Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy

Transcription

1 Materials Transactions, Vol. 48, No. 6 (2007) pp to 1372 Special Issue on Materials Science of Bulk Metallic Glasses-VII #2007 The Japan Institute of Metals Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy Yoshihiko Yokoyama 1, Hasse Fredriksson 2, Hideyuki Yasuda 3, Masahiko Nishijima 1 and Akihisa Inoue 1 1 Institute for Materials Research, Tohoku University, Sendai , Japan 2 Department of Materials Processing Casting of Metals, SE Stockholm, Addr. Brinellvagen 23, Sweden 3 Department of Adaptive Machine Systems, Osaka University, Osaka , Japan We examined the solidification morphology and structure of arc-melted Zr 50 Cu 40 Al 10 glass-forming alloys in order to determine the impurity influences using two grades of Zr metals: sponge Zr purified somewhat but not highly by the Kroll method, and highly purified crystal Zr. When crystal Zr is used, arc-melted Zr 50 Cu 40 Al 10 alloy exhibits superior glass-forming ability in forming glassy phase, even in a 40-g master alloy. When sponge Zr is used, on the other hand, we can see distinct a chain reaction of exothermic heat due to crystallization after vitrification during solidification. We conclude that the origin of the crystallization in arc-melted Zr 50 Cu 40 Al 10 alloy with sponge Zr is probably chlorine as an impurity in sponge Zr metals. Furthermore, vitrification in front of the solidification interface of arc-melted Zr 50 Cu 40 Al 10 alloy with crystal Zr can occur when the crystalline growth phase is an Al-supersaturated B2 (B19 )-type ZrCu phase. [doi: /matertrans.mf200624] (Received March 20, 2007; Accepted May 1, 2007; Published May 25, 2007) Keywords: zirconium-copper-aluminum alloy system, ternary eutectic point, sponge and crystal zirconium, solidification structure 1. Introduction Bulk glassy alloys are regarded as new materials with good glass-forming ability and excellent mechanical properties. Since the glassy alloys are characterized by aperiodic random structure and flexible metallic bonding, we can expect the development of ductile, high-toughness and high-strength bulk glassy alloys in various alloy systems. 1) However, the superior mechanical properties of bulk glassy alloys degrade significantly when crystalline inclusions are present. 2) The purity 3) and complete melting technique 4) are important in realizing the innate glass-forming ability of such alloys. In order to avoid the formation of crystalline inclusions in bulk metallic glasses, the preparation process is the most important factor, since it is this process must accomplish a complete melting state before casting. We have succeeded in developing melting and casting processes to avoid the formation of crystalline inclusions before casting, specifically, the cold copper nozzle-type 5) and ladle hearth-type 6) arc-casting methods. In these processes, all crucibles and nozzles are made of water-cooled copper. Additionally, it is important to control the purity of the Zr-Cu-Al alloy in order to avoid the formation of crystalline inclusions in cast Zr-Cu- Al bulk glassy alloys. By using crystal Zr metal, we can prepare glassy cast rods (12 mm 60 mm) in ternary Zr- Cu-Al alloys. 7) Crystal Zr material is characterized by its high purity compared with ordinary sponge Zr materials. Zr materials are usually used in sheathe material of nuclear reactors because of their high neutron transmission ability. Since the Zr material is embrittled when the oxygen concentration is over 800 ppm, 8) the oxygen concentrations of crystal Zr in the present study were maintained at less than 100 ppm. However, the effect of impurities in sponge Zr materials on the glass-forming ability of Zr-Cu-Al bulk glassy alloys remains unclear. Increases oxygen concentrations in Zrbased bulk glassy alloys enhances embrittlement and crystallization 9,10) due to the formation of oxide inclusions 11) in Zr-based glassy alloys. It has been reported that a small amount of rare earth metal, which promotes specific oxide inclusions, has a significant scavenging effect in iron-based bulk glassy alloys. 12) However, it is quite difficult to reduce the oxygen concentration after alloying, and once formed, oxide inclusions in Zr-based bulk glassy alloys can never be fully removed. Crystalline inclusions are often seen at the initiation site of a fracture surface after dynamic and static mechanical testing, indicating that they may cause the crack. In order to avoid the formation of crystalline inclusions in Zr- Cu-Al bulk glassy alloys, highly purified crystal Zr should be used, and it is important to pay attention to any impurity effects that may arise with respect to the degradation of the glass-forming ability. The purpose of the present study was to determine the effect of impurities in sponge Zr materials on the glassforming ability of Zr 50 Cu 40 Al 10 bulk glassy alloys, whose composition is considered to be a ternary eutectic point in a Zr-Cu-Al alloy system. 13) We also examined the cast morphology of arc-melted master alloys including glassy phase. Furthermore, in order to clarify the vitrification mechanism of arc-melted ingots, the solidification interface was cautiously observed by transmission electron microscopic (TEM) and high-resolution TEM (HRTEM) techniques. 2. Experimental Procedure Ternary Zr 50 Cu 40 Al 10 alloy ingots were prepared by arc melting pure Zr, Cu and Al elements in an argon atmosphere. Two different types of Zr were used as base material: crystal Zr and sponge Zr. Table 1 shows the componential analysis data of crystal Zr and sponge Zr. In addition, we performed a series of experiments with small amounts of iron, magnesium, oxygen and magnesium dichloride added to Zr 50 Cu 40 Al 10 ternary eutectics to determine the effect of impurities on the glass-forming ability of the tested alloys. Oxygen was introduced by Zr-O alloy, which was produces

2 1364 Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima and A. Inoue Table 1 Impurities and their concentrations in sponge and crystal Zr materials. element Sponge-Zr contents/ppm Crystal-Zr contents/ppm Mg Fe O Cl C Cr Si Hf Al 75 1 Ni W N Ti 50 5 Mo 50 1 Mn 50 1 Cu 30 3 Co 20 5 U 3 0 B Cd Nb Ta 0 40 H 0 15 S 0 10 Ca 0 10 Pd 0 1 Ag 0 1 Sn 0 1 (a) (b) (c) (d) (e) 5 Al (at%) K 1101 K 1125 K 1126 K 1150 K 1151 K 1175 K 1176 K 1200 K Cu (at%) Liquid (matrix) Zr 50Cu40Al10 Liquid (matrix) Solid e f d c b Zr49Cu38Al13 Zr35Cu34Al31 Liquid (matrix) Zr 54Cu41Al5 Solid Zr66Cu33Al1 (Zr2Cu ) Liquid (matrix) Zr 56Cu34Al K 1225 K 1226 K 1250 K 1251 K 1275 K 1276 K (Zr(CuAl)2) Solid Zr 64Cu31Al5 (Zr2Cu) by preliminary melting of pure Zr in Ar-O 2 (5 vol %) atmosphere. The concentration of oxygen in arc-melted alloys was measured by fusion in helium gas-infrared absorption method. In order to obtain the glassy phase in top half of the arc-melted ingots, we melted the ingots carefully for 5 minutes to restrain thermal convection in the molten alloy. The cast structure was examined by optical microscopy (OM) after being etched by nitric-hydrofluoric acid (30 vol % HNO3 + 5 vol % HF aqua) at room temperature. Microscopic cast structures were examined by TEM and HRTEM techniques using a JEOL 4000 FX Nanobeam energy dispersed spectrometry (EDS) was also performed to measure the compositional distribution in a few nanoscales with a JEOL 300 EX. TEM samples were prepared using an electrolytic polishing machine with nitric methanol (30 vol % HNO 3 ) at approximately 250 K. The ternary phase diagram was investigated 13) and is presented in Fig. 1. It is of particular interest for the present investigation that the phase diagram shows a ternary eutectic point, whose liquid during solidification goes through the following three phases; ZrCu, Zr 2 Cu and ZrCuAl compound as a primary phase. In compositions slightly different from the eutectic one, the following three phases are observed: the growth morphology of ZrCuAl phase is a coarse dendrite and is denoted as 5 in phase diagrams; the Zr 2 Cu phase has a plate-like morphology; and the 3 phase has a cubic symmetrical dendrite structure. In our further analysis, we found (f) all of these three phases in our samples except for a martensite ZrCu structure in TEM observation. 3. Results Liquid (matrix) Zr51Cu35Al14 Solid Zr50Cu25Al25 (τ 3) Fig. 1 Liquidus surface of Zr-Cu-Al alloy (a) and cast (10 60 mm) structure of Zr 50 Cu 40 Al 10 (b), Zr 55 Cu 30 Al 15 (c), Zr 55 Cu 40 Al 5 (d), Zr 55 - Cu 35 Al 10 (e) and Zr 50 Cu 35 Al 15 (f). The surface of the solidified alloy, which is composed of glassy phase, is smooth, metallic and shiny like a mirror. However, it becomes rough like smoky glass when the solidified surface is crystallized in a Zr-Cu-Al alloy system. Inhomogeneous crystallization during vitrification is probably promoted by an excessive oxygen concentration to form crystalline inclusions before casting. For the purposes of the present study, we designed an arc-melting furnace specifically to inhibit gas leakage and complete the mixing process in order to restrict oxidization during the melting and alloying processes. Using the arc-melting furnace, we can fabricate large Zr 50 Cu 40 Al 10 ingots over 40 g with a glassy shining surface.

3 Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy 1365 (c) (f) (e) interface 100µ m (a) ` (b) Chilled region (d) columnar ` 5mm Fig. 2 OM images of an arc-melted 20-g Zr 50 Cu 40 Al 10 ingot with crystal Zr (a) and magnified partial images (be). (c) (e) 100µ m (a) (b) Chilled region ` (d) 5mm Fig. 3 OM images of an arc-melted 20-g Zr 50 Cu 40 Al 10 ingot with sponge Zr (a) and magnified partial images (be). 3.1 Crystal Zr and Sponge Zr Figure 2(a) shows a cross-sectional OM image of an arcmelted 20-g Zr 50 Cu 40 Al 10 ingot with crystal Zr, and (b)(f) show magnified images of the chilled region, the interface between chilled and columnar regions, the columnar region, the interface between the columnar region and glass, and single glassy region, respectively. In the arc-melted Zr 50 - Cu 40 Al 10 ingot, a morphological transition from crystalline phase to glass phase was observed. The glass phase formed after the crystalline phase grew from the bottom of the ingot. The transition is closely related to a high glass forming ability (GFA), actually, the formation of glassy phase in arc-melted ingots may require a certain amount of skill in arc melting in order to restrict active thermal convection in the molten alloy. Approximately 1 mm from the lower chilled region, the columnar structure becomes a featureless structure. X-ray analysis shows that the columnar structure is composed of ZrCu phase with a CsCl type B2 structure and a monoclinic B19 structure, 14) while the analyzed composition of the ZrCu phase was Zr 50 Cu 40 Al 10. Therefore, the present ZrCu phase is considered to be an Al-supersaturated ZrCu intermetallic compound with mixed complex crystal structures of B2 and B19. The transition from the featureless ZrCu phase to a glassy phase occurs at a distance of approximately 2 mm from the lower chilled region. The distinct interface can be seen in Fig. 2(f); this interface indicates that vitrification occurred immediately. Figure 3(a) shows a cross-sectional OM image of an arc-melted 20-g Zr 50 Cu 40 Al 10 ingot with sponge Zr, and (b)(e) show magnified images of the chilled region, the starting point of a unidirectional dendrite, a unidirectional dendrite, and dendritic crystalline inclusions, respectively. No glassy phase can be seen in Fig. 3. The matrix phase of the final region in Fig. 3(e) was probably glassy phase immediately following solidification. The structure of the matrix is too tiny to be observed by OM because the crystalline phase is formed after vitrification. The change in morphology during solidification, which usually brings about the vitrification without any recalescence, can be seen through a window in the arc-melting chamber. Figure 4 shows photographs of the solidification process of an arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr in the arc furnace. After the arc melting is stopped, the redheated molten alloy gradually becomes dark without any

4 1366 Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima and A. Inoue 0 s 8 s 13 s 0 s 13 s 18 s Crystallization 2 s 9 s 14 s 3 s 14 s 19 s 4 s 10 s 15 s 6 s 15 s 20 s 6 s 11 s 16 s 9 s 16 s 21 s 7 s 12 s 17 s 11 s 17 s 22 s Fig. 4 Photographic images of the solidification process of a Zr 50 Cu 40 Al 10 ingot with crystal Zr after arc melting. Fig. 5 Photographic images of the solidification process of a Zr 50 Cu 40 Al 10 ingot with sponge Zr after arc melting. recalescence, indicating the vitrification of the arc-melted ingot. Figure 5 shows photographs of the solidification process of an arc-melted Zr 50 Cu 40 Al 10 ingot with sponge Zr in the arc furnace. At 13 s, recalescence due to crystallization was seen from the bottom side of the ingot, and it propagated gradually toward the top side until 19 s. These photographs show that the crystallization occurred as a result of the chain reaction from the crystallized regions at the bottom of the arc-melted ingot. We thus concluded that the impurities, which promote inhomogeneous crystallization, act as a trigger of chain reaction crystallization immediately following vitrification during the solidification process. The remarkable latent heat of crystallization at the solidification interface will reheat significantly, causing a chain reaction of crystallization. We also measured the cooling curve of Zr 50 Cu 40 Al 10 ingots with sponge Zr and crystal Zr, seeing distinct recalescence on the cooling curve of the Zr 50 Cu 40 Al 10 ingot with sponge Zr, as shown in Fig. 6. When crystal Zr is used, there is little recalescence due to crystallization on the cooling curve. Slippage of cooling curves was observed from a much higher temperature than the melting point (1100 K), and it emerged that the exothermic chain reaction, which completely interrupts glassy solidification, had occurred during the early stage of solidification of the arc-melted Zr 50 Cu 40 Al 10 ingot with sponge Zr. Consequently, highly purified Zr 50 Cu 40 Al 10 bulk glassy alloys exhibit an extremely high glass-forming ability that makes it possible to obtain glassy phase even in a 40-g arc-melted ingot, as shown in Fig. 7. A 40-g arc-melted Temperature, T / K Crystal Zr Zr 50Cu 40Al 10 20g ingot recalescence Time, t / s Sponge Zr Fig. 6 Cooling curves of arc-melted Zr 50 Cu 40 Al 10 ingots with crystal Zr and sponge Zr. Zr 50 Cu 40 Al 10 ingot with crystal Zr showed solidification morphology similar to that of a 20-g arc-melted ingot, as shown in Fig. 2. Specifically, although the glassy solidified interface seems smooth, significant amounts of globular crystalline inclusion can be seen, as shown in Fig. 7(c). The crystallized region on the bottom of the arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr is shown in Fig. 8. The crystallized region is characterized by its periodic layered structure with binary and ternary eutectic structures. Such a lamellar fluctuated structure has its origins in fluctuations in the Al concentration. Figure 9 shows the partial projection diagram of a ternary Zr-Cu-Al system with a tie triangle of ternary eutectic and binary eutectic lines. The structure of the crystallized region shows the repetition of binary and ternary eutectic reactions. Since the crystallization of an Al-enriched

5 Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy 1367 (a) (g) (d) (f) 50 µ m 100µ m (b) 50 µ m (c) (e) 50 µ m 5mm 50 µ m Fig. 7 Outer appearance of an arc-melted Zr 50 Cu 40 Al 10 ingot (a), a cross-sectional image (b) and magnified partial images (cg). 5 intermetallic compound requires a sufficient driving force, it might be difficult for ternary eutectics to occur under the peculiarly supercooled solidification conditions. Moreover, binary eutectics with Zr 2 Cu and ZrCu occur easily because no significant Al redistribution is required. Therefore, Al redistribution during solidification determines the periodic structure of binary and ternary eutectic layered structures. 3.2 Impurities In the present study, we concluded that crystal Zr is better for the fabrication of good glassy alloys, while harmful elements in sponge Zr cannot be clarified. Table 1 shows the types of impurities and their concentrations in sponge Zr and crystal Zr. Note that there are great differences in the quantities of Mg, Fe, oxygen and chlorine. The Fe impurity originates from the stainless steel chamber in the Kroll reduction process, Mg and chlorine are considered to be byproducts of MgCl 2, which is produced by the Mg reducing process for ZrCl 4 in the Kroll method, and oxygen is introduced into sponge Zr primarily by MgCl 2, which exhibits strong deliquescence. In the Kroll method, byproduct MgCl 2 is usually removed in a vacuum state at high temperatures (over 1000 K) during a long period of time (about 3 days). With the exception of oxygen in sponge Zr, impurities should be removed by modifying the conventional Kroll method. Therefore, we used a highly purified Zr metal designated as crystal Zr, which is produced by thermal decomposition of zirconium iodide gas to completely remove the byproduct MgCl 2. Figures 10(ad) shows cross-sectional scanning electron microscopic (SEM) images of a Zr 50 Cu 40 Al 10 arc-melted ingot with small additives: 1000 ppm oxygen, 0.5 at% Fe, 0.5 at% Mg and 0.03 at% MgCl 2. A glassy solidified region can be seen on the top of three ingots (Figs. 10(ac), and the interface between the crystalline and glassy regions is indicated by the pair of white arrows. All of the crystallized regions were composed of intermetallic compounds: ZrCu, Zr 2 Cu and 5 on the tie triangle. Peculiarly, the arc-melted Zr 50 Cu 40 Al 10 ingot with 0.03 at% MgCl 2 exhibited whole crystallized regions with many globular crystalline inclusions with a complex structure (not identified). The difference in the solidification morphology in a Zr 50 Cu 40 Al 10 arc-melted ingot with impurities was also examined by phase characterization, because the added impurities act as inoculants to

6 1368 Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima and A. Inoue (a) (b) Binary Eutectic Ternary Eutectic Binary Eutectic (c) Ternary Eutectic (d) Binary Eutectic Fig. 8 SEM images of the chilled region of an arc-melted 20-g Zr 50 Cu 40 Al 10 ingot with crystal Zr (a), a magnified image of the layer structure of binary and ternary eutectics (b), a magnified image of ternary eutectics (c) and binary eutectics (d). Al at% Zr 2 Cu Zr at% ZrCu 50 Fig. 9 Phase diagram showing phase relations with broken lines and binary eutectic lines with solid lines of a Zr-Cu-Al alloy system. change the type of solidification phase. Figure 11 shows the X-ray diffraction patterns of several Zr 50 Cu 40 Al 10 arc-melted ingots with impurities, specifically, patterns taken from the cross-section of the ingot shown in Fig. 10. The X-ray diffraction patterns of Zr 50 Cu 40 Al 10 arc-melted ingots to which 1000 ppm oxygen, 0.5 at% Fe or 0.5 at% Mg have been added exhibit ZrCu (B2 and B19 ), Zr 2 Cu and 5 phases that can be recognized by the phase diagram with tie triangle in Fig. 9. However, the MgCl 2 -added Zr 50 Cu 40 Al 10 arc-melted τ 3 τ 5 ingot was almost composed of 3 phase and ZrCu B19. The solidified phase change in the 0.03 at% MgCl 2 -added Zr 50 Cu 40 Al 10 ingot probably originates from the disappearance of the glassy phase in the solidified structure, indicating a significant degradation in the glass-forming ability. We conclude that the fatal impurity, that is, the one which significantly disturbs vitrification, is chlorine in sponge Zr metal. Consequently, the preparation of glassy alloys without crystalline inclusions using sponge Zr can be accomplished by controlling the concentration of byproduct MgCl 2 in sponge Zr as a raw material. 3.3 TEM observation of the glassy solidification interface It has been reported 15) that the solidification process of glass-forming Zr-Cu-Al alloys is characterized by its extremely low crystal growth rate. In the present study, the purity of Zr metal was found to have a great influence on vitrification, and we conclude that vitrification is significantly degraded by chlorine, which promotes inhomogeneous crystallization during solidification. On the other hand, in order to achieve partial vitrification of an arc-melted ingot, the restriction of inhomogeneous crystallization and a sufficiently low crystal growth rate are required simultaneously. In the present study, we examined the reasons for the vitrification in front of the solidified interface through observation of the crystal/glass interface of arc-melted Zr 50 Cu 40 Al 10 ingots using TEM and HRTEM techniques. The change in composition at the crystal/glass interface was measured in nanoscale, and compositional distributions at the solidified interface were clarified.

7 Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy 1369 (a) +O (b) +Fe (c) +Mg (d) +MgCl 2 500µm 500µm 500µm 500µm Fig. 10 SEM images of partial cross-section from the bottom to the top of an arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr with an oxygen concentration of approximately 1000 mass ppm (a), 0.5 at% iron added (b), 0.5 at% magnesium added (c) and 0.03 mol% MgCl 2 added (d). The arrows indicate the interface between the glassy and crystalline phases.

8 1370 Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima and A. Inoue +O ZrCu B2 ZrCu B19 Zr2Cu τ5 Growth direction Crystal Interface Glass +Fe Intensity (arb. units) +Mg +MgCl 2 τ3 ZrCu B θ (degree) Fig. 11 X-ray diffraction patterns of arc-melted Zr 50 Cu 40 Al 10 ingots with crystal Zr with an oxygen concentration of approximately 1000 mass ppm (a), 0.5 at% iron added (b), 0.5 at% magnesium added (c) and 0.03 mol% MgCl 2 added (d). 200nm Fig. 12 TEM images of the interface between the glassy and crystalline phases of an arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr. The bright field image of the crystal/glass interface of a 20-g arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr is shown in Fig. 12. Although the crystallized region at the crystal/ glass interface had a columnar structure, OM observation showed that the structure was featureless inside the column. TEM observation indicated that the column was composed of martensitic 10-nm fine nanocrystals, which was confirmed by X-ray diffraction showing a mixture of B2 and B19 phases. 19Þ We believe that the martensite transformation occurred after solidification because the nanocrystal structure shows a random crystal direction which is not related to the growth direction of the columnar crystals. A HRTEM image of the crystal/glass interface is shown in Fig. 13. No nanocrystalline inclusions were seen in front of the solidified interface; moreover, the solid/liquid interface is flat. Both of these morphologies suggest that the temperature gradient at the crystal/glass interface was not very steep. Careful observation of the HRTEM image shown in Fig. 13 shows onion-like contrast, 16,17) which suggests the existence of a cluster in liquid in front of the crystallized interface. Additionally, a little Zr-diluted and Cu-enriched compositional fluctuation, whose absolute values were less than a few at%, was observed at the crystal/glass interface by nanobeam EDS. No change in the Al concentration was observed around the crystal/glass interface, however, a significant fluctuation in the Al concentration in the early stage of solidification causes the unique structural change of the binary and ternary eutectic regions with the layered structure, as shown in Fig Discussion High glass-forming ability is one of the features of Zrbased bulk glassy alloys, and vitrification of the arc-melted master alloy can be seen around ternary eutectic composition Zr 50 Cu 40 Al 10. In order to achieve sufficient glass-forming ability, it is important to pay attention to the purity of materials because a given impurity may promote inhomogeneous crystallization or interrupt vitrification. Chlorine was found to be an impurity which significantly prevents vitrification in a Zr-Cu-Al alloy system. By removing impurities in Zr metal, we can fabricate a glassy arc-melted Zr 50 Cu 40 Al 10 master alloy with sufficient weight up to 40 g. It has been reported 19Þ that the origin of the high glassforming ability of Zr-based glassy alloys is not the high stability of their liquid structure but their extremely low crystal growth rate, which probably originates in the characteristic liquid structure. Vitrification of arc-melted Zr-Cu-Al alloys is seen only near the ternary eutectic composition, suggesting that the solidification process may differ from that of other alloy compositions. The compositional ratio of Zr:(Cu and Al) = 1:1, which has the advantage of enhancing the mixing entropy effect to maximize a pseudo-binary system, can be seen in the ternary eutectic

9 Glassy Solidification Criterion of Zr 50 Cu 40 Al 10 Alloy 1371 B2 (B19 ) Growth direction Interface Onion like contrast 5nm Fig. 13 HRTEM images of the interface between the glassy and crystalline phases of an arc-melted Zr 50 Cu 40 Al 10 ingot with crystal Zr. composition Zr 50 Cu 40 Al 10. It is also considered to be the origin of the low temperature at the ternary eutectic point. Furthermore, the ternary eutectic composition Zr 50 Cu 40 Al 10 is close to the B2 type ZrCu intermetallic compound composition Zr 50 Cu 44 Al 6, whose structure is partially martensitic transformed to B19. 14) Therefore, the combination conditions of the unique structural change in the primary crystal and the peculiar features of the phase diagram may cause the extremely low crystal growth rate, which is considered to be the basis of the high glass-forming ability. The origin of the solidification interface stoppage, which allows the vitrification of the remaining liquid phase, is an extremely interesting problem. Supercooling is essential in order for the vitrification in the middle of the solidification process of an arc-melted Zr 50 Cu 40 Al 10 ingot to occur. The arc-melted Zr 50 Cu 40 Al 10 alloy with crystal Zr starts solidification with the typical layered structure of the binary and ternary eutectics due to fluctuations in the Al concentration, as seen in Fig. 8. The growth velocities of Zr 2 Cu in the binary Zr 66:7 Cu 33:3 alloy, Zr 2 Cu, ZrCu or 5 (not identified) in the ternary Zr 50 Cu 40 Al 10 alloy were examined using an electromagnetic levitation furnace. 15) The velocity of the Zr 2 Cu compound ranged from 20 to 900 mm/s in the undercooling range of approximately 100 K. The growth velocity of the crystalline region in the Zr 50 Cu 40 Al 10 alloy was approximately 0.1 to 0.8 mm/s. On the other hand, the velocity of the 5 phase in the Zr 45 Cu 40 Al 15 alloy could not be measured by the levitation technique using alternating and static magnetic fields since it was too low to be measured by this technique. The solidified 5 microstructure, whose grain size was less than 20 mm, showed that the crystallization of the 5 phase was controlled by nucleation. The velocity of the 5 phase, which could be estimated by the grain size and solidification time, was at most 100 mm/s even at an undercooling value exceeding 100 K. Thus, the velocity of the 5 phase was much lower than those of the Zr 2 Cu and ZrCu compounds. In the arc-melted Zr 50 Cu 40 Al 10 alloy, a binary Zr 2 Cu-ZrCu eutectic in the bottom region, a eutectic Zr 2 Cu-ZrCu- 5 in the bottom middle region, a columnar structure ZrCu in the middle region, and a glass phase in the top were observed. This morphological transition is explained by considering the competition between the growth of the coupled growth (the binary and ternary eutectics) and the formation of the glass phase, as shown in Fig. 8. We observed the competition between the binary coupled growth of the Zr 2 Cu and ZrCu and the ternary coupled growth of the Zr 2 Cu, the ZrCu and 5 phase. In the ternary eutectic composition (Zr 50 Cu 40 Al 10 ), the ternary coupled growth is selected in equilibrium. In the coupled growth, the constituent phases have to have the same interface temperature and the same growth velocity. The growth velocity of the 5 phase was much lower than those of Zr 2 Cu and ZrCu. Thus, a certain undercooling is required so that the 5 phase will grow cooperatively with the Zr 2 Cu and ZrCu. The undercooling required for coupled growth shifts the coupled growth zone to the Al-rich region. Thus, binary Zr 2 Cu-ZrCu eutectic solidification occurs at the beginning of solidification. Since the growth velocities of the Zr 2 Cu and ZrCu were relatively low, the interface temperature of the binary coupled growth decreased as solidification proceeded. In addition, the Al element was rejected by the solidifying front. When the undercooling from the liquidus temperature of the 5 phase reached a certain temperature, the 5 phase could grow cooperatively with the Zr 2 Cu and ZrCu, and consequently ternary coupled growth occurred. Even during the ternary coupled growth, the interface temperature continued to decrease. When the interface temperature decreased to the glass transition temperature, the glass phase formed ahead of the coupled growth interface. Thus, the observed transition can be understood by taking the competition into account. The low growth velocities of the constituent crystalline phases in the Zr-Cu-Al alloys significantly contribute to the high GFA. After the formation of layered structure, fluctuations in the Al concentration decrease with the increase in the supercooling effect, and the growing crystal changes to the columnar structure of the Al-supersaturated ZrCu intermetallic compound. However, the microscopic structure of the Al-supersaturated ZrCu intermetallic compound shows a martensitic nanocrystal structure in TEM observation, a structure that shows no marks that indicate the original direction of solidification. Nanocrystallization occurred after solidification. The martensitic nanocrystallization in the columnar crystal, which is probably caused by the huge strain immediately following solidification, is caused by the

10 1372 Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima and A. Inoue significant structural differences between the precursor of the B2 structure and the local structure of the liquid. In the HRTEM image shown in Fig. 13, a great amount of onion like contrast is seen in front of the solidified interface with red circles. This contrast indicates the existence of an icosahedral-like cluster 16,17) in the remaining Zr 50 Cu 40 Al 10 liquid phase. In general, the existence of clusters in the liquid usually act as a precursor for crystallization, however, the structure of an icosahedral-like cluster in the liquid is probably quite different from that of the precursor for a ZrCu B2 structure. Huge strain was probably introduced into the Al-supersaturated ZrCu intermetallic compound by the remaining typical local structure of the Zr 50 Cu 40 Al 10 liquid phase under the supercooled slow solidification just before the liquid froze. Furthermore, since solidification is usually accompanied by latent heat due to crystallization, it is difficult to realize sufficient supercooling for vitrification in front of the solidification interface. As shown in Fig. 12, the crystal/ glass interface in the present study was smooth, and compositional distributions of less than a few percent were detected by nanobeam EDS measurement. We thus conclude that the remarkable reduction in the crystal growth rate, which enables vitrification in front of the solidification interface, was not caused by solute segregation in front of the solidified interface. Moreover, the densities of the glassy Zr 50 Cu 40 Al 10 alloy and the B2 type Zr 50 Cu 40 Al 10 compound were almost equivalent at approximately 6.8 g/cm 3, suggesting that the large solidification latent heat is seldom present when the solidified crystal is a B2 type Zr 50 Cu 40 Al 10 compound. That is, it is possible that the Al-supersaturated ZrCu intermetallic compound as a primary crystal is not accompanied by large solidification latent heat. Thus, the formation of an Al-supersaturated ZrCu B2 type compound, which shows less latent heat, occurs in Zr 50 Cu 40 Al 10 ternary eutectic composition. The precursor structure of the B2 structure is quite different from the local structure (cluster) of liquid. Consequently, supercooling is promoted by less latent heat, and the crystallization rate decreases significantly in Zr 50 Cu 40 Al 10 ternary eutectic composition. It is therefore possible to obtain a glassy arcmelted ingot in Zr 50 Cu 40 Al Summary Differences in the solidification morphology and structure of arc-melted Zr 50 Cu 40 Al 10 ingots were examined using Zr metals of different purity grades. A solidified interface after vitrification under supercooled solidification was observed by TEM and HRTEM methods, clarifying the unique vitrification mechanism of solidification. The results are summarized as follows: (1) By controlling the purity of Zr metal, the arc-melting process can produce an arc-melted Zr 50 Cu 40 Al 10 alloy as a glassy phase with total weight of up to 40 g. (2) A clear solidification interface between the crystal and glass was seen in an arc-melted Zr 50 Cu 40 Al 10 alloy with crystal Zr, however, it was not observed in an arcmelted Zr 50 Cu 40 Al 10 ingot with sponge Zr. (3) The fatal impurity, that is, the impurity which completely interrupts vitrification, was determined to be chlorine in sponge Zr. (4) The cross-sectional solidification morphology of arcmelted Zr 50 Cu 40 Al 10 ingots was found to consist of a multi-layered structure with binary and ternary eutectic structures, a columnar crystal of Al-supersaturated ZrCu intermetallic compound and a glassy phase, in that order, from the bottom up. (5) The extremely low crystal growth rate of the Alsupersaturated ZrCu intermetallic compound, which is considered to be the basis of vitrification, is probably caused by the unique liquid structure of Zr 50 Cu 40 Al 10, which has icosahedral-like clusters. Acknowledgements This research was funded in part by Grant-in-Aid for Scientific Research on Priority Area (Materials Science of Bulk Metallic Glasses) from the ministry of Education, Culture, Sports, Science and Technology and NEDO (the New Energy and Industrial Technology Development Organization). 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