane compression (here L or T directions) is also absent in these plots. The nominal response typically exhibits a maximum followed by a gradual, albeit short, decrease in the load before the specimen fails by splitting in two parts (shear failure).

Fig. 3. Scanning electron micrograph of as-received WE43 alloy revealing second phase 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–S plane. The boundaries of these contrast bands are decorated with second phase particles. 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

600

600

500

500

400

400

300

F/A0

F/A0 (MPa)

78

200

0

200

compression-L compression-T compression-S

100 0

2

4

6

300

RN2-L RN4-L RN10-L tension-L

100

8 10 12 14 16 18

0

e (%)

0

500

600

400

500

300

100 0

ΔΦS/Φ0

400

200

F/A0

σ (MPa)

600

0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

compression-L tension-L

0

2

4

6

8

300 200

10 12 14 16

RN2-T RN4-T RN10-T tension-T

100

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

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

Direction

L

T

S

Yield strength (MPa)

Tension Compression

231.5 264.0

255.0 281.0

– 216.5

Ultimate strength (MPa)

Tension Compression

418.4 438.2

430.6 450.3

– 454.0

3.3. Fracture behavior To investigate the effect of stress state triaxiality on the deformation and fracture of WE43, the uniaxial tension and compression results are now supplemented with the data of notched bar experiments. Curves of applied load, divided by initial minimum cross-section area, versus reduction of diameter are presented in Fig. 6. Two sets of curves are shown. One set corresponds to loading along the rolling direction L (Fig. 6a), the other to loading along the transverse direction T (Fig. 6b). Diameter reduction was consistently recorded along the S direction for all tensile specimens, including some smooth tensile bars. In situ measurement of lateral diameter during compression of pins was not possible; thus, the figure only includes experiments in tension. Subsequent to yielding, each specimen deforms until it fractures catastrophically. The uniaxial bar breaks after shallow necking while the RN10 and RN4 specimens break at the limit load or thereabout. On the other hand, the RN2 specimen (sharpest notch) breaks before the limit load is attained. This behavior contrasts with that of alloy AZ31 previously investigated by the authors [31]. As depicted in Fig. 6, the specimens that experience higher triaxiality require higher (axial) loads for similar reduction in

0

0

0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

ΔΦS/Φ0

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

diameter. For instance, at 0.01 diameter reduction, the normalized load increases from 300 MPa to 350, 400 and 450 MPa starting from the uniaxial bar to RN10, RN4 and RN2, respectively. In some strain intervals, a nearly constant load was measured whereas in others, there were load fluctuations to sustain the applied nominal strain rate. Also evident in Fig. 6 is the fact that notched specimens with lower triaxialities accommodate higher strains to fracture. Incidentally, the diameter reduction to fracture does not change significantly between the RN10 and uniaxial bars. Postmortem measurements confirm that the strain to fracture ε¯f is strongly dependent on the stress state triaxiality. The results are reported for both loading orientations in Fig. 7a. Note that the total strain to fracture is the sum of two contributions (see Eq. (4)): one due to diameter reduction along S (shown in Fig. 6) and another due to diameter reduction along T or L for loading along L and T. Although stress triaxiality is a field, it is useful to associate a nominal value of triaxiality to each RN specimen based on prevalent values over the deformation history, as inferred from finite element calculations. Hence, approximate figures are reported in Fig. 7a (top abscissa). Clearly, the lowest fracture strains ( ∼0.04 ) are realized at the highest triaxiality (i.e., in RN2 specimens). By reducing stress triaxiality (i.e., going from RN2 to RN10 specimens) ε¯f increases to over 0.15 (L loading). The trend of increasing ductility with decreasing triaxiality does not continue upon further reduction of triaxiality, as the ε¯f of uniaxial bars is lower than that of RN10 bars. This is in contrast with the behavior of most metallic alloys where ductility is usually higher under uniaxial loading. The above mentioned trends persist for all in-plane specimens. Interestingly, the T orientation is less ductile than the L orientation at low triaxialities. However, this trend is reversed at higher triaxialities although the fracture anisotropy is noticeably small. For

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Nominal Triaxiality -0.66

0.2

-0.33

0

0.33

0.66

0.99

1.32

1.65

0.16

– εf

0.12 0.08

79

hot-rolled AZ31 from [31] are shown in Fig. 7b. Although the two alloys exhibit comparable uniaxial ductility, the strain to failure of WE43 under triaxial loading is significantly lower than that of AZ31. For instance, AZ31-RN2 specimens have an average strain to failure four times higher than their WE43 counterparts. This figure shows the significantly different sensitivity of WE43 and AZ31 to stress triaxiality. Note that AZ31 used in this comparison is an alloy with strong anisotropy, tension–compression asymmetry and strong basal texture whereas WE43 exhibits no tension–compression asymmetry and has weakened texture and isotropic behavior in tension.

0.04 L-direction T-direction

0

comp ression

3.4. Fractography tension

RN10 RN4

RN2

Specimen Type Nominal Triaxiality

-0.66

0.4

-0.33

0

0.33

0.66

0.99

1.32

1.65

0.35 0.3 – εf

0.25 0.2 0.15 0.1 0.05 0

WE43-L AZ31-L comp ression

tension

RN10 RN4

RN2

Specimen Type

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].

The fracture surfaces of various WE43 specimens were analyzed in SEM to infer the microscopic mechanisms controlling the overall fracture response. Unlike in AZ31 [31], it was not possible to interrupt the experiments just prior to fracture so that fractography is limited to post-mortem examinations. Under uniaxial loading, fracture was slightly slanted (much less than in AZ31) and the fracture surface exhibited mixed characteristics of trans- and intergranular fracture, TGF and IGF, respectively, as shown in Fig. 8. The observation of IGF is in accord with the literature [26]. Occasionally, some large dimples are seen. In addition, the presence of shallow dimples and ductile ridges covering a significant portion of the studied area points to the extent of plasticity prior to failure and indicates that void growth in this alloy is truncated by early coalescence of voids/microcracks. An increase in triaxiality enhances the IGF features on the fracture surface, as shown in Figs. 9 and 10. Smoother facets are observed more frequently, rendering a brittle-like appearance to the fracture surface. The presence of grain pull-out provides clear evidence for IGF. Traces of particles along grain boundaries (not shown for brevity) show their involvement in the IGF process. In addition to grain boundaries (GBs), twin boundaries (TBs) may also be preferred locations for damage initiation [22,27,31] and ideal path for macroscopic crack advancement [23]. When brittle-like microcracks encounter such boundaries as GBs or TBs, the crossing of the latter may lead to crack front segmentation as has been reported in the literature. This is illustrated in Fig. 11, which presents a high resolution micrograph of the central region in Fig. 10c. Cleavage facets are observed, although sporadically on the fracture surface. It is clear on the same micrograph of an RN2 specimen that ductile ridges and a series of parallel cracks in their vicinity coexist. This clearly demonstrates the mixed nature of the fracture process in this alloy. It is worth noting that second phase particles are frequently observed on the fracture surface, located on flat and faceted features and, sometimes, at the center of shallow dimples. In many occasions, outlines of grains on the fracture surface are decorated with second phase particles. Presence of second phase particles on the fracture surface suggests that these particles are actively involved in damage initiation, notably favoring IGF (also refer to Fig. 3).

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

completeness, the fracture strains of compressed pins are also reported in Fig. 7a. The strain to failure in compression is consistently higher than in tension; its value is close to that obtained in RN10 specimens. In order to compare the fracture locus of WE43 with that of common magnesium alloys, results of the same experiments on

In this work, the mechanics and mechanisms of fracture in hotrolled WE43-T5 alloy were studied. No attempt has been made to evaluate fracture toughness or investigate crack propagation. Instead, focus was laid on crack initiation in nominally crack-free specimens. To this end, the stress-state dependence of the fracture strain, which characterizes ductility, was determined. An important parameter that influences the fracture strain of materials is stress state triaxiality [30,31]. The effect is generally associated with the strong dependence of void growth rates on triaxiality

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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, brittle features and grain-boundary precipitates.

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

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

[39]. In a notched bar, a triaxial stress state develops in the notched region, the intensity of which depends on the notch geometry. One advantage of using round notched bars in this regard is that triaxiality variations upon straining are minimized, thus leading to nearly proportional stressing paths at failure locations even if the material is anisotropic [40]. Another advantage of using these specimens in the context of Mg alloys is that the axisymmetric notch reduces the propensity for shear localization, thus enabling a thorough analysis of damage progression mechanisms. From a macroscopic viewpoint, the main finding of this work is captured by the qualitative and quantitative differences between the 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 tensile ductility that is close to that of AZ31. On the other hand, the notch ductility of WE43 is much lower than that of AZ31, irrespective of notch acuity, Fig. 7b. In AZ31, the notch ductility is consistently larger than uniaxial ductility. Kondori and Benzerga [31] attributed the enhanced notch ductility of AZ31 to the activation of void growth to coalescence mechanisms (macroscopically normal fracture). By way of contrast, the round tensile unnotched specimens of that material exhibited macroscopic shear failure. Diffuse plastic flow, which is needed for void growth, was rationalized in the notched bars by the activation of multiple deformation systems, including extension twinning, favored by the multiaxial stress state prevailing therein. In WE43, however, the void growth and coalescence mechanisms are not operative, at least not to the same extent as in AZ31. Fractography reveals a mix of ductile intergranular failure, ductile tearing, cleavage, and possibly twinsized cracks, Figs. 9 and 10. Since twin-sized cracks and cleavage were also observed in AZ31, it is hypothesized that it is the increased propensity for intergranular failure (IGF) which adversely affects ductility in WE43. The observation of IGF is consistent with a large concentration of GB particles and precipitates, as observed in this alloy (Fig. 3) as well as some Al alloys [41] and possibly others. Note that the decrease in WE43 fracture strain with increasing notch severity is consistent with the effect of stress triaxiality in ductile IGF [41]. Similarly, the decrease in AZ31 fracture strain is in keeping with triaxiality effects in ductile fracture by cavitation [39]. The relatively low notch ductility of WE43 is an obvious concern. In order to mitigate it, either in WE43-like compositions or

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other RE alloys, it is important to understand its origins. The experiments reported herein show that the plastic and fracture behavior of Mg alloys is ultimately determined by intricate factors. It is of utmost importance to combine concepts from the mechanics of ductile and brittle fracture with fundamental materials science notions of alloying effects on plastic deformation in HCP metals in order to put forth a consistent theory of fracture in Mg alloys. Thus, it would be of interest to develop a broader database on the notch sensitivity of other RE alloys. Unfortunately, it is not easy to remove second phase particles from GBs of Mg-RE alloys. Aging at 210 °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 alternative heat treatment could dissolve all GB precipitates remains to be seen. In fact, Hadorn et al. [12] showed that annealing at 673 K does not “desegregate” the RE elements from grain boundaries; also see [42]. Note that Kumar et al. [20] have shown evidence of a significant increase in ductility when the sheets were processed by friction stirring, probably due to the dissolution/redistribution of GB precipitates. These authors, however, have not examined triaxiality effects. It is worth noting that while the notch ductility of WE43 is much lower than in AZ31, their tensile ductilities are comparable. There are two aspects to this issue: (i) why is the WE43 uniaxial ductility 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 as it affects its notch ductility? We argue that the answers are essentially rooted in mechanics with details associated with the complex deformation mechanisms entering indirectly through their net effect on strain-hardening, plastic anisotropy and/or nonSchmid behavior. Consider first the increased ductility of magnesium alloys containing RE elements in dilute concentrations. It is believed to originate from a weakened texture. Concurrent activation of multiple deformation systems (i.e., contraction twinning, prismatic and pyramidal 〈c + a〉 slip) is proposed to contribute to this improved ductility. In addition, homogeneously distributed shear/ slip bands are observed in RE-containing Mg alloys with more frequency compared to other Mg alloys [15]. Presumably, this distributed shearing delays failure by macroscopic shear localization. Higher activity of contraction and secondary twins is observed in RE-containing alloy which can be related to the change in CRSS of these deformation mechanisms by alloying elements. Ease of activation of prismatic slip and its effect on easier accommodation of strain caused by deformation twinning has also been invoked as the reason for increased contraction twinning activity [12,43]. Clearly, alloy WE43 has the same weakened texture as dilute alloys; yet its tensile ductility is not as large. Note that the ductility of dilute alloys continues to increase even after texture weakening by RE elements has saturated [11,44]. This suggests that other factors contribute to the improved ductility, in addition to texture weakening. Therefore, a simpler explanation on pure mechanistic grounds is as follows. As a measure of tensile ductility, the fracture strain in uniaxial tension is the sum of two terms:

εf = εu + εpn

(5)

where εu is the true strain at the onset of necking and εpn is the post-necking strain. Typical values for εu are 0.12 (this work) and 0.17 (Ref. [15]) for WE43 and Mg–3 wt%Y, respectively. Typical values for εpn are on the other hand <0.005 (this work) and >0.07 [15]. The relative values of uniform elongation scale with the strain hardening capacity: the Mg–Y alloy has a low strength (UTS <

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250 MPa) and large hardening exponent whereas WE43 has a high strength (UTS ¼420 MPa) and relatively low hardening. Part of the decrease in ductility with increasing RE element concentration is therefore due to a decrease in hardening capacity. It is unclear, however, why the activation of multiple deformation systems in WE43 does not reflect on post-necking ductility. With the above as basis, it becomes clear that (i) the tensile ductilities of WE43 and AZ31 are close because they are both determined by uniform elongation. The ductility of AZ31 is larger because the post-necking strain is somewhat higher ∼0.03 [31]; (ii) IGF does not affect the uniaxial ductility of WE43 as much as it affects its notch ductility because its uniaxial ductility is basically limited by uniform elongation, which is controlled to first order by hardening. Another fact that hints at mechanistic aspects of fracture is that uniaxial axisymmetric bars of both materials exhibit shear failure. In AZ31, the fracture surface (whether planar or conical) is slanted at about 45° to the loading direction. In WE43 the slanting is much less (about 20° to the plane normal to the axial load; also see [20] who used rectangular-prismatic specimens). The reduction in slant angle is presumably due to the much reduced plastic anisotropy of WE43. Therefore the fundamental question is why are neither AZ31 nor WE43 capable of large postnecking deformation? In AZ31 there is a tendency for enhanced ductility upon increasing triaxiality (which is retrieved in notched bars). It is unclear why such tendency does not reflect on higher post-necking deformation in neither alloy. Possible reasons include the isolated or concurrent effects of plastic anisotropy (mostly for AZ31), propensity for IGF (only for WE43), exhaustion in hardening capacity, and strong deviations from Schmid's law and associated dilatation. Regarding the latter, significant volume change was measured in AZ31 in the absence of significant damage [31]. At present, a complete theory of fracture in Mg alloys is still lacking. In addition to the proposed rationale for fracture in WE43, the presence of flat and facet-like surfaces on the fracture surface suggest the possibility of cleavage fracture in this alloy. Features similar to what was observed on the fracture surfaces of notched bars and especially those depicted in Fig. 11 have been referred to as cleavage planes in the literature [45,46]. This might require more in-depth investigation similar to that in Ref. [27], with attention focused on thermal and alloying effects. Wu and Curtin [27] showed that at 0 K, cleavage in pure magnesium is dominant for most crack orientations whereas crack blunting via dislocation emission from crack tip is very limited. This finding shows the importance of accounting for brittle fracture in Mg alloys. The above-mentioned study, however, does not consider the effects of temperature 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 RE elements was determined for two loading orientations at ambient temperature under quasi-static loading conditions. The material is endowed with quasi-isotropic plastic flow properties, due to a weak non-basal texture, higher strength and similar uniaxial ductility 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 of AZ31 Mg alloy was greater than its uniaxial ductility over a wide range of notch geometries. The same is not true for WE43. For small amounts of superimposed hydrostatic tension, the same positive trend is found for both loading orientations. However, upon increasing further the hydrostatic tension a sharp







decrease in fracture strain is measured, irrespective of loading orientation. The decrease in ductility with increasing triaxiality is associated with an increased propensity for intergranular fracture at high levels of hydrostatic tension. This mode of fracture is favored by the formation of particles and precipitates at grain-boundaries, most likely due to strong propensity for segregation of Y and other rare earth elements to these boundaries. This propensity is believed to originate from a combination of high atomic misfit between these elements and Mg atoms and their high bulk solubility in the matrix. Although the damage mechanisms in AZ31 and WE43 are quite different at high stress triaxiality, the present findings illustrate that mere texture weakening does not lead to enhanced ductility. The results provide the groundwork for understanding the effects of microstructural and loading variables on damage and fracture in magnesium alloys, in particular their RE containing family.

Acknowledgments This research was supported by NPRP Grant no 4-1411-2-555 from the Qatar National Research Fund (a member of Qatar Foundation). The statements made herein are solely the responsibility of the authors. B.K. also gratefully acknowledges a Texas A&M University Dissertation Fellowship. The authors thank Magnesium Elektron for supplying the material and Dr. E. Dogan for his assistance with texture measurements.

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