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Characterisation of precipitates in an aged Mg–Zn–Ti alloy
Journal of Alloys and Compounds 472 (2009) 171–177
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Characterisation of precipitates in an aged Mg–Zn–Ti alloy J. Buha National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 11 April 2008 Received in revised form 3 May 2008 Accepted 6 May 2008 Available online 18 June 2008 Keywords: Magnesium alloys Precipitation TEM SEM
a b s t r a c t Up to about 0.2 at.% of Ti may be dissolved in the solid magnesium in the presence of Zn (in Mg–2.4 at.% Zn alloy) which is orders of magnitude higher than the solubility of this element in the pure solid magnesium. The presence of Ti does not modify significantly the constituent particles which are compositionally and morphologically similar to the eutectic phase and contain up to about 0.1 at.% of Ti. Ti is exceptionally effective in promoting the nucleation of the precipitates during ageing. The precipitates formed during ageing in Mg–Zn–Ti alloy were generally similar to those forming in the binary alloy, although the results indicate that Ti promotes favourable nucleation of some precipitate species thus changes the relative fractions of the precipitates present in the microstructure. The greatest level of hardening in Mg–Zn–Ti alloy can be produced by ageing at about 70 ◦ C due to the formation of a dense dispersion of at least five types of precipitates observed also in the binary alloy. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The use of magnesium and magnesium alloys offers a significant weight reduction when compared to aluminium alloys and steels. For this reason magnesium alloy are becoming increasingly attractive for the potential light-weight applications in the automotive and aerospace industries. Cast magnesium alloys already have some application in the automotive industry (seat frames, steering wheel, etc.). Only a small number of wrought magnesium alloys, mainly those processed by extrusion, are currently available. This is primarily due to a number of challenges associated with the mechanical processing of magnesium alloys. Most magnesium alloys, both cast and wrought, exhibit mechanical properties that are still inadequate for a wider application. The mechanical properties of some magnesium alloys may be controlled and improved by age hardening, however the increment in strength/hardness produced by ageing is considerably lower than what is commonly observed in aluminium alloys. This is primarily due to a much lower number density of the precipitates formed, which are then easily by-passed by dislocations during deformation. Alloys based on the Mg–Zn system are particularly interesting for further development as wrought alloys. These alloys also respond well to aged hardening, typically conducted at temperatures above 150 ◦ C (T6 heat treatment). Recent studies show that a favourable combination of the mechanical properties, such as high ductility and hardness combined with an appreciably high yield
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strength, can be produced by ageing at much lower temperatures, in particular at ambient temperature [1] and most likely also at intermediate temperatures (up to about 100 ◦ C) [2]. The precipitation in the Mg–Zn based alloys generally takes place through the formation of a number of intermediate phases [2–10]: SSSS → solute clusters → GP zones
→ 1 (rods and blocky precipitates ⊥ {0 0 0 1}Mg ; possibly Mg4 Zn7 ) →
2 (coarse
plates||{0 0 0 1}Mg and laths ⊥ {0001}Mg;
MgZn2 ) →  equilibrium phase (MgZn or Mg2 Zn3 ) The decomposition of the supersaturated solid solution (SSSS) initially occurs through clustering of solute atoms [2,10]. Depending on the ageing temperature and alloy composition, three types of coherent Guinier-Preston (GP) zones may form: GP1 zones as plates on {1 1 2¯ 0}Mg [11], GP2 zones as oblate spheroids on {0 0 0 1}Mg [11], and GP zones as discs on {0 0 0 1}Mg [12]. The 1 phase has possibly a monoclinic crystal structure and may exhibit two morphologies and corresponding orientation relationships with the magnesium lattice: long rod and block (cube), both forming perpendicular to basal plane of magnesium. The rod-shaped form of the 1 phase, also referred to as [0 0 0 1]Mg rod, is typical for the peak-aged T6 condition and has been considered to be the major strengthening precipitate in the Mg–Zn based alloys. A smaller fraction of the block-shaped 1 phase may also be present in the
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Table 1 Chemical composition of Mg–Zn and Mg–Zn–Ti alloys and the heat treatments applied
peak-aged condition [3]. The 2 phase has a hexagonal crystal structure and may also exhibit two morphologies and respective orientation relationships with the magnesium lattice: coarse plate parallel to basal plane and long lath perpendicular to basal plane. The lath-shaped form of the 2 phase, also referred to as [0 0 0 1]Mg lath, may be present in the peak-aged T6 condition [3]. These laths are often indistinguishable from the 1 rods if viewed edge-on, however in the cross section, i.e. from the [0 0 0 1]Mg direction, the two types of precipitates can be easily distinguished. In addition to [0 0 0 1]Mg rods and laths, thin plates on basal planes, some of which are GP zones [13–15], may also be present in the peak-aged T6 condition. The GP zones on basal planes are only a single magnesium atomic layer thick. The precipitates here referred to as (0 0 0 1)Mg plates are slightly thicker (a few magnesium atomic planes), however considerably smaller and thinner than the coarse plates of the 2 phase that typically form upon overageing. The formation of the latter precipitates generally leads to a loss of mechanical properties due to their very sparse distribution. Alloys aged at ambient temperature are strengthened primarily by the combination of the GP1 zones, solute (co)clusters and prismatic precipitates of an unknown phase [1,2,10,13]. At intermediate temperatures, a large number of precipitates typical for both artificial and natural ageing may form [2], many of which are yet to be fully characterised. It has been well established that the precipitates in the T6 condition of a binary Mg–Zn alloy are generally very coarse, widely spaced and inhomogeneously distributed [6] thus give rise to limited strengthening. The number density of the precipitates, thus also the mechanical properties achieved by ageing, may be increased by additional alloying with some elements. Alloying elements such as Ag [4], Au [16], rare earth (RE) elements [17], Cu [18,19] and Ca [13] have been known to be effective in this regard, although sometimes at the expense of alloy’s corrosion resistance (e.g. in the case of Cu, and to some extent also Ag and Ca [20]), or at the expense of alloy cost (e.g. in the case of alloying with Au, Ag and RE). In the more recent studies, a number of novel and uncommon alloying elements have been introduced, some of which have never been used in magnesium alloys before and many of which had been presumed insoluble in the magnesium lattice. It has been shown that small amounts (trace or micro-alloying additions) of Ba [13], Cr [15], V [14] and Ti [21] have a number of disproportionately beneficial effects (relative to their content in the alloy) on the microstructure and mechanical properties of Mg–Zn alloy. Alloying with Ti [21] proves to be particularly beneficial. Addition of this element simultaneously results in a significant grain refinement, considerably increased number density of the precipitates in T6 condition and rapid natural ageing. Ti therefore represents a particularly suitable novel alloying addition for magnesium alloys. The solubility of Ti in the solid magnesium is unknown and may possibly be about 0.001 at.% at the peritectic temperature [22]. The natural ageing response of Mg–Zn–Ti alloy was reported earlier [21] and the comprehensive study of the mechanism of grain refinement by Ti will be a subject of separate publication. The current article describes in detail the effect of Ti on the precipitation and precip-
itates formed primarily during artificial ageing (T6) and ageing at intermediate temperatures. 2. Experimental procedures The ternary Mg–Zn–Ti and binary Mg–Zn alloys of the compositions given in Table 1 were prepared from the pure Mg, pure Ti and an Mg–Zn pre-alloy (prepared in the same manner) using an induction melting furnace under protective Ar atmosphere, and then cast into a permanent mould. Both alloys were homogenized and then solution heat treated, as described in Table 1, prior to quenching and ageing at ambient temperature (∼22 ◦ C) in air and at 70, 95 and 160 ◦ C in oil baths. Age hardening was monitored by microhardness measurements made using a load of 50 g and hardness values reported here were averaged form 12 measurements. Scanning electron microscopy (SEM) of the as-cast and as-homogenized alloy specimens was performed using JEOL JSM 5400 scanning electron microscope equipped with JEOL JED 2140 energy dispersive X-ray (EDX) microanalyzer. Transmission electron microscopy (TEM) observations were performed using Phillips CM200 microscope operated at 200 kV. High resolution TEM (HRTEM) observations were conduced using FEI Tecnai G2 F30 microscope operated at 300 kV. The specimens for TEM and HRTEM observations were prepared from aged alloys by electropolishing in a solution of 10.6 g LiCl and 22.32 g Mg(ClO4 )2 in 1000 ml of methanol and 200 ml of 2-butoxi-ethanol at about −45 ◦ C and 90 V. The model diffraction patterns (Fig. 4) were constructed with the help of EMS software [23].
3. Results and discussion 3.1. The constituent particles and solubility of Ti in Mg–Zn alloy The SEM images of the as-cast and as-homogenized alloys are shown in Fig. 1. The precipitates observed generally along grain boundaries in the cast binary alloy (Fig. 1a) were identified as the eutectic phase Mg7 Zn3 [13]. Since the alloy’s composition and the homogenization temperature fall within the two-phase region of the Mg–Zn phase diagram between the eutectic and eutectoid temperatures [24], the eutectic phase was also present after homogenization (Fig. 1b). Morphologically similar particles were observed in Mg–Zn–Ti alloy. The particles in the cast Mg–Zn–Ti alloy (Fig. 1c) contained 23.4–30.1 at.% Zn and about half of the particles analysed also contained 0.01–0.12 at.% Ti, with Mg being the reminder. It is possible that these particles were also the eutectic phase containing a trace amount of Ti, although a detailed crystallographic analysis is required in order to confirm this. The EDX scans of the magnesium matrix in the precipitate-free regions showed presence of Ti in the amount between 0.02 and 0.2 at.% in some areas, while in other areas of the specimen no Ti could be detected. This value is slightly higher than the Ti content in the particles. It is suspected that fine particles of the pure Ti might have been present in the microstructure (in the grain interiors away from the grain boundary) and these would have contributed to the EDX spectra even if embedded under the surface of the specimen. In the homogenized Mg–Zn–Ti alloy (Fig. 1d) the particles observed contained 26–28.1 at.% Zn and 0.01–0.03 at.% Ti. The magnesium matrix between the particles contained between 0.01 and 0.05 at.% Ti in some regions. The lower level of Ti in the magnesium matrix after homogenization is suspected to be a result of annealing which led to possibly more homogeneous distribution of Ti and dissolu-
J. Buha / Journal of Alloys and Compounds 472 (2009) 171–177
173
Fig. 1. SEM images of the as-cast (a and c) and as-homogenized (b and d) microstructures of Mg–Zn–Ti (c and d) and Mg–Zn (a and b) alloys.
tion of the Ti, or Ti-rich, aggregates in the same way in which the distribution of Zn is affected by the annealing of a cast alloy. In addition to the previously published findings [21], the current results undoubtedly show that Ti exhibits a considerable solubility in the magnesium lattice in the presence of Zn, which is at least an order of magnitude greater than the solubility of this element in the pure solid magnesium.
[1] alloys. This indicates that different alloying elements added to an Mg–Zn alloy promote or favour nucleation of different types of precipitates. The microstructures produced by ageing are described in detail in the following sections.
3.2. Age hardening Hardness curves in Fig. 2 show that the addition or a small amount of Ti to Mg–Zn alloy significantly improves both the magnitude and the kinetics of hardening observed during artificial (T6) and natural (T4) ageing [21]. The hardness increment produced by artificial ageing is nearly doubled compared to that of the binary alloy and peak hardness was reached after a considerably shorter ageing time. As observed also in a number of other Mg–Zn based alloys, the greatest level of hardening in Mg–Zn–Ti alloy was produced by ageing at intermediate temperatures. Hardness increment of 39 VHN was produced after 208 h (8.7 days) of ageing at 95 ◦ C with very little change in the hardness being observed on prolonged ageing. The greatest level of hardening was produced by ageing at 70 ◦ C, which was 44 VHN after 408 h (17 days). Hardness value reached nearly a plateau after about 2 months of ageing (1344 h) and a hardness increment of 47 VHN, and then remained on a similar level for the duration of the experiment. Although the hardening behaviour of Mg–Zn–Ti alloy during artificial ageing and ageing at intermediate temperatures seems generally comparable to that of other Mg–Zn based alloys [1,2,13,15], the microstructure or more precisely the relative fractions of the precipitates formed are visibly different from those observed in the binary [2] and Mg–Zn–Cu–Mn
Fig. 2. Hardness curves of Mg–Zn–Ti alloy (circles) aged at 160 ◦ C (T6), 95 ◦ C, 70 ◦ C and at ambient temperature (∼22 ◦ C; T4) as indicated in the chart. The T6 and T4 hardness curves of the binary Mg–Zn alloy (diamonds) are also included for comparison (after [21]).
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3.3. Precipitates formed by ageing at 160 ◦ C The [2 1¯ 1¯ 0]Mg TEM image of the microstructure of the binary alloy produced by ageing at 160 ◦ C shows elongated precipitates aligned with the [0 0 0 1]Mg direction and also a number of precipitates having blocky and plate-shaped morphology, believed to be the variants of the 1 and 2 phases, respectively (Fig. 3a). As reported earlier, the number density of the precipitates in the Ticontaining alloy was significantly higher than in the binary alloy [5], which is also evident from the comparison of Fig. 3a and b. This indicates that Ti is very effective in promoting the nucleation of the precipitates. In comparison to the binary alloy and other Mg–Zn based alloys [2,10,13–15], Mg–Zn–Ti alloy exhibited a noticeably higher fraction of the blocky precipitates, believed to correspond to the 1 phase (Fig. 3b). This suggests that Ti eases the nucleation of this morphological and orientational variant of the 1 phase. ¯ ¯ Based on the [2 1 1 0]Mg HRTEM observations, fine (0 0 0 1)Mg plates and GP zones commonly observed in other Mg–Zn based alloys [1,2,13–15] did not represent a large fraction of the precipitates in the T6 peak-aged microstructure (two rods and a blocky precipitate are visible in Fig. 3c). The GP zones, which are a single magnesium atomic layer in thickness, and thin (0 0 0 1)Mg plates, which are several magnesium atomic layers thick, are among the precipitates that make the greatest contribution to strengthening in alloys such as Mg–Zn–Cu–Mn [1], Mg–Zn–Cr [15], Mg–Zn–Ba [13] or Mg–Zn–V [14]. These precipitates were however rarely observed in the T6 microstructure of Mg–Zn–Ti alloy. One plate-like precipitate is shown in [0 0 0 1]Mg HRTEM image in Fig. 4a. The major contribution to the hardening was made primarily by the [0 0 0 1]Mg rods similar to 1 and to a smaller extent also by the [0 0 0 1]Mg laths similar to 2 . The [0 0 0 1]Mg HRTEM images of the major strengthening precipitates in the T6 condition are shown in Fig. 4. The end-on [0 0 0 1]Mg rod similar to 1 is shown in Fig. 4a along with a (0 0 0 1)Mg plate, while the end-on [0 0 0 1]Mg lath similar to 2 is shown in Fig. 4d. The Fourier transform (FFT) of the HRTEM image corresponding to the [0 0 0 1]Mg rod (Fig. 4e) may be indexed according to the [0 0 1] zone axis of a base centred monoclinic lat tice [25] reported for the 1 phase in a binary alloy [4]. However, closer examination of the FFTs from several rods showed that the inter-planar angles, and possibly also the spot distances, do not match those reported for the 1 phase in a binary alloy in the [0 0 1] precipitate orientation. This indicates that the lattice parameters of the rods in Mg–Zn–Ti alloy are slightly different to those of the equivalent precipitates in the binary alloy. Similar observations were made on rod-like precipitates in Mg–Zn–Ba [13] and Mg–Zn–V [14] alloys and it is suspected that the subtle differences in the lattice parameters may be due to the presence of atoms of the ternary alloying elements within the precipitates. The FFT from the (0 0 0 1)Mg plate (Fig. 4b) indicates that its crystal structure is very similar to that of the 2 phase having the orientation relationship with the magnesium matrix such that [0 0 0 1]p ||[0 0 0 1]Mg and (0 1¯ 0 1)p ||(1¯ 1¯ 2 0)Mg (corresponds to a plate-shaped morphology of the 2 phase), thus very similar also to the crystal structure of fine (0 0 0 1)Mg plates observed in a number of other Mg–Zn based alloys aged to T6 temper or at intermediate temperatures [1,2,13,15]. The corresponding model diffraction pattern showing the magnesium and precipitate reflections is given in Fig. 4c. The satellite reflections around each precipitate spot, as well as around the magnesium reflections and the central spot, arise from double diffraction within the precipitate. The [0 0 0 1]Mg laths were found to be very similar to those observed in the binary alloy based on the analysis of the corresponding FFT (Fig. 4f compared to a model diffraction pattern in Fig. 4g) thus very similar to the lath-shaped form of the 2 phase having the orientation relationship with the magnesium
Fig. 3. [2 1¯ 1¯ 0]Mg TEM image of Mg–Zn alloy aged to peak-hardness at 160 ◦ C (a); [2 1¯ 1¯ 0]Mg TEM (b) and HRTEM (c) images of Mg–Zn–Ti alloy also aged to peakhardness at 160 ◦ C. Insets in (a) and (b) are corresponding diffraction patterns (DP).
J. Buha / Journal of Alloys and Compounds 472 (2009) 171–177
175
Fig. 5. [2 1¯ 1¯ 0]Mg TEM (a) and HRTEM (b) images of the microstructure aged 4 weeks at 70 ◦ C. Inset in (a) is corresponding FFT.
matrix such that [2 1¯ 1¯ 0]p ||[0 0 0 1]Mg and (0 0 0 1)p ||(1¯ 2 1¯ 0)Mg . The exact matching of the lattice parameters could not be confirmed due to inadequate quality of the FFT. These observations indicate that the precipitates in the T6 condition of Mg–Zn–Ti alloy are generally very similar to those forming in the binary alloy and other Mg–Zn based alloys, although the lattice parameters are likely slightly different. In contrast to the binary and other Mg–Zn based alloys, the presence of Ti does not seem to be favouring the formation of GP zones and fine (0 0 0 1)Mg plates on basal planes in alloy aged at 160 ◦ C. 3.4. Precipitates formed by ageing at 70 ◦ C Fig. 4. [0 0 0 1]Mg HRTEM images (a and d) of the strengthening precipitates in Mg–Zn–Ti alloy aged to peak-hardness at 160 ◦ C; FFTs from the precipitates (b, e and f) with corresponding model DPs (c is a model DP for FFT in b; g is a model DP for FFT in f; large circles correspond to Mg reflections and small circles to precipitate reflections).
The [2 1¯ 1¯ 0]Mg TEM and HRTEM images of the microstructure produced by ageing at 70 ◦ C are shown in Fig. 5a and b, respectively. Based on these images, a considerable fraction of the precipitates present in the microstructure were fine (0 0 0 1)Mg plates which gave rise to characteristic dark strain field contrast (Fig. 5a) typically present either side of a plate with a line of no contrast (bright line) in the middle. The bright line parallel to basal plane indicates that a plate-like precipitate is coherent with the magnesium lattice
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alent of the 2 phase. The precipitates visible in HRTEM image in Fig. 6b are believed to most likely be equivalent to the [0 0 0 1]Mg rods, [0 0 0 1]Mg laths and prismatic precipitates observed earlier in the binary alloy aged at the same temperature [2]. Compared to the binary alloy [2], a considerably higher fraction of thin plates on {2 1¯ 1¯ 0}Mg planes were also observed (arrowed in Fig. 6b). These precipitates correspond well to GP1 zones, observed also in a number of Mg–Zn based alloys aged at ambient and intermediate temperatures. Although the types of the precipitates observed in Mg–Zn–Ti alloy aged at 70 ◦ C are generally similar to those observed in the binary alloy [2] and in Mg–Zn–Cu–Mn alloy [1], their relative fractions are noticeably different. The Ti-containing alloy exhibited a considerably higher fraction of the fine (0 0 0 1)Mg plates, while the GP zones were not present in any substantial amount. The fraction of GP1 zones on {2 1¯ 1¯ 0}Mg planes was also noticeably higher in the Ti-containing alloy. Based on the hardness measurements, the hardness levels and possibly also the tensile properties of Mg–Zn based alloys containing different ternary alloying additions are most likely very similar. However, it is interesting to note that even a trace amount of these ternary additions affect the type of the precipitates that from in the microstructure as well as the temperature at which these precipitates form. In this regard, Ti seems to be favouring the formation of GP1 zones and fine (0 0 0 1)Mg plates (but not GP zones) at intermediate temperatures, such as 70 ◦ C, however (0 0 0 1)Mg plates were not observed in the T6 condition. This suggest that, in comparison to the binary alloy, Ti may be lowering the solvus line of the (0 0 0 1)Mg plates and reducing their thermal stability, as these precipitates were not readily observed in alloy aged at 160 ◦ C. On the other hand, Ti promotes the nucleation of the (0 0 0 1)Mg plates, which can be deduced from their increased fraction in alloy aged at 70 ◦ C.
4. Summary Fig. 6. [0 0 0 1]Mg TEM (a) and HRTEM (b) images of the microstructure aged 4 weeks at 70 ◦ C. The orientation of the images is indicated by the inset FFT in (a).
along the [0 0 0 1]Mg direction. Some of these plates are indicated by black arrows in Fig. 5b (one is about 2 nm wide). Some distortion of the magnesium lattice in the vicinity of these precipitates is also visible. Elongated precipitates aligned with the [0 0 0 1]Mg direction were also observed from the [2 1¯ 1¯ 0]Mg zone axis (Fig. 5a) although these precipitates that produced almost no elastic strains in the surrounding lattice were mainly obscured by the strain field contrast around the plate-like precipitates. Some clearly visible elongated precipitates are indicated by white arrows in Fig. 5a and b. In addition, precipitates of possibly prismatic morphology were also observed by HRTEM (indicated by star-like symbols in Fig. 5b). The [0 0 0 1]Mg TEM and HRTEM images in Fig. 6 show far more clearly the different types of the precipitate species present in the microstructure and also give a better indication of their relative fractions. Fig. 6a shows that in addition to the fine plates on basal planes, visible as honeycomb-looking moire´ fringes form this direction (some are circled), a large number of other precipitates viewed end-on were also present in the microstructure. No usable diffraction information could be obtained from most of these precipitates. The FFT from the precipitate arrowed in Fig. 6a indicates that its crystal structure and the orientation relationship with the magnesium lattice are very similar to that of the [0 0 0 1]Mg laths observed in the T6 condition and suspected to be the equiv-
• Ti is soluble in the solid magnesium in the presence of Zn in the amount that is at least an order of magnitude higher than the solubility of Ti in the pure solid magnesium. • The constituent particles in Mg–Zn–Ti alloy are compositionally and morphologically similar to the eutectic Mg7 Zn3 phase and may contain up to about 0.1 at.% Ti. • The highest level of hardening in Mg–Zn–Ti alloy can be achieved by ageing at 70 ◦ C for less than 3 weeks, which is generally consistent with the age hardening response of the binary and other Mg–Zn based alloys. • The microstructure produced by ageing at 70 ◦ C consisted of at least five different types of precipitates: [0 0 0 1]Mg rods, [0 0 0 1]Mg laths, [0 0 0 1]Mg prismatic precipitates, (0 0 0 1)Mg plates and GP1 zones. Similar precipitates were observed earlier also in the binary alloy, although the relative fractions of these precipitate species in Mg–Zn–Ti alloy were different as a result of the presence of Ti. The results suggest that Ti promotes the formation of the (0 0 0 1)Mg plates and GP1 zones at intermediate temperatures. • The microstructure of Mg–Zn–Ti alloy aged to T6 temper is strengthened primarily by the [0 0 0 1]Mg rods (equivalent to 1 in the binary alloy) and to a smaller extent by the [0 0 0 1]Mg laths (equivalent to 2 ), as well as blocky precipitates (possibly also 1 ). The crystal structures of these precipitates are most likely identical to those of the equivalent precipitates forming in the binary alloy although the lattice parameters may be slightly different, presumably as a result of the presence of Ti atoms within the precipitates.
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Acknowledgment The author wishes to thank the Japan Society for the Promotion of Science (JSPS) for providing financial support for this work in the form of JSPS Postdoctoral Fellowship. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
J. Buha, Mater. Sci. Eng. A, in press. doi:10.1016/j.msea.2008.12.006. J. Buha, Mater. Sci. Eng. A 492 (2008) 11–19. X. Gao, J.F. Nie, Scripta Mater. 56 (2007) 645–648. G. Mima, Y. Tanaka, Trans. JIM 12 (1971) 76–81. L. Sturkey, J.B. Clark, J. Inst. Metals 88 (1959–1960) 177–181. J.B. Clark, Acta Metall. 13 (1965) 1281–1289. J. Gallot, Thesis, University of Rouen, 1966. M. Bernole, J. Gallot, R. Graf, J. Microscopie 4 (1965) 787–792. G. Mima, Y. Tanaka, Trans. JIM 12 (1971) 71–75. J. Buha, T. Ohkubo, Metall. Mater. Trans. A 39 (2008) 2259–2273.
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[11] T. Takahashi, Y. Kojima, K. Takanishi, Jpn. J. Inst. Light Metals 23 (1973) 376–382. [12] I.J. Polmear, Light Alloys, Metallurgy of the Light Metals, 3rd ed., Arnold, London, 1995, p. 204. [13] J. Buha, Mater. Sci. Eng. A 491 (2008) 70–79. [14] J. Buha, Acta Mater 56 (2008) 3533–3542. [15] J. Buha, Mater. Sci. Eng. A 492 (2008) 293–299. [16] E.O. Hall, J. Inst. Metals 96 (1968) 21–27. [17] N.K. Suseelan, M.C. Mittal, Mater. Sci. Forum 30 (1988) 89–94. [18] W. Unsworth, J.F. King, Magnesium Technology, Institute of Metals, London, 1987, pp. 25–31. [19] G.W. Lorimer, Magnesium Technology, Institute of Metals, London, 1986, pp. 47–53. [20] C.J. Bettles, M.A. Gibson, K. Venkatesan, Scripta Mater. 51 (2003) 193–198. [21] J. Buha, J. Mater. Sci. 43 (2008) 1220–1227. [22] M. Hansen, Constitution of Binary Alloys, 2nd ed., McGraw-Hill Book Co., New York, 1958, p. 924 (written in cooperation with K. Anderko). [23] P.A. Stadelmann, Ultramicroscopy 21 (1987) 131–146. [24] A.A. Nayeb-Hashemi, J.B. Clark (Eds.), Phase Diagrams of Binary Magnesium Alloys, ASM International, 1988, p. 353. [25] L.L. Rokhlin, A.A. Oreshkina, Fiz. Metal. Metalloved. 66 (1988) 559–566.
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Characterisation of precipitates in an aged Mg–Zn–Ti alloy J. Buha National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan
a r t i c l e
i n f o
Article history: Received 11 April 2008 Received in revised form 3 May 2008 Accepted 6 May 2008 Available online 18 June 2008 Keywords: Magnesium alloys Precipitation TEM SEM
a b s t r a c t Up to about 0.2 at.% of Ti may be dissolved in the solid magnesium in the presence of Zn (in Mg–2.4 at.% Zn alloy) which is orders of magnitude higher than the solubility of this element in the pure solid magnesium. The presence of Ti does not modify significantly the constituent particles which are compositionally and morphologically similar to the eutectic phase and contain up to about 0.1 at.% of Ti. Ti is exceptionally effective in promoting the nucleation of the precipitates during ageing. The precipitates formed during ageing in Mg–Zn–Ti alloy were generally similar to those forming in the binary alloy, although the results indicate that Ti promotes favourable nucleation of some precipitate species thus changes the relative fractions of the precipitates present in the microstructure. The greatest level of hardening in Mg–Zn–Ti alloy can be produced by ageing at about 70 ◦ C due to the formation of a dense dispersion of at least five types of precipitates observed also in the binary alloy. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The use of magnesium and magnesium alloys offers a significant weight reduction when compared to aluminium alloys and steels. For this reason magnesium alloy are becoming increasingly attractive for the potential light-weight applications in the automotive and aerospace industries. Cast magnesium alloys already have some application in the automotive industry (seat frames, steering wheel, etc.). Only a small number of wrought magnesium alloys, mainly those processed by extrusion, are currently available. This is primarily due to a number of challenges associated with the mechanical processing of magnesium alloys. Most magnesium alloys, both cast and wrought, exhibit mechanical properties that are still inadequate for a wider application. The mechanical properties of some magnesium alloys may be controlled and improved by age hardening, however the increment in strength/hardness produced by ageing is considerably lower than what is commonly observed in aluminium alloys. This is primarily due to a much lower number density of the precipitates formed, which are then easily by-passed by dislocations during deformation. Alloys based on the Mg–Zn system are particularly interesting for further development as wrought alloys. These alloys also respond well to aged hardening, typically conducted at temperatures above 150 ◦ C (T6 heat treatment). Recent studies show that a favourable combination of the mechanical properties, such as high ductility and hardness combined with an appreciably high yield
E-mail address: [email protected]. 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.05.019
strength, can be produced by ageing at much lower temperatures, in particular at ambient temperature [1] and most likely also at intermediate temperatures (up to about 100 ◦ C) [2]. The precipitation in the Mg–Zn based alloys generally takes place through the formation of a number of intermediate phases [2–10]: SSSS → solute clusters → GP zones
→ 1 (rods and blocky precipitates ⊥ {0 0 0 1}Mg ; possibly Mg4 Zn7 ) →
2 (coarse
plates||{0 0 0 1}Mg and laths ⊥ {0001}Mg;
MgZn2 ) →  equilibrium phase (MgZn or Mg2 Zn3 ) The decomposition of the supersaturated solid solution (SSSS) initially occurs through clustering of solute atoms [2,10]. Depending on the ageing temperature and alloy composition, three types of coherent Guinier-Preston (GP) zones may form: GP1 zones as plates on {1 1 2¯ 0}Mg [11], GP2 zones as oblate spheroids on {0 0 0 1}Mg [11], and GP zones as discs on {0 0 0 1}Mg [12]. The 1 phase has possibly a monoclinic crystal structure and may exhibit two morphologies and corresponding orientation relationships with the magnesium lattice: long rod and block (cube), both forming perpendicular to basal plane of magnesium. The rod-shaped form of the 1 phase, also referred to as [0 0 0 1]Mg rod, is typical for the peak-aged T6 condition and has been considered to be the major strengthening precipitate in the Mg–Zn based alloys. A smaller fraction of the block-shaped 1 phase may also be present in the
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Table 1 Chemical composition of Mg–Zn and Mg–Zn–Ti alloys and the heat treatments applied
peak-aged condition [3]. The 2 phase has a hexagonal crystal structure and may also exhibit two morphologies and respective orientation relationships with the magnesium lattice: coarse plate parallel to basal plane and long lath perpendicular to basal plane. The lath-shaped form of the 2 phase, also referred to as [0 0 0 1]Mg lath, may be present in the peak-aged T6 condition [3]. These laths are often indistinguishable from the 1 rods if viewed edge-on, however in the cross section, i.e. from the [0 0 0 1]Mg direction, the two types of precipitates can be easily distinguished. In addition to [0 0 0 1]Mg rods and laths, thin plates on basal planes, some of which are GP zones [13–15], may also be present in the peak-aged T6 condition. The GP zones on basal planes are only a single magnesium atomic layer thick. The precipitates here referred to as (0 0 0 1)Mg plates are slightly thicker (a few magnesium atomic planes), however considerably smaller and thinner than the coarse plates of the 2 phase that typically form upon overageing. The formation of the latter precipitates generally leads to a loss of mechanical properties due to their very sparse distribution. Alloys aged at ambient temperature are strengthened primarily by the combination of the GP1 zones, solute (co)clusters and prismatic precipitates of an unknown phase [1,2,10,13]. At intermediate temperatures, a large number of precipitates typical for both artificial and natural ageing may form [2], many of which are yet to be fully characterised. It has been well established that the precipitates in the T6 condition of a binary Mg–Zn alloy are generally very coarse, widely spaced and inhomogeneously distributed [6] thus give rise to limited strengthening. The number density of the precipitates, thus also the mechanical properties achieved by ageing, may be increased by additional alloying with some elements. Alloying elements such as Ag [4], Au [16], rare earth (RE) elements [17], Cu [18,19] and Ca [13] have been known to be effective in this regard, although sometimes at the expense of alloy’s corrosion resistance (e.g. in the case of Cu, and to some extent also Ag and Ca [20]), or at the expense of alloy cost (e.g. in the case of alloying with Au, Ag and RE). In the more recent studies, a number of novel and uncommon alloying elements have been introduced, some of which have never been used in magnesium alloys before and many of which had been presumed insoluble in the magnesium lattice. It has been shown that small amounts (trace or micro-alloying additions) of Ba [13], Cr [15], V [14] and Ti [21] have a number of disproportionately beneficial effects (relative to their content in the alloy) on the microstructure and mechanical properties of Mg–Zn alloy. Alloying with Ti [21] proves to be particularly beneficial. Addition of this element simultaneously results in a significant grain refinement, considerably increased number density of the precipitates in T6 condition and rapid natural ageing. Ti therefore represents a particularly suitable novel alloying addition for magnesium alloys. The solubility of Ti in the solid magnesium is unknown and may possibly be about 0.001 at.% at the peritectic temperature [22]. The natural ageing response of Mg–Zn–Ti alloy was reported earlier [21] and the comprehensive study of the mechanism of grain refinement by Ti will be a subject of separate publication. The current article describes in detail the effect of Ti on the precipitation and precip-
itates formed primarily during artificial ageing (T6) and ageing at intermediate temperatures. 2. Experimental procedures The ternary Mg–Zn–Ti and binary Mg–Zn alloys of the compositions given in Table 1 were prepared from the pure Mg, pure Ti and an Mg–Zn pre-alloy (prepared in the same manner) using an induction melting furnace under protective Ar atmosphere, and then cast into a permanent mould. Both alloys were homogenized and then solution heat treated, as described in Table 1, prior to quenching and ageing at ambient temperature (∼22 ◦ C) in air and at 70, 95 and 160 ◦ C in oil baths. Age hardening was monitored by microhardness measurements made using a load of 50 g and hardness values reported here were averaged form 12 measurements. Scanning electron microscopy (SEM) of the as-cast and as-homogenized alloy specimens was performed using JEOL JSM 5400 scanning electron microscope equipped with JEOL JED 2140 energy dispersive X-ray (EDX) microanalyzer. Transmission electron microscopy (TEM) observations were performed using Phillips CM200 microscope operated at 200 kV. High resolution TEM (HRTEM) observations were conduced using FEI Tecnai G2 F30 microscope operated at 300 kV. The specimens for TEM and HRTEM observations were prepared from aged alloys by electropolishing in a solution of 10.6 g LiCl and 22.32 g Mg(ClO4 )2 in 1000 ml of methanol and 200 ml of 2-butoxi-ethanol at about −45 ◦ C and 90 V. The model diffraction patterns (Fig. 4) were constructed with the help of EMS software [23].
3. Results and discussion 3.1. The constituent particles and solubility of Ti in Mg–Zn alloy The SEM images of the as-cast and as-homogenized alloys are shown in Fig. 1. The precipitates observed generally along grain boundaries in the cast binary alloy (Fig. 1a) were identified as the eutectic phase Mg7 Zn3 [13]. Since the alloy’s composition and the homogenization temperature fall within the two-phase region of the Mg–Zn phase diagram between the eutectic and eutectoid temperatures [24], the eutectic phase was also present after homogenization (Fig. 1b). Morphologically similar particles were observed in Mg–Zn–Ti alloy. The particles in the cast Mg–Zn–Ti alloy (Fig. 1c) contained 23.4–30.1 at.% Zn and about half of the particles analysed also contained 0.01–0.12 at.% Ti, with Mg being the reminder. It is possible that these particles were also the eutectic phase containing a trace amount of Ti, although a detailed crystallographic analysis is required in order to confirm this. The EDX scans of the magnesium matrix in the precipitate-free regions showed presence of Ti in the amount between 0.02 and 0.2 at.% in some areas, while in other areas of the specimen no Ti could be detected. This value is slightly higher than the Ti content in the particles. It is suspected that fine particles of the pure Ti might have been present in the microstructure (in the grain interiors away from the grain boundary) and these would have contributed to the EDX spectra even if embedded under the surface of the specimen. In the homogenized Mg–Zn–Ti alloy (Fig. 1d) the particles observed contained 26–28.1 at.% Zn and 0.01–0.03 at.% Ti. The magnesium matrix between the particles contained between 0.01 and 0.05 at.% Ti in some regions. The lower level of Ti in the magnesium matrix after homogenization is suspected to be a result of annealing which led to possibly more homogeneous distribution of Ti and dissolu-
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Fig. 1. SEM images of the as-cast (a and c) and as-homogenized (b and d) microstructures of Mg–Zn–Ti (c and d) and Mg–Zn (a and b) alloys.
tion of the Ti, or Ti-rich, aggregates in the same way in which the distribution of Zn is affected by the annealing of a cast alloy. In addition to the previously published findings [21], the current results undoubtedly show that Ti exhibits a considerable solubility in the magnesium lattice in the presence of Zn, which is at least an order of magnitude greater than the solubility of this element in the pure solid magnesium.
[1] alloys. This indicates that different alloying elements added to an Mg–Zn alloy promote or favour nucleation of different types of precipitates. The microstructures produced by ageing are described in detail in the following sections.
3.2. Age hardening Hardness curves in Fig. 2 show that the addition or a small amount of Ti to Mg–Zn alloy significantly improves both the magnitude and the kinetics of hardening observed during artificial (T6) and natural (T4) ageing [21]. The hardness increment produced by artificial ageing is nearly doubled compared to that of the binary alloy and peak hardness was reached after a considerably shorter ageing time. As observed also in a number of other Mg–Zn based alloys, the greatest level of hardening in Mg–Zn–Ti alloy was produced by ageing at intermediate temperatures. Hardness increment of 39 VHN was produced after 208 h (8.7 days) of ageing at 95 ◦ C with very little change in the hardness being observed on prolonged ageing. The greatest level of hardening was produced by ageing at 70 ◦ C, which was 44 VHN after 408 h (17 days). Hardness value reached nearly a plateau after about 2 months of ageing (1344 h) and a hardness increment of 47 VHN, and then remained on a similar level for the duration of the experiment. Although the hardening behaviour of Mg–Zn–Ti alloy during artificial ageing and ageing at intermediate temperatures seems generally comparable to that of other Mg–Zn based alloys [1,2,13,15], the microstructure or more precisely the relative fractions of the precipitates formed are visibly different from those observed in the binary [2] and Mg–Zn–Cu–Mn
Fig. 2. Hardness curves of Mg–Zn–Ti alloy (circles) aged at 160 ◦ C (T6), 95 ◦ C, 70 ◦ C and at ambient temperature (∼22 ◦ C; T4) as indicated in the chart. The T6 and T4 hardness curves of the binary Mg–Zn alloy (diamonds) are also included for comparison (after [21]).
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3.3. Precipitates formed by ageing at 160 ◦ C The [2 1¯ 1¯ 0]Mg TEM image of the microstructure of the binary alloy produced by ageing at 160 ◦ C shows elongated precipitates aligned with the [0 0 0 1]Mg direction and also a number of precipitates having blocky and plate-shaped morphology, believed to be the variants of the 1 and 2 phases, respectively (Fig. 3a). As reported earlier, the number density of the precipitates in the Ticontaining alloy was significantly higher than in the binary alloy [5], which is also evident from the comparison of Fig. 3a and b. This indicates that Ti is very effective in promoting the nucleation of the precipitates. In comparison to the binary alloy and other Mg–Zn based alloys [2,10,13–15], Mg–Zn–Ti alloy exhibited a noticeably higher fraction of the blocky precipitates, believed to correspond to the 1 phase (Fig. 3b). This suggests that Ti eases the nucleation of this morphological and orientational variant of the 1 phase. ¯ ¯ Based on the [2 1 1 0]Mg HRTEM observations, fine (0 0 0 1)Mg plates and GP zones commonly observed in other Mg–Zn based alloys [1,2,13–15] did not represent a large fraction of the precipitates in the T6 peak-aged microstructure (two rods and a blocky precipitate are visible in Fig. 3c). The GP zones, which are a single magnesium atomic layer in thickness, and thin (0 0 0 1)Mg plates, which are several magnesium atomic layers thick, are among the precipitates that make the greatest contribution to strengthening in alloys such as Mg–Zn–Cu–Mn [1], Mg–Zn–Cr [15], Mg–Zn–Ba [13] or Mg–Zn–V [14]. These precipitates were however rarely observed in the T6 microstructure of Mg–Zn–Ti alloy. One plate-like precipitate is shown in [0 0 0 1]Mg HRTEM image in Fig. 4a. The major contribution to the hardening was made primarily by the [0 0 0 1]Mg rods similar to 1 and to a smaller extent also by the [0 0 0 1]Mg laths similar to 2 . The [0 0 0 1]Mg HRTEM images of the major strengthening precipitates in the T6 condition are shown in Fig. 4. The end-on [0 0 0 1]Mg rod similar to 1 is shown in Fig. 4a along with a (0 0 0 1)Mg plate, while the end-on [0 0 0 1]Mg lath similar to 2 is shown in Fig. 4d. The Fourier transform (FFT) of the HRTEM image corresponding to the [0 0 0 1]Mg rod (Fig. 4e) may be indexed according to the [0 0 1] zone axis of a base centred monoclinic lat tice [25] reported for the 1 phase in a binary alloy [4]. However, closer examination of the FFTs from several rods showed that the inter-planar angles, and possibly also the spot distances, do not match those reported for the 1 phase in a binary alloy in the [0 0 1] precipitate orientation. This indicates that the lattice parameters of the rods in Mg–Zn–Ti alloy are slightly different to those of the equivalent precipitates in the binary alloy. Similar observations were made on rod-like precipitates in Mg–Zn–Ba [13] and Mg–Zn–V [14] alloys and it is suspected that the subtle differences in the lattice parameters may be due to the presence of atoms of the ternary alloying elements within the precipitates. The FFT from the (0 0 0 1)Mg plate (Fig. 4b) indicates that its crystal structure is very similar to that of the 2 phase having the orientation relationship with the magnesium matrix such that [0 0 0 1]p ||[0 0 0 1]Mg and (0 1¯ 0 1)p ||(1¯ 1¯ 2 0)Mg (corresponds to a plate-shaped morphology of the 2 phase), thus very similar also to the crystal structure of fine (0 0 0 1)Mg plates observed in a number of other Mg–Zn based alloys aged to T6 temper or at intermediate temperatures [1,2,13,15]. The corresponding model diffraction pattern showing the magnesium and precipitate reflections is given in Fig. 4c. The satellite reflections around each precipitate spot, as well as around the magnesium reflections and the central spot, arise from double diffraction within the precipitate. The [0 0 0 1]Mg laths were found to be very similar to those observed in the binary alloy based on the analysis of the corresponding FFT (Fig. 4f compared to a model diffraction pattern in Fig. 4g) thus very similar to the lath-shaped form of the 2 phase having the orientation relationship with the magnesium
Fig. 3. [2 1¯ 1¯ 0]Mg TEM image of Mg–Zn alloy aged to peak-hardness at 160 ◦ C (a); [2 1¯ 1¯ 0]Mg TEM (b) and HRTEM (c) images of Mg–Zn–Ti alloy also aged to peakhardness at 160 ◦ C. Insets in (a) and (b) are corresponding diffraction patterns (DP).
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Fig. 5. [2 1¯ 1¯ 0]Mg TEM (a) and HRTEM (b) images of the microstructure aged 4 weeks at 70 ◦ C. Inset in (a) is corresponding FFT.
matrix such that [2 1¯ 1¯ 0]p ||[0 0 0 1]Mg and (0 0 0 1)p ||(1¯ 2 1¯ 0)Mg . The exact matching of the lattice parameters could not be confirmed due to inadequate quality of the FFT. These observations indicate that the precipitates in the T6 condition of Mg–Zn–Ti alloy are generally very similar to those forming in the binary alloy and other Mg–Zn based alloys, although the lattice parameters are likely slightly different. In contrast to the binary and other Mg–Zn based alloys, the presence of Ti does not seem to be favouring the formation of GP zones and fine (0 0 0 1)Mg plates on basal planes in alloy aged at 160 ◦ C. 3.4. Precipitates formed by ageing at 70 ◦ C Fig. 4. [0 0 0 1]Mg HRTEM images (a and d) of the strengthening precipitates in Mg–Zn–Ti alloy aged to peak-hardness at 160 ◦ C; FFTs from the precipitates (b, e and f) with corresponding model DPs (c is a model DP for FFT in b; g is a model DP for FFT in f; large circles correspond to Mg reflections and small circles to precipitate reflections).
The [2 1¯ 1¯ 0]Mg TEM and HRTEM images of the microstructure produced by ageing at 70 ◦ C are shown in Fig. 5a and b, respectively. Based on these images, a considerable fraction of the precipitates present in the microstructure were fine (0 0 0 1)Mg plates which gave rise to characteristic dark strain field contrast (Fig. 5a) typically present either side of a plate with a line of no contrast (bright line) in the middle. The bright line parallel to basal plane indicates that a plate-like precipitate is coherent with the magnesium lattice
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alent of the 2 phase. The precipitates visible in HRTEM image in Fig. 6b are believed to most likely be equivalent to the [0 0 0 1]Mg rods, [0 0 0 1]Mg laths and prismatic precipitates observed earlier in the binary alloy aged at the same temperature [2]. Compared to the binary alloy [2], a considerably higher fraction of thin plates on {2 1¯ 1¯ 0}Mg planes were also observed (arrowed in Fig. 6b). These precipitates correspond well to GP1 zones, observed also in a number of Mg–Zn based alloys aged at ambient and intermediate temperatures. Although the types of the precipitates observed in Mg–Zn–Ti alloy aged at 70 ◦ C are generally similar to those observed in the binary alloy [2] and in Mg–Zn–Cu–Mn alloy [1], their relative fractions are noticeably different. The Ti-containing alloy exhibited a considerably higher fraction of the fine (0 0 0 1)Mg plates, while the GP zones were not present in any substantial amount. The fraction of GP1 zones on {2 1¯ 1¯ 0}Mg planes was also noticeably higher in the Ti-containing alloy. Based on the hardness measurements, the hardness levels and possibly also the tensile properties of Mg–Zn based alloys containing different ternary alloying additions are most likely very similar. However, it is interesting to note that even a trace amount of these ternary additions affect the type of the precipitates that from in the microstructure as well as the temperature at which these precipitates form. In this regard, Ti seems to be favouring the formation of GP1 zones and fine (0 0 0 1)Mg plates (but not GP zones) at intermediate temperatures, such as 70 ◦ C, however (0 0 0 1)Mg plates were not observed in the T6 condition. This suggest that, in comparison to the binary alloy, Ti may be lowering the solvus line of the (0 0 0 1)Mg plates and reducing their thermal stability, as these precipitates were not readily observed in alloy aged at 160 ◦ C. On the other hand, Ti promotes the nucleation of the (0 0 0 1)Mg plates, which can be deduced from their increased fraction in alloy aged at 70 ◦ C.
4. Summary Fig. 6. [0 0 0 1]Mg TEM (a) and HRTEM (b) images of the microstructure aged 4 weeks at 70 ◦ C. The orientation of the images is indicated by the inset FFT in (a).
along the [0 0 0 1]Mg direction. Some of these plates are indicated by black arrows in Fig. 5b (one is about 2 nm wide). Some distortion of the magnesium lattice in the vicinity of these precipitates is also visible. Elongated precipitates aligned with the [0 0 0 1]Mg direction were also observed from the [2 1¯ 1¯ 0]Mg zone axis (Fig. 5a) although these precipitates that produced almost no elastic strains in the surrounding lattice were mainly obscured by the strain field contrast around the plate-like precipitates. Some clearly visible elongated precipitates are indicated by white arrows in Fig. 5a and b. In addition, precipitates of possibly prismatic morphology were also observed by HRTEM (indicated by star-like symbols in Fig. 5b). The [0 0 0 1]Mg TEM and HRTEM images in Fig. 6 show far more clearly the different types of the precipitate species present in the microstructure and also give a better indication of their relative fractions. Fig. 6a shows that in addition to the fine plates on basal planes, visible as honeycomb-looking moire´ fringes form this direction (some are circled), a large number of other precipitates viewed end-on were also present in the microstructure. No usable diffraction information could be obtained from most of these precipitates. The FFT from the precipitate arrowed in Fig. 6a indicates that its crystal structure and the orientation relationship with the magnesium lattice are very similar to that of the [0 0 0 1]Mg laths observed in the T6 condition and suspected to be the equiv-
• Ti is soluble in the solid magnesium in the presence of Zn in the amount that is at least an order of magnitude higher than the solubility of Ti in the pure solid magnesium. • The constituent particles in Mg–Zn–Ti alloy are compositionally and morphologically similar to the eutectic Mg7 Zn3 phase and may contain up to about 0.1 at.% Ti. • The highest level of hardening in Mg–Zn–Ti alloy can be achieved by ageing at 70 ◦ C for less than 3 weeks, which is generally consistent with the age hardening response of the binary and other Mg–Zn based alloys. • The microstructure produced by ageing at 70 ◦ C consisted of at least five different types of precipitates: [0 0 0 1]Mg rods, [0 0 0 1]Mg laths, [0 0 0 1]Mg prismatic precipitates, (0 0 0 1)Mg plates and GP1 zones. Similar precipitates were observed earlier also in the binary alloy, although the relative fractions of these precipitate species in Mg–Zn–Ti alloy were different as a result of the presence of Ti. The results suggest that Ti promotes the formation of the (0 0 0 1)Mg plates and GP1 zones at intermediate temperatures. • The microstructure of Mg–Zn–Ti alloy aged to T6 temper is strengthened primarily by the [0 0 0 1]Mg rods (equivalent to 1 in the binary alloy) and to a smaller extent by the [0 0 0 1]Mg laths (equivalent to 2 ), as well as blocky precipitates (possibly also 1 ). The crystal structures of these precipitates are most likely identical to those of the equivalent precipitates forming in the binary alloy although the lattice parameters may be slightly different, presumably as a result of the presence of Ti atoms within the precipitates.
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Acknowledgment The author wishes to thank the Japan Society for the Promotion of Science (JSPS) for providing financial support for this work in the form of JSPS Postdoctoral Fellowship. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
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