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Glass formation and crystallization behavior in Mg65Cu25Y10−xGdx (x=0, 5 and 10) alloys
Journal of Non-Crystalline Solids 337 (2004) 29–35 www.elsevier.com/locate/jnoncrysol
Glass formation and crystallization behavior in Mg65Cu25Y10xGdx (x ¼ 0, 5 and 10) alloys H. Men a, W.T. Kim b, D.H. Kim a
a,*
Department of Metallurgical Engineering, Center for Non-Crystalline Materials, Yonsei University, 134 Shinchon-dong, Seodaemun-ku, Seoul 120-749, South Korea b Department of Physics, Center for Non-Crystalline Materials, Chongju University, Chongju 360-764, South Korea Received 20 May 2003; received in revised form 8 March 2004
Abstract The glass forming ability and crystallization behavior of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys have been investigated. The glass forming ability (GFA) is significantly improved when Y in Mg65 Cu25 Y10 is substituted with Gd. Ternary Mg65 Cu25 Gd10 bulk metallic glass (BMG) with diameter of at least 8 mm was successfully fabricated by conventional Cu-mold casting method in air atmosphere. Mg2 (Y, Gd) is the first competing crystalline phase against the glass formation in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys, while Mg2 Cu and Cu2 Gd are the competing crystalline phases in the Mg65 Cu25 Gd10 alloy. Therefore, the suppression of the formation of Mg2 (Y, Gd) during cooling from the liquid improves the GFA significantly. Ó 2004 Elsevier B.V. All rights reserved. PACS: 81.05.K; 61.43; 64.70.p
1. Introduction The increased demand for light and strong materials able to withstand severe environmental conditions has stimulated considerable research into bulk metallic glasses (BMGs), due to their significantly improved properties, such as high strength, toughness and good corrosion resistance [1,2]. Among the BMGs developed so far, Mg–TM–RE (TM: transition metal such as Ni, Cu, Zr; RE: rare-earth metal such as Y) alloys are reported to show a large supercooled liquid region, DTx ( ¼ Tx Tg , Tx : crystallization onset temperature, Tg : glass transition temperature) and a high glass forming ability (GFA) [3–8]. Bulk amorphous specimens with a diameter of 4 mm were successfully produced by injection casting Mg65 Cu25 Y10 alloy into a Cu-mold [7]. The critical cooling rate for the amorphous phase formation is about 102 K/s for this alloy system [7]. Further improvement of GFA has been reported in Mg–Cu–Y *
Corresponding author. Tel.: +82-2 361 4255/2 123 4255; fax: +82-2 312 8281. E-mail address: [email protected] (D.H. Kim). 0022-3093/$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2004.03.110
alloy systems where Cu is partially substituted with M (M ¼ Ag, Pd, or Zn). For example, Mg65 Cu15 Ag10 Y10 [9,10], Mg65 Cu20 Zn5 Y10 [11] and Mg65 Cu15 Ag5 Pd5 Y10 [12,13] alloys exhibit high GFA, enabling the fabrication of metallic glass rods with diameters of 6, 6 and 7 mm, respectively, by a Cu-mold injection casting method. In these alloy systems, M has larger negative heat of mixing against other component elements and larger Goldschmidt atomic radius than Cu, satisfying the empirical rules for the formation of BMG [14]. Since the DTx decreases with partial substitution of Cu with M, the improvement of GFA is attributed to the decrease of melting temperature, thus larger value of reduced glass transition temperature, Trg ( ¼ Tg =Tl , Tl : liquidus temperature). In this study we investigated the effect of substitution of Y in Mg65 Cu25 Y10 alloy with Gd on the BMG forming tendency and crystallization behavior. Gd and Y have similar Goldschmidt atomic radius (Gd: 0.180 nm; Y: 0.181 nm) and similar heat of mixing against Mg and Cu (Mg–Gd: )6 J/mol; Cu–Gd: )22 J/mol; Mg–Y: )6 J/mol; Cu–Y: )22 J/mol). The variations of Tg , Tx and melting temperature, Tm with Gd content were
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
studied by the thermal analysis of the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys. Bulk glassy samples were prepared by injection casting into cylinder-shaped Cumolds under controlled inert atmosphere in an evacuated closed chamber or by casting into a cone-shaped Cu-mold in air atmosphere. The competing crystalline phases against the glass formation were investigated by continuous and isothermal heating experiments.
(a)
3. Results Fig. 1(a) and (b) show the XRD patterns taken from injection-cast Mg65 Cu25 Y10 alloy specimens with
u Mg2Y
d=3 mm d=4 mm u v
20
2. Experimental procedure
oo
30
(b) v Mg2Cu
v
v
d=5 mm
40
2 θ, deg.
50
60
u Mg2(Y, Gd)
o Cu2(Y,Gd)
Intensity, a.u.
ribbon d=4 mm d=5 mm u v
20
30
v
v
oo
d=6 mm
40
2 θ, deg.
50
v Mg2Cu
(c)
60
o Cu2 G d
ribbon
Intensity, a.u.
Cu–(Y, Gd) master alloy was prepared by arc melting Cu(99.99%), Y(99.9%), Gd(99.9%) under a Ti-gettered argon atmosphere in a water-cooled copper crucible. The master alloy was then alloyed with Mg(99.9%) in a graphite crucible using an induction furnace. After complete melting, the liquid alloy was poured into a Cumold in air atmosphere. The Cu-mold is cone-shaped with 45 mm in height, 12 mm in diameter at the top and 6 mm in diameter at the bottom. Rapidly solidified ribbon specimens were prepared by re-melting the alloys in quartz tubes, and ejected with an over-pressure of 50 kPa through a nozzle onto a Cu wheel rotating with a surface velocity of 40 m/s. The resulting ribbons have a thickness of about 45 lm and a width of about 2 mm. The injection casting was performed to make bulk samples. Appropriate amounts of each alloy were re-melted in quartz tubes under a purified inert atmosphere and injected through a nozzle into Cu-molds, having cylindrical cavities of varying diameters from 2 to 6 mm. X-ray diffraction (XRD) experiments were performed to identify the formation of the amorphous and crystalline phases in ribbon and bulk samples by using a monochromatic Cu Ka radiation. In particular, for the XRD analysis of the cone-shaped Mg65 Cu25 Gd10 alloy ingot, thin slices were cut from the cast ingot, and the transverse cross sections of about 7.5, 8, 8.5 and 9 mm diameters were examined. Thermal analysis of the ribbon and bulk samples was carried out by differential scanning calorimetry (DSC) and differential thermal analysis (DTA). The DSC and DTA traces were monitored during continuous heating from 323 to 623 K and from 323 to 823 K, respectively, using a constant heating rate of 0.333 K/s. For the isothermal calorimetry, the samples were heated up to 20 K below the target temperature with the highest possible heating rate of 1.67 K/s, and then heated to the isothermal holding temperature with a heating rate of 0.333 K/s. The isothermal holding temperature was maintained with a precision of ±1 K.
o Cu2Y
v Mg2Cu
ribbon
Intensity, a.u.
30
d=7.5 mm d=8 mm oo v
20
30
40
v
2 θ, deg.
d=9 mm
50
60
Fig. 1. XRD patterns taken from: (a) as-melt-spun ribbon and injection cast samples of Mg65 Cu25 Y10 alloy; (b) as-melt-spun ribbon and injection cast samples of Mg65 Cu25 Y5 Gd5 alloy; and (c) as-melt-spun ribbon and conventionally cast cone-shaped samples of Mg65 Cu25 Gd10 alloy.
diameters of 3, 4 and 5 mm and from the injection-cast Mg65 Cu25 Y5 Gd5 alloy specimens with diameters of 4, 5 and 6 mm. For comparison, results from the melt-spun ribbons are also included. Except for the samples with diameter of 5 mm in Fig. 1(a) and of 6 mm in Fig. 1(b), the XRD patterns from the bulk and ribbon specimens showed a broad diffraction peak characteristic of an amorphous structure, indicating that the critical diameters for BMG formation in Mg65 Cu25 Y10 and Mg65 Cu25 Y5 Gd5 are 4 and 5 mm, respectively. The samples with diameter of 5 mm in Fig. 1(a) and of 6 mm in Fig. 1(b) showed sharp diffraction peaks from the crystalline phases superimposed on a broad halo peak, indicating coexistence of crystalline and amorphous phases. Fig. 1(c) shows the XRD patterns taken from
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
the transverse cross-sections of the cone-shaped Mg65 Cu25 Gd10 alloy ingot with diameters of 7.5, 8 and 9 mm, along with the result of melt-spun ribbons. The XRD patterns from the transverse cross-sections up to diameter of 8 mm exhibited a broad diffraction peak, with no evidence of any crystalline peaks, indicating that the Mg65 Cu25 Gd10 alloy has a significantly improved GFA enabling the fabrication of the BMG ingot by conventional Cu-mold casting in air atmosphere. The transverse section with diameter of 9 mm showed sharp diffraction peaks from the crystalline phases superimposed on a broad halo peak, indicating the coexistence of crystalline and amorphous phases. In the partially crystallized samples of Mg65 Cu25 Y10x Gdx (x ¼ 0 and 5) alloys, the crystalline phases were identified as Mg2 Cu, Mg2 (Y, Gd) and Cu2 (Y, Gd) (Fig. 1(a) and (b)). For the partially crystallized Mg65 Cu25 Gd10 alloy, the peaks from Mg2 Cu and Cu2 Gd were present in the XRD pattern (Fig. 1(c)), indicating that the competing crystalline phases against the glass formation are different from those of the Mg65 Cu25 Y10x Gdx (x ¼ 0 and 5) alloys. Fig. 2(a) shows the DSC traces obtained from asmelt-spun ribbons of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys during continuous heating with a heating rate of 0.333 K/s. All the alloys exhibited a distinct glass transition, followed by a broad supercooled liquid re-
TX
Heat flow, a.u.
(a)
x=0
Tg
x=5
x=5 exo.
exo.
Heat flow, a.u.
x=10
525
400
450
550
500
575
550
Temperature, K (b)
T
x=0
sol m liq m
x=5 endo.
∆ T, K
T
600
x=10
650
700
750
800
Temperature, K Fig. 2. (a) DSC; (b) DTA traces obtained from as-melt-spun Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) ribbons during heating with a heating rate of 0.333 K/s.
31
gion and then exothermic reaction due to crystallization. The Tg and Tx are marked by arrows in the DSC traces. With increasing x from 0 to 5, the Tg remained constant, and then decreased slightly from 413 to 408 K with further increase of x from 5 to 10. The Tx increased from 473 K at x ¼ 0 to 486 K at x ¼ 5, and then decreased to 478 K at x ¼ 10. Thus, the DTx increased from about 60 K at x ¼ 0 to 73 K at x ¼ 5 and was nearly constant up to x ¼ 10. Also it can be noticed that the x ¼ 0 alloy exhibited the second exothermic reaction with much less heat compared with that of the first exothermic reaction. The second exothermic reaction became very weak in the x ¼ 5 alloy, and disappeared in the x ¼ 10 alloy. Fig. 2(b) shows the DTA traces obtained from as-melt-spun ribbons of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys during continuous heating with a heating rate of 0.333 K/s. The x ¼ 0 alloy exhibited a single endothermic peak with the melting range of about 30 K. The onset and finishing temperatures of the melting endotherms in the DTA traces were designated by Tmsol and Tmliq , respectively. For the x ¼ 0 alloy, the Tmliq and Tmsol were 760 and 728 K, respectively. It is well known that this alloy consists of a ternary Mg–Cu–Y eutectic composition [15]. Compared with the x ¼ 0 alloy, the Tmliq of the x ¼ 5 and 10 alloys decreased slightly to about 755 K, but Tmsol decreased more significantly, reaching 703 and 681 K, respectively, resulting in a large melting range of about 70 K for the x ¼ 10 alloy. The results of the thermal analysis are summarized in Table 1. It was reported that primary crystallization of ternary Mg65 Cu25 Y10 alloy corresponds to the transformation from the supercooled liquid into a mixture of nanocrystalline Mg2 Cu and a supercooled liquid matrix [15] or polymorphous transformation into Mg2 Cu crystalline phase while Y is soluble in Mg2 Cu [16]. Inoue et al., however, reported that the equilibrium structure of Mg65 Cu25 Y10 alloy consists of mixed phases of Mg2 Y, Mg2 Cu and Cu3 Y [7]. In order to identify the crystallization products of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys, as-melt-spun amorphous ribbons were heated up to various temperatures in the DSC. Fig. 3(a)–(c) show the XRD patterns of the x ¼ 0, 5 and 10 alloys after heat treatment. The diffraction peaks in the XRD pattern (Fig. 3(a)) of the x ¼ 0 alloy heated up to 493 K, i.e. just above the end temperature of the first exothermic reaction, were indexed as Mg2 Cu and Mg2 Y. A weak and broad peak was observed around 2h ¼ 31–40°, indicating the existence of a small amount of nanometer-scale crystalline particles. After heating up to 578 K, i.e. just above the end temperature of the second exothermic reaction, diffraction peaks from Cu2 Y were present together with those from Mg2 Cu and Mg2 Y. For the x ¼ 5 alloy (Fig. 3(b)), Mg2 Cu, Mg2 (Y, Gd) and Cu2 (Y, Gd) were present after heating up to 503 K, i.e. just above the end temperature of first exothermic reaction. After heating up to 603 K, i.e. just above the end temperature
32
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
Table 1 The results of thermal analysis and GFA parameters (DTx and Trg ) for Mg65 Cu25 Y10 Gdx (x ¼ 0, 5, 10) alloys Alloys
Tg (T/K)
Tx (T/K)
Tmsol (T/K)
Tmliq (T/K)
DTx (T/K)
Trg
Mg65 Cu25 Y10 Mg65 Cu25 Y5 Gd5 Mg65 Cu25 Gd10
413 413 408
473 486 478
727 703 681
765 703 681
60 73 70
0.54 0.55 0.54
v
493K
4
v Mg 2Cu o Cu 2Y u Mg 2Y
(a) v
(a)
v u 578K
v
v (b) u
v
503K
o
603K
oo v
v Mg 2Cu o Cu 2(Y, Gd) u Mg 2(Y, Gd)
v
453K
3 2
443K
1
433K 0 0
250 500 750 1000 Time (without incubation time), s
v u v
0.6
u
v Mg 2Cu o Cu Gd
(c)
2
497 K v 603 K un
o ov
u Mg 2Gd n unknown
v
o v v o n
Heat flow, mW/mg
Intensity, a.u.
vu
oo
Heat flow, mW/mg
v u
(b) 453K
0.5 0.4 0.3
443K
0.2 433K
0.1 0.0
20
30
40
50
60
0
70
500 1000 1500 2000 Time (without incubation time), s
2θ, deg.
of the second exothermic reaction, the intensity of the diffraction peaks from Cu2 (Y, Gd) increased, while those from Mg2 Cu and Mg2 (Y, Gd) remained almost same. For the x ¼ 10 alloy (Fig. 3(c)), the supercooled liquid transformed into a mixture of Mg2 Cu and Cu2 Gd after heating up to 497 K, i.e. just above the end temperature of the exothermic reaction. After heating up to higher temperature of 603 K, the XRD pattern showed the diffraction peaks from Mg2 Cu, Cu2 Gd, Mg2 Gd and unidentified phase.Fig. 4(a)–(c) show the results of the isothermal calorimetry experiments taken at the temperatures of 433, 443 and 453 K for Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys. For clarity, the incubation time prior to crystallization is excluded from the plot. The incubation time at the annealing temperatures of 433 K for the x ¼ 0, 5 and 10 alloys were 680, 1677 and 845 s, respectively. The bell-shaped exothermic traces reveal that the first exothermic reaction of all the studied alloys is due to nucleation and growth of crystalline phases [17]. But the traces of the x ¼ 0 alloy exhibit a weak
0.6 Heat flow, mW/mg
Fig. 3. XRD patterns obtained from: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys heated up to the end temperatures of the first and second exothermic peaks.
(c) 453 K
0.5 0.4 0.3 0.2
443 K
0.1
433 K
0.0 0
1000
2000
3000
Time (without incubation time), s Fig. 4. Isothermal DSC traces taken at the holding temperatures of 433, 443 and 453 K for: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys. For clarity, incubation period is excluded in the plot.
exothermic shoulder peak before the main exothermic reaction peak. The enlarged traces (Fig. 5) show that the shape of the exothermic shoulder peaks at different annealing temperatures actually are of no difference, indicating that the crystallization behavior is not influenced by the variation of the annealing temperature. Fig. 6(a) and (b) show the results of the isothermal calorimetry at 433 K for as-melt-spun and as-cast bulk
v Mg 2Cu
u Mg 2Y
(a)
2
u
0.9 ks
u
1.3 ks
1 0 1.2
453K 125
150
175
200
225
0.8
v u
as-spun
v v
2.4 ks
0.4 443K 0.0 0.3
300
20 400
500
50
1000
1250
1500
1750
Fig. 5. Enlarged isothermal DSC traces of melt-spun Mg65 Cu25 Y10 alloy shown in Fig. 4 (a).
v
(c)
30
3.6 ks 40 50 2θ, deg. u Mg 2Gd
v Mg 2Cu
60
70
o Cu 2Gd as-spun
ribbon
0.5
(b)
d=3 mm 1500
Intensity, a. u.
exo.
v oo
(a)
1000
o Cu2 (Y,Gd) as-spun
2.7 ks
u v u
20 1.0
70
1.8 ks Intensity, a. u.
750
60
2θ. deg.
433K
Time, s
0.0 500
40
(b) v Mg 2Cu u Mg 2(Y,Gd)
0.1 0.0
30
600
0.2
Heat flow, mW/mg
33
3
Intensity, a.u.
Heat flow, mW/mg
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
oo v
v
1.5 ks
v u
oo v
v
2.1 ks
v u
oo v
v
3.9 ks
2000
Time, s Fig. 6. Isothermal DSC traces taken at the holding temperature of 433 K for the Mg65 Cu25 Y10 alloy: (a) as-melt-spun ribbon; and (b) injection-cast sample with d ¼ 3 mm.
(d ¼ 3 mm) x ¼ 0 alloy samples, respectively. The two traces exhibit almost same shapes, indicating that the crystallization behavior is not influenced by the change of cooling rate. The DSC result also showed that the Tg and Tx from the two samples were almost same within the experimental error. With the increase of x from 0 to 5, the heat for the exothermic shoulder peak decreased significantly, but the traces were still asymmetric (Fig. 4(b)). For the x ¼ 10 alloy, no shoulder peak was observed, and the exothermic peak exhibited a symmetric bell-shape (Fig. 4(c)). Fig. 7(a)–(c) show the XRD patterns of x ¼ 0, 5 and 10 alloys heat treated isothermally at 433 K, along with the XRD pattern of the as-melt-spun ribbons. For the x ¼ 0 alloy (Fig. 7(a)), Mg2 Y formed first in the amorphous matrix, as can be seen in the XRD patterns obtained after isothermal holding for 0.9 and 1.3 ks. After
20
30
40
50
60
70
2θ, deg. Fig. 7. XRD patterns of: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys heat treated isothermally at 433 K.
primary crystallization of Mg2 Y, Mg2 Cu formed in the remaining matrix with further isothermal heat treatment. Formation of small amount of nanometer-scale Cu2 Y also could be noticed from the XRD pattern after isothermal holding for 2.4 ks. For the x ¼ 5 alloy (Fig. 7(b)), no sharp diffraction peaks from the crystalline phase was observed after isothermal holding up to 2.7 ks, due to higher crystallization temperature than those of the x ¼ 0 and 10 alloys. After isothermal holding for 3.6 ks, the supercooled liquid crystallized into Mg2 (Y, Gd), Mg2 Cu and Cu2 (Y, Gd). The relative intensity of the diffraction peaks was similar to that of x ¼ 0 alloy. For the x ¼ 10 alloy (Fig. 7(c)), the crystallization behavior during isothermal heat treatment was different from those of the x ¼ 0 and 5 alloys. After isothermal
34
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
holding for 1.5 ks, Mg2 Cu and Cu2 Gd formed first in the amorphous matrix. With further isothermal holding at 433 K, Mg2 Gd formed, as can be seen in the XRD patterns obtained after isothermal holding for 2.1 and 3.9 ks.
4. Discussion Present study shows that the GFA is significantly improved by substituting Y in Mg65 Cu25 Y10 with Gd. Ternary Mg65 Cu25 Gd10 BMG with diameter of at least 8 mm is successfully fabricated by conventional Cumold casting method in air atmosphere (Fig. 1). The primary crystallization of the amorphous Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys exhibits a sharp exothermic peak during continuous heating in the DSC (Fig. 2). However, during isothermal annealing of the Y containing Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys in the temperature region between Tg and Tx (Figs. 4(a), (b) and 7(a), (b)), Mg2 (Y, Gd) precipitates first in the amorphous matrix after the incubation period. Due to the primary crystallization of Mg2 (Y, Gd), the exothermic shoulder peak appears in the DSC trace. The shoulder peak becomes weak with increasing Gd content from x ¼ 0 to 5. Therefore, Mg2 (Y, Gd) is the competing crystalline phase against the glass formation in Y containing Mg65 Cu25 Y10x Gdx (x ¼ 0) alloy. For the x ¼ 5 alloy, formation of Mg2 (Y, Gd) is partially suppressed, as evidenced by the decrease of the shoulder peak. For the Y-free Mg65 Cu25 Gd10 alloy (Figs. 4(c) and 7(c)), the shoulder peak corresponding to the formation of Mg2 Gd does not appear in the isothermal DSC traces. Fig. 3(c) confirms that the diffraction peak corresponding to Mg2 Gd does not appear in the XRD pattern obtained after continuous heating up to the end temperature of the exothermic reaction, indicating that the supercooled liquid of the Mg65 Cu25 Gd10 alloy crystallizes in a different path from that of the Mg65 Cu25 Y10 alloy, i.e. by a simultaneous formation of Mg2 Cu and Cu2 Gd. Thus, Mg2 Cu and Cu2 Gd are the competing crystalline phases against the glass formation in the Mg65 Cu25 Gd10 alloy. Consequently, with increasing x in Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys the GFA gradually increases by the suppression of Mg2 (Y, Gd) formation during cooling from the liquid. As mentioned above, ternary Mg65 Cu25 Gd10 alloy shows a surprisingly high GFA. From an alloy chemistry point of view, Gd and Y are nearly identical. The negative heat of mixing in liquid and the difference of atomic size between component elements in Mg–Cu–Y and Mg–Cu–Gd alloy systems are nearly the same, and thus the degree of dense random packed structure in the undercooled liquid should remain almost unchanged. Generally, the nucleation of crystalline phases from the undercooled liquid phase requires substantial
atomic diffusion and composition redistribution. Gd and Y have nearly same Goldschmidt atomic radius, and thus the diffusivity of Gd and Y in the liquid should not show significant difference. Therefore, Mg65 Cu25 Gd10 alloy exhibits only a slightly larger DTx than Mg65 Cu25 Y10 alloy. The reduced glass transition Trg , defined as Tg =Tl , has been suggested as an indicator of the GFA. For Mg65 Cu25 Y10x Gdx alloy system, Trg increases slightly from 0.54 to 0.55 with the increase of x from 0 to 5 and then again decreases to 0.54 with further increasing x to 10, i.e. the Mg65 Cu25 Gd10 and Mg65 Cu25 Y10 alloys have almost same value of Trg , although there is a much difference in GFA. The critical cooling rate for glass formation (Rc ) in the Mg65 Cu25 Gd10 alloy is estimated to be about 1 K/s [18], which is several orders of magnitude smaller than that of the Mg65 Cu25 Y10 alloy. It has been pointed out that the nucleation rate should be vanishingly small for Trg ¼ 2=3, which causes kinetically very sluggish crystallization process [19]. In fact, the other BMG alloys with Rc of 1 K/s or below exhibit Trg values near 2/3 [20,21]. The electronic band structure of metallic glasses is known to be dependent on the type of short-range order [22]. In Ni–Fe–B metallic glasses, it has been shown theoretically [23] and experimentally [24–26] that the chemical ordering occurs by the preferential bonding interaction between B and Ni due to the difference in the electronic structure of Ni (3d8 4s2 ) and Fe (3d6 4s2 ), although Ni and Fe have similar Goldschmidt atomic radius (Ni: 0.125 nm, Fe: 0.128 nm) and heat of mixing with B (Ni–B: )9 kJ/mol, Fe–B: )11 kJ/mol). Thus, the difference in electronic configuration between Y (4d1 5s2 ) and Gd (4f7 5d1 6s2 ) may result in certain change of the short-range order in undercooled liquid of Mg65 Cu25 Y10x Gdx alloys, which causes the different crystallization route, and favors the glass formation.
5. Conclusion Ternary Mg65 Cu25 Gd10 BMG with diameter of at least 8 mm was successfully fabricated by conventional Cu-mold casting method in air atmosphere. When compared with the GFA of Mg65 Cu25 Y10 alloy, the significant improvement of the GFA of Mg65 Cu25 Gd10 alloy cannot be interpreted by the DTx and Trg values. The isothermal DSC experiments shows that Mg2 (Y, Gd) is the competing crystalline phase against the glass formation in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys, while Mg2 Cu and Cu2 Gd are the competing crystalline phases against the glass formation in the Mg65 Cu25 Gd10 alloy. Therefore, in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys, the suppression of the formation of Mg2 (Y, Gd) during cooling from the liquid improves the GFA significantly.
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
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Glass formation and crystallization behavior in Mg65Cu25Y10xGdx (x ¼ 0, 5 and 10) alloys H. Men a, W.T. Kim b, D.H. Kim a
a,*
Department of Metallurgical Engineering, Center for Non-Crystalline Materials, Yonsei University, 134 Shinchon-dong, Seodaemun-ku, Seoul 120-749, South Korea b Department of Physics, Center for Non-Crystalline Materials, Chongju University, Chongju 360-764, South Korea Received 20 May 2003; received in revised form 8 March 2004
Abstract The glass forming ability and crystallization behavior of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys have been investigated. The glass forming ability (GFA) is significantly improved when Y in Mg65 Cu25 Y10 is substituted with Gd. Ternary Mg65 Cu25 Gd10 bulk metallic glass (BMG) with diameter of at least 8 mm was successfully fabricated by conventional Cu-mold casting method in air atmosphere. Mg2 (Y, Gd) is the first competing crystalline phase against the glass formation in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys, while Mg2 Cu and Cu2 Gd are the competing crystalline phases in the Mg65 Cu25 Gd10 alloy. Therefore, the suppression of the formation of Mg2 (Y, Gd) during cooling from the liquid improves the GFA significantly. Ó 2004 Elsevier B.V. All rights reserved. PACS: 81.05.K; 61.43; 64.70.p
1. Introduction The increased demand for light and strong materials able to withstand severe environmental conditions has stimulated considerable research into bulk metallic glasses (BMGs), due to their significantly improved properties, such as high strength, toughness and good corrosion resistance [1,2]. Among the BMGs developed so far, Mg–TM–RE (TM: transition metal such as Ni, Cu, Zr; RE: rare-earth metal such as Y) alloys are reported to show a large supercooled liquid region, DTx ( ¼ Tx Tg , Tx : crystallization onset temperature, Tg : glass transition temperature) and a high glass forming ability (GFA) [3–8]. Bulk amorphous specimens with a diameter of 4 mm were successfully produced by injection casting Mg65 Cu25 Y10 alloy into a Cu-mold [7]. The critical cooling rate for the amorphous phase formation is about 102 K/s for this alloy system [7]. Further improvement of GFA has been reported in Mg–Cu–Y *
Corresponding author. Tel.: +82-2 361 4255/2 123 4255; fax: +82-2 312 8281. E-mail address: [email protected] (D.H. Kim). 0022-3093/$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2004.03.110
alloy systems where Cu is partially substituted with M (M ¼ Ag, Pd, or Zn). For example, Mg65 Cu15 Ag10 Y10 [9,10], Mg65 Cu20 Zn5 Y10 [11] and Mg65 Cu15 Ag5 Pd5 Y10 [12,13] alloys exhibit high GFA, enabling the fabrication of metallic glass rods with diameters of 6, 6 and 7 mm, respectively, by a Cu-mold injection casting method. In these alloy systems, M has larger negative heat of mixing against other component elements and larger Goldschmidt atomic radius than Cu, satisfying the empirical rules for the formation of BMG [14]. Since the DTx decreases with partial substitution of Cu with M, the improvement of GFA is attributed to the decrease of melting temperature, thus larger value of reduced glass transition temperature, Trg ( ¼ Tg =Tl , Tl : liquidus temperature). In this study we investigated the effect of substitution of Y in Mg65 Cu25 Y10 alloy with Gd on the BMG forming tendency and crystallization behavior. Gd and Y have similar Goldschmidt atomic radius (Gd: 0.180 nm; Y: 0.181 nm) and similar heat of mixing against Mg and Cu (Mg–Gd: )6 J/mol; Cu–Gd: )22 J/mol; Mg–Y: )6 J/mol; Cu–Y: )22 J/mol). The variations of Tg , Tx and melting temperature, Tm with Gd content were
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
studied by the thermal analysis of the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys. Bulk glassy samples were prepared by injection casting into cylinder-shaped Cumolds under controlled inert atmosphere in an evacuated closed chamber or by casting into a cone-shaped Cu-mold in air atmosphere. The competing crystalline phases against the glass formation were investigated by continuous and isothermal heating experiments.
(a)
3. Results Fig. 1(a) and (b) show the XRD patterns taken from injection-cast Mg65 Cu25 Y10 alloy specimens with
u Mg2Y
d=3 mm d=4 mm u v
20
2. Experimental procedure
oo
30
(b) v Mg2Cu
v
v
d=5 mm
40
2 θ, deg.
50
60
u Mg2(Y, Gd)
o Cu2(Y,Gd)
Intensity, a.u.
ribbon d=4 mm d=5 mm u v
20
30
v
v
oo
d=6 mm
40
2 θ, deg.
50
v Mg2Cu
(c)
60
o Cu2 G d
ribbon
Intensity, a.u.
Cu–(Y, Gd) master alloy was prepared by arc melting Cu(99.99%), Y(99.9%), Gd(99.9%) under a Ti-gettered argon atmosphere in a water-cooled copper crucible. The master alloy was then alloyed with Mg(99.9%) in a graphite crucible using an induction furnace. After complete melting, the liquid alloy was poured into a Cumold in air atmosphere. The Cu-mold is cone-shaped with 45 mm in height, 12 mm in diameter at the top and 6 mm in diameter at the bottom. Rapidly solidified ribbon specimens were prepared by re-melting the alloys in quartz tubes, and ejected with an over-pressure of 50 kPa through a nozzle onto a Cu wheel rotating with a surface velocity of 40 m/s. The resulting ribbons have a thickness of about 45 lm and a width of about 2 mm. The injection casting was performed to make bulk samples. Appropriate amounts of each alloy were re-melted in quartz tubes under a purified inert atmosphere and injected through a nozzle into Cu-molds, having cylindrical cavities of varying diameters from 2 to 6 mm. X-ray diffraction (XRD) experiments were performed to identify the formation of the amorphous and crystalline phases in ribbon and bulk samples by using a monochromatic Cu Ka radiation. In particular, for the XRD analysis of the cone-shaped Mg65 Cu25 Gd10 alloy ingot, thin slices were cut from the cast ingot, and the transverse cross sections of about 7.5, 8, 8.5 and 9 mm diameters were examined. Thermal analysis of the ribbon and bulk samples was carried out by differential scanning calorimetry (DSC) and differential thermal analysis (DTA). The DSC and DTA traces were monitored during continuous heating from 323 to 623 K and from 323 to 823 K, respectively, using a constant heating rate of 0.333 K/s. For the isothermal calorimetry, the samples were heated up to 20 K below the target temperature with the highest possible heating rate of 1.67 K/s, and then heated to the isothermal holding temperature with a heating rate of 0.333 K/s. The isothermal holding temperature was maintained with a precision of ±1 K.
o Cu2Y
v Mg2Cu
ribbon
Intensity, a.u.
30
d=7.5 mm d=8 mm oo v
20
30
40
v
2 θ, deg.
d=9 mm
50
60
Fig. 1. XRD patterns taken from: (a) as-melt-spun ribbon and injection cast samples of Mg65 Cu25 Y10 alloy; (b) as-melt-spun ribbon and injection cast samples of Mg65 Cu25 Y5 Gd5 alloy; and (c) as-melt-spun ribbon and conventionally cast cone-shaped samples of Mg65 Cu25 Gd10 alloy.
diameters of 3, 4 and 5 mm and from the injection-cast Mg65 Cu25 Y5 Gd5 alloy specimens with diameters of 4, 5 and 6 mm. For comparison, results from the melt-spun ribbons are also included. Except for the samples with diameter of 5 mm in Fig. 1(a) and of 6 mm in Fig. 1(b), the XRD patterns from the bulk and ribbon specimens showed a broad diffraction peak characteristic of an amorphous structure, indicating that the critical diameters for BMG formation in Mg65 Cu25 Y10 and Mg65 Cu25 Y5 Gd5 are 4 and 5 mm, respectively. The samples with diameter of 5 mm in Fig. 1(a) and of 6 mm in Fig. 1(b) showed sharp diffraction peaks from the crystalline phases superimposed on a broad halo peak, indicating coexistence of crystalline and amorphous phases. Fig. 1(c) shows the XRD patterns taken from
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
the transverse cross-sections of the cone-shaped Mg65 Cu25 Gd10 alloy ingot with diameters of 7.5, 8 and 9 mm, along with the result of melt-spun ribbons. The XRD patterns from the transverse cross-sections up to diameter of 8 mm exhibited a broad diffraction peak, with no evidence of any crystalline peaks, indicating that the Mg65 Cu25 Gd10 alloy has a significantly improved GFA enabling the fabrication of the BMG ingot by conventional Cu-mold casting in air atmosphere. The transverse section with diameter of 9 mm showed sharp diffraction peaks from the crystalline phases superimposed on a broad halo peak, indicating the coexistence of crystalline and amorphous phases. In the partially crystallized samples of Mg65 Cu25 Y10x Gdx (x ¼ 0 and 5) alloys, the crystalline phases were identified as Mg2 Cu, Mg2 (Y, Gd) and Cu2 (Y, Gd) (Fig. 1(a) and (b)). For the partially crystallized Mg65 Cu25 Gd10 alloy, the peaks from Mg2 Cu and Cu2 Gd were present in the XRD pattern (Fig. 1(c)), indicating that the competing crystalline phases against the glass formation are different from those of the Mg65 Cu25 Y10x Gdx (x ¼ 0 and 5) alloys. Fig. 2(a) shows the DSC traces obtained from asmelt-spun ribbons of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys during continuous heating with a heating rate of 0.333 K/s. All the alloys exhibited a distinct glass transition, followed by a broad supercooled liquid re-
TX
Heat flow, a.u.
(a)
x=0
Tg
x=5
x=5 exo.
exo.
Heat flow, a.u.
x=10
525
400
450
550
500
575
550
Temperature, K (b)
T
x=0
sol m liq m
x=5 endo.
∆ T, K
T
600
x=10
650
700
750
800
Temperature, K Fig. 2. (a) DSC; (b) DTA traces obtained from as-melt-spun Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) ribbons during heating with a heating rate of 0.333 K/s.
31
gion and then exothermic reaction due to crystallization. The Tg and Tx are marked by arrows in the DSC traces. With increasing x from 0 to 5, the Tg remained constant, and then decreased slightly from 413 to 408 K with further increase of x from 5 to 10. The Tx increased from 473 K at x ¼ 0 to 486 K at x ¼ 5, and then decreased to 478 K at x ¼ 10. Thus, the DTx increased from about 60 K at x ¼ 0 to 73 K at x ¼ 5 and was nearly constant up to x ¼ 10. Also it can be noticed that the x ¼ 0 alloy exhibited the second exothermic reaction with much less heat compared with that of the first exothermic reaction. The second exothermic reaction became very weak in the x ¼ 5 alloy, and disappeared in the x ¼ 10 alloy. Fig. 2(b) shows the DTA traces obtained from as-melt-spun ribbons of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys during continuous heating with a heating rate of 0.333 K/s. The x ¼ 0 alloy exhibited a single endothermic peak with the melting range of about 30 K. The onset and finishing temperatures of the melting endotherms in the DTA traces were designated by Tmsol and Tmliq , respectively. For the x ¼ 0 alloy, the Tmliq and Tmsol were 760 and 728 K, respectively. It is well known that this alloy consists of a ternary Mg–Cu–Y eutectic composition [15]. Compared with the x ¼ 0 alloy, the Tmliq of the x ¼ 5 and 10 alloys decreased slightly to about 755 K, but Tmsol decreased more significantly, reaching 703 and 681 K, respectively, resulting in a large melting range of about 70 K for the x ¼ 10 alloy. The results of the thermal analysis are summarized in Table 1. It was reported that primary crystallization of ternary Mg65 Cu25 Y10 alloy corresponds to the transformation from the supercooled liquid into a mixture of nanocrystalline Mg2 Cu and a supercooled liquid matrix [15] or polymorphous transformation into Mg2 Cu crystalline phase while Y is soluble in Mg2 Cu [16]. Inoue et al., however, reported that the equilibrium structure of Mg65 Cu25 Y10 alloy consists of mixed phases of Mg2 Y, Mg2 Cu and Cu3 Y [7]. In order to identify the crystallization products of Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys, as-melt-spun amorphous ribbons were heated up to various temperatures in the DSC. Fig. 3(a)–(c) show the XRD patterns of the x ¼ 0, 5 and 10 alloys after heat treatment. The diffraction peaks in the XRD pattern (Fig. 3(a)) of the x ¼ 0 alloy heated up to 493 K, i.e. just above the end temperature of the first exothermic reaction, were indexed as Mg2 Cu and Mg2 Y. A weak and broad peak was observed around 2h ¼ 31–40°, indicating the existence of a small amount of nanometer-scale crystalline particles. After heating up to 578 K, i.e. just above the end temperature of the second exothermic reaction, diffraction peaks from Cu2 Y were present together with those from Mg2 Cu and Mg2 Y. For the x ¼ 5 alloy (Fig. 3(b)), Mg2 Cu, Mg2 (Y, Gd) and Cu2 (Y, Gd) were present after heating up to 503 K, i.e. just above the end temperature of first exothermic reaction. After heating up to 603 K, i.e. just above the end temperature
32
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
Table 1 The results of thermal analysis and GFA parameters (DTx and Trg ) for Mg65 Cu25 Y10 Gdx (x ¼ 0, 5, 10) alloys Alloys
Tg (T/K)
Tx (T/K)
Tmsol (T/K)
Tmliq (T/K)
DTx (T/K)
Trg
Mg65 Cu25 Y10 Mg65 Cu25 Y5 Gd5 Mg65 Cu25 Gd10
413 413 408
473 486 478
727 703 681
765 703 681
60 73 70
0.54 0.55 0.54
v
493K
4
v Mg 2Cu o Cu 2Y u Mg 2Y
(a) v
(a)
v u 578K
v
v (b) u
v
503K
o
603K
oo v
v Mg 2Cu o Cu 2(Y, Gd) u Mg 2(Y, Gd)
v
453K
3 2
443K
1
433K 0 0
250 500 750 1000 Time (without incubation time), s
v u v
0.6
u
v Mg 2Cu o Cu Gd
(c)
2
497 K v 603 K un
o ov
u Mg 2Gd n unknown
v
o v v o n
Heat flow, mW/mg
Intensity, a.u.
vu
oo
Heat flow, mW/mg
v u
(b) 453K
0.5 0.4 0.3
443K
0.2 433K
0.1 0.0
20
30
40
50
60
0
70
500 1000 1500 2000 Time (without incubation time), s
2θ, deg.
of the second exothermic reaction, the intensity of the diffraction peaks from Cu2 (Y, Gd) increased, while those from Mg2 Cu and Mg2 (Y, Gd) remained almost same. For the x ¼ 10 alloy (Fig. 3(c)), the supercooled liquid transformed into a mixture of Mg2 Cu and Cu2 Gd after heating up to 497 K, i.e. just above the end temperature of the exothermic reaction. After heating up to higher temperature of 603 K, the XRD pattern showed the diffraction peaks from Mg2 Cu, Cu2 Gd, Mg2 Gd and unidentified phase.Fig. 4(a)–(c) show the results of the isothermal calorimetry experiments taken at the temperatures of 433, 443 and 453 K for Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys. For clarity, the incubation time prior to crystallization is excluded from the plot. The incubation time at the annealing temperatures of 433 K for the x ¼ 0, 5 and 10 alloys were 680, 1677 and 845 s, respectively. The bell-shaped exothermic traces reveal that the first exothermic reaction of all the studied alloys is due to nucleation and growth of crystalline phases [17]. But the traces of the x ¼ 0 alloy exhibit a weak
0.6 Heat flow, mW/mg
Fig. 3. XRD patterns obtained from: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys heated up to the end temperatures of the first and second exothermic peaks.
(c) 453 K
0.5 0.4 0.3 0.2
443 K
0.1
433 K
0.0 0
1000
2000
3000
Time (without incubation time), s Fig. 4. Isothermal DSC traces taken at the holding temperatures of 433, 443 and 453 K for: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys. For clarity, incubation period is excluded in the plot.
exothermic shoulder peak before the main exothermic reaction peak. The enlarged traces (Fig. 5) show that the shape of the exothermic shoulder peaks at different annealing temperatures actually are of no difference, indicating that the crystallization behavior is not influenced by the variation of the annealing temperature. Fig. 6(a) and (b) show the results of the isothermal calorimetry at 433 K for as-melt-spun and as-cast bulk
v Mg 2Cu
u Mg 2Y
(a)
2
u
0.9 ks
u
1.3 ks
1 0 1.2
453K 125
150
175
200
225
0.8
v u
as-spun
v v
2.4 ks
0.4 443K 0.0 0.3
300
20 400
500
50
1000
1250
1500
1750
Fig. 5. Enlarged isothermal DSC traces of melt-spun Mg65 Cu25 Y10 alloy shown in Fig. 4 (a).
v
(c)
30
3.6 ks 40 50 2θ, deg. u Mg 2Gd
v Mg 2Cu
60
70
o Cu 2Gd as-spun
ribbon
0.5
(b)
d=3 mm 1500
Intensity, a. u.
exo.
v oo
(a)
1000
o Cu2 (Y,Gd) as-spun
2.7 ks
u v u
20 1.0
70
1.8 ks Intensity, a. u.
750
60
2θ. deg.
433K
Time, s
0.0 500
40
(b) v Mg 2Cu u Mg 2(Y,Gd)
0.1 0.0
30
600
0.2
Heat flow, mW/mg
33
3
Intensity, a.u.
Heat flow, mW/mg
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
oo v
v
1.5 ks
v u
oo v
v
2.1 ks
v u
oo v
v
3.9 ks
2000
Time, s Fig. 6. Isothermal DSC traces taken at the holding temperature of 433 K for the Mg65 Cu25 Y10 alloy: (a) as-melt-spun ribbon; and (b) injection-cast sample with d ¼ 3 mm.
(d ¼ 3 mm) x ¼ 0 alloy samples, respectively. The two traces exhibit almost same shapes, indicating that the crystallization behavior is not influenced by the change of cooling rate. The DSC result also showed that the Tg and Tx from the two samples were almost same within the experimental error. With the increase of x from 0 to 5, the heat for the exothermic shoulder peak decreased significantly, but the traces were still asymmetric (Fig. 4(b)). For the x ¼ 10 alloy, no shoulder peak was observed, and the exothermic peak exhibited a symmetric bell-shape (Fig. 4(c)). Fig. 7(a)–(c) show the XRD patterns of x ¼ 0, 5 and 10 alloys heat treated isothermally at 433 K, along with the XRD pattern of the as-melt-spun ribbons. For the x ¼ 0 alloy (Fig. 7(a)), Mg2 Y formed first in the amorphous matrix, as can be seen in the XRD patterns obtained after isothermal holding for 0.9 and 1.3 ks. After
20
30
40
50
60
70
2θ, deg. Fig. 7. XRD patterns of: (a) melt-spun Mg65 Cu25 Y10 ; (b) melt-spun Mg65 Cu25 Y5 Gd5 ; and (c) melt-spun Mg65 Cu25 Gd10 alloys heat treated isothermally at 433 K.
primary crystallization of Mg2 Y, Mg2 Cu formed in the remaining matrix with further isothermal heat treatment. Formation of small amount of nanometer-scale Cu2 Y also could be noticed from the XRD pattern after isothermal holding for 2.4 ks. For the x ¼ 5 alloy (Fig. 7(b)), no sharp diffraction peaks from the crystalline phase was observed after isothermal holding up to 2.7 ks, due to higher crystallization temperature than those of the x ¼ 0 and 10 alloys. After isothermal holding for 3.6 ks, the supercooled liquid crystallized into Mg2 (Y, Gd), Mg2 Cu and Cu2 (Y, Gd). The relative intensity of the diffraction peaks was similar to that of x ¼ 0 alloy. For the x ¼ 10 alloy (Fig. 7(c)), the crystallization behavior during isothermal heat treatment was different from those of the x ¼ 0 and 5 alloys. After isothermal
34
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
holding for 1.5 ks, Mg2 Cu and Cu2 Gd formed first in the amorphous matrix. With further isothermal holding at 433 K, Mg2 Gd formed, as can be seen in the XRD patterns obtained after isothermal holding for 2.1 and 3.9 ks.
4. Discussion Present study shows that the GFA is significantly improved by substituting Y in Mg65 Cu25 Y10 with Gd. Ternary Mg65 Cu25 Gd10 BMG with diameter of at least 8 mm is successfully fabricated by conventional Cumold casting method in air atmosphere (Fig. 1). The primary crystallization of the amorphous Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys exhibits a sharp exothermic peak during continuous heating in the DSC (Fig. 2). However, during isothermal annealing of the Y containing Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys in the temperature region between Tg and Tx (Figs. 4(a), (b) and 7(a), (b)), Mg2 (Y, Gd) precipitates first in the amorphous matrix after the incubation period. Due to the primary crystallization of Mg2 (Y, Gd), the exothermic shoulder peak appears in the DSC trace. The shoulder peak becomes weak with increasing Gd content from x ¼ 0 to 5. Therefore, Mg2 (Y, Gd) is the competing crystalline phase against the glass formation in Y containing Mg65 Cu25 Y10x Gdx (x ¼ 0) alloy. For the x ¼ 5 alloy, formation of Mg2 (Y, Gd) is partially suppressed, as evidenced by the decrease of the shoulder peak. For the Y-free Mg65 Cu25 Gd10 alloy (Figs. 4(c) and 7(c)), the shoulder peak corresponding to the formation of Mg2 Gd does not appear in the isothermal DSC traces. Fig. 3(c) confirms that the diffraction peak corresponding to Mg2 Gd does not appear in the XRD pattern obtained after continuous heating up to the end temperature of the exothermic reaction, indicating that the supercooled liquid of the Mg65 Cu25 Gd10 alloy crystallizes in a different path from that of the Mg65 Cu25 Y10 alloy, i.e. by a simultaneous formation of Mg2 Cu and Cu2 Gd. Thus, Mg2 Cu and Cu2 Gd are the competing crystalline phases against the glass formation in the Mg65 Cu25 Gd10 alloy. Consequently, with increasing x in Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys the GFA gradually increases by the suppression of Mg2 (Y, Gd) formation during cooling from the liquid. As mentioned above, ternary Mg65 Cu25 Gd10 alloy shows a surprisingly high GFA. From an alloy chemistry point of view, Gd and Y are nearly identical. The negative heat of mixing in liquid and the difference of atomic size between component elements in Mg–Cu–Y and Mg–Cu–Gd alloy systems are nearly the same, and thus the degree of dense random packed structure in the undercooled liquid should remain almost unchanged. Generally, the nucleation of crystalline phases from the undercooled liquid phase requires substantial
atomic diffusion and composition redistribution. Gd and Y have nearly same Goldschmidt atomic radius, and thus the diffusivity of Gd and Y in the liquid should not show significant difference. Therefore, Mg65 Cu25 Gd10 alloy exhibits only a slightly larger DTx than Mg65 Cu25 Y10 alloy. The reduced glass transition Trg , defined as Tg =Tl , has been suggested as an indicator of the GFA. For Mg65 Cu25 Y10x Gdx alloy system, Trg increases slightly from 0.54 to 0.55 with the increase of x from 0 to 5 and then again decreases to 0.54 with further increasing x to 10, i.e. the Mg65 Cu25 Gd10 and Mg65 Cu25 Y10 alloys have almost same value of Trg , although there is a much difference in GFA. The critical cooling rate for glass formation (Rc ) in the Mg65 Cu25 Gd10 alloy is estimated to be about 1 K/s [18], which is several orders of magnitude smaller than that of the Mg65 Cu25 Y10 alloy. It has been pointed out that the nucleation rate should be vanishingly small for Trg ¼ 2=3, which causes kinetically very sluggish crystallization process [19]. In fact, the other BMG alloys with Rc of 1 K/s or below exhibit Trg values near 2/3 [20,21]. The electronic band structure of metallic glasses is known to be dependent on the type of short-range order [22]. In Ni–Fe–B metallic glasses, it has been shown theoretically [23] and experimentally [24–26] that the chemical ordering occurs by the preferential bonding interaction between B and Ni due to the difference in the electronic structure of Ni (3d8 4s2 ) and Fe (3d6 4s2 ), although Ni and Fe have similar Goldschmidt atomic radius (Ni: 0.125 nm, Fe: 0.128 nm) and heat of mixing with B (Ni–B: )9 kJ/mol, Fe–B: )11 kJ/mol). Thus, the difference in electronic configuration between Y (4d1 5s2 ) and Gd (4f7 5d1 6s2 ) may result in certain change of the short-range order in undercooled liquid of Mg65 Cu25 Y10x Gdx alloys, which causes the different crystallization route, and favors the glass formation.
5. Conclusion Ternary Mg65 Cu25 Gd10 BMG with diameter of at least 8 mm was successfully fabricated by conventional Cu-mold casting method in air atmosphere. When compared with the GFA of Mg65 Cu25 Y10 alloy, the significant improvement of the GFA of Mg65 Cu25 Gd10 alloy cannot be interpreted by the DTx and Trg values. The isothermal DSC experiments shows that Mg2 (Y, Gd) is the competing crystalline phase against the glass formation in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5) alloys, while Mg2 Cu and Cu2 Gd are the competing crystalline phases against the glass formation in the Mg65 Cu25 Gd10 alloy. Therefore, in the Mg65 Cu25 Y10x Gdx (x ¼ 0, 5 and 10) alloys, the suppression of the formation of Mg2 (Y, Gd) during cooling from the liquid improves the GFA significantly.
H. Men et al. / Journal of Non-Crystalline Solids 337 (2004) 29–35
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