High cycle fatigue improvement by heat-treatment for semi-continuous casting Mg96.34Gd2.5Zn1Zr0.16 alloy

Materials Science & Engineering A 604 (2014) 78–85

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High cycle fatigue improvement by heat-treatment for semi-continuous casting Mg96.34Gd2.5Zn1Zr0.16 alloy Zongling He, Liming Peng, Penghuai Fu n, Yingxin Wang, Xiaoyu Hu, Wenjiang Ding National Engineering Research Center of Light Alloy Net Forming and State Key Laboratory of Metal Matrix Composite, Shanghai Jiao Tong University, Shanghai 200240, PR China

art ic l e i nf o

a b s t r a c t

Article history: Received 18 December 2013 Received in revised form 1 March 2014 Accepted 5 March 2014 Available online 12 March 2014

High cycle fatigue behavior of casting Mg96.34Gd2.5Zn1Zr0.16 alloy was investigated for its improvement by heat-treatment. After solution treatment (T4, 10 h@773 K) or solution treatment plus artificial aging (T6, 10 h@773 K þ128 h@473 K), fatigue strength of this alloy was found to be enhanced. The T6-treated alloy achieved the highest fatigue strength, 130 MPa, being 25 MPa and 18 MPa greater than those of the as-cast and T4-treated alloy, respectively. The average fatigue life of the heat-treated alloys is longer than that of the as-cast alloy a given stress amplitude. For the distribution of fatigue life, a fatigue life gap spanned from 105 to 107 can be observed in the as-cast and T4-treated alloy. Such a gap is absent after the alloy received artificial aging. The mechanism for the high cycle fatigue behavior of the casting alloy after heat-treatment was also discussed. & 2014 Elsevier B.V. All rights reserved.

Keywords: Semi-continuous casting Mg–Gd–Zn alloy Heat-treatment Fatigue properties

1. Introduction Fatigue properties of cast Mg alloys are strongly deteriorated by their casting defects. Existence of casting defects, e.g., gas pores, shrinkages on or below the surfaces, can easily trigger initiation of fatigue cracks in the alloys [1,2], especially at high cycle fatigue (HCF) conditions. Any practice involving reduction of casting defects can improve the fatigue properties of the alloys. The fatigue properties of the alloys also depend on their microstructures, e.g., grain size [3], second-phase particles [4,5], and precipitates [3], determined by their chemical compositions, casting processes, and heat-treatment methods. For a certain cast Mg alloy, when its chemical composition and casting technique are fixed, the mechanical properties of the alloy are mainly influenced by the applied heat-treatment. It was reported that application of appropriate heat-treatment can promote both tensile and fatigue properties of the cast Mg alloys. Bag and Zhou [6] pointed out that a solution-treated (T4) AZ91D alloy has the highest fatigue strength at any given stress level compared with that of the as-cast and solution plus ageing (T6) alloy. Their work also suggested that the aging treatment can reduce the fatigue crack propagation (FCP) rate of the AZ91D alloy. Recent work on a casting Mg–3.0Nd–0.2Zn–Zr (wt%, NZ30K) alloy [3] showed that the fatigue strength of the alloy was promoted by heat-treatment. The peak-aged alloy (T6) assumed the highest fatigue strength of 100 MPa compared with the as-cast and T4-treated (90 MPa) alloy. For the application of Mg alloys, they are

usually used in heat-treated conditions as structural materials. A better understanding of how the heat-treatment changes the microstructure and fatigue behavior is necessary for the wider usage of Mg alloys. Mg–Gd alloy, with high strength and creep resistance, is one of the most promising new structural materials. Addition of Zn to the Mg–Gd alloy can bring evident age-hardening effect [7,8]. It was reported that a high-strength cast Mg–3.2Gd–0.5Zn (at%) alloy developed by Ozaki et al. [9] was carrying an ultimate tensile strength of 410 MPa and elongation of 2%. It is the highest ultimate tensile strength for a cast Mg alloy obtained at present stage. Our previous study [10] on an as-cast Mg96.34Gd2.5Zn1Zr0.16 alloy indicated that this alloy had a relatively high tension-compression (R¼  1) fatigue strength of 10578 MPa, obviously higher than commercial AZ91 [11] and AM60 [12] alloys. It was also higher than the high-strength Mg–10Gd–3Y–Zr (wt%) alloy in T6 condition [13]. It could be deduced that the fatigue strength of the Mg96.34Gd2.5Zn1Zr0.16 alloy can be further improved if this aging harden-able alloy received appropriate heat-treatment [7,8]. In this study, the influence of heat-treatment on the fatigue behavior of semi-continuous casting Mg96.34Gd2.5Zn1Zr0.16 alloy was examined. The heat-treatment included solution and aging treatment, which is commonly seen in Mg alloys.

2. Experimental procedure 2.1. Alloy and heat-treatment

n

Corresponding author. Tel.: þ 86 21 54742911; fax: þ 86 21 34202794. E-mail address: [email protected] (P. Fu).

http://dx.doi.org/10.1016/j.msea.2014.03.017 0921-5093/& 2014 Elsevier B.V. All rights reserved.

Alloy ingots with nominal compositions (at%) of Mg96.34Gd2.5 Zn1Zr0.16 was melt from pure Mg, pure Zn, Mg–30Gd and Mg–30Zr

Z. He et al. / Materials Science & Engineering A 604 (2014) 78–85

(wt%) master alloys in an electric resistance furnace. Under protection of a mixture gas of CO2 and SF6, the ingots were prepared by a semicontinuous casting process, in which the casting speed, melt temperature, and water flow were maintained constantly at 100 mm/min, 983 K, and 25 L/min, respectively. The dimension of the ingot was 70 mm in diameter and 2300 mm in length. The actual chemical composition of the ingot was Mg–2.56Gd–0.9Zn–0.16Zr (at%), measured by an inductively coupled plasma atomic emission spectroscopy (ICP-AES) analyzer (Perkin-Elmer, Plasma 400). After casting, the ingots were solution treated at 773 K for 10 h, followed by water quenching (designated as ‘T4’) and aged at 473 K for 128 h (designated as ‘T6’). The samples of the as-cast alloy (designated as ‘as-cast’) are also tested for comparison.

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samples were observed using an optical microscope (OM) and FEI scanning electric microscopy (SEM, Sirion 200) equipped with Energy dispersive X-ray (EDX) analysis. The intermetallic phase distribution on the fracture surface of the samples was obtained using back-scatter electron imaging (BEI) technique. The area fractions of second-phase particles were calculated by OM images. At least 10 OM images (387 μm in length and 300 μm in width) were used in each condition. The area fraction is determined by dividing the total area of second-phase particles by the total area of the whole OM image. The average grain size of the alloy was measured from OM images using a linear intercept method.

3. Results 2.2. Tensile and fatigue test 3.1. Microstructure Rectangle tensile samples, 4 mm in width, 2 mm in thickness, and 15 mm as gauge length, were tested with a crosshead speed of 1 mm/min at ambient temperature on a Zwick/Roell-20 kN testing machine. For fatigue testing, plate specimens with a gauge length of 10 mm, a cross sectional area of 5 mm  4 mm, and a radius of 40 mm between the gauge length and grip ends were adopted. The gage surfaces of all fatigue specimens were polished before fatigue testing to avoid machining influence. The fatigue tests were performed under a tension-compression sinusoidal waveform (frequency of 20 Hz, stress ratio of R¼  1) in ambient air (temperature of 293–298 K and relative humidity of 40–50%) on a BOSE Electronforce 3550 fatigue testing machine. The tests were conducted in a staircase mode. That is, if a specimen failed less than 107 cycles, the next stress amplitude was decreased by 5 MPa. If no failure occurred within 107 cycles, the stress amplitude of the next specimen was increased by 5 MPa until the specimen failure happened. The calculation method of the fatigue strength from the staircase testing results was given in Refs. [10,14]. 2.3. Microstructural analysis Phase analysis was determined by X-ray diffraction (XRD) using a copper target on a Rigaku D/max 2550 V diffractometer. The microstructures and fatigue fracture surfaces of the investigated

Fig. 1 shows SEM images of the semi-continuous cast Mg96.34Gd2.5Zn1Zr0.16 alloy in different conditions, i.e., as-cast, T4, and T6. The as-cast alloy consists of α-Mg grains separated by a network of dendritic eutectic compounds, verified as (Mg,Zn)3Gd phase (DO3, a ¼0.72 nm) [15], at grain boundary. The shapes of the compounds are irregular and in sharp closed angles, forming many interfaces between α-Mg grains. Some small lamellar phases, identified as 14H-type long period stacking ordered (LPSO) phases [16], are located near the compounds and extend into the interior of the α-Mg matrix (Fig. 1a). The average grain size of the as-cast alloy is 12.2 μm, finer than that of convention cast Mg alloys (30–50 μm). Zr addition as grain refiner and quick cooling rate may help to form the fine grains in the semi-continuous casting process. After solution treatment, the interfaces between the (Mg, Zn)3Gd phases and the α-Mg matrix become smooth without shape angles. The volume fraction of the second-phase particles decreases from 14.5% (the as-cast alloy) to 10.5% (T4). 14H LPSO structured X phases are observed with flat interfaces to the α-Mg matrix, as indicated in Fig. 1b (indicated by arrows). Similar morphologies of X phases were reported in the Mg97Gd1Zn1 alloy [15] and Mg96.5Gd2.5Zn1 alloy [17]. Compared with the as-cast alloy, the 14H LPSO structured X phases become larger in size and easier to distinguish. Fig. 1b and c shows some bright-white fine

Fig. 1. SEM images of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions: (a) as-cast; (b) and (c) T4-treated; (d) T6-treated.

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precipitates forming during solution treatment near the grain boundaries. The similar precipitates were often observed in Zn– Zr containing Mg alloys, e.g., Mg–Nd–Zn–Zr [18,19], AM-SC1 [4,20], Mg–Ca–Zn–Zr [21], and Mg–Zn–Zr alloys [22]. They are determined as zinc–zirconium particles [21]. Together with the residual (Mg,Zn)3Gd compounds, the Zn–Zr particles, located at or near grain boundaries in Fig. 1c, usually observed at grain interiors in other alloys [18–21], can act as grain growth inhibitors. Subsequent aging after solution treatment does not lead to obvious difference in SEM images, as shown in Fig. 1d. The residual (Mg,Zn)3Gd compounds, X phases, and the Zn–Zr particles were kept unchanged. According to Yamashiki's research [15], aging treatment at 473 K for 128 h would lead to precipitation of metastable β0 and β1 phase. The β0 phase (Mg7Gd) has an orthorhombic structure [15], as the main strengthening phase of Mg–Gd–Zn and Mg–Gd alloys. The β1 phase has an fcc structure [15], as the next meta-stable phase transformed from the β0 phase. The chemical compositions of different phases in the three conditions were detected by EDX, listed in Table 1. The volume fraction of the (Mg,Zn)3Gd compounds decreases after solution treatment, and the solution content of the α-Mg matrix only increases a little. The matrix of the as-cast alloy contains 10.09 wt% Gd and 0.77 wt% Zn, and that of the T4-treated alloy contains 10.88 wt% Gd and 0.94 wt% Zn. The Gd content of the (Mg,Zn)3Gd phase increases evidently, from 26.33 wt% of the as-cast alloy to 42.2 wt% of the T4 alloy. The Zn content drops a little. After the solution treatment, the volume fraction of the LPSO structured phase becomes larger compared with that of the as-cast alloy. EDX result shows that the X phase contains less Gd and Zn content than the (Mg,Zn)3Gd phase. The exact chemical compositions of the X phase need to be further verified. It can be deduced from the distribution of the solute elements in the three conditions that the solution treatment obviously changes the morphologies of the (Mg,Zn)3Gd and X phase, with no much alternation of the α-Mg matrix composition. The average grain sizes of the T4-treated and T6-treated alloy are 19.3 μm and 19.7 μm, respectively, larger (460%) than that of the as-cast alloy (12.2 μm). Fig. 2 shows the XRD patterns of the cast alloy in three conditions. Both α-Mg and (Mg,Zn)3Gd phase are present in all conditions, confirming the above analysis. The X phase and precipitation are not detected. The microstructure characteristics are summarized in Table 2. 3.2. Tensile properties The tensile properties of the cast alloy in three conditions at room temperature are listed in Table 3. It can be seen that the as-cast and T4-treated alloy have nearly the same tensile properties. After the T6 treatment, the alloy exhibits a significant increase in its yield strength ( þ90 MPa) and ultimate tensile strength (þ 95 MPa), and an obvious drop of elongation (  6.8%). The average yield strength (s0.2) and ultimate tensile strength (sb) of the T6-treated alloy are 260.3 MPa and 366.8 MPa, with an elongation of 3.3%. 3.3. Fatigue strength The S–N curves of the cast alloy in three conditions are presented in Fig. 3. The fatigue strength (sf) at 107 cycles, ultimate tensile strength (sb), and corresponding fatigue ratio (sf/sb) are listed in Table 4. The results show that the fatigue strength increases after the heat-treatments, but not as markedly as the tensile strength does. Among the three conditions, the T6-treated alloy shows the highest fatigue strength of 130 MPa, which is 25 MPa and 18 MPa greater than that of the as-cast and T4-treated

Table 1 EDX analysis of different phases in the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions. Thermal condition

Position

Wt%

At%

Mg

Gd

Zn

Mg

Gd

Zn

As-cast

α-Mg matrix (Mg,Zn)3Gd

89.14 58.45

10.09 26.33

0.77 15.22

97.97 85.73

1.71 5.97

0.31 8.30

T4

α-Mg matrix (Mg,Zn)3Gd X phase

88.89 43.71 62.20

10.88 42.20 28.50

0.94 14.09 9.32

98.05 78.79 88.72

1.86 11.76 6.32

0.39 9.45 4.97

T6

α-Mg matrix (Mg,Zn)3Gd X phase

88.28 45.85 77.00

11.10 40.20 18.25

0.93 13.96 4.75

97.84 79.57 94.38

1.90 11.18 3.46

0.39 9.25 2.16

• • α

• •

•

•

•

••

Fig. 2. XRD patterns of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions.

alloys. The values of the fatigue ratio (sf/sb) for the investigated alloy in the as-cast, T4 and T6 conditions are 0.38, 0.41, and 0.35, respectively. It is in good agreement with the reported value for magnesium alloys (0.25–0.5) [23]. As seen from Fig. 3, the increase in the fatigue strength also leads to the improvement of the fatigue life. At any given stress amplitude, the average fatigue life of the heat treated alloys is longer than that of the as-cast alloy. One thing to be noted here is the life gap phenomenon of the S–N curves, where no failures happening within 105 to 107 cycles are observed in both the as-cast and T4-treated alloys, being not in the case of T6-treated alloy. This interesting phenomenon of no-lifegap after aging treatment will be discussed in detail later.

3.4. Fatigue fracture behavior To clearly understand the behavior of fatigue crack initiation and propagation, the fracture surfaces of the failed specimens were examined. As shown in Fig. 4a, the overall fracture surface can be divided into three regions, i.e., crack initiation region (Region 1), crack propagation region (Region 2), and final rupture region (Region 3). The fatigue cracks in the present study all initiate from the outer surface of the specimen, as indicated in Fig. 4b–d. The average height of the fatigue crack initiation site is evidently lower than other areas of Region 1, similar to the previous as-cast alloys [10]. The EDX analysis shows that the oxygen content of the fatigue crack initiation site is significantly higher (60 wt%) than other locations (18 wt%). As discussed in the previous study [10], the fatigue cracks initiate from oxide films, formed from the fatigue test. Similar phenomenon was also reported in the fatigue properties of the forged Mg–Zn–Y–Zr alloys [23], where the oxide

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Table 2 Microstructure characristics of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions. Thermal conditions

α-Mg matrix (wt%)

As-cast

Solute atoms (10.09 Gd, 0.77 Zn)

Eutectic compounds

14.5%, dentritic morphologies Solute atoms (10.88 Gd, 0.94 Zn) 10.5%, elliptic morphologies Precipitation (β0 & β1 phases); Solution atoms 11.0%, elliptic (at least 3.82 Gd). morphologies

T4 T6

Table 3 Tensile properties of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy at room temperature. Thermal condition

Yield strength s0.2 (MPa)

Ultimate tensile strength sb Elongation (MPa) (%)

As-cast T4 T6

172.5 169.2 260.3

274.8 272.1 366.8

9.5 10.1 3.3

Stress Amplitude (MPa)

200 As-cast T4 T6

180 160

runout

140 120 Life gap

100 80 103

104

105

106

107

108

Number of cycle Fig. 3. Stress–life curves of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions.

Table 4 Fatigue strength of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy at room temperature. Thermal conditions

sf (MPa)

Fatigue strength

Deviation (MPa)

Ultimate tensile strength sb (MPa)

sf/sb

As-cast T4 T6

105 112 130

7 8.8 7 12.1 7 4.1

274.8 272.1 366.8

0.38 0.41 0.35

LPSO phases (X phases)

Zn–Zr particles

Average grain size (μm)

Little amount at grain inteiors

None

12.2

Small amount, increased based on ascast condition. Small amount, increased based on ascast condition.

Located near the grain boundaries. Located near the grain boundaries.

19.3 19.7

A similar observation was also found in the casting AZ91E-T4 [1] and Mg–12Zn–1.2Y–0.4Zr (wt%) [5] alloys. There, as indicated in Fig. 4a, exists a transition region between Region 1 and Region 2 in all of the three conditions, different from the casting AZ91E-T4 [1] and Mg–12Zn–1.2Y–0.4Zr (wt%) [5] alloys. The surface in this region is rougher than that of Region 1. It is mainly featured by some unique steps inside the grains and fracture eutectics can be observed around them. Fig. 6 shows the typical morphologies of these steps. The spacing of these steps, like branches, varies from less than 1 μm to several micrometers, pointed out by the double-head arrow in Fig. 6. Some of the Zn–Zr particles are also observed in the BEI image (Fig. 6b). It confirms that the unique morphologies mainly locate at grain interiors. As stated in the as-cast alloy [10], they may be attributed by the interaction between the unique LPSO phase and dislocations during the crack propagation. Fig. 7 shows typical images of fatigue fracture surface in Region 2. More serrated fatigue striations are observed in all samples, indicating a typical cyclic crack growth and retardation during fatigue damage. In all three conditions, the fracture surfaces of Region 2 are characterized by fractured grains around by the fractured eutectic compounds, as shown in Fig. 7b, d, and f. The tear ridges, i.e., white ribbons in the images, of the T4 and T6-treated alloys are more evident than that of the as-cast alloy, as indicated in Fig. 7c and e. It may indicate that trans-granular fracture of the T4 and T6-treated alloys are larger than that of the as-cast alloy. The fracture surface of the T6-treated alloy, unlike the as-cast and T4-treated alloys, has some large steps (Fig. 7e and f), in addition to the fatigue striations. The outline of these steps looks like twinning, being consistent with Yang et al.'s [24] and S. M. Yin et al.'s [25] observation in Region 1. Fig. 8 shows the final rupture region (Region 3) of the fracture surface. Compared with Regions 1 and 2, Region 3 plays a minor role in determining the fatigue life. In all the samples, this region consists of a high volume fraction of second-phase particles along grain boundaries, leading to inter-granular failures. 3.5. Surface deformation morphology of the fatigue samples

films were formed by the interaction between the single-slip and environment near the sites of the crack initiation. Close to the fatigue crack initiation site, the fracture surface of the relatively flat Region 1 is characterized by two kinds of typical fracture morphologies. Type I morphology is shown in Fig. 5a, characterized by part of extreme flat areas nearly with no striations, like cleavage fracture planes. Type II morphology is characterized by flat areas with fine fatigue striations, as shown in Fig. 5b. From Fig. 4b–d, it can be seen that the percentage of the Type I morphology (large flat planes without striations) is lower than that of the Type II morphology (large flat planes with striations). The Type I morphology in the T6 condition is higher than other two conditions. Most of the fracture surfaces are covered by the Type II morphology in Region 1. The fatigue cracks in this region tend to run across the grains, leaving with flat fracture surfaces and fractured eutectic compounds (Fig. 5).

The sample surfaces (polished before fatigue tests) after fatigue tests were investigated in the present study to reveal the deformation mechanisms during the test, shown in Fig. 9. In the as-cast samples, slip bands are clearly observed at the interiors of some grains in Fig. 9a and b. Those bands are parallel to each other and believed to be the basal slips. The grains with the bands are optimally oriented for the basal slips. Similar morphologies were observed in Marrow et al.'s EBSD results [26] and F. Yang et al.'s TEM work [27]. After solution treatment, the X phase appears. It becomes more difficult to distinguish the parallel basal slip lines with the X phases. The X phase is less in density and thicker compared with the slip bands, seen in Fig. 9c and d. The highdensity thinner parallel lines at grain interiors are the basal slip bands. Compared with the as-cast alloy, the density of the bands in the T4-treated alloy is greatly lower. It indicates that the T4

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Fig. 4. Fractopography of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions: (a) and (b) as-cast, (c) T4-treated, and (d) T6-treated.

Fig. 5. Fatigue fracture images of Region 1 in the casting Mg96.34Gd2.5Zn1Zr0.16 alloy: (a) extremely flat regions with little striations (Type I) in the T6-treated alloy; (b) flat fracture surface with fine fatigue striations at grain interiors (Type II) in the T6-treated alloy.

Fig. 6. Morphologies of unique steps in the transition region between Region 1 and Region 2. These two images come from the fatigue fracture surface of the T4-treated alloy. They are secondary electron image and back scatter electron image at the same position.

treatment improves the threshold amplitude for the formation of persistent slip bands. In the T6-treated samples, no such slip bands can be observed, as shown in Fig. 9e and f. It is hard for the basal slips to take place in the T6-treated alloy than those in the as-cast and T4-treated alloys. As known as a fact, the basal slips are much easier to be triggered than other slip systems during the plastic deformation at room temperature. The present study confirms that the basal slips are the main deformation mechanism during the

high cycle fatigue tests, though the basal slip lines are not observed on the surfaces of the T6-treated samples.

4. Discussion Through modifications on the microstructures, heat-treatment can enhance both tensile and fatigue properties of the alloy.

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Fig. 7. Fatigue fracture surface images of Region 2 in the casting Mg96.34Gd2.5Zn1Zr0.16 alloy in three conditions: (a) and (b) as-cast, (c) and (d) T4-treated alloy, and (e) and (f) T6-treated alloy. Image (b), (d), and (f) are high-magnification observation of the areas indicated by squares in images (a), (c), and (e), respectively.

Fig. 8. Fatigue fracture surface images of Region 3 in the casting Mg96.34Gd2.5Zn1Zr0.16 alloy. These two images come from the fatigue fracture surface of the T6-treated alloy. They are secondary electron image and back scatter electron image at the same position.

The tensile properties and fatigue strength of the as-cast and T4-treated alloys are comparable, corresponding to the similar microstructures, as shown in Table 2. After subsequent aging, the meta-stable β0 and β1 phases precipitate from the α-Mg matrix [15]. The yield strength of the alloy obviously increases as high as 90 MPa, and the elongation decreases from 10.1% to 3.3%. The fatigue strength of the T6-treated alloy also has an improvement of 25 MPa compared with the as-cast alloy and 18 MPa to the T4-treated alloy. The precipitation strengthening can help elevation of the mechanical properties in the Mg–Gd–Zn alloy. It was also reported in the cast Mg–3.0Nd–0.2Zn–Zr (wt%) alloy

(Mg-RE based alloy) [3], with less amount (ca. 12 MPa). All the MgRE alloys probably can acquire improvement on the fatigue strength by the aging after the solution treatment. It is different from the AZ91 alloy (Mg–Al based alloy) [6], whose fatigue strength decreased after aging treatments. In the present study, an interesting phenomenon was observed in Fig. 3. A fatigue life gap from 105 to 107 where no failure happens exists in both the as-cast and T4-treated alloys. It, however, does not exist in the T6-treated alloy. The reason for the existing life gap in the as-cast alloy was discussed in the previous work [10]. The main reason contributing to the life gap is

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Fig. 9. Surface morphologies of the fatigue samples of the casting Mg96.34Gd2.5Zn1Zr0.16 alloy before and after the fatigue tests in three conditions: (a) and (b) as-cast alloy (s ¼ 115 MPa, Nf ¼ 31,976), (c) and (d) T4-treated alloy (s ¼ 120 MPa, Nf ¼74,704), and (e) and (f) T6-treated alloy (s ¼ 140 MPa, Nf ¼ 37,073).

the finer grains compared with the plastic zone and the weak incoherent interface between the eutectics and α-Mg matrix. Similar consideration on the comparison of the plastic zone size and the average grain size can explain the difference about the life gap in Fig. 3 among the three conditions. The size of the plastic zone, ry, can be estimated from the stress-filed equation under the plane strain condition as [5] ry ¼(ΔK/sy)2/3π, where ΔK is stress intensity factor and sy is the yield strength. According to the Ksolution [5], the stress intensity between Region 1 and Region 2, ΔK, of the as-cast, T4-treated and T6-treated alloys are calculated to be 2.4 MPa  m1/2, 2.56 MPa  m1/2, and 3 MPa  m1/2, respectively. For the as-cast alloy, the yield strength is 172.5 MPa and ry ¼20.5 μm. It indicates that the size of the plastic zone at the crack tip is larger than the average grain size (12.2 μm). The T4treated alloy can obtain the same result. The plastic zone size, ry ¼24.3 μm, is also larger than the average grain size (19.3 μm). After the aging treatment, the precipitation results in a sharp increase in the yield strength of the alloy. Considering ΔK ¼3 MPa  m1/2 and sy ¼260.3 MPa, the plastic zone size of the T6-treated alloy is approximately 14.1 μm, less than the average grain size as 19.7 μm. The larger plastic zone size permits the crack to interact with several grains simultaneously during the cyclic loading. This will accumulate large damages, leading to a faster propagation rate of the fatigue cracks, and vice versa. On the contrary, for the smaller plastic zone and relatively larger grain

size, the damage is much smaller. For this reason, the propagation rate of the T6-treated alloy is significantly lower than those of the as-cast and T4-treated alloys. It explains the disappearance of the life gap in the S–N curves of the T6-treated alloy (Fig. 3). The weak incoherent interface between the eutectic compounds and α-Mg matrix [10] is not the reason discussed here. In Fig. 9, different deformation morphologies were observed in the three conditions. The as-cast alloy has the most significant amount of the basal slip lines (Fig. 9b), and the T4-treated alloy (Fig. 9d) shows obvious less amount of the basal slip lines. The T6treated alloy assumes no deformation morphologies at all. Those morphologies are intimately related to the extent of plastic deformation of the alloy before fatigue fracture. It illustrates that the T6-treated alloy does not experience much plastic deformation before the initiation of localized fatigue cracks, unlike the as-cast and T4-treated alloys. The ductility of the T6-treated alloy is too low to introduce the plastic deformation before the failure. The fatigue properties probably can be further improved by the ductility increase at an expense of the strength of the alloy.

5. Conclusion The influence of solution treatment (T4) and solution plus aging treatment (T6) on the microstructure, tensile properties,

Z. He et al. / Materials Science & Engineering A 604 (2014) 78–85

and fatigue properties at room temperature of the semicontinuous casting Mg96.34Gd2.5Zn1Zr0.16 alloy was investigated in the present study. The following conclusions can be made. (1) Solution treatment leads to modifications of eutectic (Mg, Zn)3Gd compounds and the LPSO structured X phase, increase of the average grain size, and the precipitation of the Zn–Zr phase. The shape of the eutectic compounds changes from dendritic to elliptic, with a little decrease in content. After the solution treatment, the X phase becomes larger in size, and the solution content of the α-Mg matrix only varies a little. Most of the new Zn–Zr precipitates are located near grain boundaries, inhibiting grain growth, together with the residual (Mg,Zn)3Gd compounds. The modification effect of the solution treatment on the tensile and fatigue properties is not noticeable. (2) The subsequent aging does not leads to the modification of the eutectic compounds, LPSO structured X phase, the average grain size, and the Zn–Zr phase. The meta-stable β0 phase with orthorhombic structure and β1 phase with fcc structure are precipitated when aged at 200 1C for 128 h. The precipitates can greatly improve both the yield strength by 90 MPa and the fatigue strength by 18 MPa, for the alloy. (3) The basal slip lines are observed on the surface of the as-cast and T4-treated fatigued samples, believed to be the main deformation mechanism during the high cycle fatigue. (4) There are life gaps in the as-cast and T4-treated alloys, i.e., there is no failure between 105 and 107 cycles. It does not exist in the T6-treated alloy. The plastic zone sizes of the as-cast and T4-treated alloys are much larger than their average grain sizes, leading to a faster propagation rate of fatigue cracks. This can explain the phenomenon of the life gaps.

Acknowledgment This work was supported by National Natural Science Foundation of China project (51201103) and Specialized Research Fund for the Doctoral Program of Higher Education (20110073120008).

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