Microstructures and creep behavior of as-cast and annealed heat-resistant Mg–4Al–2Sr–1Ca alloy

Materials Science and Engineering A 531 (2012) 130–140

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Microstructures and creep behavior of as-cast and annealed heat-resistant Mg–4Al–2Sr–1Ca alloy Jing Bai ∗ , Yangshan Sun, Feng Xue, Jian Zhou School of Materials Science and Engineering, Southeast University, Jiangning, 211189, Nanjing, China

a r t i c l e

i n f o

Article history: Received 24 May 2011 Received in revised form 27 September 2011 Accepted 17 October 2011 Available online 21 October 2011 Keywords: Magnesium alloys Strontium Calcium Creep Microstructure

a b s t r a c t Microstructures, mechanical and creep properties of Mg–4Al–2Sr–1Ca alloy were investigated. As-cast microstructure of the experimental alloy consists of dendritic ␣-Mg and grain boundary intermetallics, predominantly lamellar eutectic C14–Mg2 Ca and bulky Mg–Al–Sr ternary phase. In addition, small amounts of C36–(Mg, Al)2 Ca and Mg17 Sr2 phases were also observed. Annealing at 400 ◦ C leads to the transformation of Laves phases from C14 (or C14 + C36) to C15. Meanwhile, the lamellar eutectic tends to be spheroidised and the continuous intermetallic network is broken up with prolongation of annealing time. The as-cast alloy shows a very high creep resistance at the temperatures between 150 and 200 ◦ C and applied stresses between 50 and 80 MPa. Annealing at 400 ◦ C results in a remarkable decrease of creep properties due to the morphological modification of the grain boundary intermetallics. It is proposed that both dislocation motion and grain boundary sliding contribute to the creep deformation of the present alloy. © 2011 Elsevier B.V. All rights reserved.

1. Introduction As the lightest metallic engineering structural material, Mg alloys, especially Mg–Al based alloys, such as AZ91D (Mg–9Al–1Zn, wt%) and AM60B (Mg–6Al–0.4Mn, wt%), offer superior die castability and a good balance of strength and ductility, thereby has been increasingly used in the automotive industry in view of the worldwide growing environmental concerns for energy conservation and recycle promotion in the past several years. Mg–Al alloys, however, currently faces a challenge in meeting the performance requirements of most critical automotive components, such as transmission and engine parts, due to the restriction of strength and creep resistance at elevated temperatures above 150 ◦ C [1]. ␤-Mg17 Al12 phase is generally considered to be a major factor deteriorating the mechanical properties of Mg–Al alloys at elevated temperature [2]. Some Mg–Al series alloys, such as Mg–Al–RE (RE: rare earth misch metal, which is usually composed of Ce, La, Nd and Pr etc.) system [3,4] (AE alloys) and Mg–Al–Si system [5] (AS alloys), have been developed for improving elevated-temperature performance. These alloys, however, have a high cost or other disadvantages, e.g. poor diecastability, high oxidation rate, low fatigue properties and even decomposition of intermetallics at elevated temperatures. It was recently reported [6,7] that alkaline earth

∗ Corresponding author. Tel.: +86 25 52090689. E-mail address: [email protected] (J. Bai). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.10.045

elements Sr and Ca are effective alloying elements on improving the high temperature properties of Mg–Al based alloys. Some main intermetallics, Al4 Sr [8], Mg2 Ca [9], Al2 Ca [10] or (Mg, Al)2 Ca [11] have been observed in the microstructure of Mg–Al based alloys with addition of Sr or Ca. All these intermetallics have high melting point and thermal stability, significantly resisting the creep deformation. The microstructures and mechanical properties of the alloys based on Mg–Al–Ca and Mg–Al–Sr ternary systems were systematically investigated in the previous work [12], however, there is a lack of research concerning the effect of combined additions of Sr and Ca on the microstructures and creep behavior of Mg–Al based alloys. In the present work alloy AJX421 (where J represents Sr and X represents Ca) were prepared and the creep tests performed on this alloy reveal that it has excellent creep resistance at temperatures up to 200 ◦ C. In this paper the microstructures and mechanical properties of alloy AJX421 were studied, and the mechanisms responsible for creep deformation were also investigated. According to the Mg–Al binary phase diagram, the highest solubility of Al in the ␣-Mg is 12.6 wt% at the eutectic temperature of 437 ◦ C and decreases rapidly with the reduction of temperature. The ␤-Mg17 Al12 phase will dissolve into the ␣-Mg matrix with an increase of temperature for most commercial Mg–Al alloys. Accordingly, the conventional Mg–Al alloys can be solid solution and aging strengthened by T4 and T6 heat treatment. For the present alloy, some new intermetallics which have high thermal stability at elevated temperatures replaced the Mg17 Al12 with the addition of Sr and Ca. Thereby, it is of interests investigating the microstructural change of these intermetallics as well as the effect on mechanical

J. Bai et al. / Materials Science and Engineering A 531 (2012) 130–140 Table 1 Chemical compositions (wt%) of AJX421 alloy. Alloying elements

Al

Sr

Ca

Mn

Impurity

Mg

Designed compositions Analyzed compositions

4.0 3.89

2.0 2.17

1.0 1.01

0.3 0.32

– <0.05

Bal. Bal.

properties after annealing at high temperature. In the present paper, therefore, the microstructures, tensile and creep properties of AJX421 alloy after annealing at 400 ◦ C for 5–200 h were investigated. In addition, the transition of three calcium containing Laves phases and their thermodynamic and structural stability were also discussed. 2. Experimental procedure Mg–4Al–2Sr–1Ca alloy with small amount of Mn addition was prepared and its designed compositions are listed in Table 1. To fabricate this alloy, strontium, calcium and manganese were conducted by adding master alloys of Mg–27 (wt%) Sr, Mg–30 (wt%) Ca and Al–10 (wt%) Mn, respectively. Pure Mg and Al stocks were used to achieve the target compositions. The alloy was melted in a mildsteel crucible under the protection of a mixed gas atmosphere of SF6 (1 v/v) and CO2 (Bal.). The melt was held at 740 ◦ C for 10 min then poured into a cylindrical water cooled mold made of cast copper with 60 mm diameter and 250 mm height. The chemical composition of the ingots was determined by inductively coupled plasma analyzer (ICP) and the results are listed in Table 1. Homogenization annealing heat treatment was performed at the temperature of 400 ◦ C, close to the solutionizing temperature in order to accelerate atomic diffusion, for time ranges from 5 h to 200 h, and followed by water quenching. The Vickers hardness was measured in an HVI-10A Vickers hardness tester with a load of 2.5 kg and a dwell time of 15 s. The tensile strength was performed using a CMT5105 Electronic Universal Testing Machine on a flat specimen with gauge section size of 18 mm in length, 3.2 mm in width and 1.8 mm in thickness. Cylindrical specimens of 100 mm gauge length and 10 mm diameter cross section was used for the creep testing by RD2 -3 high-temperature creep testing machine with constant load and the temperature maintained within ±2 ◦ C. Based on the present works, the creep curves of all alloys studied have been gone through the steady state level under the experimental conditions within 100 h. The creep tests in this paper, therefore, were stopped at 100 h and it is reliable to obtain the minimum creep rate by these creep curves. The microstructures of the alloy studied were examined by optical microscopy (OM), scanning electron microscopy (SEM) equipped with energy dispersive X-ray spectroscopy (EDS) and

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transmission electron microscopy (TEM). For TEM studies, discs with 0.1 mm in thickness and 3 mm in diameter were cut and mechanically polished, and then subjected to twin-jet electropolishing in a solution of ethanol with 3 vol.% perchloric acid. The crystal structure of the intermetallics was characterized using selected area diffraction (SAD) patterns and X-ray diffractometer (XRD).

3. Results 3.1. Microstructures 3.1.1. As-cast alloy The optical micrograph of as-cast sample is shown in Fig. 1(a), from which it can be seen that the as-cast microstructures of this alloy consist of dendritic ␣-Mg matrix and intermetallics which form an almost continuous network surrounding the matrix grains. SEM observations on as-cast specimen reveal two predominant morphologies of these second-phases: lamellar eutectic with bright contrast and bulk-shaped phase mainly residing at triple junctions with intermediate contrast, as shown in Fig. 1(b). The SAD analysis was conducted on thin foil specimens to identify the crystalline structure of the intermetallics. Fig. 2(a) shows the TEM bright field image of the eutectic compound with its corresponding SAD pattern taken along the [1¯ 2 1¯ 0] zone axis, which is indexed as arising from a hexagonal Laves structure C14 consistent with the structure of Mg2 Ca phase. In some intergranular eutectic region two kinds of intermetallics with different contrast can be observed, as shown in Fig. 2(b). The corresponding SAD pattern (on the top right corner of Fig. 2(b)) taken from the particles with dark contrast exhibits extra reflections (small spots) besides the C14 reflections (big spots). This is consistent with the reflections of dihexagonal Laves C36 structure with the same incident beam directions [1¯ 2 1¯ 0] in the SAD pattern of Fig. 2(a). This C36 phase is commonly represented as (Mg, Al)2 Ca in Mg–Al–Ca system. The atom arrangement on the close-packed plane of the Laves hexagonal structure C14 and C36 is very similar, but the existence of reflections between 0 0 0 0 and 0 0 0 2 in the SAD pattern of Fig. 2(b) indicates that the unit cell of the C36 is about twice that of the C14 along the [0 0 0 1] direction. In the previous investigations both the C14 and C36 phases were observed in Mg–Al–Ca based alloys with different content of Al and Ca [9]. From SEM observations, Suzuki et al. [9] reported that the eutectic compounds C14 (Mg2 Ca) and C36 ((Mg, Al)2 Ca) showed different morphologies, fine lamella for C14 and coarse irregular shape for C36, respectively. In the present study, however, the C36 phase cannot be distinguished from C14 phase by morphology resulting from the volume fraction of the C36 ((Mg, Al)2 Ca) is very small compared with C14 (Mg2 Ca) phase.

Fig. 1. Microstructure of as-cast AJX421 alloy: (a) optical micrograph and (b) SEM micrograph.

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Fig. 2. (a) TEM image of Mg2 Ca and its corresponding SAD pattern (B = 1¯ 2 1¯ 0); (b) TEM image of the region of (Mg, Al)2 Ca and Mg2 Ca and their corresponding SAD pattern (B = 1¯ 2 1¯ 0).

TEM observations additionally reveal a small amount of intermetallic with light contrast surrounding the lamellar Mg2 Ca in some eutectic region, as indicated by arrowhead in Fig. 3(a). This intermetallic is identified as Mg17 Sr2 with a hexagonal structure ˚ c = 10.356 A). ˚ Fig. 3(b) is a (space group P63 /mmc, a = 10.535 A, SAD pattern taken from both Mg2 Ca and Mg17 Sr2 phases along a ¯ zone axis of [1 1¯ 0 2]Mg 2 Ca. By indexing the pattern it can be concluded that there is an orientation relationship between Mg17 Sr2 and Mg2 Ca as: (0 0 0 1)Mg17 Sr2 //(1 1¯ 0 1)Mg2 Ca, (1 1¯ 0 1)Mg17 Sr2 //(1 1 2¯ 0)Mg2 Ca ¯ [1¯ 2 1¯ 0]Mg17 Sr2 //[1 1¯ 0 2]Mg 2 Ca.

The chemical compositions of the bulky phase shown in Fig. 1(b) with intermediate contrast is approximately 17.5 wt% Al–7.9 wt% Sr–Mg measured by EDS. However, the electron diffraction patterns taken from this compound cannot be identified due to the lack of information in regard to the crystal structure of Mg–Al–Sr ternary intermetallics. In order to further confirm the crystalline structure of the intermetallics existing in the as-cast alloy studied, XRD was conducted on as-cast specimen and the result is shown in Fig. 4, from which two main phases, ␣-Mg and Mg2 Ca were positively identified. No peaks in the XRD pattern can be indexed as arising from the phases of (Mg, Al)2 Ca and Mg17 Sr2 because the volume fractions of these phases in the microstructure are too small. Some distinct peaks in

Fig. 3. (a) TEM image and (b) SAD pattern taken from both of Mg17 Sr2 and Mg2 Ca (B = 1¯ 2 1¯ 0Mg17 Sr2 //1 1¯ 0 2¯ Mg2 Ca ).

J. Bai et al. / Materials Science and Engineering A 531 (2012) 130–140

Fig. 4. X-ray diffraction pattern of as-cast AJX421 with descriptions of phases detected.

this XRD pattern cannot be indexed according to the PDF cards. Based on the results of microstructural observation mentioned above it is proposed that these peaks are from the bulky ternary phase shown in Fig. 1(b). Similar ternary phases were reported by the previous investigations [13], however the locations of the peaks appearing in the XRD patter of the present paper do not match well with that reported in the previous work. Therefore, this phase is probably a new ternary compound and is tentatively designated as ␶ phase in the XRD pattern (Fig. 4). 3.1.2. Microstructures of annealed alloy Some samples of the alloy studied were heat treated at 400 ◦ C for times of 5 h, 15 h, 50 h, 100 h and 200 h, respectively, followed by water quenching. The optical micrograph taken from the sample annealed at 400 ◦ C for 200 h is shown in Fig. 5. In comparison with as-cast microstructures (Fig. 1(a)), the difference between the as-cast and heat treated microstructure is unnoticeable. The microstructures of the annealed sample also consist of the ␣Mg matrix and intermetallic compounds surrounding the matrix grains, suggesting that the Sr and Ca containing intermetallics have high thermodynamic stability and hardly dissolved into the ␣-Mg matrix after heat treatment at elevated temperature of 400 ◦ C for time up to 200 h. However, the intermetallic network seen in the ascast microstructures (Fig. 1(a)) becomes discontinuous along grain boundaries. High magnification SEM observations performed on the samples annealed at 400 ◦ C for different times reveal the change in microstructures during annealing and the results are shown in Fig. 6. Compared with the as-cast microstructures (Fig. 1(b)), the

Fig. 5. Optical micrograph of 200 h annealed AJX421 sample.

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eutectic found at grain boundaries in the sample annealed for 15 h is no longer connected lamellar, as shown in Fig. 6(a). With prolongation of annealing time, the eutectic structure gradually breaks up, and then changes to irregular blocks or plate shaped particles after 100 h, finally, becomes discrete spherical particles after 200 h (Fig. 6(b)). It is believed that the intermetallics transform to an increasingly favored shape with increasing annealing time to reduce the interface energy. Hence, the eutectic phases tend to reduce excess surface area and spheroidise with slow diffusion of Mg and Al atoms during the annealing process. Compared with the as-cast microstructures (Fig. 1(b)), there are hardly remarkable change in the volume fraction and morphology of the bulky Mg–Al–Sr ternary phase after annealing at 400 ◦ C, even for 200 h (Fig. 6(b)), suggesting that the bulky phase has relatively higher thermodynamic stability than eutectic Ca containing phases. TEM observations were also conducted on annealed samples. Fig. 7(a) shows the TEM bright field image of eutectic phase in the annealed sample and corresponding SAD pattern, which can be indexed as arising from Al2 Ca of cubic structure C15 with B = 5 2 1. In addition, observations on some eutectic regions reveal a high density of staking faults with a spacing of several nanometer (Fig. 7(b)). The diffraction pattern (down-left corner) taken from this region is consistent with the twinned C15 structure (type of 1 1 2 {1 1 1}) with B = 3¯ 2 1 and the stacking faults are parallel to the close-packed plane of the cubic C15 structure. There is no C14 or C36 eutectic found by TEM observations in the sample annealed for 200 h, indicating that both C14 (Mg2 Ca) and C36 ((Mg, Al)2 Ca) have completely transformed to the C15 (Al2 Ca) during long time annealing. 3.2. Mechanical properties In order to gain an insight into the relationship between microstructures and mechanical behavior, the hardness, tensile and creep properties of the alloy studied on both as-cast and annealed conditions were tested. 3.2.1. Hardness and tensile properties The Vickers hardness (HV) measurement of the experimental alloy as a function of the annealing time is shown in Fig. 8(a). The first point on the hardness curve corresponds to the hardness of the as-cast sample. The hardness value increases with prolongation of annealing time and reaches a maximum value after annealing for 15 h, then decreases with increase of annealing time gradually. After 100 h annealing, the effect of annealing time on hardness is not visible and the hardness values does not show obvious variation when the annealing treatment is further prolonged. Tensile properties of as-cast and annealed samples were tested and the results are based on the average of three specimens for each sample on each condition. The variation trend of tensile properties at ambient temperature with increase of annealing time is similar to that of hardness, as shown in Fig. 8(b). The as-cast sample has relatively lower strength, including ultimate strength ( b ), yield strength ( 0.2 ), and ductility compared to that of annealed samples. The highest yield strength correlating with ductility is obtained from the 15 h annealed sample. Both strength and ductility of annealed sample slightly decrease if the annealing time is further prolonged. 3.2.2. Creep properties For heat-resistance magnesium alloys, the creep property plays a more important role on selection and application of alloys for high temperature components. The tensile creep tests in the present investigation were carried out under the conditions of temperature between 150 and 200 ◦ C and applied stress between 50 and 80 MPa. Two sets of typical creep curves are shown in Fig. 9(a) and

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Fig. 6. The microstructural change in AJX421 alloy with annealing at 400 ◦ C for: (a) 15 h, (b) 200 h.

Fig. 7. (a) TEM image and SAD pattern of eutectic region (B = 5 2¯ 1); (b) TEM image and SAD pattern of fine twin of Al2 Ca (B = 3¯ 2 1).

Fig. 8. (a) Vickers hardness of the alloy as a function of annealing time at 400 ◦ C and (b) tensile strength and elongation of as-cast and annealed AJX421 alloy.

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Fig. 9. (a) Creep curves of as-cast AJX421 at 175 ◦ C and applied stress between 50 and 80 MPa, and (b) creep curves of as-cast AJX421 at 70 MPa and temperature between 150 and 200 ◦ C.

(b). The average minimum creep rates obtained from several different records are also exhibited in the figures. The first set of curves, as shown in Fig. 9(a), was obtained at the temperature of 175 ◦ C and applied stress in the range of 50–80 MPa. From these curves it can be seen that the creep rate increases with the increase of applied stress. In Fig. 9(b), the second set of curves was obtained under the applied stress of 70 MPa and temperatures range between 150 and 200 ◦ C. In this case the creep rate also increases with a rise of temperature. In both figures all creep curves present an obvious primary region, which occurs within the first 10 h and results in about 60–80% of the creep elongation during the total 100 h, and is then followed by a long region of almost constant creep rate. However, the tertiary-stage creep cannot be observed within a 100 h under any conditions. Creep resistance is usually evaluated by the minimum creep rate and creep extension. In comparison with commercial heat-resistant magnesium alloys, such as AS21 and AE42 [1,2,13,14], the AJX421 alloy shows much lower minimum creep rate and elongation. The creep behavior of polycrystalline materials is now understood reasonably well. The important characteristic of high˙ generally temperature creep is that the steady-state creep rate, ε, varies with the applied stress, , the absolute temperature, T, through a relationship described by the following equation [15]:

ε˙ =

ADGb kT

 b p   n d

G

(1)

where D is the diffusion coefficient, defined as D0 exp(−Q/RT) (where D0 is a frequency factor, Q is the activation energy for creep, R is the gas constant, T is the absolute temperature), G is the shear modulus (=1.92 × 104 − 8.6T MPa) for magnesium [16], b is the Burgers vector, k is Boltzmann’s constant, d is the grain size, p and n are the exponents of the inverse grain size and the stress, respectively, and A is a dimensionless constant. Fig. 10(a) shows the all experimental data logarithmically plot˙ versus the normalized stress, at the 175 ◦ C and ted as strain rate, ε, the applied stress between 50 and 80 MPa. The results of samples tested can be fitted by a straight line with a stress exponent n of about 3.5.

According to Eq. (1), the creep activation energy, Q, can be calculated based on the following relation: Q =

˙ n−1 T ) ∂ ln(εG ∂(−1/RT )

(2)

˙ n−1 T ) against The value of Q may be determined by plotting ln(εG 1/T for selected , as shown in Fig. 10(b), in which Q is estimated as ∼73 kJ/mol at the temperatures ranging from 150 to 200 ◦ C on the condition of  = 70 MPa with n = 3.5. In order to study the effect of annealing treatment on the creep properties, creep tests were also performed for some annealed samples under the condition of 175 ◦ C and 70 MPa. The 100 h creep curves are shown in Fig. 11(a) and the creep curve of an as-cast sample obtained on the same condition is also plotted in the figure. At the same time, the minimum creep rates obtained from the steady-state stage of creep curves and 100 h creep extension for selected samples are plotted in Fig. 11(b). Compared with the development of hardness and tensile properties during annealing process, the influence of heat treatment on creep resistance is more remarkable. The creep properties significantly decrease with the proceeding of annealing. At as-cast state, the alloy exhibits excellent creep property under the conditions of 175 ◦ C/70 MPa. The minimum creep rate of as-cast samples is as low as 3.5 × 10−10 s–1 . The minimum creep rate of the sample annealed for 15 h, however, increases to 5.2 × 10−9 s–1 , one order of magnitude higher than that of as-cast sample. Further deterioration of creep resistance is found with the increase of annealing time, and the highest minimum creep rate obtained from the sample annealed for 200 h is 5.2 × 10−8 s–1 , beyond two orders of magnitude higher than that of as-cast sample. None of these samples reach the point to fracture after 100 h creep test. The change of creep extension with annealing time exhibits a similar trend to the minimum creep rate, as shown in Fig. 11(b). 3.3. Post-crept microstructure In order to better understand the mechanism responsible for the change of creep behavior caused by annealing the post-crept microstructures along the longitudinal direction of creep specimens have been investigated. After creep tests at 175 ◦ C and 70 MPa for 100 h, no change can be found in comparison with the as-cast microstructures. The morphology and size of the intermetallics in

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Fig. 10. (a) Minimum creep rate as a function of the stress for AJX421 at 175 ◦ C and (b) minimum creep rates as a function of the inverse temperature for AJX421 at 70 MPa.

Fig. 11. (a) Creep curves and (b) minimum creep rate and creep elongation of as-cast and annealed samples at 175 ◦ C/70 MPa during 100 h.

the post-crept microstructure is basically the same as that found in as-cast microstructures. When testing temperature increases to 200 ◦ C, small amounts of microcracks approximately normal to the tensile stress axis were observed inside the intermatellics after 100 h creep, as indicated by arrowheads in Fig. 12(a). Microcracks were also observed in the annealed samples, however, the location of the cracks is different from that in as-cast samples. SEM

observations for the 100 h post-crept sample annealed for 200 h reveal that microcracks initiate at the interfaces between Al2 Ca and Mg–Al–Sr ternary phase, and propagate along grain boundaries, as shown in Fig. 12(b) and (c), respectively. More detailed observations were performed for the sample after creep at 175 ◦ C/70 MPa for 100 h using TEM, as shown in Fig. 13(a) and (c), revealing a kinds of new plate shaped precipitates which

Fig. 12. SEM micrographs of (a) microcracks within intermetallics of as-cast alloy after creep at 200 ◦ C/70 MPa for 100 h, and (b) microcracks at the interface of Al2 Ca and Mg–Al–Sr phase and (c) cracks propagating along grain boundary in 200 h annealed sample after creep at 175 ◦ C/70 MPa for 100 h.

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Fig. 13. After creep at 175 ◦ C/70 MPa for 100 h in the as-cast AJX421 alloy, (a) TEM image and (b) SAD pattern of the plate shaped Al2 Ca particles parallel to basal plane of ␣-Mg (B = [1¯ 1 0]Al2 Ca //[1¯ 1 0 0]␣-Mg ); (c) TEM image and (d) SAD pattern of the plated shaped Al2 Ca particles parallel to prismatic plane of ␣-Mg (B = [1¯ 1 1]Al2 Ca //[1¯ 1 0 0]␣-Mg ); (e) interaction between dislocations and Al2 Ca particles.

form inside grains during creep and present two different directions, parallel to the basal planes and the prismatic planes of ␣-Mg, respectively. This particle intermetallic is idendified by SAD pattern (Fig. 13(b) and (d)) as cubic Laves structure C15–Al2 Ca phase, same as the crystal structure of the eutectic region after annealing, and appears to be two kinds of orientation relationship with the ␣-Mg matrix as: (1 1 1)Al2 Ca//(0 0 0 1)␣, [1¯ 1 0]Al2 Ca//[1¯ 1 0 0]␣

(Fig. 13(b)),

(1 1 0)Al2 Ca//(0 0 0 1)␣-Mg, [1¯ 1 1]Al2 Ca//[1¯ 1 0 0]␣-Mg (Fig. 13(d))

Simultaneously, large numbers of dislocation, as well as their interaction with Al2 Ca particles were detected by TEM, as shown in Fig. 13(e), where a series of dislocations are confined on the slip plane and pile up against the particle precipitates. This demonstrates that the Al2 Ca precipitates are effective on restricting the dislocation motion within the matrix grains during the creep process and further reducing creep deformation.

4. Discussion 4.1. Microstructures For Mg–Al series alloys, it is reasonable to anticipate that reduction or elimination of Mg17 Al12 , which facilitates grain boundary migration in elevated temperature creep process [2], will lead to improvement in creep resistance. In this study, the major intermetallics, Mg2 Ca, (Mg, Al)2 Ca and the Mg–Al–Sr ternary phase, substitute the formation of Mg17 Al12 and have high thermodynamic stability. Their morphology and distribution hardly show obvious change after creep at the temperature of 200 ◦ C. Furthermore, the interaction between dislocation and Al2 Ca plate shaped precipitates revealed by TEM observations (Fig. 13(e)) indicates that these particles have the effect of impeding dislocation motion in the interior of the grains and thus contribute good creep property to the alloy studied. This is consistent with the previous reports [17,18]. The as-cast AJX421, consequently, can provide superior creep resistance in comparison with conventional Mg–Al based alloys. Since Laves phases C14 and C36 were found as mainly eutectic phase forming along grain boundaries in alloy studied, the crystal structure of these phases are also of interest in this study. The

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Table 2 Lattice parameters of C14 and C36 phases identified in AJX421 alloy. Phase

Mg2 Ca from JCPDS standard Mg2 Ca (Mg, Al)2 Ca

Structure

Hexagonal Hexagonal Hexagonal

Table 3 Parameters for the Mg2 Ca and Al2 Ca phase [25].

Lattice parameters (Å)

Phase

Terms

a (J/mol)

b (J/mol)

a

c

6.24 6.21 6.17

10.14 9.94 19.39

Mg2 Ca Al2 Ca

Gf Gf

−12704.4 −29281.7

1.80939 5.38176

Laves phase is a type of intermetallic compound with the AB2 stoichiometry that can exist with crystalline structures of cubic C15, hexagonal C14 or dihexagonal C36 [19]. In this notation, C14 and C15 are hcp and fcc based structures with hexagonality of 100% and 0%, respectively. C36 is an intermediate structure between C14 and C15 with the hexagonality of 50% [20] and the lattice parameter of C36 structure along [0 0 0 1] direction is twice that of C14 structure in terms of the stacking sequence of Laves structure. The precise lattice parameters of the both phases in this study have been calculated according to their electron diffraction patterns and the results are listed in Table 2. The values of a and c of Mg2 Ca phase ˚ smaller than those of in AJX421 alloy are a = 6.21 A˚ and c = 9.94 A, data of Mg2 Ca from JCPDS standard because the solution of a small amount of Al atoms with 1.43 A˚ in radius substitute Mg atoms with 1.60 A˚ in radius in the unit cell. In the unit cell of (Mg, Al)2 Ca, more Al atoms enter and substitute Mg atoms than that in Mg2 Ca, therefore, the a value of (Mg, Al)2 Ca is smaller than that of Mg2 Ca and the c value of (Mg, Al)2 Ca is smaller than twice that of Mg2 Ca, as listed in Table 2. Based on the stacking sequence of close-packed planes in Laves phases, the transformation between these Laves phases is geometrically possible by a shear mechanism involving synchro-shockley dislocations within the interior of the individual intermetallic compound region [21,22] and dose not require separate nucleation of the product phase. Twinning, therefore, is thought to be a consequence of the C14 to C15 transformation based on a number of experimental observations [23] and a suggested synchroshear mechanism in the process of forming twinning in Laves phase has been reported by Kumar and Hazzledine [19]. In this investigation, the C15 twinning resulted from the high density stacking faults observed in annealed AJX421 alloy (Fig. 8(b)) confirms the transition of C14 (Mg2 Ca) (or C14 + C36 (Mg, Al)2 Ca) to C15 (Al2 Ca) during annealing process. The thermodynamic stability of the Laves phases can be indicated by comparison of their Gibbs free energy. The Gibbs energy 0, function Gi (T ) = Gi0 (T ) − HiSER for the pure element i in the ϕ phase is described by the equation: 0,

Gi

(T ) = a + bT + cT ln T + d T 2 + eT 3 + fT −1 + gT 7 + hT −9

For the phases of Mg2 Ca and Al2 Ca, the values of Gf obtained by calculations based on Eqs. (3)–(5) as a function of the temperature from 25 to 450 ◦ C is shown in Fig. 14. It can be seen that the free energy of Al2 Ca is smaller than that of Mg2 Ca, indicating that Al2 Ca has relatively higher thermodynamic stability at the calculational temperature ranging. This is consistent with the direction of phase transition between the Laves, from metastable C14 to stable C15, during annealing. The bond energy of Al2 Ca and Mg2 Ca, on the other hand, also reflects the stability of crystal structure and were calculated in our previous paper [26] by means of the bond length difference (BLD) method based on the empirical electron theory (EET) of solids and molecules formulated. The bond energy of three main bonds in Al2 Ca structure are 56.88 kJ/mol for Ca–Ca bond, 22.83 kJ/mol for Al–Al bond and 17.63 kJ/mol for Ca–Al bond, respectively. In the unit cell of Mg2 Ca, however, the strongest bond energy is only 20.60 kJ/mol between Mg (II) and Mg (II). Especially, there are several obvious weak bonds, such as Ca–Mg (I) and Mg (I)–Mg (II) with bond energy of 5.82 kJ/mol and 9.89 kJ/mol, respectively. Al2 Ca, therefore, shows higher structural stability and higher melting point, 1079 ◦ C, in comparison with Mg2 Ca with a melting point of 717 ◦ C. As an intermediate structure in the pseudo-binary system between Mg2 Ca and Al2 Ca, the C36 phase is considered to have intermediate stability between Mg2 Ca and Al2 Ca. 4.2. Creep behavior ˙ of polycrystalline The steady-state secondary creep rate (ε) materials is generally described by a relationship of Eq. (1). The n and Q parameters are used to infer the dominant creep deformation mechanisms for a material in specific ranges of stress and temperature. Therefore, a preliminary assessment of the likely creep deformation mechanisms for the alloys studied can be proposed by determining the stress exponent (n) and the apparent activation energy (Q), as shown in Fig. 10(a) and Fig. 10(b). There is a wide range of stresses at intermediate stresses and a single testing temperature where the datum points usually fall along a line having a slope within the range of ∼3–6 [14,16,27]. This region is usually related with the power-law creep and associated

(3)

where HiSER is the molar enthalpy of the stable element reference (SER) at 298.15 k and 1 bar, and T is the absolute temperature. The value of the coefficient a to h for Mg, Al and Ca are taken form the SGTE compilation by Dinsdale [24]. The Gibbs energy for binary stoichiometric compounds is described by the following equation: G = xi Gi0,1 + xj Gj0,2 + Gf

(4)

Gf = a + bT

(5)

where xi , xj are the mole fraction of component i and j, respectively, and Gi0,1 , Gj0,2 represent the Gibbs energy of a component in its standard state. However the Gibbs energy of the compound phase may refer to a different crystal structure than those of the pure elements, ϕ1 andϕ2 . Gf is the Gibbs energy for the formation of per mole stoichiometric compound. The parameters a and b in Eq. (5) for Mg2 Ca and Al2 Ca are obtained by optimization in the phase equilibria and thermodynamic data from [25] and shown in Table 3.

Fig. 14. The Gibbs free energy as a function of the temperature.

J. Bai et al. / Materials Science and Engineering A 531 (2012) 130–140

unambiguously with some form of dislocation creep in which the creep strain accrued from the intragranular movement of dislocation through the processes of glide and climb. Generally, the rate-controlling mechanism may be a dislocation glide with n = 3 and is defined as class A (alloy-type) behavior, which involves viscous glide motion of dislocations dragging solute atmospheres (Al atoms in Mg–Al based magnesium alloys) along their glide path. And the n-value in the range of ∼4–6 are indicative of class M (metal-type) behavior where creep is believed to be controlled by a recovery process of dislocation climb by vacancy diffusion. Additionally, at higher stresses, there is commonly a rapid increase in the value of n with n > 7 due either to a breakaway of dislocations from their solute atmospheres in class A behavior or power-law breakdown. The creep results of studied alloy AJX421 exhibit a stress exponent of n = 3.5 at 175 ◦ C between 50 and 80 MPa, and fall in the range of 3–6, associating with dislocation controlled creep. At the same time, it is also believed that both the viscous glide motion of dislocations drag aluminum atom atmospheres and dislocation climb affect the creep deformation process at the investigated conditions. According to above theory, if the creep is dominantly controlled by solute drag creep related with n = 3, the apparent activation energy, Q, should be approximately equal to activation energy for inter-diffusion of Al in the Mg lattice (143 kJ/mol) [28]. For dislocation climb creep with n between 4 and 6, on the other hand, the Q-value can be closely evaluated as activation energy for lattice self-diffusion in pure magnesium (135 kJ/mol) [14]. However, the value of Q obtained for as-cast AJX421 at applied stress 70 MPa and the temperature range of 150–200 ◦ C, is 73 kJ/mol, much lower than 143 kJ/mol or 135 kJ/mol, but very close to the activation energy for grain boundary diffusion in pure magnesium (82 kJ/mol) [14]. Similar values of Q were also found in binary Mg–Sc alloys after T6 heat treatment at 40 MPa and low-temperature range by Mordike et al. [29]. This suggests that grain boundary plays an important role in creep process of the studied alloy and grain boundary migration or sliding can still make strong contributions to the overall creep strain. The as-cast AJX421 alloy offers excellent creep resistance with elongation of 0.1% after 100 h creep test at 175 ◦ C and 70 MPa. The grain boundary intermetallics in the as-cast microstructures forms a continuous network which is effective on inhibiting grain boundary migration and sliding so that the alloy shows very high creep resistance in comparison with that of commercial Mg–Al based alloys. However, the creep property, measured as minimum creep rate and 100 h creep elongation, significantly decreases after annealing. The decrease of creep resistance may be accounted for by the following factors. Firstly, the grain boundary intermetallics are no longer effective at inhibiting grain boundary sliding during creep deformation due to the morphology changes and the disconnection of intermetallic compound network occurring during annealing. Secondly, the effect of solid solution hardening is weakened due to annealing. In the as-cast microstructure the ␣Mg matrix is a supersaturated solid solution since the cooling rate of the melt in the water cooled mold is very high. The phase transition from Mg2 Ca to Al2 Ca during annealing causes the reduction of Al concentration in the matrix grain, which reduces the effect of solid solution strengthening. In as-cast samples, microcracks usually form in grain boundary intermetallics during the creep process because of their high hardness and brittleness. With the prolongation of annealing time, disconnection and spheroidisation of grain boundary intermetallics occurs so that more microcraks initiate at the interfaces between the ␣-Mg matrix and second-phases or between the Al2 Ca and the ternary Mg–Al–Sr phase during creep deformation due to the pileup of dislocation and stress concentration. Therefore, it is evident that the continuity of grain boundary intermetallics is a dominant

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factor affecting creep resistance of Mg–Al–Sr–Ca alloys. The results on creep tests obtained in the present investigation also verify that the contributions to creep by grain boundary sliding cannot be ignored. Also, the mechanism responsible for the creep deformation in the studied alloy is a mixed one composed of both dislocation motion and grain boundary sliding. Therefore, controlling of grain boundary microstructures is critical in the design and application of alkaline containing Mg–Al based alloys. It is also important to consider additional matrix strengthening methods. 5. Conclusions (1) The as-cast microstructure of alloy AJX421 consists of a dendritic ␣-Mg and intermetallics, which are the lamellar eutectic C14–Mg2 Ca and a bulky type Mg–Al–Sr ternary phase forming a continuous network along grain boundaries. Meanwhile, a small amount of C36–(Mg, Al)2 Ca and Mg17 Sr2 along the grain boundary and plate shaped C15–Al2 Ca precipitates in the matrix grains were also observed and they have determinate orientation relationships with Mg2 Ca and ␣-Mg, respectively. (2) The distribution of grain boundary phases does not show obvious changes after annealing at 400 ◦ C for 200 h. However, the lamellar eutectic tended to be spheroidised and the continuous intermetallic network is broken up with prolongation of annealing time. Annealing at 400 ◦ C causes the transformation of Laves phases from C14 (or C14 + C36) to C15. (3) As-cast AJX421 alloy shows very high creep resistance at temperature of 175 ◦ C and applied stress range from 50 to 80 MPa and at applied stress of 70 MPa and temperatures range from 150 to 200 ◦ C. Under the above conditions the stress exponent n and creep activation energy Q are 3.5 and 73 kJ/mol, respectively. Annealing at 400 ◦ C results in remarkable decrease of creep properties due to the morphology modification of the grain boundary intermetallics. From the result of the creep tests, it is proposed that both dislocation motion and grain boundary sliding contribute to the creep deformation. Acknowledgements This research was supported by the Natural Science Foundation of Jiangsu Province of China (No. BK2010392) and the Innovation Foundation of Southeast University (No. 3212000502). The authors are grateful to Nanjing Welbow Metals Co., Ltd. for providing Mg–Ca and Mg–Sr master alloys. References [1] A.A. Luo, Int. Mater. Rev. 49 (2004) 13–30. [2] W. Blum, B. Watzinger, P. Zhang, Adv. Eng. Mater. 2 (2000) 349–355. [3] J. Zhang, Z. Leng, S. Liu, M. Zhang, J. Meng, R. Wu, J. Alloys Compd. 509 (2011) L187–L193. [4] M. Sumida, Recent Pat. Mater. Sci. 4 (2011) 94–105. [5] W. Blum, P. Zhang, B. Watzinger, B.V. Grossmann, H.G. Haldenwanger, Mater. Sci. Eng. A 319–321 (2001) 735–740. [6] T. Sato, M.V. Kral, Mater. Sci. Eng. A 498 (2008) 369–376. [7] Y. Nakaura, A. Watanabe, K. Ohori, Mater. Trans. 47 (2006) 1031–1039. [8] B. Jing, S. Yangshan, X. Shan, X. Feng, Z. Tianbai, Mater. Sci. Eng. A 419 (2006) 181–188. [9] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Acta Mater. 53 (2005) 2823–2834. [10] A. Suzuki, N.D. Saddock, J.R. Terbush, B.R. Powell, J. Wayne Jones, T.M. Pollock, Magnesium Technology 2007: TMS 2007 Annual Meeting and Exhibition, Orlando, FL, United States, 2007, pp. 375–380. [11] J. Bai, Y. Sun, F. Xue, S. Xue, J. Qiang, T. Zhu, J. Alloys Compd. 437 (2007) 247–253. [12] M.O. Pekguleryuz, E. Baril, Mater. Trans. 42 (2001) 1258–1267. [13] E. Baril, P. Labelle, M.O. Pekguleryuz, JOM 55 (2003) 34–39. [14] M.O. Pekguleryuz, A.A. Kaya, Adv. Eng. Mater. 5 (2003) 866–878. [15] T.G. Langdon, Metall. Mater. Trans. A: Phys. Metall. Mater. Sci. 33 (2002) 249–259. [16] S.S. Vagarali, T.G. Langdon, Acta Metall. 29 (1981) 1969–1982.

140

J. Bai et al. / Materials Science and Engineering A 531 (2012) 130–140

[17] B. Smola, I. Stulikova, J. Pelcova, B.L. Mordike, J. Alloys Compd. 378 (2004) 196–201. [18] J.F. Nie, Scripta Mater. 48 (2003) 1009–1015. [19] K.S. Kumar, P.M. Hazzledine, Intermetallics, vol. 12, Elsevier Ltd, 2004, pp. 763–770. [20] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Scripta Mater. 51 (2004) 1005–1010. [21] P.M. Hazzledine, P. Pirouz, Scripta Metall. Mater. 28 (1993) 1277–1282. [22] M. Grujicic, S. Tangrila, O.B. Cavin, W.D. Porter, C.R. Hubbard, Mater. Sci. Eng. A A160 (1993) 37–48. [23] T. Takasugi, S. Hanada, M. Yoshida, Mater. Sci. Eng. A A192 (19) (1995) 805–810.

[24] A.T. Dinsdale, CALPHAD: Comput. Coupling Phase Diag. Thermochem. 15 (1991) 317–425. [25] F. Islam, M. Medraj, Proceedings of the 2004 CSME Forum, London, Ontario, Canada, 2004. [26] X. Min, Y. Sun, F. Xue, W. Du, D. Wu, Mater. Chem. Phys. 78 (2003) 88–93. [27] M. Regev, E. Aghion, A. Rosen, Mater. Sci. Eng. A: Struct. Mater.: Properties, Microstruct. Process. A234 (23) (1997) 123–126. [28] W.J. Kim, S.W. Chung, C.S. Chung, D. Kum, Acta Mater. 49 (2001) 3337–3345. [29] B.L. Mordike, I. Stulikova, B. Smola, Metall. Mater. Trans. A: Phys. Metall. Mater. Sci. 36 (2005) 1729–1736.