structural evolution in mg-al alloys

Title of Publication Edited by TMS (The Minerals, Metals & Materials Society), Year

STRUCTURAL EVOLUTION IN MG-AL ALLOYS Shaul Avraham, Yael Maoz and Menachem Bamberger Technion – Israel Institute of Technology, Haifa 32000, Israel Keywords: alloy development, microstructure, Mg-Al, Mg-Sn Abstract The development of a new microstructuraly stable Mg-Al alloys was performed via the CALPHAD method. It was found that precipitation of Mg2Sn takes place as a result of the Sn presence in the alloy. The improved microhardness of the new alloy results from the combined effect of solution strengthening elements and stable phase at the grain boundaries. Al-Nd and Al-Mn-Nd intermetallics are expected to constrain grain boundary sliding during high temperature creep. The fine precipitation of Mg2Sn can hinder dislocation creep. The CALPHAD method was found to be consistent with experimental results. Introduction Mg alloys are highly attractive for engineering proposes due to their high specific mechanical properties [1, 2, 3]. The reduction of engineering parts weight can result in reduced energy consumption and beneficial environmental implications. Currently, a relative small number of traditional Mg alloys are available. The possible benefits resulting from the application of Mg alloys motivates scientific and engineering interest in the development of new creep resistant Mg alloy for elevated temperature (T>200°C) [4]. Mg-Al based alloys present excellent castability [4]. AE42 is a commercial Mg alloy; it presents high mechanical properties at temperatures lower than 150 ºC [5]. The development of a new alloy is governed by the possible alloy constituents, alloy composition, and processing parameters. Computational thermochemistry (CT) based on the CALPHAD method is applied in a successful manner and saves time and resources [6, 7]. This work is aimed at development of new microstructuraly stable MgAl alloys. The stability of the alloys is expected to result in high creep resistant at elevated temperatures. Experimental Two alloys were considered during this work: B+Nd and AE42 alloys. The Base alloy was determined as a AZ50 alloy modified with Sn [8, 9]. The B+Nd alloy is a derivative of the Base alloy with the addition of 1 wt.% Nd, the addition of Nd in Mg-Al alloys results with the formation of intermetallics with high melting temperature [10]. AE42 serves as a commercial reference alloy. The samples were produced by melting of 200 gr of the alloys under protective atmosphere (CO2+ Freon 134) and casting in a steel mould. The composition of the alloys was analyzed and confirmed by spark emission spectrometer (BAIRD, Model DV5). The prediction of the precipitation of the different phases during solidification was based upon CT simulations performed using the

Thermo-Calc software. The simulation results are based on thermodynamic considerations and a specially developed software database [7, 11]. Microstructural stability tests were composed of thermal exposure (TE) of as-cast samples and aging of solution treated samples. The solution treatment was conducted according to the CT simulation results. The solution treatment was designed as-follows: heating of the as-cast samples (120 °C/hour), a 20 hours dwell at 400°C, steady state heating (1 °C/hour) to 480°C, 120 hours dwell and water quench to room temperature. The TE and aging of the samples were conducted by the suspension of the samples in a hot (200°C) molten salt (sodium nitrate 50%+ potassium nitrate 50%) bath for different time periods (1, 2, 4, 8, 12, 16, 20, 24, 28 and 32 days). The samples designation is as follows: initially the alloy type is noted, As-cast or solutionized samples are denoted by additional a or s, respectively. The sample dwell time (days) is separated by a _ mark. Micro hardness measurements were conducted in order to estimate the thermal stability of the alloys during prolonged aging at 200°C. The samples for micro hardness measurements were sliced and cold mounted in an epoxy resin. The sample (102•5 [mm]) was polished by series of SiC polishing cloth and a final polishing step using an Al2O3 particle (0.05 µm, Buehler) suspension in water. The Vickers micro hardness measurements were conducted using a DMH-2 micro hardness tester (MATSUZAWA SEIKS Co. LTD Japan). The load level was set at 50 grf and the load time was 15 sec. The microstructure of the samples was investigated using X-ray diffraction (XRD) and scanning electron microscopy (SEM). XRD is used for phase identification. XRD measurements were conducted by a PW-3020, Philips X-ray automatic powder diffractometer. SEM was conducted in order to evaluate the microstructure of the different alloys, the different phases that are present and the influence of the thermal treatments on the microstructure evolution. The composition of the different phases was evaluated by energy dispersive spectroscopy (EDS). SEM was conducted using a FEI Quanta 200 equipped with an INCA 350 EDS system for elemental microanalysis of the sample. Results Figure 1 presents the results of the CT simulation for equilibrium and non-equilibrium phase evolution during cooling of the B+Nd alloy. During equilibrium solidification (Figure 1.a.) α-Mg starts to solidify during cooling of the melt at the liquidus temperature (620°C), Al8Mn5 starts to precipitate at 630°C. Solidification terminates at 490°C. The presence of Sn in the base alloy results with the formation of Mg2Sn, the addition of Nd results in the formation of Nd containing phases (Al2Nd and Mg41Nd5). The

precipitation of γ-Mg17Al12 from the super saturated α-Mg matrix starts above 350°C, the weight percent of γ-Mg17Al12 at 0°C is 8.7%. Figure 1.b presents the results of the CT simulation for nonequilibrium phase evolution during cooling of the B+Nd alloy. αMg solidified at 620°C. Al8Mn5 precipitate in a negligible amount at 634°C. Al2Nd and Mg2Sn are expected at 470°C and 440°C, respectively. The eutectic solidification takes place at 425°.

during TE of the as-cast sample, apparently there is an additional maxima (12-16 days) but its confirmation is still underway. The microhardness of B+Nd alloy appears to stabilize at ≈80 [Hv]. In the case of as-cast AE42 alloy (AE42a_0d) the initial micro hardness is ≈70 [Hv], two maxima (≈75 [Hv]) takes place after 12-16 days and 28 days, the microhardness of the AE42a alloy stabilize at ≈70 [Hv]. The initial micro hardness of the assolutionized B+Nds_0d alloy is ≈70 [Hv], aging of the sample do not results with a major variation, one maxima can be observed (16 days) during aging of B+Nds. The initial micro hardness of the as-solutionized AE42s_0d alloy is ≈65 [Hv], aging of the sample do not results with a major variation. The microhardness and stability of the as-cast B+Nda alloy is higher then that of the as-cast AE42a, the microhardness results of the aged samples appears to be equivalent.

Figure 1. (a) Equilibrium and (b) non-equilibrium phase evolution during cooling of the B+Nd alloy. Figure 2 presents the results of the CT simulation for equilibrium and non-equilibrium phase evolution during cooling of alloy AE42. During equilibrium solidification (Figure 2.a.) α-Mg starts to solidify during cooling of the melt at the liquidus temperature (621°C), Al8Mn5 starts to precipitate at 600°C. Solidification terminates at 515°C. The presence of RE elements results in the formation of RE containing phases (Al11Ce3, Al2Nd and Mg41Nd5). The precipitation of γ-Mg17Al12 from the super saturated α-Mg matrix starts above 320°C, the weight percent of γ-Mg17Al12 at 0°C is 6.8%. Figure 2.b presents the results of the CT simulation for nonequilibrium phase evolution during cooling of alloy AE42. α-Mg solidified at 620°C. Al8Mn5 starts to precipitate at 600°C but its amount is negligible. Al11Ce3, Al4CeMg4, and Al2Nd are expected at 515°C, 500°C, and 470°C, respectively. Eutectic solidification takes place at 425°C. The liquidus temperature according to the equilibrium and non-equilibrium CT simulations (620°C) is in good agreements with experimental results (625°C) [12].

Figure 2. (a) Equilibrium and (b) non-equilibrium phase evolution during cooling of the AE42 alloy. Figure 3 presents micro hardness results of B+Nd and the AE42 alloys during TE and aging (200 ºC, for 32 days). The initial micro hardness of the as-cast B+Nd alloy (B+Nda_0d) is ≈70 [Hv], two maxima (≈80 [Hv]) can be observed (1-2 & 28 days)

Figure 3. Micro hardness results of the B+Nd and AE42 during TE and aging (200 ºC, for 32 days). Figure 4 presents a SEM BSE micrograph of the B+Nd alloy in the (a) as-cast state (B+Nda_0d), (b) as-cast B+Nd after TE (200 ºC & 32 days, B+Nda_32d), (c) solution treated as-cast B+Nd alloy (B+Nds_0d) and (d) solutionized and aged base alloy (200 ºC & 32 days, B+Nds_32d). The microstructure of the B+Nda_0d sample (Figure 4.a) consists of α-Mg matrix, the matrix phase dissolved 1.5 at.% Al and 0.4 at.% Sn. The presence of a thin bright layer is apparent near the α-Mg grain boundary (GB), the layer is enriched with Al, Zn and Sn (4.5 at.%, 0.7 at.% and 1.3 at.%, respectively). The presence of Sn is preferred in the form of an alloying elements even thou some Mg2Sn was detected. The coring effect results from the rejection of solute elements from the solidifying layer to the interdendritic residual liquid [13]. Two types of isolated bright phases are apparent at the GB: elongated particles (various Al-Nd intermetallics) and irregularly shaped (Al-Mn-Nd intermetallics). EDS results confirmed the presence of small amounts of the γ-Mg17Al12 phase. SEM analysis of the B+Nda_32d sample (Figure 4.b) shows that solid state precipitation of Mg2Sn and small amounts of γMg17Al12 are present at the GB. The thin light enriched layer at the GB transforms to a strip of globular Mg2Sn at the GB and very fine Mg2Sn precipitates in the vicinity of the GB. The Sn content in the α-Mg matrix phase was reduced (0.2 at.%) in comparison to the as-cast sample but Zn (0.3 at.%) was added as an alloying element and the Al content increase (2.2 at.%). The two types of

precipitates at the GB (Al-Mn-Nd and Al-Nd intermetallics) are apparent with no changes. Solution treatment of the as-cast B+Nd alloy, B+Nds_0d , (Figure 4.c) results with homogenization of the sample, the α-Mg matrix phase was found to contain Al, Zn and Sn (2.8 at.%, 0.4 at.% and 0.5 at.%, respectively). The two type of isolated bright phase are apparent at the GB. Figure 4.d presents the microstructure of B+Nds_32d sample. The increased content of alloying elements in the matrix resulted with fine and homogenous precipitation of Mg2Sn. The two types of precipitates at the GB (Al-Mn-Nd and Al-Nd intermetallics) are apparent with no changes. A comparison of SEM results of the as-cast samples (B+Nda) and the solution

treated samples (B+Nds) shows that the distributing of the GB intermetallics (Al-Mn-Nd and Al-Nd) is lowered, this is a direct result of the grain growth that resulted from the solution treatment procedure. According to EDS results various types of Al-Nd intermetallics are present in the B+Nd alloy. SEM results (Figure 4.b, Figure 4.d) suggest that Al2Nd is more abundant in B+Nda_32d and B+Nds_32d than in B+Nda_0d and B+Nds_0d. According to CT results (Figure 1.a) the Al2Nd phase is the equilibrium phase in the B+Nd alloy. Apparently the Al2Nd is the equilibrium phase and the other Al-Nd intermetallics are intermediate phases.

Figure 4. SEM BSE micrograph of the B+Nd alloy in the (a) as-cast state (B+Nda_0d), (b) as-cast B+Nd after TE (200 ºC & 32 days, B+Nda_32d), (c) solution treated as-cast B+Nd alloy (B+Nds_0d) and (d) solutionized and aged base alloy (200 ºC & 32 days, B+Nds_32d). In the B+Nda sample, SEM analysis (B+Nda_0d & B+Nda_4d) showed that the first maxima (1-2 days) is related to the formation of very fine Mg2Sn precipitates in the thin bright layer near the α-

Mg GB. The second maxima in the B+Nda sample (after 28 days) is related to further growth of the Mg2Sn precipitates. The confirmation of the maxima after 12-16 days requires further

transmission electron microscopy. SEM analysis shows that the maxima in the B+Nds sample (after 16 days) is related to the growth process of the fine Mg2Sn precipitates. SEM analysis of the AE42 alloys shows that the microstructure consists of α-Mg matrix and fine lamellar Al-RE eutectic. Solution treatment of the AE42 alloy did not resulted in significant microstructural changes. High magnification SEM analysis showed that TE and aging at 200 ºC for 32 days resulted with the formation of the γ-Mg17Al12 phase between the fine lamellas structure. Table I presents a summary of the constitutional phases detected by XRD after the different thermal treatments in the B+Nd alloy and AE42 alloy samples (α-Mg included), Mg2Sn is absent at the B+Nda_0d sample. It was confirmed by SEM analysis that Sn is solutionized in the α-Mg matrix. TE and aging of the B+Nd alloy at 200 ºC for 32 days results in precipitation of Mg2Sn and the formation of Al78Mn22. The Al78Mn22 phase appears to be the AlMn-Nd phase that was observed via SEM analysis, this phase apparently contains significant amounts of Nd. AE42a_0d contains Al4RE and Al11RE3, the formation of the γ-Mg17Al12 phase is apparent only after TE at 200°C for 32 days (AE42a_32d). The AE42s_0d sample contains Al4RE and Al11RE3, in addition to the two former phases the γ-Mg17Al12 phase was detected in AE42s_32d. The AE42 CT, SEM and XRD results are consistent with the findings made by Powell et. Al. [5]. In both alloys the CT based solution treatment temperature was determined as 480 ºC (see Figure 2) to eliminate the γ-Mg17Al12 phase. Table I. The constitutional phases detected by XRD in the B+Nd and AE42 alloy samples (α-Mg included). As-cast & solutionized TE at & aged at Alloy As-cast Solutionized 200°C for 200°C for 32 days 32 days Mg2Sn, Mg2Sn, Mg2Sn, γB+Nd Al78Mn22 Al78Mn22 Al78Mn22 Mg17Al12 Al4RE, Al4RE, Al11RE3, Al4RE, Al4RE, AE42 Al11RE3, γAl11RE3 Al11RE3 γMg17Al12 Mg17Al12 Discussion The development of a new alloy is a time and resource consuming procedure. The CALPHAD method computes the equilibrium state of a multi component alloy by minimization of the total Gibbs energy at a constant temperature, pressure, and composition. CT simulations of the AE42 alloy successfully predicts the identity and relative amounts of the phases (α-Mg, Al11Ce3, γ-Mg17Al12 and Al-Mn) that forms during solidification. CT simulations are a useful tool in the estimation of new alloys (B+Nd). The addition of the alloying elements (Sn and Nd) resulted with the predicted formation of Al-Nd and Mg2Sn. Zn serves as a solution strengthening element in AZ alloys and It was found that Zn addition to Mg– 2.2 at.% Sn alloy increases the age hardening response, Zn refines and homogenize the precipitation of Mg2Sn [9, 14].

An aging procedure requires a single phase region for the solution of all alloying elements. The equilibrium CT results (Figure 1.a & Figure 2.a) reveals that at the appropriate temperature range for solution treatment ranges between of 460-490°C. In this temperature range the alloys contains the smallest amount of intermetallics, and thus is appropriate for solution treatment. The eutectic temperature of the alloys is estimated at 425°C. A high rate heating of the samples to a temperature in the solution treatment range can result with the melting of the eutectic phase, this can be avoided by a slow heating rate (1 °C/hour) to a temperature that is higher than the eutectic temperature, this procedure allows sufficient time for elemental diffusion. The information derived from the CT results allows the formulation of a compositional related solution treatment for the new alloys. The improved microhardness of the B+Nd in reference to AE42 results from the combined solution strengthening effect of Al, Zn and Sn that are present in increased amount in the B + Nd alloy. SEM results show that the Sn and Zn are solutionized in the vicinity of the α-Mg GB (B+Nda_0d), Sn is present at the core of the α-Mg phase. In accordance to the CT simulations SEM results shows that the Sn presences in B+Nd alloy results with the formation of Mg2Sn precipitates. Al-Nd and Al-Mn-Nd intermetallics were found at the GB. The solution treatment of the B+Nd samples resulted with grain growth and reduced GB intermetallics distribution, this can explain the reduced microhardness of the B+Nds samples in comparison to the B+Nda samples. The stability of microhardness results is attributed to the morphological stability of the microstructure, the Mg2Sn precipitation process is the cause for the small variations in microhardness. Further TEM analysis is needed for the analysis of the precipitation process. XRD results confirmed the expected formation of various phases that were predicted according to the CT simulations. The fact that Al-Nd intermetallics were not detected by XRD in the B+Nd alloy results from the small amount of Nd that was used (1 wt%) in the alloy. Improved elevated temperature creep resistance in magnesium alloys can be achieved by the suppression of γ-Mg17All2 by appropriate additives, grain boundary pining via secondary phase formation at the GB, and dispersion of fine, stable precipitates or solute elements that hinder the dislocations motion [4]. The B+Nd alloy appear to be stable during prolonged exposure at 200 °C, the micro hardness appeared to be stable during the aging and TE period. Thermodynamic simulations, XRD results and SEM analysis showed that the formation of γ-Mg17Al12 was negligible. The presence of Al-Nd, Mg2Sn and Al-Mn-Nd intermetallics at the GB can serve as a source for grain boundary pinning. The distribution of fine Mg2Sn precipitates at the α-Mg matrix can serve as obstacles for dislocation motion [9]. The HCP crystallographic structure of Magnesium results in reduced deformability. Room temperature deformation takes place via basal slip. In AZ91 it was found that dislocation exist on nonbasel planes in the temperature range of 120-180 °C [15]. At higher temperature (T>250°C) slip takes place in Mg on the prismatic planes and on the pyramidal planes [16, 17]. The density of plate and rod-like precipitates on non-basal planes increases with Zn addition in Mg-Sn alloys [9]. Rod precipitates on the

pyramidal planes and prismatic planes have increased influence on strengthening than precipitation on the basal plane and can serve as obstacle for dislocation creep [17, 18]. Summary The development of a new microstructuraly stable Mg-Al alloys was performed via the CALPHAD method, The simulations were used in refinement of the solution treatment. Fine precipitation of Mg2Sn was detected as a result of the Sn presence in the alloy. The improved microhardness of the new alloy results from the combined effect of solution strengthening elements and stable phases at the grain boundaries. The presence of Al-Nd and AlMn-Nd intermetallics is expected to constrain grain boundary sliding during creep. The fine precipitation of Mg2Sn is expected to hinder dislocation creep. The CT simulation was found to be consistent with experimental results. Acknowledgments The authors wish to thank Dr. Larry Kaufman, Brooklin, MA, USA. This study was supported by the EC under contract number FP6-503826 and partially by the fund for the promotion of research at the Technion. References 1 M.O. Pekguleryuz and A.A. Kaya, Creep Resistant Magnesium Alloys for Powertrain Applications, Advanced Engineering Materials, 5[12]: 866-878, 2003. 2 A. Wendt, K. Weiss, A. Ben-Dov, M. Bamberger and B. Bronfin. Magnesium Castings in Aeronautics Applications Special Requirements. in Magnesium Technology. 2005. 3 C.J. Smithells, E.A. Brandes and G.B. Brook, Smithells Light Metals Handbook. 1998, Oxford [England] ; Boston: Butterworth-Heinemann. vi, 194 p. 4 A.A. Luo, Recent Magnesium Alloy Development for Elevated Temperature Applications, International Materials Reviews, 49[1]: 13-30, 2004. 5 B.R. Powell, V. Rezhets, M.P. Balogh and R.A. Waldo, Microstructure and Creep Behavior in Ae42 Magnesium DieCasting Alloy, JOM, 54[8]: 34-38, 2002. 6 R. Schmid Fetzer and J. Grobner, Focused Development of Magnesium Alloys Using the Calphad Approach, Advanced Engineering Materials, 3[12]: 947-61, 2001. 7 M. Bamberger, Phase Formation in Mg-Sn-Zn Alloys Thermodynamic Calculations Vs Experimental Verification, Journal of Materials Science, 41[10]: 2821-2829, 2006. 8 S. Avraham, Y. Maoz and and M. Bamberger, submitted to Calphad 9 T.T. Sasaki, K. Oh-ishi, T. Ohkubo and K. Hono, Enhanced Age Hardening Response by the Addition of Zn in MgSn Alloys, Scripta Materialia, 55[3]: 251-254, 2006. 10 American Society for Metals, ASM Handbook Online. 2005, American Society for Metals: Ohio. 11 N. Saunders and A.P. Miodownik, Calphad (Calculation of Phase Diagrams): A Comprehensive Guide. Pergamon Materials Series ; V. 1. 1998, Oxford ; New York: Pergamon. xvi, 479 p. 12 International Magnesium association, http://www.intlmag.org/phys07.aspx

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