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Significance of stable and metastable phases in high temperature creep resistant magnesium–rare earth base alloys
Journal of Alloys and Compounds 378 (2004) 196–201
Significance of stable and metastable phases in high temperature creep resistant magnesium–rare earth base alloys B. Smola a,b,∗ , I. Stul´ıková a,b , J. Pelcová a , B.L. Mordike b,c a
c
Faculty of Mathematics and Physics, Charles University, Ke Karlovu 5, CZ-121 16 Prague 2, Czech Republic b Zentrum für Funktionswerkstoffe gGmbH, Sachsenweg 8, D-38678 Clausthal-Zellerfeld, Germany Institut für Werkstoffkunde und Werkstofftechnik, TU Clausthal, Agricolastr. 6, D-38678 Clausthal-Zellerfeld, Germany Received 1 September 2003; accepted 14 October 2003
Abstract The relationship between microstructure and creep properties is discussed for a number of Mg–rare earth cast alloys with reference to the minimum creep rate. Plate shaped precipitates, which form on the prismatic planes of the matrix in a dense triangular arrangement, provide not only most effective barriers to the motion of basal dislocations motion but are also very effective against creep deformation. Shear-resistant small discs or plates parallel to the basal planes of the ␣-matrix are the least effective barriers to the motion of basal dislocations but are strong obstacles to cross slip of basal dislocations and to non-basal slip. The influence of stacking fault energy decrease on the minimum creep rate is also mentioned. © 2004 Elsevier B.V. All rights reserved. Keywords: Metals; Precipitation; Scanning and transmission electron microscopy
1. Introduction Magnesium alloys are attractive for space, aeronautical, automobile and leisure applications because of their low density, high specific strength, good machinability and availability. The use of low cost Mg alloys is limited due to their poor mechanical and creep properties at elevated and high temperatures. The commercially available highly creep resistant alloys for applications at elevated and high temperatures, e.g. QE22 (MgAgNdZr) and WE54, WE43 (MgYNdZr), respectively, often fulfil the specifications [1–3] but not at an economical price. A similar situation exists for the experimental high temperature creep resistant alloys developed by alloying Mg with various rare earths and combinations thereof, e.g. Mg–Gd [4–6], Mg–Gd–Y, Mg–Gd–Nd [7], or with the addition of Sc and Mn, e.g. Mg–Sc–Mn [8], Mg–Sc–Gd (or Y or Ce)–Mn [9,10]. In particular, the knowledge acquired about the precipitate structure and mor∗ Corresponding author. Tel.: +420-2-21911458; fax: +420-2-21911490. E-mail address: [email protected] (B. Smola).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.10.099
phology leading to a very high creep resistance in these Mg alloys is especially valuable for further development of the low cost highly creep resistant Mg alloys for elevated temperature applications. The steady-state creep of Mg alloys is realized, at a desired applied stress, by dislocation climb or cross slip of basal dislocations and/or by non-basal slip at elevated temperatures. Over a wide temperature range, it is often possible to obtain two values of apparent creep activation energies for Mg–rare earth (RE) alloys (Y, Nd, Gd) (e.g. [2,11,12]) corresponding to two different ranges of temperature. The conclusions about the creep mechanisms are not unambiguous even for very similar conditions [13]. The influence of precipitate morphology on the strength and creep resistance is insufficiently understood due to a lack of knowledge of the interactions between precipitates and dislocations [14]. The predominantly non-spherical particles, precipitating as transient or equilibrium phases in Mg–RE alloys, are usually rationally oriented in the ␣-Mg matrix [15]. Contemporary knowledge concerning effect of precipitate morphology on creep resistance of Mg–RE base alloys is briefly summarized here.
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2. Influence of precipitate morphology on creep behavior 2.1. Plate shaped precipitates parallel to prismatic planes Plate shaped precipitates are the most common in Mg alloys. The precipitates with habit planes lying in the prismatic planes of the ␣-Mg matrix and which form the dense interlocking triangular arrangement are the most effective obstacles to basal dislocation slip [1,2,15]. The desirable dense triangular arrangement of prismatic plates of the metastable hexagonal D019 or stable Mg23 Th6 (fcc) phases was observed in the highly creep resistant Mg alloys with thorium [1,16]. Most of the transient and equilibrium phases with a decomposition sequence of Mg–RE (Gd or Y type), namely coherent metastable hexagonal phase ( , D019 structure a = 2.aMg , c = cMg , Mg3 Gd), metastable semicoherent c base centered orthorhombic (cbco) phase ( , a = 2.aMg , b ≈ 8.dMg (1 1¯ 0 0), c = cMg , Mg3 Gd), transient cubic phase (1 , fcc, a = 0.74 nm, Mg3 Nd) and equilibrium cubic phase (, fcc, a ≈ 8.dMg (1 1¯ 0 0), Mg5 Gd), precipitate as prismatic plates in the triangular arrangement. The transient 1 phase was observed only in WE, Mg–Gd–Nd–Zr and Mg–Dy–Nd–Zr alloys during isothermal annealing at 250 ◦ C and nucleated on the globular particles of the  cbco phase [17,18]. In binary Mg–Gd, Mg–Gd–Nd–Zr, Mg–Dy–Nd–Zr, Mg–Gd–Y–Zr, Mg–Sc–RE–Mn (RE = Gd or Y) and WE alloys the decomposition sequence without the 1 phase was observed during isochronal annealing and isothermal treatment leading to the peak hardness [1,4,7,9,19]. The occurrence of precipitation of the  and 1 phases in Mg alloys with a combination of RE of Y and Ce subgroups is probably dependent on the treatment conditions. A very dense triangular arrangement of cbco phase in the MgGd15 alloy (size of plates ∼15 × 5 nm, volume fraction ∼0.018 [6]) developed during the T6 treatment leads to the outstanding mechanical properties of this alloy (yield stress ∼263 MPa at 200 ◦ C, as well as to the excellent creep resistance at 200 ◦ C—see in Table 1 a comparison with the WE43 alloy. The same temperature dependence of the minimum creep rate is observed for WE43 and MgGd alloys with two different values of activation energies for temperatures below 275 and above 300 ◦ C. The triangular arrangement of the cbco phase transforms into a triangular arrangement of stable Mg5 Gd plates during
Table 1 Minimum creep rates at 200 ◦ C and 60 MPa for as cast and peak age hardened alloys Alloy
Minimum creep rate (s−1 )
WE43 as cast 7.3 × 10−10 WE43 T6 1.2 × 10−10 MgGd10 as cast 1.3 × 10−10
Alloy
Minimum creep rate (s−1 )
MgGd10 T6 1.1 × 10−10 MgGd15 as cast 7 × 10−11 MgGd15 T6 2 × 10−11
Fig. 1. Dislocations passing between Mg5 Gd prismatic plates. MgGd15 after creep at 300 ◦ C and 60 MPa. Bright field in (1 1¯ 0 0) reflection near [1 1 2¯ 3] pole.
creep at higher temperatures [20]. The volume fraction and the size of the plates increase [6]. Fig. 1 shows dislocations passing between large Mg5 Gd plates oriented parallel to equivalent {1 0 1¯ 0} prismatic planes in MgGd15 crept to fracture at 300 ◦ C and 60 MPa. The creep curves of MgGd15 at 350 ◦ C and 30 MPa are compared with those of WE43 in Fig. 2 [5,21]. This arrangement of coarse prismatic plates causes a decrease in the minimum creep rate also in the temperature range where most authors assume cross slip to be the mechanism controlling creep in WE type alloys [22]. 2.2. Spherical particles For a given volume fraction and distribution of precipitates, the increase in the Orowan flow stress due to spherical particles is moderate compared to other regularly shaped precipitates in magnesium alloys [15]. Fine spherical particles of Mn2 Sc (hexagonal structure, P63 /mmc, a = 0.5033 nm, c = 0.828 nm, diameter ∼5 nm) develop in the MgSc15Mn1 alloy during a T5 heat treatment. They exhibit high-temperature stability and do not dissolve up to 570 ◦ C. Fig. 3 shows the influence on the minimum creep rate compared to the binary T6 treated MgSc12 alloy. The microstructure of MgSc12 after a T6 treatment consists only of a small volume fraction of irregularly (nest-like) distributed plate- or needle-like shaped precipitates ori¯ ented in 1120 directions in the Mg basal plane. The creep resistance of the MgSc15Mn1 alloy is two or three
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B. Smola et al. / Journal of Alloys and Compounds 378 (2004) 196–201
Fig. 2. Creep curves of MgGd15 and WE43 alloys at 350 ◦ C and 30 MPa.
orders of magnitude better than that of the MgSc12 alloy with at higher temperatures a shift in the transition zone between two temperature ranges with different activation energies. The total hardening effect of Mn2 Sc spherical particles in MgSc15Mn1 alloy is moderate, but it is relatively high (35% of the initial hardness after a T5 treatment), even for the volume fraction only ∼0.001 due to the high volume density. 2.3. Plate shaped precipitates parallel to basal planes Plates, which precipitate on basal planes in the Mg-matrix, provide the least effective barriers to basal dislocations in
Mg alloys with non-shearable plates. The shape of the small discs on the basal planes is typical for Mn2 Sc precipitates in Mg–RE–Sc–Mn (RE = Gd, Y, Ce) alloys (diameter ∼10–20 nm, thickness ∼3–5 nm). The increase in hardness is low (16% of initial hardness) due to the T5 treatment, which causes precipitation of Mn2 Sc basal discs in the MgSc6Ce4Mn1 alloy, although the volume fraction is approximately three times higher than that of the spherical form in the MgSc15Mn1 alloy. The discs exhibit a coherent strain in the [0 0 0 1]Mg direction (see Fig. 4) and do not coarsen during isochronal annealing up to 480 ◦ C or long creep exposure at 350 ◦ C [10,19]. The minimum creep rates of Mg–Sc–RE–Mn (RE = Y or Gd) alloys are lower than or comparable to those of WE and Mg–Gd alloys in the high
Fig. 3. Temperature dependence of minimum creep rates at 30 MPa for binary MgSc and ternary MgScMn alloy in peak age hardened condition.
B. Smola et al. / Journal of Alloys and Compounds 378 (2004) 196–201
Fig. 4. The [0 0 0 1] direction strain contrast of Mn2 Sc discs parallel to basal plane. MgCe4Sc1Mn1 crept at 350 ◦ C and 30 MPa. Dark field weak beam image in (0 0 0 2) reflection near [1 0 1¯ 0] pole.
199
Fig. 5. Thin hexagonal plates of Mn and Y containing phase parallel to basal plane. MgY4Sc1Mn1 alloy after T5 temper. Bright field, [4 1 5¯ 0] pole, (0 0 0 2) reflection. Notice also Mn2 Sc discs.
temperature range (see Table 2: 300 ◦ C) for lower RE contents and consequently for lower volume fractions of coarse prismatic plates of stable Mg–RE system phases. This is due to the effect of basal Mn2 Sc discs as strong obstacles to the cross slip of basal dislocations and to non-basal slip. The complex precipitate structure in low Sc content Mg–Sc–RE–Mn (RE = Y or Gd) alloys includes after a T5 temper, besides Mn2 Sc, also very thin plates of transient hexagonal phase parallel to the basal plane (a = 2.dMg (1 1¯ 0 0), c = cMg diameter ∼30–100 nm, thickness ∼1.5 nm) (Fig. 5). However, these plates containing Mn and RE dissolve at temperatures higher than that of the corresponding T5 treatment [19]. Their aspect ratio is high (20–70), and therefore they can be very strong obstacles to cross slip of basal dislocations at elevated temperatures. The third constituent of the complex precipitate structure is the triangular arrangement of prismatic plates of the transient Mg–RE phase (cbco). If basal plates of the RE–Mn phase co-exist with other constituents, for example, as is
Table 2 Minimum creep rates at 300 ◦ C and 40 MPa for various peak age hardened alloys Alloy
Minimum creep rate (s−1 )
MgGd15 1.5 × 10−8 WE43 4.2 × 10−8 MgGd5Sc1Mn1 9 × 10−9
Alloy
Minimum creep rate (s−1 )
MgGd10Sc1Mn1 5 × 10−9 MgY4Sc1Mn1 1.8 × 10−9 MgCe3Sc1Mn1 2.4 × 10−9
Fig. 6. Stacking faults and partial dislocations in MgGd10 alloy over-aged after T6 treatment. Dark field weak beam in (1 1¯ 0 0) reflection near [1 1 2¯ 3] pole.
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Fig. 7. Creep curves of MgYScMn and MgYZnMn alloys at 350 ◦ C and 30 MPa in comparison to WE43.
the case in the MgY4Sc1Mn1 alloy at 250 ◦ C, they influence the creep resistance very positively—minimum creep rate at 250 ◦ C and 30 MPa ∼5 × 10−11 s−1 . In the low temperature range, where the climb of basal dislocations is expected to be the controlling mechanism for secondary creep, the basal plates of the RE–Mn phase can be also effective in inhibiting relatively long climb of basal dislocations over prismatic plates of the transient Mg–RE phase. Basal plates (GP zones) with the same structure were reported in Mg–RE–Zn–Zr alloy after a T6 treatment [23].
3. Influence of other factors It is not only the precipitates in the grain interior, which affect the creep behavior. In all Mg–RE base alloys the grain boundaries are decorated by coarse particles of stable phases. In Ce containing alloys the grain boundary eutectic persists during heat treatment up to the eutectic temperature due to the low Ce solubility in Mg. In the operating temperature range of Mg–RE base alloys grain boundary sliding is not the main factor contributing to creep deformation in squeeze cast alloys. Another factor, which can influence the creep rates at higher temperatures, is the stacking fault energy. Rare earth elements in solid solution in Mg decrease the stacking fault energy. In Fig. 6 stacking faults are imaged in the MgGd10 binary alloy, which was over-aged after a T6 temper. Splitting of basal dislocations into partials has been recently observed in Mg–0.9 at.% Y–0.04 at.% Zn [24]. The influence of a low value of stacking fault energy on the minimum creep rate is demonstrated in Fig. 7 for the MgY4Zn1Mn1 alloy, where contrary to the MgY4Sc1Mn1 alloy partial dislocations and a fringe contrast of stacking faults were observed.
Acknowledgements The supports by German Research Foundation (DFG), by the Czech Grant Agency (GACR project 106/03/0903) and by Ministry of Education of Czech Republic within the framework of the research program MSM 113200002 are gratefully acknowledged.
References [1] G.W. Lorimer, in: H. Baker (Ed.), Proceedings of the London Conference on Magnesium Technology, Institute of Metals, London, 1986, p. 47. [2] B.L. Mordike, I. Stul´ıková, in: The Metallurgy of Light Alloys, Institute of Metals, London, 1983, p. 145. [3] M. Ahmed, G.W. Lorimer, P. Lyon, R. Pilkington, in: B.L. Mordike, F. Hehman (Eds.), Magnesium Alloys and Their Applications, DGM Informationsgeselschaft, Oberusel, 1992, p. 301. [4] P. Vostrý, I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Phys. Stat. Sol. A 175 (1999) 491. [5] I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Mater. wiss. u. Werkstofftech. 32 (2001) 20. [6] B. Smola, I. Stul´ıková, F. von Buch, B.L. Mordike, Mater. Sci. Eng. A 324 (2002) 113. [7] S. Kamado, Y. Kojima, R. Ninomyia, K. Kubota, in: G.W. Lorimer (Ed.), Proceedings of the 3rd International Magnesium Conference, The Institute of Materials, Cambridge, 1997, p. 327. [8] F. von Buch, J. Lietzau, B.L. Mordike, A. Pisch, R. Schmid-Fetzer, Mater. Sci. Eng. A 263 (1999) 1. [9] B. Smola, I. Stul´ıková, J. Pelcová, F. von Buch, B.L. Mordike, Phys. Stat. Sol. A 191 (2002) 305. [10] I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Mater. wiss. u. Werkstofftech. 34 (2003) 102. [11] J.G. Wang, L.M. Hsiung, T.G. Nieh, M. Mabuchi, Mater. Sci. Eng. A 315 (2001) 81. [12] M. Suzuki, H. Sato, K. Maruyama, H. Oikawa, Mater. Sci. Eng. A 315–319 (2001) 751. [13] B.L. Mordike, Mater. Sci. Eng. A 324 (2002) 103. [14] J.F. Nie, Scripta Mater. 48 (2003) 981. [15] J.F. Nie, Scripta Mater. 48 (2003) 1009.
B. Smola et al. / Journal of Alloys and Compounds 378 (2004) 196–201 [16] I.J. Polmear, Mater. Sci. Tech. 10 (1994) 1. [17] J.F. Nie, B.C. Muddle, Acta Mater. 48 (2000) 1691. [18] P.J. Apps, H. Karimzadeh, J.F. King, G.W. Lorimer, Scripta Mater. 48 (2003) 1023. [19] B. Smola, I. Stul´ıková, J. Pelcová, F. von Buch, B.L. Mordike, Z. Metallkde. 94 (2003) 553. [20] M. Svoboda, M. Pahutova, K. Kucharova, V. Sklenicka, in: J.D. Parker (Ed.), Creep and Fracture of Engineering Materials and Structures, The Institute of Materials, London, 2001, pp. 173–181.
201
[21] F. von Buch, Ph.D. thesis, TU Clausthal, Clausthal-Zellerfeld, 1999. [22] B.L. Mordike, P. Lukáˇc, in: G.W. Lorimer (Ed.), Proceedings of the Third International Magnesium Conference, Institute of Materials, London, 1997, pp. 419–429. [23] D.-H. Ping, K. Hono, J.F. Nie, Scripta Mater. 48 (2003) 1017. [24] M. Suzuki, T. Kimura, J. Koike, K. Maruyama, Scripta Mater. 48 (2003) 997.
Significance of stable and metastable phases in high temperature creep resistant magnesium–rare earth base alloys B. Smola a,b,∗ , I. Stul´ıková a,b , J. Pelcová a , B.L. Mordike b,c a
c
Faculty of Mathematics and Physics, Charles University, Ke Karlovu 5, CZ-121 16 Prague 2, Czech Republic b Zentrum für Funktionswerkstoffe gGmbH, Sachsenweg 8, D-38678 Clausthal-Zellerfeld, Germany Institut für Werkstoffkunde und Werkstofftechnik, TU Clausthal, Agricolastr. 6, D-38678 Clausthal-Zellerfeld, Germany Received 1 September 2003; accepted 14 October 2003
Abstract The relationship between microstructure and creep properties is discussed for a number of Mg–rare earth cast alloys with reference to the minimum creep rate. Plate shaped precipitates, which form on the prismatic planes of the matrix in a dense triangular arrangement, provide not only most effective barriers to the motion of basal dislocations motion but are also very effective against creep deformation. Shear-resistant small discs or plates parallel to the basal planes of the ␣-matrix are the least effective barriers to the motion of basal dislocations but are strong obstacles to cross slip of basal dislocations and to non-basal slip. The influence of stacking fault energy decrease on the minimum creep rate is also mentioned. © 2004 Elsevier B.V. All rights reserved. Keywords: Metals; Precipitation; Scanning and transmission electron microscopy
1. Introduction Magnesium alloys are attractive for space, aeronautical, automobile and leisure applications because of their low density, high specific strength, good machinability and availability. The use of low cost Mg alloys is limited due to their poor mechanical and creep properties at elevated and high temperatures. The commercially available highly creep resistant alloys for applications at elevated and high temperatures, e.g. QE22 (MgAgNdZr) and WE54, WE43 (MgYNdZr), respectively, often fulfil the specifications [1–3] but not at an economical price. A similar situation exists for the experimental high temperature creep resistant alloys developed by alloying Mg with various rare earths and combinations thereof, e.g. Mg–Gd [4–6], Mg–Gd–Y, Mg–Gd–Nd [7], or with the addition of Sc and Mn, e.g. Mg–Sc–Mn [8], Mg–Sc–Gd (or Y or Ce)–Mn [9,10]. In particular, the knowledge acquired about the precipitate structure and mor∗ Corresponding author. Tel.: +420-2-21911458; fax: +420-2-21911490. E-mail address: [email protected] (B. Smola).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2003.10.099
phology leading to a very high creep resistance in these Mg alloys is especially valuable for further development of the low cost highly creep resistant Mg alloys for elevated temperature applications. The steady-state creep of Mg alloys is realized, at a desired applied stress, by dislocation climb or cross slip of basal dislocations and/or by non-basal slip at elevated temperatures. Over a wide temperature range, it is often possible to obtain two values of apparent creep activation energies for Mg–rare earth (RE) alloys (Y, Nd, Gd) (e.g. [2,11,12]) corresponding to two different ranges of temperature. The conclusions about the creep mechanisms are not unambiguous even for very similar conditions [13]. The influence of precipitate morphology on the strength and creep resistance is insufficiently understood due to a lack of knowledge of the interactions between precipitates and dislocations [14]. The predominantly non-spherical particles, precipitating as transient or equilibrium phases in Mg–RE alloys, are usually rationally oriented in the ␣-Mg matrix [15]. Contemporary knowledge concerning effect of precipitate morphology on creep resistance of Mg–RE base alloys is briefly summarized here.
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197
2. Influence of precipitate morphology on creep behavior 2.1. Plate shaped precipitates parallel to prismatic planes Plate shaped precipitates are the most common in Mg alloys. The precipitates with habit planes lying in the prismatic planes of the ␣-Mg matrix and which form the dense interlocking triangular arrangement are the most effective obstacles to basal dislocation slip [1,2,15]. The desirable dense triangular arrangement of prismatic plates of the metastable hexagonal D019 or stable Mg23 Th6 (fcc) phases was observed in the highly creep resistant Mg alloys with thorium [1,16]. Most of the transient and equilibrium phases with a decomposition sequence of Mg–RE (Gd or Y type), namely coherent metastable hexagonal phase ( , D019 structure a = 2.aMg , c = cMg , Mg3 Gd), metastable semicoherent c base centered orthorhombic (cbco) phase ( , a = 2.aMg , b ≈ 8.dMg (1 1¯ 0 0), c = cMg , Mg3 Gd), transient cubic phase (1 , fcc, a = 0.74 nm, Mg3 Nd) and equilibrium cubic phase (, fcc, a ≈ 8.dMg (1 1¯ 0 0), Mg5 Gd), precipitate as prismatic plates in the triangular arrangement. The transient 1 phase was observed only in WE, Mg–Gd–Nd–Zr and Mg–Dy–Nd–Zr alloys during isothermal annealing at 250 ◦ C and nucleated on the globular particles of the  cbco phase [17,18]. In binary Mg–Gd, Mg–Gd–Nd–Zr, Mg–Dy–Nd–Zr, Mg–Gd–Y–Zr, Mg–Sc–RE–Mn (RE = Gd or Y) and WE alloys the decomposition sequence without the 1 phase was observed during isochronal annealing and isothermal treatment leading to the peak hardness [1,4,7,9,19]. The occurrence of precipitation of the  and 1 phases in Mg alloys with a combination of RE of Y and Ce subgroups is probably dependent on the treatment conditions. A very dense triangular arrangement of cbco phase in the MgGd15 alloy (size of plates ∼15 × 5 nm, volume fraction ∼0.018 [6]) developed during the T6 treatment leads to the outstanding mechanical properties of this alloy (yield stress ∼263 MPa at 200 ◦ C, as well as to the excellent creep resistance at 200 ◦ C—see in Table 1 a comparison with the WE43 alloy. The same temperature dependence of the minimum creep rate is observed for WE43 and MgGd alloys with two different values of activation energies for temperatures below 275 and above 300 ◦ C. The triangular arrangement of the cbco phase transforms into a triangular arrangement of stable Mg5 Gd plates during
Table 1 Minimum creep rates at 200 ◦ C and 60 MPa for as cast and peak age hardened alloys Alloy
Minimum creep rate (s−1 )
WE43 as cast 7.3 × 10−10 WE43 T6 1.2 × 10−10 MgGd10 as cast 1.3 × 10−10
Alloy
Minimum creep rate (s−1 )
MgGd10 T6 1.1 × 10−10 MgGd15 as cast 7 × 10−11 MgGd15 T6 2 × 10−11
Fig. 1. Dislocations passing between Mg5 Gd prismatic plates. MgGd15 after creep at 300 ◦ C and 60 MPa. Bright field in (1 1¯ 0 0) reflection near [1 1 2¯ 3] pole.
creep at higher temperatures [20]. The volume fraction and the size of the plates increase [6]. Fig. 1 shows dislocations passing between large Mg5 Gd plates oriented parallel to equivalent {1 0 1¯ 0} prismatic planes in MgGd15 crept to fracture at 300 ◦ C and 60 MPa. The creep curves of MgGd15 at 350 ◦ C and 30 MPa are compared with those of WE43 in Fig. 2 [5,21]. This arrangement of coarse prismatic plates causes a decrease in the minimum creep rate also in the temperature range where most authors assume cross slip to be the mechanism controlling creep in WE type alloys [22]. 2.2. Spherical particles For a given volume fraction and distribution of precipitates, the increase in the Orowan flow stress due to spherical particles is moderate compared to other regularly shaped precipitates in magnesium alloys [15]. Fine spherical particles of Mn2 Sc (hexagonal structure, P63 /mmc, a = 0.5033 nm, c = 0.828 nm, diameter ∼5 nm) develop in the MgSc15Mn1 alloy during a T5 heat treatment. They exhibit high-temperature stability and do not dissolve up to 570 ◦ C. Fig. 3 shows the influence on the minimum creep rate compared to the binary T6 treated MgSc12 alloy. The microstructure of MgSc12 after a T6 treatment consists only of a small volume fraction of irregularly (nest-like) distributed plate- or needle-like shaped precipitates ori¯ ented in 1120 directions in the Mg basal plane. The creep resistance of the MgSc15Mn1 alloy is two or three
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Fig. 2. Creep curves of MgGd15 and WE43 alloys at 350 ◦ C and 30 MPa.
orders of magnitude better than that of the MgSc12 alloy with at higher temperatures a shift in the transition zone between two temperature ranges with different activation energies. The total hardening effect of Mn2 Sc spherical particles in MgSc15Mn1 alloy is moderate, but it is relatively high (35% of the initial hardness after a T5 treatment), even for the volume fraction only ∼0.001 due to the high volume density. 2.3. Plate shaped precipitates parallel to basal planes Plates, which precipitate on basal planes in the Mg-matrix, provide the least effective barriers to basal dislocations in
Mg alloys with non-shearable plates. The shape of the small discs on the basal planes is typical for Mn2 Sc precipitates in Mg–RE–Sc–Mn (RE = Gd, Y, Ce) alloys (diameter ∼10–20 nm, thickness ∼3–5 nm). The increase in hardness is low (16% of initial hardness) due to the T5 treatment, which causes precipitation of Mn2 Sc basal discs in the MgSc6Ce4Mn1 alloy, although the volume fraction is approximately three times higher than that of the spherical form in the MgSc15Mn1 alloy. The discs exhibit a coherent strain in the [0 0 0 1]Mg direction (see Fig. 4) and do not coarsen during isochronal annealing up to 480 ◦ C or long creep exposure at 350 ◦ C [10,19]. The minimum creep rates of Mg–Sc–RE–Mn (RE = Y or Gd) alloys are lower than or comparable to those of WE and Mg–Gd alloys in the high
Fig. 3. Temperature dependence of minimum creep rates at 30 MPa for binary MgSc and ternary MgScMn alloy in peak age hardened condition.
B. Smola et al. / Journal of Alloys and Compounds 378 (2004) 196–201
Fig. 4. The [0 0 0 1] direction strain contrast of Mn2 Sc discs parallel to basal plane. MgCe4Sc1Mn1 crept at 350 ◦ C and 30 MPa. Dark field weak beam image in (0 0 0 2) reflection near [1 0 1¯ 0] pole.
199
Fig. 5. Thin hexagonal plates of Mn and Y containing phase parallel to basal plane. MgY4Sc1Mn1 alloy after T5 temper. Bright field, [4 1 5¯ 0] pole, (0 0 0 2) reflection. Notice also Mn2 Sc discs.
temperature range (see Table 2: 300 ◦ C) for lower RE contents and consequently for lower volume fractions of coarse prismatic plates of stable Mg–RE system phases. This is due to the effect of basal Mn2 Sc discs as strong obstacles to the cross slip of basal dislocations and to non-basal slip. The complex precipitate structure in low Sc content Mg–Sc–RE–Mn (RE = Y or Gd) alloys includes after a T5 temper, besides Mn2 Sc, also very thin plates of transient hexagonal phase parallel to the basal plane (a = 2.dMg (1 1¯ 0 0), c = cMg diameter ∼30–100 nm, thickness ∼1.5 nm) (Fig. 5). However, these plates containing Mn and RE dissolve at temperatures higher than that of the corresponding T5 treatment [19]. Their aspect ratio is high (20–70), and therefore they can be very strong obstacles to cross slip of basal dislocations at elevated temperatures. The third constituent of the complex precipitate structure is the triangular arrangement of prismatic plates of the transient Mg–RE phase (cbco). If basal plates of the RE–Mn phase co-exist with other constituents, for example, as is
Table 2 Minimum creep rates at 300 ◦ C and 40 MPa for various peak age hardened alloys Alloy
Minimum creep rate (s−1 )
MgGd15 1.5 × 10−8 WE43 4.2 × 10−8 MgGd5Sc1Mn1 9 × 10−9
Alloy
Minimum creep rate (s−1 )
MgGd10Sc1Mn1 5 × 10−9 MgY4Sc1Mn1 1.8 × 10−9 MgCe3Sc1Mn1 2.4 × 10−9
Fig. 6. Stacking faults and partial dislocations in MgGd10 alloy over-aged after T6 treatment. Dark field weak beam in (1 1¯ 0 0) reflection near [1 1 2¯ 3] pole.
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Fig. 7. Creep curves of MgYScMn and MgYZnMn alloys at 350 ◦ C and 30 MPa in comparison to WE43.
the case in the MgY4Sc1Mn1 alloy at 250 ◦ C, they influence the creep resistance very positively—minimum creep rate at 250 ◦ C and 30 MPa ∼5 × 10−11 s−1 . In the low temperature range, where the climb of basal dislocations is expected to be the controlling mechanism for secondary creep, the basal plates of the RE–Mn phase can be also effective in inhibiting relatively long climb of basal dislocations over prismatic plates of the transient Mg–RE phase. Basal plates (GP zones) with the same structure were reported in Mg–RE–Zn–Zr alloy after a T6 treatment [23].
3. Influence of other factors It is not only the precipitates in the grain interior, which affect the creep behavior. In all Mg–RE base alloys the grain boundaries are decorated by coarse particles of stable phases. In Ce containing alloys the grain boundary eutectic persists during heat treatment up to the eutectic temperature due to the low Ce solubility in Mg. In the operating temperature range of Mg–RE base alloys grain boundary sliding is not the main factor contributing to creep deformation in squeeze cast alloys. Another factor, which can influence the creep rates at higher temperatures, is the stacking fault energy. Rare earth elements in solid solution in Mg decrease the stacking fault energy. In Fig. 6 stacking faults are imaged in the MgGd10 binary alloy, which was over-aged after a T6 temper. Splitting of basal dislocations into partials has been recently observed in Mg–0.9 at.% Y–0.04 at.% Zn [24]. The influence of a low value of stacking fault energy on the minimum creep rate is demonstrated in Fig. 7 for the MgY4Zn1Mn1 alloy, where contrary to the MgY4Sc1Mn1 alloy partial dislocations and a fringe contrast of stacking faults were observed.
Acknowledgements The supports by German Research Foundation (DFG), by the Czech Grant Agency (GACR project 106/03/0903) and by Ministry of Education of Czech Republic within the framework of the research program MSM 113200002 are gratefully acknowledged.
References [1] G.W. Lorimer, in: H. Baker (Ed.), Proceedings of the London Conference on Magnesium Technology, Institute of Metals, London, 1986, p. 47. [2] B.L. Mordike, I. Stul´ıková, in: The Metallurgy of Light Alloys, Institute of Metals, London, 1983, p. 145. [3] M. Ahmed, G.W. Lorimer, P. Lyon, R. Pilkington, in: B.L. Mordike, F. Hehman (Eds.), Magnesium Alloys and Their Applications, DGM Informationsgeselschaft, Oberusel, 1992, p. 301. [4] P. Vostrý, I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Phys. Stat. Sol. A 175 (1999) 491. [5] I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Mater. wiss. u. Werkstofftech. 32 (2001) 20. [6] B. Smola, I. Stul´ıková, F. von Buch, B.L. Mordike, Mater. Sci. Eng. A 324 (2002) 113. [7] S. Kamado, Y. Kojima, R. Ninomyia, K. Kubota, in: G.W. Lorimer (Ed.), Proceedings of the 3rd International Magnesium Conference, The Institute of Materials, Cambridge, 1997, p. 327. [8] F. von Buch, J. Lietzau, B.L. Mordike, A. Pisch, R. Schmid-Fetzer, Mater. Sci. Eng. A 263 (1999) 1. [9] B. Smola, I. Stul´ıková, J. Pelcová, F. von Buch, B.L. Mordike, Phys. Stat. Sol. A 191 (2002) 305. [10] I. Stul´ıková, B. Smola, F. von Buch, B.L. Mordike, Mater. wiss. u. Werkstofftech. 34 (2003) 102. [11] J.G. Wang, L.M. Hsiung, T.G. Nieh, M. Mabuchi, Mater. Sci. Eng. A 315 (2001) 81. [12] M. Suzuki, H. Sato, K. Maruyama, H. Oikawa, Mater. Sci. Eng. A 315–319 (2001) 751. [13] B.L. Mordike, Mater. Sci. Eng. A 324 (2002) 103. [14] J.F. Nie, Scripta Mater. 48 (2003) 981. [15] J.F. Nie, Scripta Mater. 48 (2003) 1009.
B. Smola et al. / Journal of Alloys and Compounds 378 (2004) 196–201 [16] I.J. Polmear, Mater. Sci. Tech. 10 (1994) 1. [17] J.F. Nie, B.C. Muddle, Acta Mater. 48 (2000) 1691. [18] P.J. Apps, H. Karimzadeh, J.F. King, G.W. Lorimer, Scripta Mater. 48 (2003) 1023. [19] B. Smola, I. Stul´ıková, J. Pelcová, F. von Buch, B.L. Mordike, Z. Metallkde. 94 (2003) 553. [20] M. Svoboda, M. Pahutova, K. Kucharova, V. Sklenicka, in: J.D. Parker (Ed.), Creep and Fracture of Engineering Materials and Structures, The Institute of Materials, London, 2001, pp. 173–181.
201
[21] F. von Buch, Ph.D. thesis, TU Clausthal, Clausthal-Zellerfeld, 1999. [22] B.L. Mordike, P. Lukáˇc, in: G.W. Lorimer (Ed.), Proceedings of the Third International Magnesium Conference, Institute of Materials, London, 1997, pp. 419–429. [23] D.-H. Ping, K. Hono, J.F. Nie, Scripta Mater. 48 (2003) 1017. [24] M. Suzuki, T. Kimura, J. Koike, K. Maruyama, Scripta Mater. 48 (2003) 997.