The Charpy impact results for Mg–10Gd–3Y–0.5Zr and Mg–11Y–5Gd–2Zn–0.5Zr alloys at various temperatures showed that Mg–10Gd–3Y–0.5Zr was more sensitive to temperature. The increase in impact toughness with temperature was related to the blunt crack-tip at high temperatures. The delamination and local melt of matrix were responsible for the brittle-to-ductile transition of GW103 alloy. The branch and bridging of cracks resulting from ordered phases played an import role in the change in fracture mode from cleavage fracture to quasi-cleavage and dimple-fracture for WGZ1152 alloy

Anyanwu

D.D. Yin

Janik

Liao

M.L. Hu

Q.D. Wang

Somekawa

V. Janik

Y.L. Mu

English

Crossref

Scripta Materialia

Elevated-temperature impact toughness of Mg–(Gd, Y)–Zr alloy

Mu, Y. L.

Wang, Q. D.

Hu, M. L.

Janik, Vit

Yin, D. D.

Warwick Research Archives Portal Repository

Available online at www.sciencedirect.comScripta Materialia 68 (2013) 885–888www.elsevier.com/locate/scriptamatElevated-temperature impact toughness of Mg–(Gd, Y)–Zr alloyY.L. Mu,a Q.D. Wang,a,⇑ M.L. Hu,a V. Janikb and D.D. YinaaNational Engineering Research Center of Light Alloy Net Forming, State Key Laboratory of Metal Matrix Composites,Shanghai Jiao Tong University , Dongchuan Road 800, Shanghai 200240, ChinabCzech Technical University in Prague, Karlovo nam. 13, Prague 12135, Czech RepublicReceived 18 December 2012; revised 8 February 2013; accepted 9 February 2013Available online 21 February 2013The Charpy impact results for Mg–10Gd–3Y–0.5Zr and Mg–11Y–5Gd–2Zn–0.5Zr alloys at various temperatures showed thatMg–10Gd–3Y–0.5Zr was more sensitive to temperature. The increase in impact toughness with temperature was related to the bluntcrack-tip at high temperatures. The delamination and local melt of matrix were responsible for the brittle-to-ductile transition ofGW103 alloy. The branch and bridging of cracks resulting from ordered phases played an import role in the change in fracture modefrom cleavage fracture to quasi-cleavage and dimple-fracture for WGZ1152 alloy. 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.Keywords: Magnesium alloys; Toughness; Impact behaviour; Fracture; Long-period stacking orderedIt has been demonstrated that magnesium (Mg)alloys containing gadolinium (Gd) and yttrium (Y) pos-sess high strength [1–4], creep-resistance [5–7] and wear-resistance [8,9]. These alloys thus have great potential tobe used as heat-resistant alloys in the automotive andaerospace industries. However, most research has fo-cused on the static mechanistic behavior of Mg alloys.When the materials are subjected to impulsive loadssuch as impact, explosion or thermomechanical shockduring service, the dynamic response and fracture pro-cess exhibits special characteristics that diﬀer from thoseof static response, owing to the presence of high loadingrates. Therefore, it is necessary for a proper assessmentof the dynamic fracture behavior of Mg–(Gd, Y) alloysfor structural applications, especially at elevated temper-atures. One method is to investigate their impact tough-ness at various temperatures. There have been manystudies on the fracture toughness of wrought Mg/Mg al-loy at room temperature [10–16]. Recently, Liao et al.[17,18] reported the room temperature impact toughnessof extrude Mg alloy. Results demonstrated that theroom temperature fracture/impact toughness ofwrought Mg/Mg alloy is aﬀected by microstructural fea-tures, such as texture [10,11], solid-solution/precipita-tion strengthening [12–15], grain size [13,16–19] andoxygen content [17]. However, few reports focus on1359-6462/$ - see front matter  2013 Acta Materialia Inc. Published by Elhttp://dx.doi.org/10.1016/j.scriptamat.2013.02.031⇑Corresponding author. Tel.: +86 21 54742715; fax: +86 2134202794; e-mail: wangqudong@sjtu.edu.cnthe elevated-temperature impact toughness of castheat-resistant Mg–RE alloy, which is important for esti-mating the service performance of Mg–RE alloy underhigh temperature and transient dynamic load condi-tions. However, it is diﬃcult to obtain more details nearthe crack tips throughout the test process, owing to theextremely short impact time.In the present work, the impact toughness and frac-ture feature of two heat-resistant Mg–RE alloys, peak-aged Mg–10Gd–3Y–0.5Zr (wt.%) (GW103) and Mg–11Y–5Gd–2Zn–0.5Zr (wt.%) (WGZ1152), were studiedat a temperature range from room temperature to350 C. It needs to be emphasized that the local stressstates of impact samples are quite diﬀerent from theplane-strain fracture toughness test. The impact loadand impact energy, therefore, cannot truly reﬂect thecritical stress-intensity factor KIc of samples. The impacttoughness represents the whole energy consumption ofthe specimen from elastic deformation to crack initia-tion until fracture, while KIc reﬂects the energy con-sumption before unstable extension of the crack.However, the results still have reference value for assess-ing the elevated-temperature fracture responses underimpact loading.GW103 alloy and WGZ1152 alloy were preparedfrom high purity (P99.9 wt.%) Mg, Zn and master al-loys of Mg–25 wt.% Gd, Mg–25 wt.% Y and Mg–30 wt.% Zr. Further details are available in Refs. [2–4,6]. The GW103 alloy was solution heat-treated atsevier Ltd. All rights reserved.0 3 6 9 12 1502468100 3 6 9 12 1520406080WGZ115(b)T=350Impact energy/JDisplacement/mmLoad/kNIIGW103II2II10 1 2 3 402468101214161 2 3 4510152025WGZ115IIIIImpact energy/JDisplacement/mmLoad/kNGW103WGZ115IIT=25(a)GW103Figure 1. Typical Charpy unnotched impact load and energy vs.displacement curves of GW103 alloy and WGZ1152 alloy at (a) roomtemperature and (b) 350 C. Macro-fractographs of GW103 alloy andWGZ1152 alloy impact at RT and 350 C (insets).0 2 4 6 8 10 12 14 16 180246810 (a)Displacement/mmDisplacement/mmLoad/kN RT 100 150 200 250 300 350GW1030 1 2 3 4 5 6 702468 (b)Load/kN RT 100 150 200 250 300 350WGZ115Figure 2. LD curves for (a) GW103 and (b) WGZ1152 at varioustemperatures.886 Y. L. Mu et al. / Scripta Materialia 68 (2013) 885–888500 C for 6 h and quenched in water. The T6 peak-aging heat treatment involved aging in a hot oil-bathat 225 C for 16 h followed by air cooling. TheWGZ1152 alloy was solution treated at 535 C for20 h, quenched in water at 25 C, and then aged at225 C for 24 h in an oil-bath. All the impact tests wereperformed on material that had undergone the peak-aged (T6) heat treatment.Impact toughness was determined in terms of theCharpy impact testing instrument IMP450J at tempera-tures ranging from room temperature (RT) to 350 C torecord the load and displacement of the hammer. Thespecimens for the impact test were Charpy unnotchedtest specimens with dimensions 10  10  55 mm. Allspecimens were heated to the testing temperature andheld for 20 min. An average of three measurements wereused to evaluate the impact toughness. Scanning elec-tron microscopy (SEM; JEOL JSM-5200) was employedfor observation of the fractured surface.The microstructures of peak-aged GW103 andWGZ1152 alloy have been described previously [2–4].The peak-aged microstructure evolution of GW103 al-loy involves mainly a-Mg solid solution, ﬁne and denserod-like b0 precipitates inside a-Mg matrix. The averageequiaxed a-phase grain size of the peak-aged alloy was40 lm. The peak-aged WGZ1152 alloy consists mainlyof a ﬁne-lamellar long-period stacking ordered (LPSO)structure distributed in the grain interior, strip-shapedLPSO Mg12YZn distributed along the grain boundaries,dense spheroidal b0 precipitates and the a-Mg solid solu-tion. The average grain size of WGZ1152 alloy was120 lm. In addition, Gd/Y enriched cuboidal phase(face-centered cubic) and Zr cores exist in both alloys.Typical impact load–displacement (LD) curves andimpact energy (IE) curves (insets) at RT and 350 Cfor both alloys are shown in Figure 1a and b. The LDcurve for GW103 alloy can be divided into three phasesat RT: the elastic phase (I), the hardening phase (II) andthe collapse phase (III). Here, the absorbed IE duringphases I and II includes crack initiation and growth en-ergy, while phase III represents crack instability propa-gation energy. The LD curve for WGZ1152 alloy,however, shows a distinct peak and the absence of phaseII, indicating the rapid extension of the crack the mo-ment the crack initiates under impact loading. It is clearthat the crack initiation and growth energy is muchgreater than the crack instability propagation energyfor GW103 alloy, demonstrating higher ductility andfracture resistance (toughness) than the WGZ1152 alloy.As shown in Figure 1b, the most apparent diﬀerence forimpact response at temperatures up to 350 C is themuch higher displacement and total IE (inset) for bothalloys, which increases by 300–400% compared with thatof RT (Fig. 1a). However, it should be noted that notonly does the impact load decrease for both alloys, butnotable changes at phase II of LD curves have occurred.First, a hardening phase appeared on the WGZ1152 LDcurve, while phase II of GW103 LD curve can be subdi-vided into a sustained load rise zone (II1) and a constantload zone (II2). Second, the load in phase II forWGZ1152 alloy is greater.Figure 2a and b shows the impact LD curves ofGW103 and WGZ1152 alloy at RT of 100, 150, 200,250, 300 and 350 C, respectively. The maximum loadfor GW103 alloy decreases gradually from 10 kN atRT to 5.5 kN at 350 C, but ranges between 7 andY. L. Mu et al. / Scripta Materialia 68 (2013) 885–888 8879 kN for WGZ1152 alloy. At temperatures >250 C, themaximum load for WGZ1152 alloy exceeds that forGW103 alloy, but the value of total energy forGW103 alloy is still higher, owing to the longer impactdisplacement. For WGZ1152 alloy, the values of load inphase I are progressively lower with increasing tempera-ture. However, phase II, whose max load value is higherthan phase I, appears at 200 C and above. The UTS ofWGZ1152 alloy at elevated temperatures has shown asimilar response [4]. The mechanism of this unusual phe-nomenon, however, remains enigmatic.The IE as a function of temperature is shown inFigure 3a. The total IE for both alloys increases withtest temperature, indicating rising impact toughness athigh temperatures. It seems that the impact toughnessof GW103 alloy is much more sensitive to impact tem-perature, because the amount of IE increases (from24 J at RT to 82 J at 350 C) per 100 C (E100) reaches17 J, but E100 reduces to 8 J for WGZ1152 alloy(from 6 J at RT to 32 J at 350 C). The rapid rise in totalIE and appreciable reduction in load as the impact tem-perature exceeds 200 C are essential features of ele-vated-temperature impact response for both alloys. Itshould be remarked here that the increase in IE comesmainly from phase II of the LD curve (Fig. 2a and b),that is, the crack initiation and growth process. Fig-ure 3b shows the relation between the ultimate tensilestrength (UTS) [3,4] and IE of both alloys at varioustemperatures. GW103 alloy shows an improved combi-nation of strength and impact toughness compared withWGZ1152 alloy. The temperature dependence of tensileproperties of GW103 and WGZ1152 alloys has demon-strated that UTS and elongation for GW103 alloy werehigher than that for WGZ1152 alloy at 200 C and0 50 100 150 200 250 300 3500153045607590GW103WGZ115Temperature/Total energy/J(a)Figure 3. IE absorption of GW103 and WGZ1152 alloy as a functionof (a) temperature and (b) UTS [3,4] in the temperature range fromroom temperature to 350 C.below, while UTS and elongation were lower forGW103 alloy than for WGZ1152 alloy >200 C [3,4].In addition, the grain size [16,17–19] may be one ofthe reasons for higher impact toughness for GW103 al-loy (40 lm) than WGZ1152 alloy (120 lm), indicat-ing that it is possible to attain higher strength andimpact toughness at high temperatures for WGZ1152 al-loy by grain reﬁnement.The macro-fractographs and micro-fractographs ofboth alloys samples impacted at temperatures betweenRT and 350 C are shown in Figure 1 (insets) and Fig-ure 4, respectively. It is clear that both alloys fail in brit-tle transgranular mode at RT. The shear lips are barelynoticeable on the fracture surface, and apparent radialmarks cover the whole fracture surface. Low-energydecohesion, without visible plastic deformation and evi-dent secondary cracks on the micro-fractographs ofboth alloys shown in Figure 4a and b also identify themechanism of formation of RT brittle fracture on themacroscopic scale. These observations are in accordancewith the LD response of both alloys with relative appar-ent peak load (P) at low impact temperature (6150 C),which indicates that the crack propagates rapidly underthe impact load once it originates. As the impact temper-ature reaches 350 C, the quite narrow shear lips on thesides of the fracture surface of WGZ1152 specimen canbe observed, although the radiation zone still coversmost of the surface. For GW103 alloy, the ﬁbrous zoneand shear rupture zone of the specimen impacted at350 C cover the whole surface, indicating considerableplastic deformation. Moreover, the GW103 specimensexhibited minor delaminations, where cracks branchedFigure 4. Micro-fractographs of samples impact for GW103 alloy at:(a) room temperature; (c) 250 C; (e) 350 C; and for WGZ1152 alloyat: (b) room temperature; (d) 250 C; (f) 350 C.888 Y. L. Mu et al. / Scripta Materialia 68 (2013) 885–888parallel to the longitudinal direction of the impact testbar, and therefore the crack propagation in the strikingdirection were substantially blunted.The brittle-to-ductile transition occurred clearly forGW103 alloy with rising impact temperature. Thefeatures of fracture surface morphology of the alloychanged from cleavage steps (Fig. 4a) to dimples(Fig. 4c) when the impact temperature increased fromRT to 250 C. The dimples’ nucleation, growth andcoalescence are the major fracture mechanisms, and thesamples can absorb more energy during the impactfracture process at a temperature of 250 C. As the im-pact temperature reached 350 C, tear ridges and wide-spread viscous-like ﬂow shear surface in the a-Mg solidsolution (Fig. 4e) could be observed. The local detail ofviscous-like ﬂow shear surface, a ripple morphology,shown in Figure 4e indicates that local melt may occurfor GW103 alloy under elevated-temperature impactload, which can explain the cause of much greater crackextension phase (phase II) and IE at 350 C than that at250 C in Figures 2a and 3. The higher impact toughnessfor GW103 alloy as the impact temperature exceeds250 C can therefore be attributed to three factors:(I) the plastic zone size of the crack-tip increases as aresult of decreasing yield strength at high temperatures,leading to a blunt crack front to improve intrinsictoughening; (II) the minor delamination relaxes triaxialstress conditions ahead of the crack-tip and blunts thecrack-tip, thus improving the impact toughness [20];(III) crack-tip blunting/shielding due to local melt ofthe matrix.A more detailed observation of the fracture surface ofWGZ1152 alloy at 350 C (Fig. 4f) shows the change infracture mode from cleavage fracture to quasi-cleavageand dimple fracture for WGZ1152 alloy, indicatingweak energy dissipation ability over the whole rangeof impact temperature, corresponding to the similar va-lue of P at various temperatures and low E100 in Figures2b and 3a. One prominent feature for WGZ1152 alloy atthe high impact temperature is the rising load during thehardening phase (Fig. 2b). The feature resulting fromthermally stable LPSO phases strengthening [5] elevatedthe stress level at the crack-tip, which enhanced the driv-ing force of crack propagation, thus reducing the impacttoughness. The other prominent feature manifestsmainly as more secondary cracks on the fracture surface(Fig. 4f), resulting from the branch and bridging ofcracks, improving the impact toughness. It should bementioned that the crack-tip blunting due to large plas-tic zone size at high temperatures has also played an im-port role in improvement of the impact toughness ofWGZ1152 alloy in addition to the branch and bridgingof cracks. This can be conﬁrmed from the plastic defor-mation of a-Mg matrix in Figure 4f. At low impact tem-peratures, the cleavage faces and steps form rapidly oncethe fracture stress of brittle a-Mg or LPSO phase alongthe grain boundaries is reached by stress concentration.Larger and more cleavage faces can thus be observed onthe fracture surface (Fig. 4b). With the rising impacttemperature, the yield strength of a-Mg matrix de-creases, and even the brittle LPSO phases themselveshave a certain capacity for plastic deformation, alleviat-ing the stress concentration near the brittle phase. Thegrowth of the initial micro-crack on the brittle phase be-comes slow, which creates favorable conditions for theinitiation of a lot of micro-crack on other positions,further reducing the stress concentration. Therefore,the secondary cracks on the fracture surface forWGZ1152 alloy multiplied at high impact temperatures.The eﬀect of temperature on the impact response andfractography was evaluated in GW103 and WGZ1152alloys. It was found that GW103 alloy exhibited higherimpact toughness than WGZ1152 alloy. The total IE forboth alloys rose remarkably as the impact temperatureincreased, owing to the enhanced dynamic plastic defor-mation ability. Delamination and local melt of matrixunder elevated-temperature impact load led to the brit-tle-to-ductile transition of GW103 alloy. The increasein branches and bridging of cracks resulting from LPSOphases plays an import role in improving the impacttoughness of WGZ1152 alloy.The authors gratefully acknowledge ﬁnancialsupport from the National Natural Science Foundationof China (No. 51074106), the Key Hi-Tech Researchand Development Program of China (2009AA033501)and the National Key Technology R&D Program ofChina (2011BAE22B01-5, 2006BAE04B01-2).[1] I.A. Anyanwu, S. Kamado, Y. Kojima, Mater. Trans. 42(2001) 1206.[2] S.M. He, X.Q. Zeng, L.M. Peng, X. Gao, J.F. Nie, W.J.Ding, J. Alloys Comp. 427 (2007) 316.[3] V. Janik, D.D. Yin, Q.D. Wang, S.M. He, C.J. Chen, Z.Chen, C.J. Boehlert, Mater. Sci. Eng. A 528 (2011) 3105.[4] D.D. Yin, Q.D. Wang, Y. Gao, C.J. Chen, J. Zheng, J.Alloys Comp. 509 (2011) 1696.[5] I.A. Anyanwu, S. Kamado, Y. Kojima, Metall. Mater.Trans. A 42 (2001) 1212.[6] D.D. Yin, Q.D. Wang, C.J. Boehlert, V. Janik, Metall.Mater. Trans. A 43 (2012) 3338.[7] X.B. Liu, X. Guan, R.S. Chen, E.H. Han, Trans.Nonferrous Met. Soc. China 20 (2010) S545.[8] M.L. Hu, Q.D. Wang, C. Li, W.J. Ding, Trans. Nonfer-rous Met. Soc. China 22 (2012) 1918.[9] M.L. Hu, Q.D. Wang, C.J. Chen, D.D. Yin, W.J. Ding,Z.S. Ji, Mater. Design 42 (2012) 223.[10] H. Somekawa, T. Mukai, Mater. Sci. Eng. A 459 (2007)366.[11] H. Somekawa, T. Mukai, Scr. Mater. 53 (2005) 541.[12] H. Somekawa, Y. Osawa, A. Singh, T. Mukai, J. Mater.Res. 23 (2008) 1128.[13] H. Somekawa, A. Singh, T. Mukai, Scr. Mater. 56 (2007)1091.[14] H. Somekawa, A. Singh, T. Mukai, J. Mater. Res. 22(2007) 965.[15] H. Somekawa, Y. Osawa, T. Mukai, Scr. Mater. 55(2006) 593.[16] H. Somekawa, T. Mukai, Scr. Mater. 53 (2005) 1059.[17] J.S. Liao, M. Hotta, A. Koshi, Mater. Lett. 65 (2011)2995.[18] J.S. Liao, M. Hotta, K. Kaneko, K. Kondoh, Scr. Mater.61 (2009) 208.[19] A. Ma, K. Suzuki, Y. Nishida, N. Saito, I. Shigematsu,M. Takagi, H. Iwata, A. Watazu, T. Imura, Acta Mater.53 (2005) 211.[20] D.W. Kum, T. Oyama, J. Wadsworth, O.D. Sherby, J.Mech. Phys. 31 (1983) 173.

Elevated-temperature impact toughness of Mg–(Gd, Y)–Zr alloy

Abstract

Similar works

Full text

Available Versions

Crossref

Warwick Research Archives Portal Repository