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Microstructure and mechanical properties of a newly developed high strength Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy
Intermetallics 45 (2014) 84e88
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Microstructure and mechanical properties of a newly developed high strength Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy P.F. Gostin a, *, H. Wendrock a, I. Schneider b, M. Bleckmann b, M. Stoica a, U. Kühn a, J. Eckert a, c a b c
IFW Dresden, Institute for Complex Materials, P.O. Box 27 01 16, D-01171 Dresden, Germany Bundeswehr Research Institute for Materials, Fuels and Lubricants, D-85368 Erding, Germany Institute of Materials Science, TU Dresden, D-01062 Dresden, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 August 2013 Received in revised form 14 October 2013 Accepted 17 October 2013 Available online 2 November 2013
Starting from the bulk glass-forming composition Mg59.5Cu22.9Ag6.6Gd11, the development of a bulk glassy matrix composite alloy with high strength and ductility was attempted by adding a large amount of Sc to induce the formation of the MgSc intermetallic compound. However, copper-mould casting of the modified composition, Mg54.7Cu11.5Ag3.3Gd5.5Sc25, as rods with 3 and 5 mm diameter, yields besides the expected MgSc, a second (CuSc) and a third (Mg3Gd-type) intermetallic. These compounds are surrounded by a Mg-rich matrix, possibly composed of glassy regions and a Mg-based terminal solid solution. Although the obtained microstructure deviates largely from the expected one, the new alloy exhibits very high strength, i.e. 790 MPa, and visible plastic deformation, i.e. 0.9%. Shear dimples were found in MgSc. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: B. Alloy design B. Glasses, metallic B. Mechanical properties at ambient temperature D. Microstructure
1. Introduction Mg-based bulk metallic glasses (BMG) have attracted a lot of attention in recent time due to their high strength and low density, which makes them potential candidates for light-weight applications. They have strength values of typically 750e950 MPa [1e7], which are higher than those of commercial Mg alloys by a factor of 2e4 [8]. However, a major drawback for their use in structural applications is their very low ductility at room temperature. Also compared to other glassy alloy systems they are most brittle, fracturing basically without any plastic deformation [9]. A general method to circumvent the limited macroscopic ductility of monolithic metallic glasses is to realize so-called composite microstructures containing crystalline phases in a glassy matrix; these were indeed demonstrated to show much better deformation behaviour than monolithic glasses [10,11]. This principle was also applied to Mg-based alloys and several such bulk metallic composites with high strength and ductility were
* Corresponding author. Tel.: þ49 351 4659 767; fax: þ49 351 4659 541. E-mail addresses: [email protected], fl[email protected] (P. F. Gostin). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.10.015
developed. For example, Hui et al. reported on the preparation of an in-situ composite with precipitated Mg flakes (Mg solid solution) in a glassy matrix by casting, achieving a plastic strain of 18%. However, the microstructure of these alloys depends strongly on the cooling rate and therefore, on the sample thickness. To overcome this limitation, in the present paper a different strategy is proposed. Starting with a known good glass-forming composition, an element is added that forms with the main element, Mg, a binary intermetallic with high melting temperature, higher than the glass transition temperature of the starting composition. The composition is designed so that after the intermetallic is formed in the melt, the remaining melt has the starting glass-forming composition and therefore vitrifies during cooling forming a glassy matrix around the already formed intermetallic crystals. In this paper, this strategy is applied, using Mg59.5Cu22.9Ag6.6 Gd11 as a starting alloy, which has a critical diameter for glass formation of 27 mm [12] and Sc as additional element. Sc forms with Mg the MgSc intermetallic with a “melting” temperature of 1370e 1520 K, much higher than the liquidus temperature of the glassy alloy, i.e. w730 K [12]. Additionally, the MgSc intermetallic has a cubic crystalline structure, i.e. Pm-3m a ¼ 3.5921 A [13], with high symmetry and low lattice constant, indicating a high ability of plastic deformation. The desired phase composition (in apparent
P.F. Gostin et al. / Intermetallics 45 (2014) 84e88
atomic concentration) is 50 at.% MgSc and 50 at.% Mg59.5Cu22.9Ag6.6Gd11, which results after calculation in the final composition: Mg54.7Cu11.5Ag3.3Gd5.5Sc25. Both the starting bulk glassy Mg59.5Cu22.9Ag6.6Gd11 and the newly designed Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys were prepared by copper mould injection casting and characterized regarding their microstructure and mechanical properties under compression at room temperature. Although the Sc-containing alloy does not form the expected microstructure, it exhibits significant plastic deformation in addition to maintaining high fracture strength as the initial BMG. 2. Experimental Ingots (50 g in weight) of nominal composition Mg59.5Cu22.9 Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (at.%) were synthesized by the cold crucible method in a high purity argon atmosphere. To ensure chemical homogeneity, the alloy ingots were re-melted two times. From these ingots, rods with a diameter of 3 and 5 mm and a length of 70 mm were prepared by injection casting into a Cu mould. The chemical composition of both the ingot and of the cast rods of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy was determined by inductively coupled plasma optical emission spectrometry (ICPOES). The actual elemental concentration values agree well with
Fig. 1. XRD patterns of the as-cast (a) Mg59.5Cu22.9Ag6.6Gd11 Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy samples with a diameter of 5 mm.
and
(b)
85
the nominal ones, i.e. 0.65 at.%. For microstructure characterization, X-ray diffraction (XRD) with Co Ka radiation, and scanning electron microscopy (SEM) coupled with energy dispersive X-ray analysis (EDX) have been used. Further investigations were carried out by means of electron backscatter diffraction (EBSD) on a Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy sample with a diameter of 5 mm. This sample was polished by sputtering with Arþ ions with an acceleration voltage of 6 kV for 8 h. For the EBSD analysis, a ZEISS Ultra Plus SEM coupled with an EBSD system CHANNEL5 with a Nordlys F detector (Oxford Instruments) was used. Primary electrons were accelerated with a 10 kV voltage and the working distance was 8 mm, while the step size was set to 0.3 mm. Room temperature compression tests were performed on rod-shaped samples 3 and 5 mm in diameter, and 6 and 10 mm in length with an initial strain rate of 1.7 104 s1 using an Instron 5869 device. Subsequent to the compression tests, the fracture surface was analysed by SEM. 3. Results and discussion 3.1. Microstructure characterization Fig. 1 shows XRD patterns of the Mg59.5Cu22.9Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy samples. The pattern corresponding to the bulk glass-forming composition, Mg59.5Cu22.9Ag6.6Gd11, shows no sharp peaks indicating the glassy nature of the material. On the contrary, the pattern of the Sc-containing alloy exhibits a multitude of sharp peaks. As expected, high intensity Bragg peaks
Fig. 2. SEM images of an as-cast Mg54.7Cu11.5Ag3.3Gd5.5Sc25 rod sample (d ¼ 5 mm); (a) low magnification image showing the distribution of the various constituent phases, and (b) higher magnification image showing their size and morphology.
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Table 1 Atomic concentration of each alloy element in each constituent phase of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy produced by copper mould injection casting (rod-shaped sample with a diameter of 5 mm). Alloy components, at.% Mg Dendritic phase Around dendrites Polyhedral phase Matrix Nominal
60.7 62.5 7.6 73.2 54.7
Possible correspondence to structure from XRD
Sc
2.4 1.4 0.6 1.6
31.9 18.7 43.1 7.7 25
Cu
1.3 0.9 1.3 0.6
3.2 9.8 44.3 6.9 11.5
Ag
0.1 1.1 1.7 0.4
1.3 3.4 2.0 7.0 3.3
Gd
0.1 0.2 0.1 0.3
1.8 6.0 1.1 4.8 5.5
0.2 0.5 0.2 0.4
Mg65Sc35 Mg3Gd/Mg2AgGd CuSc (Mg)
Fig. 3. (a) SEM image on the cross-section of an Mg54.7Cu11.5Ag3.3Gd5.5Sc25 cast rod sample (d ¼ 5 mm). The lines mark the area investigated by EBSD shown in subfigure (b); (b) band contrast map of the area marked in subfigure (a).
corresponding to the MgSc-type structure (Pm-3m) are observed. MgSc is a non-stoichiometric intermetallic with an equilibrium Sc content varying from 35 to 55 at.% [14]. However, no diffuse background corresponding to an (expected) amorphous phase is detectable in the XRD pattern of rod samples of either 3 or 5 mm diameter. Instead, additional crystalline phases can be indexed: a second primitive cubic structure (Pm-3m) corresponding most probably to the CuSc intermetallic, a face-centred cubic structure (Fm-3m) assignable to Mg3Gd or Mg2AgGd, and possibly a hexagonal structure (P63/mmc) corresponding to a Mg- or aSc-based solid solution (Mg and Sc have close lattice parameters). Additionally, there are several other low intensity peaks (indicated in Fig. 1 by question marks), which do not correspond to any of the above mentioned phases. These peaks probably originate from superficial oxidation or oxide inclusions. The XRD pattern of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy prepared at a higher cooling rate in the shape of a rod with a diameter of 3 mm is quite similar to the one of the rod with 5 mm diameter (not shown here). As visible in the SEM image shown in Fig. 2(a), the alloy Mg54.7Cu11.5Ag3.3Gd5.5Sc25 consists of a homogeneous multiphase microstructure. In rods with diameter of 5 mm, the size of the grains is between 1 and 10 mm and somewhat smaller at the outer rim where the cooling rate was higher. In rods with diameter of 3 mm the microstructure is essentially similar with the difference that the grain size is between about 0.2 and 3 mm. The higher magnification image in Fig. 2(b) shows a more detailed view of the constituent phases. Four clearly distinguishable phases are observed. First, there is a dark phase with a dendritic shape. Second, there is a phase having the brightest contrast that closely surrounds the dendrites. Third, there is a preponderantly polyhedral-shaped phase and finally a fourth phase appearing as matrix. In order to check their compositions, EDX point analysis of each of the four phases was performed. The average measured elemental atomic concentration values of each phase are shown in Table 1. Obviously, the dendrites are the expected MgSc intermetallic also clearly detected by XRD, while the polyhedral phase is the CuSc
intermetallic. Both of them contain other elements as well, but the concentration ratio of the main elements does not deviate much from the (binary) equilibrium one, i.e. 35e55 at.% Sc in MgSc [14] and 50 at.% Sc in CuSc [15]. The light phase around the dendrites is as rich in Mg as the dendrites themselves, but contains a significantly lower Sc amount and more Cu, Ag and Gd. The matrix has the highest concentration of Mg and is, therefore, assumed to be a solid solution of Mg. This correlates with the XRD results, which revealed the presence of a hexagonal phase with P63/mmc structure. It results then, that the Mg3Gd-type structure should be associated with the phase around the dendrites. To reveal further microstructure details, EBSD analyses were carried out on the cross-section of a rod with diameter of 5 mm (rods with diameter of 3 mm have grain sizes too small to yield meaningful EBSD results). The lines in Fig. 3(a) mark the boundaries of the area which was investigated by EBSD. Fig. 3(b) represents a
Fig. 4. True stress-true strain curves under uniaxial quasistatic compressive loading of bulk glassy Mg59.5Cu22.9Ag6.6Gd11 (3 mm rod) and of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy (3 and 5 mm rods).
P.F. Gostin et al. / Intermetallics 45 (2014) 84e88
87
Table 2 Mechanical properties under compression of the bulk glassy Mg59.5Cu22.9Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys.
Mg59.5Cu22.9Ag6.6Gd11 (Ø3) Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (Ø3) Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (Ø5)
E [GPa]
sy [MPa]
smax [MPa]
εoff.
59 2 65 3 75 2
e 728 23 690 11
958 27 813 45 790 48
e 1.3 0.0 1.1 0.0
band contrast map of that area. Black regions correspond to the lowest band contrast, i.e. a massively distorted or amorphous microstructure. A proper discrimination of the cubic crystal phases was not possible, since the relative differences in the lattice constants of these phases are less than 1%. But the band contrast, which corresponds to the local amount of crystallographic information, reveals that a low but significant fraction of the material is massively distorted or amorphous. It is clear that those possibly amorphous regions are found in the matrix, but it appears that they only partially cover the matrix area. Another constituent phase of the matrix might be, as discussed earlier, a terminal solid solution of Mg. Considering the complexity of the microstructure, only a rough model of the phase formation sequence can be constructed. Firstly, it is assumed that all the constituent elements of this alloy, i.e. Mg54.7Cu11.5Ag3.3Gd5.5Sc25, are miscible in the liquid state, which means that initially during the casting process there is a homogeneous melt. Upon cooling, the first solid phases to crystallize should be the CuSc-based intermetallic and the b-Sc solid solution, since those have the highest melting point among all observed phases, i.e. 1398 K for the binary CuSc [15] and 1323 K for the liquidus temperature of binary Mg-Sc at 33 at.% Sc [14]. With further cooling, more b-Sc solid solution is formed, while at the same time the remaining liquid is enriched in Mg until a peritectic reaction starts to take place at the interface between the b-Sc dendrites and the neighbouring liquid, as inferable, from the binary Mg-Sc phase diagram. The occurrence of a peritectic reaction is also indicated by the morphology of the phase surrounding the black dendrites. By now, the liquid is severely depleted of Sc and also significantly enriched in Gd and Ag, being therefore closer to the known wide glass-forming composition region of the MgeCueAgeGd system [12]. As a consequence, amorphous regions form in the matrix, as detected by EBSD. 3.2. Mechanical properties Fig. 4 shows the stressestrain curves determined under quasistatic compression conditions for the bulk glassy
0.2
[%]
εf [%]
εpl [%]
1.6 0.1 1.9 0.3 1.9 0.5
0 0.7 0.2 0.9 0.4
Mg59.5Cu22.9Ag6.6Gd11 and the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys in form of rods with 3 and 5 mm diameter. The values of significant mechanical characteristics are summarized in Table 2. The bulk glassy alloy exhibits a very high strength of 958 MPa, but it fractures before reaching the elastic limit without any visible plastic deformation. Multiple pores were observed on the fracture surface of the resulting fragments indicating that the samples failed in an axial splitting mode typical for unconfined brittle materials [16]. The Sccontaining alloy also has a high strength, which is only about 15% lower. However, it shows significant plastic deformation and workhardening before fracture, as evident in the true stress-true strain curves (see Fig. 4). Slight differences in the mechanical parameters between the two different rods with diameter 3 and 5 mm are caused by the slightly different microstructure, which is, in turn, caused by the different cooling rate. To understand the fracture mechanism of the Mg54.7Cu11.5Ag3.3 Gd5.5Sc25 alloy, samples fractured in compression under quasistatic compression at room temperature were macro- and microscopically investigated. Upon reaching the maximum stress under compression conditions the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy fractures into several fragments. The inset in Fig. 5 shows one of the main resulting fragments. The fracture surface is typically either parallel to the compression direction or inclined at an angle of 45 5 . As Fig. 5(a) shows, the former type of fracture surface has a brittle appearance with clear river lines visible at the grain size level. This fracture surface parallel to the loading direction forms by axial splitting due to the tensile stress generated at pores [16]. The inclined surface displayed in Fig. 5(b), appears as a typical shear surface revealing shearing dimples and oriented parallel grooves. In order to relate these observations to the microstructure, imaging using both secondary electrons and back-scattered electrons in parallel was carried out. The shearing dimples appear to be located mostly at the dendrites. This indicates that, as expected, the MgSc dendrites are ductile. To the best of our knowledge, there is no data available regarding the mechanical properties of the binary MgSc intermetallic, although it recently received increased attention as secondary phase in Mg-alloys [17,18] with high creep resistance.
Fig. 5. SEM images of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy after fracture under uniaxial quasistatic compressive loading. The inset shows a macrograph of the main resulting fragment; (a) close-up on fracture plane parallel to compression direction and (b) close-up on inclined shearing plane (dual imaging: left e using secondary electrons, and right e using back-scattered electrons).
88
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According to Gschneidner et al. [19], the binary CuSc intermetallic is, contrary to expectations, ductile. This was explained by the particular electronic structure in the AB B2 CsCl-type compounds of which CuSc is one. Microscopic observations of fracture surfaces revealed that the CuSc polyhedral phase in the Mg54.7Cu11.5Ag3.3 Gd5.5Sc25 alloy does not undergo plastic deformation. This discrepancy can be explained by the presence of other elements, i.e. Mg, Ag and Gd, in the structure of the CuSc intermetallic perturbing the electronic structure and impeding dislocation movement.
4. Conclusions As expected, the MgSc intermetallic formed upon cooling of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 melt in a copper mould with diameter of 3 or 5 mm. However, competing crystallization of the CuSc intermetallic overcame the expected formation of an amorphous matrix phase. The resulting microstructure is composed of the abovementioned intermetallics plus a third phase rich in Mg that assumes the crystalline structure of Mg3Gd and a Mg-rich matrix. The alloy exhibits a high compressive strength of about 790 MPa together with a plastic deformation of 0.9%. Fracture occurs in a mixed mode: parallel to the compression direction by axial splitting, and by shearing along planes inclined at 45 5 .
Acknowledgements S. Donath is gratefully acknowledged for alloy and sample preparation, A. Voß and R. Buckan for ICP-OES tests, R. Keller for metallography, H.-J. Klauß for mechanical testing, and S. Pauly for fruitful discussions. Financial support for this work was provided by Wehrwissenschaftliches Institut für Werk- und Betriebsstoffe (WIWeB) e contact person: W. Kreuzer.
References [1] Inoue A, Kato A, Zhang T, Kim SG, Masumoto T. MgeCueY amorphous alloys with high mechanical strengths produced by a metallic mold casting method. Mater Trans JIM 1991;32:609e16. [2] Ma H, Ma E, Xu J. A new Mg65Cu7.5Ni7.5Zn5Ag5Y10 bulk metallic glass with strong glass-forming ability. J Mater Res 2003;18:2288e91. [3] Men H, Kim DH. Fabrication of ternary Mg-Cu-Gd bulk metallic glass with high glass-forming ability under air atmosphere. J Mater Res 2003;18:1502e4. [4] Park ES, Lee JY, Kim DH. Effect of Ag addition on the improvement of glassforming ability and plasticity of Mg-Cu-Gd bulk metallic glass. J Mater Res 2005;20:2379e85. [5] Zheng Q, Cheng S, Strader JH, Ma E, Xu J. Critical size and strength of the best bulk metallic glass former in the MgeCueGd ternary system. Scr Mater 2007;56:161e4. [6] Ma H, Shi LL, Xu J, Li Y, Ma E. Discovering inch-diameter metallic glasses in three-dimensional composition space. Appl Phys Lett 2005;87:181915. [7] Qin W, Li J, Kou H, Gu X, Kecskes L, Chang H, et al. Effects of Zn addition on the improvement of glass forming ability and plasticity of MgeCueTb bulk metallic glasses. J Non-cryst Solids 2008;354:5368e71. [8] Housh S, Mikucki B, Stevenson A. Selection and application of magnesium and magnesium alloys. In: ASM Handbook. 10th ed. ASM International; 1990. p. 1424e589. [9] Xu Y-K, Ma H, Xu J, Ma E. Mg-based bulk metallic glass composites with plasticity and gigapascal strength. Acta Mater 2005;53:1857e66. [10] Hays CC, Kim CP, Johnson WL. Microstructure controlled shear band pattern formation and enhanced plasticity of bulk metallic glasses containing in situ formed ductile phase dendrite dispersions. Phys Rev Lett 2000;84:2901e4. [11] Eckert J, Das J, Pauly S, Duhamel C. Mechanical properties of bulk metallic glasses and composites. J Mater Res 2007;22:285e301. [12] Zheng Q, Xu J, Ma E. High glass-forming ability correlated with fragility of MgeCu(Ag)eGd alloys. J Appl Phys 2007;102:113519. [13] Parthe E, Hohnke D, Jeitschko W, Schob O. Structure data of new intermetallic compounds. Naturwissenschaften 1965;52:155. [14] Predel B. Mg-Sc (Magnesium-Scandium). In: Madelung O, editor. SpringerMaterials e the Landolt-Börnstein Database. Springer-Verlag; 1998. [15] Predel B. Cu-Sc (Copper-Scandium). In: Madelung O, editor. SpringerMaterials e the Landolt-Börnstein Database. Springer-Verlag; 1998. [16] Meyers M, Chawla K. Mechanical behavior of materials. 2nd ed. Cambridge University Press; 2009. [17] Peng Q, Ma N, Li X, Zhang J. Age hardening response of a Mge5wt.% Sc alloy under an applied stress field. Mater Lett 2012;78:58e61. [18] Mordike BL, Stulíková I, Smola B. Mechanisms of creep deformation in Mg-Scbased alloys. Metall Mater Trans A 2005;36:1729e36. [19] Gschneidner KA, Ji M, Wang CZ, Ho KM, Russell AM, Mudryk Y, et al. Influence of the electronic structure on the ductile behavior of B2 CsCl-type AB intermetallics. Acta Mater 2009;57:5876e81.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Microstructure and mechanical properties of a newly developed high strength Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy P.F. Gostin a, *, H. Wendrock a, I. Schneider b, M. Bleckmann b, M. Stoica a, U. Kühn a, J. Eckert a, c a b c
IFW Dresden, Institute for Complex Materials, P.O. Box 27 01 16, D-01171 Dresden, Germany Bundeswehr Research Institute for Materials, Fuels and Lubricants, D-85368 Erding, Germany Institute of Materials Science, TU Dresden, D-01062 Dresden, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 August 2013 Received in revised form 14 October 2013 Accepted 17 October 2013 Available online 2 November 2013
Starting from the bulk glass-forming composition Mg59.5Cu22.9Ag6.6Gd11, the development of a bulk glassy matrix composite alloy with high strength and ductility was attempted by adding a large amount of Sc to induce the formation of the MgSc intermetallic compound. However, copper-mould casting of the modified composition, Mg54.7Cu11.5Ag3.3Gd5.5Sc25, as rods with 3 and 5 mm diameter, yields besides the expected MgSc, a second (CuSc) and a third (Mg3Gd-type) intermetallic. These compounds are surrounded by a Mg-rich matrix, possibly composed of glassy regions and a Mg-based terminal solid solution. Although the obtained microstructure deviates largely from the expected one, the new alloy exhibits very high strength, i.e. 790 MPa, and visible plastic deformation, i.e. 0.9%. Shear dimples were found in MgSc. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: B. Alloy design B. Glasses, metallic B. Mechanical properties at ambient temperature D. Microstructure
1. Introduction Mg-based bulk metallic glasses (BMG) have attracted a lot of attention in recent time due to their high strength and low density, which makes them potential candidates for light-weight applications. They have strength values of typically 750e950 MPa [1e7], which are higher than those of commercial Mg alloys by a factor of 2e4 [8]. However, a major drawback for their use in structural applications is their very low ductility at room temperature. Also compared to other glassy alloy systems they are most brittle, fracturing basically without any plastic deformation [9]. A general method to circumvent the limited macroscopic ductility of monolithic metallic glasses is to realize so-called composite microstructures containing crystalline phases in a glassy matrix; these were indeed demonstrated to show much better deformation behaviour than monolithic glasses [10,11]. This principle was also applied to Mg-based alloys and several such bulk metallic composites with high strength and ductility were
* Corresponding author. Tel.: þ49 351 4659 767; fax: þ49 351 4659 541. E-mail addresses: [email protected], fl[email protected] (P. F. Gostin). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2013.10.015
developed. For example, Hui et al. reported on the preparation of an in-situ composite with precipitated Mg flakes (Mg solid solution) in a glassy matrix by casting, achieving a plastic strain of 18%. However, the microstructure of these alloys depends strongly on the cooling rate and therefore, on the sample thickness. To overcome this limitation, in the present paper a different strategy is proposed. Starting with a known good glass-forming composition, an element is added that forms with the main element, Mg, a binary intermetallic with high melting temperature, higher than the glass transition temperature of the starting composition. The composition is designed so that after the intermetallic is formed in the melt, the remaining melt has the starting glass-forming composition and therefore vitrifies during cooling forming a glassy matrix around the already formed intermetallic crystals. In this paper, this strategy is applied, using Mg59.5Cu22.9Ag6.6 Gd11 as a starting alloy, which has a critical diameter for glass formation of 27 mm [12] and Sc as additional element. Sc forms with Mg the MgSc intermetallic with a “melting” temperature of 1370e 1520 K, much higher than the liquidus temperature of the glassy alloy, i.e. w730 K [12]. Additionally, the MgSc intermetallic has a cubic crystalline structure, i.e. Pm-3m a ¼ 3.5921 A [13], with high symmetry and low lattice constant, indicating a high ability of plastic deformation. The desired phase composition (in apparent
P.F. Gostin et al. / Intermetallics 45 (2014) 84e88
atomic concentration) is 50 at.% MgSc and 50 at.% Mg59.5Cu22.9Ag6.6Gd11, which results after calculation in the final composition: Mg54.7Cu11.5Ag3.3Gd5.5Sc25. Both the starting bulk glassy Mg59.5Cu22.9Ag6.6Gd11 and the newly designed Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys were prepared by copper mould injection casting and characterized regarding their microstructure and mechanical properties under compression at room temperature. Although the Sc-containing alloy does not form the expected microstructure, it exhibits significant plastic deformation in addition to maintaining high fracture strength as the initial BMG. 2. Experimental Ingots (50 g in weight) of nominal composition Mg59.5Cu22.9 Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (at.%) were synthesized by the cold crucible method in a high purity argon atmosphere. To ensure chemical homogeneity, the alloy ingots were re-melted two times. From these ingots, rods with a diameter of 3 and 5 mm and a length of 70 mm were prepared by injection casting into a Cu mould. The chemical composition of both the ingot and of the cast rods of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy was determined by inductively coupled plasma optical emission spectrometry (ICPOES). The actual elemental concentration values agree well with
Fig. 1. XRD patterns of the as-cast (a) Mg59.5Cu22.9Ag6.6Gd11 Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy samples with a diameter of 5 mm.
and
(b)
85
the nominal ones, i.e. 0.65 at.%. For microstructure characterization, X-ray diffraction (XRD) with Co Ka radiation, and scanning electron microscopy (SEM) coupled with energy dispersive X-ray analysis (EDX) have been used. Further investigations were carried out by means of electron backscatter diffraction (EBSD) on a Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy sample with a diameter of 5 mm. This sample was polished by sputtering with Arþ ions with an acceleration voltage of 6 kV for 8 h. For the EBSD analysis, a ZEISS Ultra Plus SEM coupled with an EBSD system CHANNEL5 with a Nordlys F detector (Oxford Instruments) was used. Primary electrons were accelerated with a 10 kV voltage and the working distance was 8 mm, while the step size was set to 0.3 mm. Room temperature compression tests were performed on rod-shaped samples 3 and 5 mm in diameter, and 6 and 10 mm in length with an initial strain rate of 1.7 104 s1 using an Instron 5869 device. Subsequent to the compression tests, the fracture surface was analysed by SEM. 3. Results and discussion 3.1. Microstructure characterization Fig. 1 shows XRD patterns of the Mg59.5Cu22.9Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy samples. The pattern corresponding to the bulk glass-forming composition, Mg59.5Cu22.9Ag6.6Gd11, shows no sharp peaks indicating the glassy nature of the material. On the contrary, the pattern of the Sc-containing alloy exhibits a multitude of sharp peaks. As expected, high intensity Bragg peaks
Fig. 2. SEM images of an as-cast Mg54.7Cu11.5Ag3.3Gd5.5Sc25 rod sample (d ¼ 5 mm); (a) low magnification image showing the distribution of the various constituent phases, and (b) higher magnification image showing their size and morphology.
86
P.F. Gostin et al. / Intermetallics 45 (2014) 84e88
Table 1 Atomic concentration of each alloy element in each constituent phase of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy produced by copper mould injection casting (rod-shaped sample with a diameter of 5 mm). Alloy components, at.% Mg Dendritic phase Around dendrites Polyhedral phase Matrix Nominal
60.7 62.5 7.6 73.2 54.7
Possible correspondence to structure from XRD
Sc
2.4 1.4 0.6 1.6
31.9 18.7 43.1 7.7 25
Cu
1.3 0.9 1.3 0.6
3.2 9.8 44.3 6.9 11.5
Ag
0.1 1.1 1.7 0.4
1.3 3.4 2.0 7.0 3.3
Gd
0.1 0.2 0.1 0.3
1.8 6.0 1.1 4.8 5.5
0.2 0.5 0.2 0.4
Mg65Sc35 Mg3Gd/Mg2AgGd CuSc (Mg)
Fig. 3. (a) SEM image on the cross-section of an Mg54.7Cu11.5Ag3.3Gd5.5Sc25 cast rod sample (d ¼ 5 mm). The lines mark the area investigated by EBSD shown in subfigure (b); (b) band contrast map of the area marked in subfigure (a).
corresponding to the MgSc-type structure (Pm-3m) are observed. MgSc is a non-stoichiometric intermetallic with an equilibrium Sc content varying from 35 to 55 at.% [14]. However, no diffuse background corresponding to an (expected) amorphous phase is detectable in the XRD pattern of rod samples of either 3 or 5 mm diameter. Instead, additional crystalline phases can be indexed: a second primitive cubic structure (Pm-3m) corresponding most probably to the CuSc intermetallic, a face-centred cubic structure (Fm-3m) assignable to Mg3Gd or Mg2AgGd, and possibly a hexagonal structure (P63/mmc) corresponding to a Mg- or aSc-based solid solution (Mg and Sc have close lattice parameters). Additionally, there are several other low intensity peaks (indicated in Fig. 1 by question marks), which do not correspond to any of the above mentioned phases. These peaks probably originate from superficial oxidation or oxide inclusions. The XRD pattern of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy prepared at a higher cooling rate in the shape of a rod with a diameter of 3 mm is quite similar to the one of the rod with 5 mm diameter (not shown here). As visible in the SEM image shown in Fig. 2(a), the alloy Mg54.7Cu11.5Ag3.3Gd5.5Sc25 consists of a homogeneous multiphase microstructure. In rods with diameter of 5 mm, the size of the grains is between 1 and 10 mm and somewhat smaller at the outer rim where the cooling rate was higher. In rods with diameter of 3 mm the microstructure is essentially similar with the difference that the grain size is between about 0.2 and 3 mm. The higher magnification image in Fig. 2(b) shows a more detailed view of the constituent phases. Four clearly distinguishable phases are observed. First, there is a dark phase with a dendritic shape. Second, there is a phase having the brightest contrast that closely surrounds the dendrites. Third, there is a preponderantly polyhedral-shaped phase and finally a fourth phase appearing as matrix. In order to check their compositions, EDX point analysis of each of the four phases was performed. The average measured elemental atomic concentration values of each phase are shown in Table 1. Obviously, the dendrites are the expected MgSc intermetallic also clearly detected by XRD, while the polyhedral phase is the CuSc
intermetallic. Both of them contain other elements as well, but the concentration ratio of the main elements does not deviate much from the (binary) equilibrium one, i.e. 35e55 at.% Sc in MgSc [14] and 50 at.% Sc in CuSc [15]. The light phase around the dendrites is as rich in Mg as the dendrites themselves, but contains a significantly lower Sc amount and more Cu, Ag and Gd. The matrix has the highest concentration of Mg and is, therefore, assumed to be a solid solution of Mg. This correlates with the XRD results, which revealed the presence of a hexagonal phase with P63/mmc structure. It results then, that the Mg3Gd-type structure should be associated with the phase around the dendrites. To reveal further microstructure details, EBSD analyses were carried out on the cross-section of a rod with diameter of 5 mm (rods with diameter of 3 mm have grain sizes too small to yield meaningful EBSD results). The lines in Fig. 3(a) mark the boundaries of the area which was investigated by EBSD. Fig. 3(b) represents a
Fig. 4. True stress-true strain curves under uniaxial quasistatic compressive loading of bulk glassy Mg59.5Cu22.9Ag6.6Gd11 (3 mm rod) and of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy (3 and 5 mm rods).
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Table 2 Mechanical properties under compression of the bulk glassy Mg59.5Cu22.9Ag6.6Gd11 and Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys.
Mg59.5Cu22.9Ag6.6Gd11 (Ø3) Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (Ø3) Mg54.7Cu11.5Ag3.3Gd5.5Sc25 (Ø5)
E [GPa]
sy [MPa]
smax [MPa]
εoff.
59 2 65 3 75 2
e 728 23 690 11
958 27 813 45 790 48
e 1.3 0.0 1.1 0.0
band contrast map of that area. Black regions correspond to the lowest band contrast, i.e. a massively distorted or amorphous microstructure. A proper discrimination of the cubic crystal phases was not possible, since the relative differences in the lattice constants of these phases are less than 1%. But the band contrast, which corresponds to the local amount of crystallographic information, reveals that a low but significant fraction of the material is massively distorted or amorphous. It is clear that those possibly amorphous regions are found in the matrix, but it appears that they only partially cover the matrix area. Another constituent phase of the matrix might be, as discussed earlier, a terminal solid solution of Mg. Considering the complexity of the microstructure, only a rough model of the phase formation sequence can be constructed. Firstly, it is assumed that all the constituent elements of this alloy, i.e. Mg54.7Cu11.5Ag3.3Gd5.5Sc25, are miscible in the liquid state, which means that initially during the casting process there is a homogeneous melt. Upon cooling, the first solid phases to crystallize should be the CuSc-based intermetallic and the b-Sc solid solution, since those have the highest melting point among all observed phases, i.e. 1398 K for the binary CuSc [15] and 1323 K for the liquidus temperature of binary Mg-Sc at 33 at.% Sc [14]. With further cooling, more b-Sc solid solution is formed, while at the same time the remaining liquid is enriched in Mg until a peritectic reaction starts to take place at the interface between the b-Sc dendrites and the neighbouring liquid, as inferable, from the binary Mg-Sc phase diagram. The occurrence of a peritectic reaction is also indicated by the morphology of the phase surrounding the black dendrites. By now, the liquid is severely depleted of Sc and also significantly enriched in Gd and Ag, being therefore closer to the known wide glass-forming composition region of the MgeCueAgeGd system [12]. As a consequence, amorphous regions form in the matrix, as detected by EBSD. 3.2. Mechanical properties Fig. 4 shows the stressestrain curves determined under quasistatic compression conditions for the bulk glassy
0.2
[%]
εf [%]
εpl [%]
1.6 0.1 1.9 0.3 1.9 0.5
0 0.7 0.2 0.9 0.4
Mg59.5Cu22.9Ag6.6Gd11 and the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloys in form of rods with 3 and 5 mm diameter. The values of significant mechanical characteristics are summarized in Table 2. The bulk glassy alloy exhibits a very high strength of 958 MPa, but it fractures before reaching the elastic limit without any visible plastic deformation. Multiple pores were observed on the fracture surface of the resulting fragments indicating that the samples failed in an axial splitting mode typical for unconfined brittle materials [16]. The Sccontaining alloy also has a high strength, which is only about 15% lower. However, it shows significant plastic deformation and workhardening before fracture, as evident in the true stress-true strain curves (see Fig. 4). Slight differences in the mechanical parameters between the two different rods with diameter 3 and 5 mm are caused by the slightly different microstructure, which is, in turn, caused by the different cooling rate. To understand the fracture mechanism of the Mg54.7Cu11.5Ag3.3 Gd5.5Sc25 alloy, samples fractured in compression under quasistatic compression at room temperature were macro- and microscopically investigated. Upon reaching the maximum stress under compression conditions the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy fractures into several fragments. The inset in Fig. 5 shows one of the main resulting fragments. The fracture surface is typically either parallel to the compression direction or inclined at an angle of 45 5 . As Fig. 5(a) shows, the former type of fracture surface has a brittle appearance with clear river lines visible at the grain size level. This fracture surface parallel to the loading direction forms by axial splitting due to the tensile stress generated at pores [16]. The inclined surface displayed in Fig. 5(b), appears as a typical shear surface revealing shearing dimples and oriented parallel grooves. In order to relate these observations to the microstructure, imaging using both secondary electrons and back-scattered electrons in parallel was carried out. The shearing dimples appear to be located mostly at the dendrites. This indicates that, as expected, the MgSc dendrites are ductile. To the best of our knowledge, there is no data available regarding the mechanical properties of the binary MgSc intermetallic, although it recently received increased attention as secondary phase in Mg-alloys [17,18] with high creep resistance.
Fig. 5. SEM images of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 alloy after fracture under uniaxial quasistatic compressive loading. The inset shows a macrograph of the main resulting fragment; (a) close-up on fracture plane parallel to compression direction and (b) close-up on inclined shearing plane (dual imaging: left e using secondary electrons, and right e using back-scattered electrons).
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According to Gschneidner et al. [19], the binary CuSc intermetallic is, contrary to expectations, ductile. This was explained by the particular electronic structure in the AB B2 CsCl-type compounds of which CuSc is one. Microscopic observations of fracture surfaces revealed that the CuSc polyhedral phase in the Mg54.7Cu11.5Ag3.3 Gd5.5Sc25 alloy does not undergo plastic deformation. This discrepancy can be explained by the presence of other elements, i.e. Mg, Ag and Gd, in the structure of the CuSc intermetallic perturbing the electronic structure and impeding dislocation movement.
4. Conclusions As expected, the MgSc intermetallic formed upon cooling of the Mg54.7Cu11.5Ag3.3Gd5.5Sc25 melt in a copper mould with diameter of 3 or 5 mm. However, competing crystallization of the CuSc intermetallic overcame the expected formation of an amorphous matrix phase. The resulting microstructure is composed of the abovementioned intermetallics plus a third phase rich in Mg that assumes the crystalline structure of Mg3Gd and a Mg-rich matrix. The alloy exhibits a high compressive strength of about 790 MPa together with a plastic deformation of 0.9%. Fracture occurs in a mixed mode: parallel to the compression direction by axial splitting, and by shearing along planes inclined at 45 5 .
Acknowledgements S. Donath is gratefully acknowledged for alloy and sample preparation, A. Voß and R. Buckan for ICP-OES tests, R. Keller for metallography, H.-J. Klauß for mechanical testing, and S. Pauly for fruitful discussions. Financial support for this work was provided by Wehrwissenschaftliches Institut für Werk- und Betriebsstoffe (WIWeB) e contact person: W. Kreuzer.
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