Effect of addition of Sn on the microstructure and mechanical properties of Mg–MM (misch-metal) alloys

Journal of Alloys and Compounds 454 (2008) 515–522

Effect of addition of Sn on the microstructure and mechanical properties of Mg–MM (misch-metal) alloys Hyun Kyu Lim a , Sung Woo Sohn a , Do Hyung Kim a , Ju Youn Lee a , Won Tae Kim b , Do Hyang Kim a,∗ a

Center for Noncrystalline Materials, Department of Metallurgical Engineering, Yonsei University, 134 Sinchon-dong, Seodaemun-gu, Seoul 120-749, South Korea b Applied Science Division, Cheongju University, 36 Naedok-dong, Sangdang-gu, Cheongju, Chongbuk 360-764, South Korea Received 9 August 2007; received in revised form 11 September 2007; accepted 17 September 2007 Available online 25 September 2007

Abstract The effects of addition of Sn on the microstructure and mechanical properties of rolled Mg–MM alloy sheets have been investigated. The secondary solidification phase in Mg–MM alloy changes dramatically with addition of Sn forming small size rod-shaped phase. Although the strength of alloy is decreased with addition of Sn, the ductility is improved when the small rod-shaped phase forms in Mg-rich Mg–MM–Sn alloys due to the sound interface without forming any void at the end of particles after tensile test at 200 ◦ C. The result indicates that the small size rod-shaped particles can enhance the ductility while maintaining the moderate level of strength in Mg–MM–Sn alloys. © 2007 Elsevier B.V. All rights reserved. Keywords: Mg–MM alloys; Sn addition; Secondary solidification phase; Ductility

1. Introduction Due to the demand for light-weight alloys for structural applications, magnesium alloys are receiving attention as structural material with high specific strength. Recently, since addition of rare earth (RE) elements in magnesium alloys has been reported to show a favorable effect on mechanical properties, in particular, at elevated temperature properties, investigation on the detailed strengthening mechanism with addition of RE elements has become an important issue [1]. Key properties of these alloys include high specific strength at room and high temperatures, good creep resistance, good castability and damping characteristics. However, the scientific understanding on the influence of RE elements on mechanical properties is not completed, due to the complexity of alloy systems [2]. Moreover, it is well-known that there is a limitation in using RE elements for industrial magnesium alloy products due to high cost of elements and difficulties in achievement of homogeneous structures. Recently, however, many researchers use misch-metal (MM) instead of

∗

Corresponding author. Tel.: +82 2 2123 4255; fax: +82 2 312 8281. E-mail address: [email protected] (D.H. Kim).

0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.09.074

individual RE elements, since the addition of MM can provide a beneficial effect in terms of the economical perspectives [3–6]. On the other hand, when Sn is alloyed in magnesium alloys, strength and creep resistance have been reported to improve by formation of thermally stable Mg2 Sn particles within the matrix and along grain boundaries [7]. The effect of Sn addition has been investigated in various aspects such as creep resistance in Mg–Sn–Si [8] and Mg–Al–Sn–Ca [9] systems, and high temperature strength in Mg–Al–Zn–Sn [10] system. Moreover, it has been shown that Sn serves to increase the ductility of the alloy and makes it better for hammer forging by reducing the tendency for the alloy to crack while being hot worked [11]. As mentioned above, RE elements and Sn separately have been used as alloying elements to enhance the strength and creep resistance of magnesium alloys. Recently, Y. Chan et al. reported that addition of didymium (neodymium–praseodymium mixed metal) in Mg–5% Sn alloy results in the formation of the small rod-shaped Snx (Nd, Pr)y phase on the grain boundaries leading to an increase in strength and elongation [12]. Considering the beneficial effect of RE and Sn addition in magnesium alloys, it is strongly required to investigate Mg–RE–Sn system for casting products as well as wrought products. Therefore, the present study aims to identify the phases in Mg–RE–Sn sys-

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Table 1 Nominal chemical compositions of Mg–MM–Sn alloys Alloy code

E3 ET43 ET33 ET34

Nominal composition (wt%) Mg

MM

Sn

Balance Balance Balance Balance

3 4 3 3

– 3 3 4

tem, and to evaluate the mechanical properties of Mg–RE–Sn rolled sheets. For cost effectiveness, MM has been used instead of individual RE elements. The main results include effect of Sn on the microstructure of Mg–Mm alloy and mechanical properties of Mg–MM–Sn alloys. To identify the phase structure unanimously, the phases in Mg–Ce–Sn alloy have also been investigated.

10 cm were prepared by pouring the melt into the preheated steel mold. The cast alloys were homogenized at 500 ◦ C for 12 h followed by water quenching. To investigate the rollability and mechanical properties, the specimens were fabricated by hot-rolling to 1 mm final thickness (reduction ∼90%). Before rolling, the rolls were preheated up to ∼100 ◦ C. The ingots preheated at 500 ◦ C for 15 min were rolled with a reduction ratio of 10–30% per pass. The rolled sheets were annealed at 300 ◦ C for 1 h in an air-circulating furnace. Phase identifications were performed by the X-ray diffraction (XRD, Rigaku CN2301) using monochromatic Cu K␣ radiation. For microstructural observations, alloy specimens for optical microscope (LEICA DMRM) were etched with a solution of 30 ml acetic acid + 15 ml water + 6 g picric acid + 100 ml ethanol or 2% nital. Thin foils for the transmission electron microscopy (TEM, JEOL 2100F) were prepared by ion-beam thinning technique at a voltage of 3.0 keV and incident angle from 8◦ to 4◦ . The chemical composition of the specimen was examined by the EDS (energy dispersive spectrum (Oxford)) with INCA computer program. The uniaxial tensile test at room temperature test and 200 ◦ C was carried out on dog-bone specimens of hot-rolled sheets (specimen gauge lengths were 25 mm and 10 mm for room temperature test and 200 ◦ C test, respectively) under a constant cross-head speed. The initial strain rate was 10−3 s−1 . Scanning electron microscope (SEM, Hitachi S-2700) was used to observe the interface between matrix and the second phase in the specimens after tensile test at 200 ◦ C.

2. Experimental procedure

3. Results Nominal chemical compositions of Mg–MM–Sn alloys used in this study are listed in Table 1. The alloys were melted in an electrical resistance furnace with high purity magnesium (99.9%), zinc (99.95%), tin (99.99%) and commercial purity misch-metal (52% Ce–26% La–16% Nd–6% Pr) in the boron nitride (BN)coated steel crucible under a dynamic protective gas (SF6 + CO2 ) atmosphere. The alloy ingots with a dimension of thickness 1.5 cm, width 6 cm, and height

3.1. Microstructure and phase identification Fig. 1(a) and (b) shows the optical microstructures of ascast E3 and ET33 alloys, while Fig. 1(c)–(f) shows the optical

Fig. 1. Optical micrographs of studied alloys: as-cast (a) E3 and (b) ET33 alloys and homogenized (c) E3, (d) ET43, (e) ET33 and (f) ET34 alloys.

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Fig. 2. TEM images of homogenized E3 alloy: (a) bright field image and (b) SADP of Mg12 MM phase.

microstructures of cast E3, ET43, ET33 and ET34 alloys after homogenization at 500 ◦ C for 12 h. The samples were etched with a 2% nital solution to outline the distribution of the interdendritic second phase particles. The as-cast structure of E3 alloy (Fig. 1(a)) exhibited a typical dendritic structure consisted of ␣-Mg dendrite and continuous interdendritic second phase. It can be noticed that with addition of Sn, there is a dramatic change in the morphology of the interdendritic second phase, i.e. from the morphology of the continuous type interdendritic phase in E3 alloy into fine feather-shaped particles embedded in ␣-Mg in ET33 alloy (Fig. 1(b)). With homogenization treatment, the continuous type interdendritic phase in E3 alloy changed into isolated- or elongated-shaped particles (Fig. 1(c)), while the fine feather-shaped particles in ET33 alloy changed into small rod-type particles (Fig. 1(e)). In the ET43 alloy with larger content of MM than Sn, most of the interdendritic phase was rod-type particles, and small volume fraction of isolated particles (marked with dotted circle in Fig. 1 (d)) which appeared in the E3 alloy was still present in the interdendritic region. If the content of Sn is larger than that of MM in the ET34 alloy, polygon-type particles (marked with dotted rectangle in Fig. 1(f)) were present in the interdendritic region, although most of the interdendritic phase was small rod-type particles embedded in ␣-Mg. TEM analysis confirmed the structure of interdendritic phase in the E3 alloy. Fig. 2(a) and (b) shows the bright field TEM image of Mg12 MM phase and corresponding selected area diffraction pattern (SADP) taken from the area marked in Fig. 2(a), respectively. The SADP shows a [0 2 1] zone of the body-centered tetragonal Mg12 MM phase (I4/mmm, a = b = 1.030 nm and c = 0.598 nm). Fig. 3(a) shows the bright field TEM image of the small rod-shaped particles obtained from the ET33 alloy. The corresponding SADPs are shown in Fig. 3(b) and (c). The SADP in Fig. 3(c) was obtained by tilting the specimen about 9.4◦ along the axis parallel to the arrow marked in Fig. 3(b). As shown in the schematic diagram in Fig. 3(d) and (e), the angle between the basic reciprocal lattice vectors in the SADPs (Fig. 3(b) and (c)) was 91◦ and 90.8◦ , respectively. The SADPs could not be identified unanimously with the structure reported in Mg–MM, Mg–Sn, MM–Sn and Mg–MM–Sn systems. The SADP result in Figs. 3 (b) and (c)

indicates that the small rod-shaped phase has lower symmetry crystal structure. EDS analysis (Table 2) showed that the atomic stoichiometry of the interdendritic phase in the E3 alloy and the small rod-shaped particles in the ET33 alloy was close to Mg12 RE and Mg3 RE1 Sn1 , respectively. The result of TEM analysis indicates that the interdendritic phase in the E3 alloy is Mg12 MM phase. However, the small rod-shaped particles appearing as a major interdendritic phase in the ET43, ET33 and ET34 alloys could not be identified unanimously in the present study. The polygon-type interdendritic phase in the ET34 alloy was not identified in detail in the present study. However, the EDS analysis (Table 2) indicates that the polygon-shaped particles has similar composition to the small rod-shaped particles, but slightly lower Mg and higher Sn. It might be thought that the formation of polygon-shaped particles was due to the excess Sn in ET34 alloy. If larger amount of Sn is added in Mg–MM alloy, it is expected that Mg2 Sn phase forms with excess Sn instead of the polygon-shaped particles in ET34 alloy. 3.2. Microstructure and mechanical properties of sheets The rollability of the E3 alloy was not enough to fabricate the rolled sheets. With the addition of Sn in the Mg–MM alloy, the rollability improved dramatically, enabling successful fabrication of rolled sheets in the ET43, ET33 and ET34 alloys. Therefore, the results for the microstructure and mechanical properties of the rolled sheets are presented for the ET43, ET33 and ET34 alloys. Fig. 4(a)–(f) shows the optical microstructures of the ET43, ET33 and ET34 rolled sheets after annealing at 300 ◦ C for 1 h. The arrows in the micrographs represent the rolling direction. The samples were etched with two different solutions; i.e. with a 2% nital solution for observation of the distribution of the interdendritic phase particles (Fig. 4(a)–(c)), and with a solution of 30 ml acetic acid + 15 ml water + 6 g picric acid + 100 ml ethanol for observation of the recrystallized grain structure. Fig. 4(a)–(c) shows the distribution of the small rod-shaped particles in the rolled sheets. When compared with the as-cast microstructures in Fig. 1(b)–(d), the distribution of the particles became more homogeneous, and most of the particles were aligned along the

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Fig. 3. TEM images of ET33 alloy: (a) bright field image and (b) and (c) SADPs of the second phase.

rolling direction. However, size and shape of the Mg12 MM phase and the polygon-type phase in the ET43 and ET34 alloys, respectively (shown in the dotted circles and rectangle in Fig. 4(a) and (c)) were almost similar to those in the as-cast microstructures (Fig. 1(b) and (d)). Fig. 4(d)–(f) shows the optical microstructures of the recrystallized grain structure in the ET43, ET33 and ET34 rolled sheets corresponding to microstructures in Fig. 4(a)–(c). The grain sizes in the ET43, ET33 and ET34 rolled sheets were about 9.8 ␮m, 14.2 ␮m and 11.1 ␮m, respectively. The Mg–MM alloys containing Sn exhibited, in general, a small grain size, due to the presence of the small rod-shaped particles in the interdendritic region. However, among the three alloys the ET43 and ET34 alloys exhibited slightly smaller grain size than the ET33 alloy possibly due to the presence of Mg12 MM and polygon-type particles providing a beneficial effect in the formation of the recrystallized grains.

The engineering stress–strain curves at room temperature obtained from the annealed ET43, ET33 and ET34 alloy sheets are shown in Fig. 5(a). The results of tensile test including yield strength, ultimate tensile strength and elongation are listed in Table 3. Among the alloys investigated, the ET43 alloy exhibited the highest yield strength (159.9 MPa), while the ET33 alloy exhibited the largest elongation-to-fracture (9.45%). The result indicates that smaller grain size and presence of Mg12 MM (and polygon-shaped particle in ET34 alloy) and small rod-shaped particles played a role in enhancing the yield strength of alloys. The ET43 alloy (highest strength) exhibited the smallest grain size of 9.8 ␮m, while the ET33 alloy (lowest strength) exhibited the largest grain size of 14.2 ␮m. Despite the highest yield strength of the ET43 alloy, its ultimate tensile strength was lowest since the ET43 alloy exhibited low strain-hardening exponent and low elongation-to-

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Table 2 EDS analysis of the second phases in E3, ET33 and ET34 alloys Phase

Elements (at.%) Mg

Interdendritic phase in E3 alloy 1 2 Average

91.58 93.76 92.67

Sn – – –

Ce

La

REa

5.57 4.25 4.91

2.85 1.99 2.42

8.42 6.24 7.33

Feather-shaped phase in ET33 alloy 1 62.05 2 60.36 3 63.53 Average 61.98

18.6 19.62 18.43 18.88

12.9 13.21 12.27 12.79

6.45 6.81 5.77 6.34

19.35 20.02 18.04 19.14

Polygon-type phase in ET34 alloy 1 2 3 Average

20.59 20.8 21.15 20.85

14.36 12.94 12.44 13.25

7.52 7.44 7.07 7.34

21.88 20.38 19.51 20.59

a

57.53 58.82 59.34 58.56

RE = Ce + La.

fracture. The true stress–strain curves at an initial strain rate of 10−3 s−1 measured at 200 ◦ C are shown in Fig. 5(b). As in the results at room temperature, the ET33 alloy containing only the small rod-shaped phase exhibited the lowest level of

yield stress but the largest elongation. While ET43 and ET34 sheet containing Mg12 MM and the polygon-type phase, respectively, showed higher yield strengths but smaller elongation-tofractures.

Fig. 4. Optical micrographs of sheets which are annealed at 300 ◦ C for 1 h: (a) ET43, (b) ET33 and (c) ET34 etched with 2% nital; and (d) ET43, (e) ET33 and (f) ET34 etched with a solution of 30 ml acetic acid + 15 ml water + 6 g picric acid + 100 ml ethanol.

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Fig. 5. Curves for tensile test of annealed ET43, ET33 and ET34 specimens: (a) engineering stress–strain curves at room temperature and (b) true stress–strain curves at 200 ◦ C.

small amount of the polygon-shaped phase (volume fraction: ∼6%) coexists with the small rod-shaped phase in the interdendritic region. It should be pointed out that the size of the small rod-shaped particles (thickness: ∼0.7 ␮m and length: ∼5.6 ␮m) is much smaller than that of the Mg12 MM (∼7 ␮m) or polygon-shaped (∼10 ␮m) particles. After rolling the small rod-shaped particles are aligned along the rolling direction and are distributed more homogeneously, while the Mg12 MM and polygon-shaped particles remains almost same as the state before rolling. The EDS result shows that the stiochiometry of the small rod-shaped particles is close to Mg3 RE1 Sn1 . The small rod-shaped phase in the present study is similar to the secondary solidification phase reported in the Mg–Sn–didymium alloys [12]. However, the composition of the small rod-shaped particles, (Snx (Nd, Pr)y ) reported in the Mg–Sn–didymium alloys is not consistent with the composition observed in the present study. Since major element in MM used in the present study is Ce (52 wt%), the microstructure of the Mg–3 wt% Ce–3 wt% Sn alloy has been investigated to confirm the formation of shaped-phase in the interdendritic region. The microstructure of homogenized Mg–3 wt% Ce–3 wt% Sn alloy is same as that in the Mg–3 wt% MM–3 wt% Sn alloy, i.e. the small rod-shaped phase forms in the interdendritic region. The EDS result indicates that the composition of the small rod-type particles in Mg–3 wt%Ce–3 wt%Sn alloy is close to Mg3 Ce1 Sn1 , i.e. the average amounts of Mg, Ce and Sn are 61.08 at.%, 19.75 at.% and 19.17 at.%, respectively. The SADP analysis shows that the structure of the small rod-shaped particles in the Mg–3 wt% Ce–3 wt% Sn alloy is same as that in the Mg–3 wt% MM–3 wt% Sn alloy, although the crystal structure can not be decided unanimously. The structure of the Mg12 MM phase is also similar to the Mg12 Ce compound responsible for the Mg–Ce eutectic reaction [3].

4. Discussion

4.2. Role of particles in mechanical properties

4.1. Microstructure

In general, elongation to failure in the alloys containing intermetallic particles is small, since geometrically necessary dislocations are formed in the region surrounding the hard particles, leading to discohesion from the matrix [13]. However, several alloy systems have been reported to exhibit large elongation to failure although particles are present in the interdendritic region. For example, the Mg–Zn–Y alloy consisted of ␣-Mg and quasicrystalline particles have been reported to show a good combination of strength and ductility due to the low energy interface between ␣-Mg matrix and quasicrystalline particle [14,15]. The present study shows that the size of the secondary phase particles is one of the important factors affecting the combination

The present result shows that the secondary solidification phase in Mg–MM alloy changes dramatically with addition of Sn, i.e. the Mg12 MM binary phase in Mg–MM alloy changes into the small rod-shaped ternary phase with addition of Sn. When the amount of Sn and MM is same (3 wt%), only the small rod-shaped phase forms in the interdendritic region. When the amount of MM (4 wt%) is larger than that of Sn (3 wt%), small amount of the Mg12 MM phase (volume fraction: ∼5%) remains together with the small rod-shaped phase, while when the amount of Sn (4 wt%) is larger than that of MM (3 wt%), Table 3 Results of tensile test at room temperature and 200 ◦ C Alloys

ET43 ET33 ET34

200 ◦ C

Room temperature YS (MPa)

UTS (MPa)

Elong. (%)

YS (MPa)

UTS (MPa)

Elong. (%)

159.9 124.8 132.2

200.9 205.4 211.4

3.94 9.45 6.66

66.9 56.0 60.1

82.6 69.9 74.2

23.2 56.2 38.5

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Fig. 6. SEM images of the specimen after tensile test at 200 ◦ C: (a) ET33 and (b) ET34.

of mechanical properties. When the small size small rod-shaped particles are present, the elongation to failure increases significantly maintaining moderate level of strength. While when the large size Mg12 MM or polygon-shaped particles are present, the elongation failure is significantly deteriorated, but the strength is improved. In particular, polygon-type phase exhibits an adverse effect on the ductility due to its size and shape, although its composition is similar to that of the small rod-shaped phase. In the case of tensile test at high temperature, the stress level changes depending on the amount of the strain. Typically, the strength of the alloys containing the second phase particles reaches a peak at small amount of strain and decreases continuously with increasing strain. However, the alloys investigated in the present study exhibit strain hardening behavior in the initial small strain range and then the flow stress decreases with increasing strain. This incubation strain is necessary for recrystallization since the nucleation site for recrystallization is limited [14]. The ET43 alloy shows the smallest incubation stain due to the existence of Mg12 MM particles, while the ET33 alloy exhibits the largest incubation strain, since the small rod-shaped phase does not act effectively as an obstacle to movement of dislocations. The effect of small rod-type particles on movement of dislocations can be noticed in the recrystallized microstructures (Fig. 4(b), (d) and (f)). Generally, the recrystallized grains are assumed to form in the vicinity of some obstacles such as grain boundaries and particles where the dislocations are heavily piled. Then the newly formed recrystallized grains grow by the movement of atoms along the matrix–grain interface [16]. If the particles strongly acted as obstacles to movement of dislocations, the recrystallized grain size becomes smaller. However, the annealed microstructure of ET33 sheet exhibits the largest grain size among the alloys studied. Therefore, the largest incubation strain and the largest grain size in ET33 alloy sheet indicate that the small rod-type particles do not act effectively as an obstacle to movement of dislocations when compared to other type of particles in ET43 and ET34 alloys. As mentioned above, when the content of Sn is larger than MM, polygon-type particles with the similar composition to the small rod-shaped phase form, and act more effectually as obstacles than small rod-shaped particles. Therefore, the grain size of the ET34 alloy was smaller than that of ET33 alloy and the strength is higher at room and high temperatures. For more

detailed examination, the interfaces between matrix and the second phase in the ET33 and ET34 specimens after tensile test at 200 ◦ C have been observed from the surface of the specimen by SEM (Fig. 6). The arrow indicated the rolling and tensile loading directions. As shown in Fig. 6(a) and (b), the large polygon-type particle in the ET34 specimen is broken and the void is observed around the broken particles, while the ET33 specimen exhibited a sound interface without forming any void around the particles. The polygon-type phase particle is crushed during hot-rolling process and the void initiates from the crack during tensile deformation. Therefore, it is clear that the small rod-shaped phase exhibits favorable effect on enhancement of ductility in the Mg-rich Mg–MM–Sn alloy system. 5. Conclusions In the present study the effects of addition of Sn on the microstructure and mechanical properties of Mg–MM alloys have been investigated. The conclusions are as follows: (1) The secondary solidification phase in Mg–MM alloy changes dramatically with addition of Sn, i.e. the Mg12 MM binary phase in Mg–MM alloy changes into the small rodshaped ternary phase with addition of Sn. The stoichiometry of the small rod-shaped particle is close to Mg3 RE1 Sn1 . (2) When the amount of Sn and MM is same (3 wt%), only the small rod-shaped phase forms in the interdendritic region. When the amount of MM (4 wt%) is larger than that of Sn (3 wt%) or when the amount of Sn (4 wt%) is larger than that of MM (3 wt%), the Mg12 MM or polygon-shaped phases coexists with the small rod-shaped phase in the interdendritic region. The size of the small rod-shaped particles (thickness: ∼0.7 ␮m and length: ∼5.6 ␮m) is much smaller than that of the Mg12 MM (∼7 ␮m) or polygon-shaped (∼10 ␮m) particles. (3) The ET43 alloy exhibited the highest yield strength (159.9 MPa), while the ET33 alloy exhibited the largest elongation-to-fracture (9.45%). The result indicates coexistence of the Mg12 MM (and polygon-shaped particle in ET34 alloy) and small rod-shaped phases and smaller grain size in the ET43 and ET34 alloys played a role in enhancing the yield strength of alloys. The result also indicates that

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the small size small rod-shaped particles can enhance the ductility while maintaining the moderate level of strength. Acknowledgments This study was supported by a grant from the Fundamental R&D Program for Core Technology of Materials funded by the Ministry of Commerce, Industry and Energy, Republic of Korea. Hyun Kyu Lim and Ju Youn Lee are grateful for the support from the Second Stage of Brain Korea 21 Project in 2007. References [1] I.J. Polmear, in: B.L. Mordike, F. Hehmann (Eds.), Magnesium Alloys and their Applications, DGM Informationsgesellschaft m.b.H, Oberursel, 1992, pp. 18–36, 201–212. [2] V. G¨artnerov´a, Z. Trojanov´a, A. J¨ager, P. Palˇcek, J. Alloy Compd. 378 (2004) 180–183. [3] L.Y. Wei, G.L. Dunlop, H. Westengen, J. Mater. Sci. 32 (1997) 3334– 3340. [4] C.J. Ma, M.P. Liu, G.H. Wu, W.J. Ding, Y.P. Zhu, Mater. Sci. Eng. A 349 (2003) 207–212.

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