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Acta Materialia 59 (2011) 2243–2267

Recent development and application products of bulk glassy alloys q

A. Inoue ⇑, A. Takeuchi

WPI-Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan

Available online 20 January 2011

Abstract

This paper reviews past developments and present understanding of the glass-forming ability, structure and physical, chemical,mechanical and magnetic properties of bulk glassy alloys (BGA) with the emphasis on recent results obtained since 1990, together withapplications of BGA, achieved mainly in Tohoku University. After introducing the fundamental concepts around glassy alloys (GA) inSections 1 and 2 describes the progress of the study of structural relaxation leading to the discovery of GA with a large supercooled liquidregion. Section 3 reviews the history of BGA development, followed by BGA systems and their features in Section 4, and features ofglassy structure in Section 5. Sections 6–9 summarize the engineering and standardization of Zr-based BGA, followed by the originsof the development of useful materials on the basis of experimental data on the compositional effect on the fundamental propertiesof basic ternary and quaternary Zr-based BGA. Sections 10 and 11 include the glass-forming ability and dynamic mechanical propertiesof Zr-based hypoeutectic BGA and Cu–Zr–Al–Ag BGA. Mechanical properties of Ni- and Zr-based BGA at low temperatures areshown in Section 12, while Section 13 describes the formation and properties of Ni-free Ti-based BGA. Sections 14 and 15 deal withporous Zr-based BGA, including spherical pores and commercialized ferromagnetic and high-strength Fe-based GA, respectively, thenSection 16 reviews supercooled liquid formation. Applications for Zr-, Ti- and Fe-based GA are described in Section 17. In conclusion,Section 18 attempts to assess the present knowledge of the structure and physical properties and identify some outstanding problems forfuture work.� 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Bulk metallic glass; Glass-forming ability; Mechanical properties; Magnetic properties; Atomic structure

1. Introduction

Even in the 21st century, metals and their alloys aresome of the most important classes of materials, used inextremely wide fields, such as structural, functional andbiomedical materials, which are essential for the develop-ment of an affluent yet sustainable society. As regards bulkmetallic materials which can be used in a three-dimensionalshape, only crystalline metallic alloys have been used by

1359-6454/$36.00 � 2010 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2010.11.027

q This Acta Materialia Gold Medal Lecture was presented at the 2010WPI-AIMR Annual Workshop at Sendai Japan where WPI (WorldPremier International research center initiative) is a prior program/projectsupported by JSPS (Japan Society for the Promotion of Science).⇑ Corresponding author. Address: 2-1-1, Katahira Aoba-ku, Sendai 980-

8577, Japan. Tel.: +81 22 217 4806.E-mail address: [email protected] (A. Inoue).

human beings for several thousand years up to about themid-1990s. Before 1990, there were no data on the forma-tion of alloys in bulk form with critical dimensions aboveseveral millimeters by the copper mold casting process,which is very popular as a practical method of mass pro-duction [1]. It was reported in 1990, before crystallizationeven by the copper mold casting method [1], that bulkglassy alloys (BGA) can be formed by choosing specialmulti-component amorphous alloys with a large super-cooled liquid region. Since the first success of formingBGA by the casting method, a large number of studieson the formation and fundamental properties of BGA havebeen carried out over the past two decades [2–6]. The max-imum diameter for glass formation achieved by coppermold casting processes reached as large as >2 cm for sys-tems based on Zr [7–9], Zr–Be [10], Pd [11], Pt [12], Mg[13], La [14,15], Ni [16] and Cu [17], and >1 cm for systems

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Fig. 1. Thermograms of (Fe0.5Ni0.5)83P17 amorphous alloy subjected toannealing at 475 K for various periods from 1 to 48 h; the solid linepresents the thermogram of the sample subjected to heating to 675 K.Reproduced with permission from Ref. [127].

2244 A. Inoue, A. Takeuchi / Acta Materialia 59 (2011) 2243–2267

based on Fe [18], Co [19] and Ti [20]. These critical dimen-sions are large enough to be used as practical materials in avariety of application fields. In addition, large BGA pos-sess nearly the same fundamental properties as those ofthe corresponding glassy alloys with millimeter-sized diam-eter. Owing to the unique advantages of high glass-formingability, net-shape castability, good nanoscale imprintabilityand useful fundamental properties, BGA have been used aspractical materials and/or tested for real application initems such as sporting goods, watch parts, electromagneticcasing, optical parts, ornamental parts, choke coils, powerinductors, magnetic field identification systems, electro-magnetic wave shielding sheet, micro-geared motor parts,pressure sensors, Coriolis flow meters, surface coatingmaterials, peening shot and medical instruments [6]. Thispaper reviews the glass-forming ability, alloy components,structure, fundamental properties, unique liquid-formingprocess and application examples of BGA developed mainlyby the present authors’ group.

2. Discovery of glassy alloys with large supercooled liquid

region through structural relaxation study

When a metallic liquid with temperatures above meltingis cooled at an ordinary slow cooling rate such as 1 K s�1,the liquid always solidifies to a crystalline structure withperiodic atomic configurations on a long-range scale. Thisuniversal phenomenon was recognized, without any excep-tional examples, and crystalline metallic alloys were used ina bulk form as practical materials for several thousandyears. It was found in 1960 [21] that, when the metallicliquid is cooled at a very high cooling rate of �106 K s�1,a disordered atomic configuration similar to liquid struc-ture can be maintained down to room temperature. Thus,the first synthesis of an amorphous phase was made forAu75Si25 alloy by Duwez’s group in 1960. Since then, alarge number of studies on the synthesis, structure andproperties of amorphous alloys were carried out in con-junction with the development of various types of rapidsolidification techniques. However, the formation of amor-phous phase always required very high cooling rates of theorder of 106 K s�1 because of the high atomic mobility inconventional metallic alloys. Therefore, the resulting amor-phous alloys were usually limited to small shape sizes, i.e.,ribbon �0.02 mm thick, wire �0.1 mm in diameter andpowder �0.02 mm in diameter. When their material sizeswere further increased, it is well known that the amorphousalloys lost their material values upon catastrophic embrit-tlement caused by the precipitation of the crystalline phaseand/or the progress of structural relaxation [22].

In addition, amorphous alloys transform directly to thecrystalline phase upon continuous heating and do not showglass transition and a supercooled liquid region in the tem-perature range before crystallization. Consequently, before1990, nobody considered or believed that it would be pos-sible to form bulk amorphous alloys with critical diametersof several millimeters by conventional casting processes,

though some data were known on the formation of Pd-based glassy alloys in rod form with diameter �2 mm bywater quenching [23] and in a massive sphere form withdiameter �10 mm by repetition flux melting with B2O3, fol-lowed by water quenching [24].

It is also known that all the fundamental properties ofamorphous alloys change significantly with proceedingstructural relaxation caused by annealing [22]. In 1982, aseries of studies on the relationship between structuralrelaxation and fundamental properties for amorphousalloys was started, using a precise calorimetric measure-ment method which is a suitable way to examine quantita-tively their structural relaxation behavior. For instance,Fig. 1 shows the temperature dependence of apparent spe-cific heat for (Fe0.5Ni0.5)83P17 amorphous alloy ribbon withthickness �15 lm produced by melt spinning [25]. The as-spun amorphous sample shows a distinct irreversible exo-thermic reaction with an onset temperature of �375 K,owing to the change from the unrelaxed amorphous tothe relaxed amorphous state. When the sample is annealed,for instance, for different periods of 1–48 h at 475 K, itshows an excess reversible endothermic peak with an onsettemperature of �475 K, owing to the disappearance ofannealing-induced relaxed atomic configurations. In thestructural relaxation study, the use of amorphous alloysexhibiting glass transition and a supercooled liquid regionis very important, because the supercooled liquid is in aninternal equilibrium state. Its state can be used as a fullyrelaxed reference state in quantitative evaluation of thechange in annealing-induced configuration enthalpy. Byobtaining specific data in the fully relaxed liquid, one candetermine the change in configuration enthalpy of amor-phous alloys upon structural relaxation. Although theamorphous Fe–Ni–P alloy did not exhibit the fully super-cooled liquid state, the onset point in the transition from

A. Inoue, A. Takeuchi / Acta Materialia 59 (2011) 2243–2267 2245

glass transition to supercooled liquid state was taken as arelaxed reference state.

Thus, systematic studies on structural relaxation of Fe-,Co-, Ni- and Cu-based amorphous alloys carried out since1982 searched for amorphous alloys exhibiting glass transi-tion and a supercooled liquid region, with the aim ofobtaining more reliable, quantitative and reversible dataon structural relaxation. As a result, for several yearsbetween 1988 and 1992, a number of glass-type alloys withdistinct glass transition and a large supercooled liquidregion >50 K before crystallization were found [26–28].Using the high stabilization of supercooled liquid againstcrystallization, glassy alloys were successfully formed inbulk form by casting processes in La- [1], Mg- [29] andZr- [7–9] based metal–metal alloy systems without any met-alloid elements. Their BGA showed significant plasticity, incontrast to the embrittlement phenomenon of amorphousalloy ribbons, by increasing the ribbon thickness from20 lm to �50 lm. Since the successes in forming BGA withgood mechanical properties by the present authors’ group,the main research target in non-equilibrium metallic mate-rials drastically changed from amorphous-type alloys tobulk glassy-type alloys.

3. History of development of BGA

As described in Sections 1 and 2, since the findings onforming BGA using the high stability of supercooled liquidagainst crystallization for metallic alloys [1,29], a largenumber of studies on the development of BGA have beencarried out over the past two decades on the basis of basicscience and engineering interests. Fig. 2 shows a schematicillustration of continuous cooling transformation curvesfor bulk glassy, ordinary amorphous and crystalline alloys.The resulting lowest critical cooling rate for glass forma-tion reaches as low as 0.067 K s�1 [30], which is �108 timessmaller than those for ordinary amorphous alloys [1–4].Such a dramatic increase in the thermal stability of super-cooled liquid has enabled the production of large-scaleBGA with various outer shapes, e.g., massy ingot shape72 mm in diameter and 75 mm high [11], and 80 mm indiameter and 85 mm high [31] for a Pd–Cu–Ni–P system,

Fig. 2. Schematic illustration showing the high stability of the supercooledliquid state for long periods reaching several thousand seconds. Repro-duced with permission from Ref. [128].

25–30 mm in diameter and 40–50 mm high for systemsbased on Zr–Al–Ni–Cu [7,8], Cu–Zr–Al–Ag [17] andNi–Pd–P–B [16], cylindrical rods 25 mm in diameter and300 mm long for a Pd–Cu–Ni–P system, and 40 mm indiameter and 100 mm long for a Pt–Pd–Cu–P system,and hollow pipes 10 mm in outer diameter and 1 mm thickfor Pd–Cu–Ni–P, Zr–Al–Ni–Cu and La–Al–Ni–Cu sys-tems, as exemplified in Fig. 3 [4,5]. The production oflarge-scale BGA with a diameter of 25 mm was also repre-sented in an early study on a Pd–Ni–P system by He et al.[32]. The largest diameter of 80 mm was obtained for thequaternary alloy composition of Pd40Cu30Ni10P20 [11,30].Considering the result that a fully glassy phase was formedover a whole circular disk plate 80 mm in diameter and5 mm high cut from the original massy alloy ingot, the trueglass-forming ability is expected to be considerably higherthan the experimental data obtained to date. In additionto the large-scale massy alloys, glassy alloy sheets with uni-form thickness �2–4 mm and large surface area up to A-4size (210 � 297 mm2) [33] as well as glassy alloy balls withdiameters ranging from several millimeters to �10 mm areavailable for Zr–Al–Ni–Cu and Cu–Zr–Al–Ag systems[34]. It is thus noticed that the critical diameter reaches sev-eral centimeters even for Zr–Al–Ni–Cu [7–9], Ni–Pd–(P,B)[16] and Cu–Zr–Al–Ag [17] alloy systems, though their val-ues are smaller than the largest value (80 mm for Pd–Cu–Ni–P alloy) [11,31] among all BGA reported to date. Basedon these recent significant developments in BGA, a newtype of bulk metallic alloy consisting of a glassy structurehas been used in commercial materials since the mid-1990s, in addition to the ordinary crystalline bulk metallicalloys which have been used for several thousand years.

4. BGA systems and their features

Table 1 summarizes typical BGA systems reported todate, together with the calendar year when the first paperor patent for each alloy system was published. The alloysystems can be classified as non-ferromagnetic and ferro-magnetic alloy groups. The non-ferromagnetic alloy groupcomprises Ln-, Mg-, Zr-, Ti-, Pd-, Ca-, Cu-, Pt- and Au-based systems, while the ferromagnetic group consists ofFe-, Co- and Ni-based alloy systems. Considering that lan-thanide (Ln) metal is composed of at least 15 elements, thetotal number of BGA systems greatly exceeds 500 types.Based on the experimental results for BGA systemsobtained during the period up to 1993, the features of alloycomponents for stabilization of supercooled liquid andhigh glass-forming ability are proposed [2–5]. These fea-tures consist of the following three factors: (1) multi-com-ponent, consisting of three elements or more; (2) significantatomic size mismatch >12% among the three main constitu-ent elements; and (3) negative heats of mixing among thethree main elements. Using the three empirical rules, thepresent authors succeeded in developing a variety ofBGA systems, summarized in Table 1 for the past 15 years.

(a) Pd-Cu-Ni-P (c) Cu-Zr-Al-Ag (d) Ni-Pd-P-B

72φx 75 mm 80φx 85 mm

Massy Ingot Shape

Cylindrical Rods(e) Pd-Cu-Ni-P

(b) Zr-Al-Ni-Cu

(f) Pt-Pd-Cu-PHollow Pipes

(g) Pd-Cu-Ni-P

Fig. 3. Pd-, Ni,- Cu- and Zr-element-based large centimeter-sized BMG. Reproduced with permission from Ref. [129].

Table 1Typical bulk glassy alloy systems reported up to date together with the calendar years when the first paper or patent of each alloy system was published.

1. Non-ferromagnetic alloy systems Year 2. Ferromagnetic alloy systems Year

Mg–Ln–M (Ln = lanthanide metal, M = Ni, Cu, Zn) 1988 Fe–(Al,Ga)–(P,C,B,Si,Ge) 1995Ln–Al–TM (TM = Fe, Co, Ni, Cu) 1989 Fe–(Nb,Mo)–(Al,Ga)–(P,B,Si) 1995Ln–Ga–TM 1989 Co–(Al,Ga)–(P,B,Si) 1996Zr–Al–TM 1990 Fe–(Zr,Hf,Nb)–B 1996Zr–Ln–Al–TM 1992 Co–(Zr,Hf,Nb)–B 1996Ti–Zr–TM 1993 Fe–Co–Ln–B 1998Zr–Ti–TM–Be 1993 Fe–Ga–(Cr,Mo)–(P,C,B) 1999Zr–(Ti,Nb,Pd)–Al–TM 1995 Fe–(Cr,Mo)–(C,B) 1999Pd–Cu–Ni–P 1996 Ni–(Nb,Cr,Mo)–(P,B) 1999Pd–Ni–Fe–P 1996 Co–Ta–B 1999Ti–Ni–Cu–Sn 1998 Fe–Ga–(P,B) 2000Ca–Cu–Ag–Mg 2000 Ni–Zr–Ti–Sn–Si 2001Cu–Zr, Cu–Hf 2001 Ni–(Nb,Ta)–Zr–Ti 2002Cu–(Zr,Hf)–Ti 2001 Fe–Si–B–Nb 2002Cu–(Zr,Hf)–Al 2003 Co–Fe–Si–B–Nb 2002Cu–(Zr,Hf)–Al–(Ag,Pd) 2004 Ni–Nb–Sn 2003Pt–Cu–Ni–P 2004 Co–Fe–Ta–B–Si 2003Ti–Cu–(Zr,Hf)–(Co,Ni) 2004 Ni–Pd–P 2004Au–Ag–Pd–Cu–Si 2005 Fe–(Cr,Mo)–(C,B)–Ln (Ln = Y, Er, Tm) 2004Ce–Cu–Al–Si–Fe 2005 Co–(Cr,Mo)–(C,B)–Ln (Ln = Y, Tm) 2005Cu–(Zr,Hf)–Ag 2005 Ni–(Nb,Ta)–Ti–Zr–Pd 2006Pd–Pt–Cu–P 2007 Ni–Pd–P–B 2009Zr–Cu–Al–Ag–Pd 2007 Fe–(Nb,Cr)–(P,B,Si) 2010Ti–Zr–Cu–Pd 2007Ti–Zr–Cu–Pd–Sn 2007


It is also important to point out that the special multi-component alloy liquid always includes a eutectic pointwith the lowest liquidus temperature. Designing the alloychoosing an alloy composition exhibiting a lower liquidustemperature in the vicinity of the eutectic point is an effec-tive way of finding an appropriate BGA composition. Oncethe appropriate alloy composition has been found, it is rel-atively easy to increase further the critical diameter forglass formation by modification of the alloy component

to a more multi-component system. The selection of ele-ments with nearly zero heat of mixing belonging to thesame group in the periodic table is effective for the minoradditional elements, presumably because of an increase inentropy of the supercooled liquid and an enhancement ofthe difficulty of atomic rearrangement through the increasein the degree of complexity of liquidus structure. Thisconcept is also consistent with the entropy effect conceptproposed by Greer [35].


5. Features of glassy structure

It has also been reported that alloys with the three empir-ical rules have new glassy and liquid structures with the fol-lowing features: (1) a high degree of dense packed atomicconfigurations; (2) new local atomic configurations whichare completely different from those for the correspondingcrystalline phases; and (3) long-range homogeneity withattractive interaction [4,36]. Fig. 4 shows a schematic illus-tration of their structural features, together with typicalglassy alloy systems. Metal–metal-type glassy alloys asexemplified for Zr–Al–Ni–Cu and Zr–Be–Ti–Ni–Cu systemsare composed of icosahedral-like ordered atomic configura-tions. Pd–transition metal–metalloid-type glassy alloys suchas the Pd–Cu–Ni–P system consist of highly dense packedconfigurations of two types of polyhedra of Pd–Cu–P andPd–Ni–P atomic pairs, while metal–metalloid-type glassyalloys of Fe–Ln–B and Fe–(Zr,Hf,Nb)–B ternary systemshave network-like atomic configurations in which a distortedtrigonal prism and an anti-Archimedean prism of Fe and Bare connected with each other in face- and edge-sharedconfiguration modes through glue atoms of Ln and earlytransition metals of Zr, Hf or Nb. These icosahedral-, poly-hedral- and network-like ordered atomic configurations caneffectively suppress the long-range rearrangements of theconstituent elements which are necessary for the progressof the crystallization reaction. Among the three types ofstructures, the second and third types are similar in pointcontaining trigonal prism structure, but are different in thatthe latter forms a well-developed connected structure of theprisms by sharing their vertices and edges. As a result, onecan obtain highly stabilized supercooled liquid, leading tothe formation of BGA even by various slow cooling solidifi-cation processes.

The above experimental data on glassy structures werereported before 2002 by the present authors’ group.

Fig. 4. Schematic illustrations of the structural features of metal–metal, Pd–permission from Ref. [128].

Recently, a number of research groups have reported moredetailed structural models of BGA by computer simulationmethods. As a typical example, Fig. 5 illustrates a topolog-ical model based on coordination polyhedra proposed byMiracle [37], indicating the development of medium-rangeorder with highly efficient packing [37–40]. This model isalso consistent with the above-mentioned structural fea-tures. The present authors’ group also proposed a structuremodel using the molecular dynamics (MD) simulationmethod based on the plastic crystal model [41] for creatinga dense random packing structure including glassy, amor-phous and liquid phases in the framework of a conceptof cluster-packed structure. This model aimed to constructthe glassy structure through the sequence of ensemble ofclusters and glue atoms, followed by random rotations ofclusters and then annealing-induced structural relaxation.The total pair distribution function of the Zr–Al–Cu glassyalloy obtained by this model agrees well with the experi-mental data, suggesting that the development of medium-range ordered atomic configurations and their effectivepacking on a longer scale play an important role in the sta-bilization of supercooled liquid for special multi-compo-nent alloy systems which satisfy the three empirical rules.In addition, experimental studies performed to determinethe short-range atomic structure and the electronic struc-ture of prototypical BGA [42–44] also greatly contributeto progress in clarifying the features of glassy structures.

6. Engineering BGA in Zr–Al–Ni–Cu system

A large amount of knowledge has been obtained from alarge number of experimental and computer simulationstudies over the past two decades. This has enabled theproduction of BGA with large critical diameters >10 mmin various types of alloy systems such as those based onZr, Ti, Pd, Pt, Mg, Ln, Fe, Cu and Ni. The easy formation

metalloid and metal–metalloid types of glassy alloys. Reproduced with

α

α

Ω

α

α

αβ

β

β

β

a b

Fig. 5. Topological model based on coordination polyhedra proposed by Miracle [37]. Reproduced with permission from Ref. [130].


of such large-scale BGA has opened up new engineeringmaterial and science fields. Here, some recent progress infundamental studies on engineering BGA, in particular,Zr-, Ti-, Pd-, Pt- and Fe-based systems, is described. Table 2summarizes some typical application fields being developedin Japan at present for BGA. Their application and fieldsof development extend to the following variety of materi-als: structural, sensor, spring, sporting goods, wear-resis-tant coating, corrosion-resistant, optical, magnetic, microor nano-technology, information data storage, ornamental,biomedical and fuel-cell separator materials. It is importantto point out here that these materials have been developedin conjunction with unique advantages of net-shape castingand simple production techniques, viscous flow formabilityand nanoscale imprint ability, in addition to useful charac-teristics resulting from the long-range disordered atomicconfigurations and multi-components.

Table 2Application fields that have been proceeded at present for BMGs in Japanwhere each term with numbers 1–17 come to have a complete implicationby supplementing affix words of “Materials” behind the term, such as“Structural Materials”.

1. Structural2. Sensor3. Precision machinery4. Optical5. Ornamental6. Spring7. Sporting goods8. Wear-resistant coating9. Precision nozzle10. Corrosion-resistant11. Magnetic12. Micro-technology13. Nano-technology14. Information data storage15. Biomedical16. Medical instrument17. Fuel-cell separator

Advantages: net-shape processing, viscous flow forming processing.

Focusing on engineering BGA, in practical uses they aremainly limited to seven types of alloy systems: (i) hostmetal-base types of Zr–Al–Ni–Cu, Fe–Cr–metalloid, Fe–Nb–metalloid and Fe–Ni–Cr–Mo–metalloid systems and(ii) pseudo-host metal-base types of Zr–Be–Ni–Cu–Ti,Ti–Zr–Cu–(Ni,Pd)–Sn and Cu–Zr–Al–Ag systems. Thus,one can see that Zr- and Fe-based BGA are the mostimportant materials for practical use.

7. Attempts at standardization of Zr-based BGA

With the aim of extending further application fields andgeneral recognition of BGA, much effort has been devotedto standardizing the alloy composition and fundamentalproperties of BGA and to establishing their availabilityon a global basis. Table 3 summarizes the standardizedspecimen diameter, maximum diameter, thermal stabilityand static and dynamic mechanical properties of four alloyseries, i.e., Z-alloy series in the Zr–Cu–Al and Zr–Cu–Al–Ni systems, ZC-alloy series in the Zr–Cu–Al–Ag systemand ZF-alloy series in the Zr-Cu-Al-Fe system developedpreviously by the present authors’ group [45]. It shouldbe noted here that the terms “static” and “dynamic”

mechanical properties indicate those measured withoutand under reputational cycles, respectively. It is recognizedthat all the alloy series exhibit large critical diameters>15 mm in the case of using the copper mold casting pro-cess, and reliable fundamental properties in the diameterrange up to at least 10 mm in the alloy compositions repre-sented in Table 3.

Thus, BGA, including Zr metal as a main constituentelement, have been recognized as engineering materials. Itis known that six types of BGA in Zr–Al–(Ni,Cu), Zr–Al–Ni, Zr–Al–Cu, Zr–Be–Ti–Ni–Cu, Zr–Al–Co and Zr–Cu–Al–Ag systems were developed in 1990 [1], 1990 [28],1991 [46], 1993 [47], 1995 [48] and 2007 [49], respectively,as basic and typical Zr-based alloy systems. Their maxi-mum diameters are 32 mm [43] for Zr–Al–Ni–Cu system,15 mm [50] for Zr–Al–Ni system, 22 mm [51] for Zr–Al–Cu system, >30 mm [10] for Zr–Be–Ti–Cu–Ni system,

Table 3Fundamental data of standardized bulk metallic glasses.

Series Alloy composition(features)

dc

(mm)ds

(mm)Tg

(K)Tx

(K)Tl

(K)E

(GPa)ey

(%)ry

(MPa)CUE(kJ/m2)

Z Z1 Zr50Cu40Al10 (ternary eutectic) 22 10 706 792 1092 88 2.1 1860 104Z2 Zr55Cu30Al10Ni5 (high glass-forming

ability)30 10 683 767 1163 90 2.0 1830 125

Z3 Zr60Cu20Al10Ni10 (high resistance toembrittlement)

20 10 662 754 1164 80 2.2 1750 87

Z4 Zr65Cu17.5Al7.5Ni10 (high stability ofsupercooled liquid)

16 10 625 750 1164 82 1.9 1528 85

ZC ZC1 Zr48Cu36Al8Ag8 (high glass forming) 25 10 683 792 1142 102 1.8 1850 –ZC2 Zr42Cu42Al8Ag8 (high strength) 14 10 705 780 1213 108 1.8 1986 –

ZF Zr62.5Cu22.5Al10Fe5 (Ni-free type) 20 – 651 770 1173 88 – 1584 –

dc: critical diameter, ds: standard diameter.


18 mm [52] for Zr–Al–Co system and 30 mm [17] for Zr–Cu–Al–Ag base system and their Zr contents for bulk glassformation are in the range �50–70 at.%, 52–70 at.%, 50–70 at.%, 42–44 at.%, 55–60 at.% and 42–48 at.%, respec-tively. The four glassy alloys in Zr–Al–Cu–Ni, Zr–Al–Ni,Zr–Al–Cu and Zr–Al–Co systems can contain high Zr con-tents ranging from 50 to 70 at.%, 52 to 70 at.%, 50 to70 at.% and 55 to 60 at.%, respectively, and can beregarded as typical Zr-based alloy systems, while the Zrcontents in the other two glassy alloys are <50 at.%. Con-sequently, it is more appropriate for the other two glassyalloys to be called Zr–Be and Zr–Cu base alloy systems.Considering the experimental result that eutectic alloycompositions in Zr–Al–Ni–Cu and Zr–Al–Cu systems arelocated around Zr55Al10Cu30Ni5 [53] and Zr50Al10Cu40

[54], the Zr–Al–Cu–Ni and Zr–Al–Cu glassy alloys areextended particularly over a wide composition rangeincluding eutectic and hypoeutectic compositions. Also,the hypoeutectic Zr–Al–Ni–Cu alloy with the highest70%Zr content has a large critical diameter reaching10 mm [55]. In contrast, the Zr–Be and Zr–Cu base BGAappear to be located in the vicinity of eutectic compositionand can be regarded as a eutectic-type glassy alloy.

8. Compositional effect

In the standardization and industrialization process ofZ-, ZC- and ZF-alloy series, the hypoeutectic Zr-basedBGA with higher Zr contents than the eutectic compositionexhibit much better dynamic mechanical properties in as-cast and annealed states accompanying higher Poissonratios in comparison with the corresponding eutectic glassyalloys [56,57]. It should be noted here that, in the latestresearch on bulk metallic glasses (BMG), Poisson’s ratiois becoming a more important factor, since it has directrelationships to mechanical properties such as shear modu-lus [58] and the fracture energy [59]. Fig. 6 shows the com-positional dependence of Poisson’s ratio and Young’smodulus of Zr–Al–Cu–Ni BGA. In the wide Zr concentra-tion range from 50 to 70 at.%, there is a distinct tendencyfor Poisson’s ratio to increase with increasing Zr content,and the 70%Zr-containing hypoeutectic alloy has a high

Poisson’s ratio of 0.387. In contrast, Young’s modulustends to decrease with increasing Zr content, and the lowestvalue of 78 GPa is recognized for the 70%Zr glassy alloys.The Zr70Al8Ni16Cu6 glassy alloy exhibited large plasticstrain exceeding 70% under uniaxial compression and didnot show final rupture, as shown in Fig. 7 [60]. A high den-sity of shear bands can be observed on the peripheral sur-face of the severely deformed alloy rod. The highly plasticnature of the BGA with high Poisson’s ratio can be recog-nized. Even for the highly Zr-rich alloys subjected to severeplastic deformation, no trace of deformation-inducednanocrystallization was observed in the transmission elec-tron microscopy images. In addition, the hypoeutectic70%Zr alloy exhibits distinct plastic elongation of 1.7%under a uniaxial tension mode at a strain rate of 1.6 �10�1 s�1 [61]. As the strain rate increases to 5 � 10�1 s�1,the tensile plastic elongation increases further to 2.8%,accompanying the gradual decrease in nominal flow stressfrom 1500 MPa at 0.2% plastic elongation to 1200 MPaat 2.8%, as exemplified in Fig. 8 [62]. The tensile plasticelongation is believed to be the largest for monolithicBGA without any trace of deformation-induced nanocrys-tallization. The tensile fracture occurred through thesequence of generation of shear bands, followed by distinctnecking, appreciable shear sliding along the maximumshear stress plane and then final catastrophic rupture [62].The angle between the tensile loading direction and thefracture plane was measured as 53�, suggesting that thedeformation and fracture occurred via the Mohr–Coulombcriterion mechanism [63].

The relationship between the shear transformation zone(STZ) and Poisson’s ratio for various types of BGA wasalso examined. The STZ are nano-scale volumes of mate-rial that undergo plastic flow, and the concentration ofstress at the STZ is believed to cause localized shear bandsto form, leading to mechanical failure and poor ductility.In reality, STZ was measured by nanoindentation [64].As summarized in Fig. 9 for Zr- and Pd-based BGA [64],there is a clear tendency for the STZ to increase almostlinearly from 1.3 to 1.9 nm with an increase in Poisson’sratio from 0.32 to 0.45. It is therefore concluded that theZr-rich hypoeutectic glassy alloys possess an inherent

8284

86

Young's Modulus

80

78

78

80

82

84 86

0.3880.386

0.384

0.382

0.380

0.378

0.376

0.374

Poisson's ratioAl = 10at%

near the eutecticcomposition with

the lowest Tl= 1103 K

Fig. 6. Compositional dependence of Poisson’s ratio and Young’s modulus of Zr–Al–Ni–Cu BGA. Reproduced with permission from Ref. [131].

Fig. 7. Large plastic strain exceeding 40% under a uniaxial compressionwithout final rupture for 70%Zr glassy alloy. Reproduced with permissionfrom Ref. [132].


potentiality of obtaining higher ductility compared withthe other BGA, being consistent with experimental resultsof the extremely large compressive plasticity and distincttensile elongation shown in Figs. 7 and 8.

φ3.5 cast

φ8 cast rod

Initial Strain rate : 1.5 X 10-1

Yield Stress : 1.5 GPaYield Strain : 2.2 %Young’s modulus : 70 GPaPoisson’s ratio : 0.39

Zr70Ni16Cu6Al8 BMG

0200400600800

10001200140016001800

0

Stre

ss, σ

/ M

Pa

φ3.5 cast

φ8 cast rod

Initial Strain rate : 1.5 X 10-1

Yield Stress : 1.5 GPaYield Strain : 2.2 %Young’s modulus : 70 GPaPoisson’s ratio : 0.39

Zr70Ni16Cu6Al8 BMG

Zr70Ni16Cu6Al8 BMG φ3.5mm tilt cast sample

1.7%

1 2 3 4Elongation, ε/ %

Fig. 8. Tensile stress–elongation curve of Zr70Al8Ni16Cu6 B

Very recently, it has also been found that the hypoeutec-tic Zr70Al8Ni16Cu6 BGA exhibits very high resistanceagainst stress corrosion failure under a tensile deformationmode in 3%NaCl solution. Fig. 10 summarizes the relation-ship between the applied tensile stress and the time to finalfailure under tensile applied stress for eutectic Zr55Al10-

Ni5Cu30 and hypoeutectic Zr70Al8Ni16Cu6 glassy alloy rodswith the same gauge dimensions of 1 mm in diameter and1 mm in length [65]. The former eutectic alloy exhibits sig-nificant degradation in tensile fracture stress, i.e., from1620 MPa in air to 680 MPa after testing for 20 min underthe tensile stress loading condition in 3%NaCl solution.The fracture mode also changes from shear type alongthe maximum shear stress plane to perpendicular type.However, it is noticed that the latter hypoeutectic alloykeeps the same tensile fracture stress level as the fracturestrength obtained in air as well as the same shear type frac-ture mode. The finding of such excellent stress corrosionresistance for the hypoeutectic BGA with a high Poisson’sratio of 0.393, low Young’s modulus of 70 GPa, and largeSTZ of 2.5 nm is promising for future use as a corrosion-resistant structural material in various corrosive environ-ments. It is also important to say that the tensile fracture

rod

s-1

rod

s-1

Tensile test view

Sample dimensions

2.8%

5 6

GA rod. Reproduced with permission from Ref. [133].

Fig. 11. Temperature dependence of viscosity for Zr70Al8Ni16Cu6 glassyalloy. The temperature corresponding to the viscosity of 4000 Pa s is alsopresented.

Fig. 9. Relationships between the STZ volume and Poisson’s ratio for Zr-,Ni-, Cu- Pt- and Pd-based BGA [64].


stress of the eutectic-type glassy alloy rods remains almostconstant in other chemical aqueous solutions such as de-ionized water, 0.5 M Na2SO4 and 0.5 M NaNO3 [65].

The appearance of the distinct tensile plastic elongationprovides a good opportunity to investigate the slidingspeed of the shear band and the temperature inside theshear band in tensile deformation mode. There have beenno experimental data on the temperature rise in the shearband under tensile plastic deformation. The presentauthors tried to measure the sliding speed of the shear bandthrough the tensile rod specimen for Zr70Al8Ni16Cu6 glassyalloy using a high-speed video camera [66]. The slidingspeed (V) and shear stress (s) of the shear band were mea-sured as 2–4 m s�1 and 640–690 MPa, respectively, and theshear band thickness (h) was assumed to be 15 nm, fromprevious experimental data [67,68]. The resulting viscosity(g) in the shear band was estimated to be 4000 Pa s fromthe well-known relation g = s � (h/V). There are

Fig. 10. SSRT (Slow Strain Rate Technique) test results obtained forZr70Cu6Al8Ni16 BGA in 0.5 M NaCl at an initial strain rate of 5 � 10�6

s�1. The results for Zr50Cu40Al10 and Zr50Cu30Al10Ni10 BGAs are alsoshown for comparison. Reproduced with permission from Ref. [134].

experimental data on the temperature dependence of vis-cosity for Zr70Al8Ni16Cu6 glassy alloy, as shown inFig. 11 [66]. The dependence data allow one to estimatethat the temperature in the shear band reaches 840 K.The rising temperature lies within some previous experi-mental values obtained under bending deformation mode[69]. It is thus concluded that the tensile shear slidingoccurs via a viscous flow mode within instantaneous time,in agreement with the compressive shear sliding mode.

It is also important to clarify the structure relaxation sen-sitivity of static and dynamic mechanical properties of thehypoeutectic-type BGA. Choosing the eutectic (Zr50Cu40

Al10) and the hypoeutectic (Zr55Cu35Al10 and Zr60Cu30Al10)glassy alloys in a simple ternary Zr–Cu–Al system, the pres-ent authors examined the difference in stability of themechanical properties upon annealing [70,71]. The 60%Zr

Fig. 12. Fracture toughness of Zr50Cu40Al10, Zr55Cu35Al10 andZr60Cu30Al10 BGA in as-cast and annealed states. Reproduced withpermission from Ref. [135].


alloy keeps nearly the same tensile fracture elongation in anannealed state before crystallization, in contrast to the signif-icant degradation of tensile elongation upon annealing forthe 50%Zr alloy. The 60%Zr alloy also shows nearly the samefatigue stress amplitude after 107 cycles in as-cast unrelaxedand annealing-induced relaxed structural states. It is furthernoticed that the hypoeutectic 60%Zr alloy exhibits signifi-cant improvement in U-notched Charpy impact fractureenergy in the structural relaxed state caused by annealing,though the eutectic 50%Zr alloy shows a gradual decreasein the impact fracture energy with progress in structure relax-ation. In addition, the hypoeutectic 60%Zr and 55%Zr alloysexhibit much higher fracture toughness values than that forthe eutectic 50%Zr alloy, as shown in Fig. 12. The higherfracture toughness values are also maintained in the anneal-ing-induced structure relaxed state. Reflecting the differencein the fracture toughness values, one can observe a signifi-cant difference in the shear band structure in the front regionof the notched pre-fatigue crack. The shear band length andthe crack tip opening displacement were measured as2700 lm and 340 lm, respectively, for the as-cast 60%Zralloy, which were much larger than those (250 lm and13 lm, respectively) for the as-cast 50%Zr alloy, being con-sistent with the fracture toughness data. It is thus noticedthat the hypoeutectic-type glassy alloys exhibit higher valuesof compressive plastic strain, tensile plastic elongation, U-notched Charpy impact fracture energy and fracture tough-ness compared with the corresponding eutectic-type glassyalloy. The similar difference in the compressive plasticity

Zr (at.%

(a)

50 60

10

20

(b)

50 60

40

30

10

20

2070

Fig. 13. Composition range in which the BGA are formed by the copper moldcast glassy alloy rods in Zr–Al–Cu, Zr–Al–Ni and Zr–Al–Co systems.

and fracture toughness between the eutectic and hypoeutec-tic BGA has also been recognized in the Zr–Al–Ni ternarysystem [61]. The better properties for the hypoeutecticBGA have been presumed to result from the increase in thenumber of the more metallic Zr–Zr atomic bonding pair aswell as the increase in the population of free volumes dueto the deviation from optimal atomic configurations whichcan be obtained around the eutectic alloy composition.

9. Fundamental ternary Zr–Al–TM (TM = Co, Ni or Cu)

BGA

One of the most important and well-recognized BGA isthe Zr–Al–Cu–Ni system, because of many features of highglass-forming ability in a wide composition range of 50–70%Zr, high static mechanical properties in conjunctionwith high corrosion resistance, good dynamic mechanicalproperties in the hypoeutectic composition range >60%Zrand the absence of special toxic elements. It is further rec-ognized that Zr–Al–Cu and Zr–Al–Ni ternary systems arefundamental for the development of Zr–Al–Ni–Cu system.An earlier study also reported BGA with a critical diameterup to 18 mm in the Zr–Al–Co ternary system [52]. It istherefore important to compare the features of glass-form-ing ability and fundamental properties between the threetypes of fundamental ternary BGA, i.e., Zr–Al–Cu, Zr–Al–Ni and Zr–Al–Co systems.

Fig. 13a–c shows the composition range in which theBGA are formed by the copper mold casting method and

)

40

30

(c)

50 60

40

30

10

20

70

7020

casting method and the composition range of the maximum diameter of

Zr48Cu36Ag8Al8

6045

2535

40

50

15

Zr (at.%)

20

45

Fig. 14. Compositional dependence of maximum diameter of Zr–Cu–Al–Ag glassy alloys produced by copper mold casting. Reproduced withpermission from Ref. [136].


the composition range of the maximum diameter of castglassy alloy rods in Zr–Al–Cu, Zr–Al–Ni and Zr–Al–Cosystems, respectively. The BGA with critical diameters>10 mm can be formed by tilt casting in a wide composi-tion range of 7.5–12.5%Al and 30–42.5%Cu for theZr–Al–Cu system, 11–21%Al and 21–27%Ni for the Zr–Al–Ni system and 14–20%Al and 25–29%Co for the Zr–Al–Co system. The maximum diameter of the glassy alloyrods is 22 mm for Zr50Al10Cu40 [51], 15 mm forZr60Al15Ni25 [50] and 18 mm for Zr56Al16Co28 [52]. Thecritical diameter decreases in the order Cu > Co > Ni. Inaddition, the alloy composition at which the largest criticaldiameter is obtained is also distinctly different among thethree systems. It is noticed that the concentration ratio ofM/Al (M = Cu, Ni, Co) is much higher for Cu and isnearly the same for M = Ni and Co. The significant differ-ence in the ratio of M/Al between M = Cu and M = Ni orCo implies that the Zr–Al–Cu BGA are based on the Zr–Cu binary system, while the Zr–Al–Ni and Zr–Al–CoBGA are the multi-component system in which Al and Melements equivalently contribute to the formation ofBGA. The difference has also been supported by manyexperimental data: that is, only Zr–Cu binary alloys amongZr–M binary alloys exhibit a distinct glass transition, fol-lowed by a supercooled liquid region and are mainly com-posed of highly dense icosahedral atomic configurations,while neither the glass transition phenomenon nor thedevelopment of icosahedral atomic configuration isobserved for the other binary Zr–Ni and Zr–Co amor-phous alloys [72]. The reason for such a significantdifference between Zr–Al–Cu and Zr–Al–M (M = Ni orCo) is under investigation.

New information on the formation of BGA with criticaldiameters >10 mm has been shown, even in Zr–Al–Co andZr–Al–Ni ternary systems, comparable with that for Zr–Al–Cu system. This new information suggests that the highglass-forming ability for the Zr–Al–Cu system does notoriginate only from the icosahedral atomic configurationin the Zr–Cu binary system, and the third element of Alalso contributes to a further increase in the critical diame-ter for glass formation. This concept is consistent with theexperimental result that the critical diameters for Zr–Cu–binary alloys are �1–2 mm, which is one order smaller thanthose for Zr–Al–Cu ternary alloys. Considering the datathat BGA with critical diameters >10 mm are obtainedeven for Zr–Al–Co and Zr–Al–Ni systems, the high glass-forming ability in Zr–Al–M (M = Co, Ni or Cu) systemsmay originate from the more dense and longer-scale devel-opment of icosahedral atomic configuration achieved bythe multi-component effect consisting of the special threeelements or more. The effect of addition of Al to the Zr–Cu binary system was discussed by Cheng et al. [73]together with the subject of atomic-level structure of a rep-resentative ternary Cu–Zr–Al BMG through MD simula-tions. In the literature, Chen et al. reported that a smallpercentage of Al in the ternary BMG leads to a dramati-cally increased population of full icosahedra and their

spatial connectivity, and that the stabilizing effect of Al isnot merely topological, but also has its origin in electronicinteractions and bond shortening. The formation of BGAin Zr–Al–Co and Zr–Al–Ni systems clearly demonstratesthe importance of the three empirical rules. More detailedstudies on the abnormal icosahedral atomic configurationstructure and the bonding nature in the three ternary alloyswill give further deep understanding of the mechanism ofthe very high glass-forming ability of Zr–Al–M ternaryalloys and contribute to further developments in the basicscience and engineering of BGA.

10. Glass-forming ability of hypoeutectic Zr-based glassy

alloys

The glass-forming ability of hypoeutectic glassy alloys inthe Zr–Al–Ni–Cu system was examined in comparison withthe eutectic-type glassy alloy, using copper mold castingmethods which are useful for the production of engineeringBGA. It was reported that the eutectic Zr55Al10Cu30Ni5 alloyrods kept a glassy phase in the diameter range up to 30 mm[7,8] in the case of copper mold suction casting and 32 mmusing copper mold cap casting [74]. Similarly, the maximumdiameter for the hypoeutectic-type glassy alloys has beenreported to be at least 20 mm for Zr60Al10Ni10Cu20,Zr61Ti2Nb2Al7.5Ni10Cu17.5, Zr60Ti2Nb2Al7.5Ni10Cu18.5 andZr65Al7.5Ni10Cu15.5Pd2 [75,76], 22 mm for Zr60Ti2N-b2Al7.5Ni10 Cu16.5Pd2 [77], 16 mm for Zr65Al7.5Ni10Cu17.5

[78] and 10 mm for Zr70Al8Ni16Cu6 [55] by copper mold tiltcasting. The addition of alloy components with nearly zeroheat of mixing to the basic ternary BGA component wasfound to be effective for the further increase in maximumdiameter from 16 to 22 mm even in the hypoeutectic compo-sition range. It was also confirmed that the thermal stabilityparameters of the glass transition temperature (Tg), crystal-lization temperature (Tx) and temperature interval of thesupercooled liquid region DTx(=Tx � Tg), the staticmechanical properties of Young’s modulus, yield strengthand plastic strain, and the deformation and fracture behav-ior via shear sliding along the maximum shear stress plane

Fig. 15. Compressive stress–strain curves of Ni60Pd20P17B3 BGA rodstested at temperatures between room temperature and 77 K. Reproducedwith permission from Ref. [137].


and then catastrophic final rupture are almost independentof both a part and an angle to the longitudinal direction ofthe cast alloy rod as well as the rod dimension in the diameterrange up to 20 mm [75]. The confirmation that reliable ther-mal and mechanical properties are obtained for the hypoeu-tectic-type BGA with a large critical diameter reaching20 mm is encouraging for further extension of the applica-tion fields for Zr–Al–Cu–Ni BGA.

11. Formation and properties of Cu- and Zr-rich BGA in

Cu–Zr–Al–Ag system

Since 2004, Cu–Zr base BGA have been developed in Zr–Cu–Ag and Cu–Zr–Al–Ag base systems. Fig. 14 shows thecompositional dependence of maximum diameter for theformation of BGA in Zr-rich Zr–Cu–Al–Ag and Cu-richCu–Zr–Al–Ag systems. The largest glassy rod diameter is25 mm for Zr48Cu36Al8Ag8 and 15 mm for Cu48Zr36 Al8Ag8

[79]. It was also reported that the critical diameter increasedfurther to 30 mm by addition of 2%Pd to the Zr-rich Zr–Cu–Al–Ag alloy [17]. The as-cast structure in the Zr

48Cu36Al8Ag8

alloy rod with a much larger diameter >35 mm was com-posed mostly of a very fine eutectic lamellar structure con-sisting of CuZr2, AlCu2Zr, AgZr2 and unknown phases[80], indicating that the eutectic point is located in the vicinityof Zr48Cu36Al8Ag8. The eutectic-type Zr–Cu–Al–Ag andZr–Cu–Al–Ag–Pd glassy alloys exhibit a relatively high yieldstrength of �1900 MPa and distinct compressive plasticstrain, accompanying the shear-sliding-type fracture mode.The fracture toughness of the Zr48Cu36Al8Ag8 alloy was�17 MPa

ffiffiffiffi

mp

[81], which was considerably lower than thosefor Zr–Al–Ni–Cu BGA (40–55 MPa

ffiffiffiffi

mp

) [82] and those forZr–Al–Cu BGA (50–110 MPa

ffiffiffiffi

mpÞ) [82]. The shear band

length and stretched zone width in the front region of thenotched pre-fatigue crack were measured as 49 lm and0.8 lm, respectively, for the Zr–Cu–Al–Ag alloy [81], beingmuch smaller than those for the hypoeutectic-type Zr–Al–Ni–Cu and Zr–Al–Cu BGA, as is evident from comparisonwith Fig. 12 [71]. It is thus concluded that the fracture tough-ness of the eutectic Zr–Cu–Al–Ag BGA is considerablylower than that for the hypoeutectic Zr–Al–Ni–Cu BGA.Similarly for the Zr–Al–Ni–Cu system, the development ofhypoeutectic Zr–Cu–Al–Ag BGA is an effective methodfor improving the dynamic mechanical properties, even inthe Zr–Cu–Al–Ag system.

12. Mechanical properties of BGA at low temperatures

As described in Section 4, BGA were formed in Mg-, La-and Zr-based alloy systems around�1990–1993 [1–4]. Sincethen, much attention has been paid to the synthesis of Ni-based BGA because their alloys are expected to exhibithighly ductile nature and good corrosion resistance. As aresult, a variety of Ni-based BGA were synthesized as exem-plified for Ni–Cr–Mo–P–B [82], Ni–Nb–Cr–Mo–P–B [83]and Ni–Nb–Zr–Ti–Cu–Co [84] systems. However, therehad been no data on the formation of Ni-based BGA

containing more than 50%Ni with critical diameters>10 mm. More recently, the formation of Ni-based BGAwith large critical diameters of 10–22 mm in Ni–Pd–P [12]and Ni–Pd–P–B [16] alloys containing 50–60%Ni have beenreported. In addition to the high glass-forming ability, theseNi-based BGA also exhibit large compressive plastic strainsexceeding 2% at room temperature. It is important to clarifythe mechanical properties in a low-temperature range fromroom temperature to 77 K (boiling temperature of liquidnitrogen) because of the possibility of applying the Ni-and Zr-based BGA to structural and functional materialswhich can be used at low temperatures. In addition, theauthors do not have any data on mechanical properties inthe low-temperature range for BGA exhibiting distinctroom temperature ductility, though some data on mechan-ical properties at low temperatures have been reported foreutectic-type Zr- and Pd-based BGA without distinct plas-tic strain under compression mode at room temperature[85,86].

The results on mechanical properties and deformationbehavior at low temperatures are introduced for a Ni60

Pd20P16B4 BGA with a critical diameter of 12 mm.Fig. 15 shows typical compressive stress–strain curvesof the Ni-based glassy alloy rods with diameter 2 mm,subjected to compression test in the temperature rangefrom room temperature to 77 K [87]. It is clearly seenthat the yield strength and plastic strain increase signifi-cantly with a decrease in temperature to 77 K, accompa-nying the reduction of the serrated flow amplitude on thestress–strain curves. In addition, these changes becomemore distinct in the lower temperature range <�200 K.Although the shear-sliding-type fracture mode was inde-pendent of testing temperature, the density of the shearbands was much higher for the rod specimen deformed


at 77 K. However, the vein pattern in the fracture planesremains unchanged over the entire testing temperaturerange. Fig. 16 shows a tensile stress–strain curve andouter appearance of the fractured part for the hypoeutec-tic Zr70Al8Ni16Cu6 glassy alloy rod with diameter 0.8 mmin the gauge part tested at 77 K [88]. Although no appre-ciable plastic elongation is seen for the Zr-based alloysubjected to tensile deformation at room temperature,one can recognize a distinct tensile plastic elongation of0.2% at 77 K accompanying the generation of a numberof shear bands on the lateral outer surface of the tensilespecimen. These results indicate clearly that the strengthand ductility increase significantly with a decrease in tem-perature to 77–123 K and their improvements are inde-pendent of the type of base metals. A monotonicincrease in all the properties of Young’s modulus, shearmodulus, bulk modulus and Lame parameter (lambda)with decreasing temperature was also reported, whichreflected the gradual increase in the longitudinal andtransverse wave velocities [87]. In addition, the Debyetemperature increases monotonically with decreasing tem-perature, in contrast to the gradual decrease in Poisson’sratio in the temperature range <�200 K. These resultsindicate that the significant increases in strength and duc-tility at low temperatures seem to be attributed to thesuppression of generation and propagation of inhomoge-neous shear bands caused by the increase in the bondingforces among the constituent elements accompanying thedecrease in their inter-atomic distances, and the decrease

Fig. 16. Tensile stress–strain curve and outer appearance of the fractured par0.8 mm in the gauge part.

in the mobility of the constituent atoms and freevolumes.

13. Formation and properties of Ni-free Ti-based BGA

It is known that one of the important application fieldsfor Ti-based BGA is the biomedical field, as exemplified inimplant materials. For application to biomedical materials,there is a need to develop a new Ti-based BGA exhibitinghigh glass-forming ability, reliable mechanical propertiesand good compatibility with living tissues, which is withoutany elements that are toxic, allergenic or carcinogenic.Recently, the present authors developed a new Ti-basedBGA with a critical diameter of 7 mm in a Ni-free Ti–Zr–Cu–Pd system [89]. The largest critical diameter of7 mm was obtained in the limited compositions of Ti36Zr14-

Cu36Pd14 and Ti36Zr14Cu34Pd16. The addition of a smallamount of Sn element was effective for further increase inthe critical diameter to 10 mm [20]. These critical diametersare large enough to be applied to a variety of biomedicalmaterials. The Ni-free Ti-based BGA in the as-cast stateexhibit a high yield strength of �2000 MPa and distinctcompressive plastic strain [90] as well as better corrosionresistance than that of SUS304 stainless steel in a 3%NaClsolution [91]. The good mechanical properties are main-tained even for an annealing-induced nanocrystalline struc-ture consisting of fine Pd3Ti particles embedded in a glassymatrix. It was also confirmed that the addition of 2–3%Nbis effective for further increase in plastic strain to 6.5–8.5%

t for the hypoeutectic Zr70Al8Ni16Cu6 glassy alloy rod with a diameter of


[92]. The high plasticity provides good machinability tovarious complex shapes, which is necessary for applicationto bio-body and teeth implants. The Ni-free Ti-based BGAalso exhibit good corrosion resistance, with higher anodicpotential and lower passive current density in comparisonwith SUS304 steel and conventional Ti–Al–V alloys. Inaddition, the Ti-based BGA possess high formation abilityof a porous hydroxyapatite phase via the special two-stageprocess of electrochemical and chemical treatments [93], asexemplified in Fig. 17. It has also been demonstrated thatthe formation of such a porous hydroxyapatite phase isessential for achieving good compatibility with living tis-sues [94]. The Ni-free Ti-based BGA with highly reliablemechanical properties in as-cast and annealed states andthe high formation ability of a porous hydroxyapatite sur-face layer in conjunction with high glass-forming ability areunder investigation for application to biomedical implantsand medical instruments.

14. Porous BGA

Much attention has been paid to porous BGA becauseof the possibility of significantly changing various

Fig. 17. (a) SEM micrograph of BMG (Ti40Zr10Cu36Pd14) surfacethrough hydrothermal–electrochemical treatment for 120 min. (b) SEMmicrograph in a cross-section view of Fig. 13a. Reproduced withpermission from Ref. [138].

fundamental properties, such as density, elastic modulus,yield strength, ductility and specific surface area. Somedata on porous BGA have been reported for Pd–Cu–Ni–P alloys, including spherical and polyhedral pores [95–98], and Zr–Al–Cu–Ni alloys, including polyhedral pores[99]. The early work performed by Schroers et al. [95] hasbeen followed up by a more thorough study in the litera-ture [100]. There had been no data on the formation ofZr-based porous BGA including spherical pores. Veryrecently, the present authors found that Zr–Cu–Al–Ag por-ous BGA including spherical pores can be formed by evac-uation and supercooling of alloy liquid melted in a high-pressure helium gas atmosphere [101]. Fig. 18 shows theouter shape, cross section and mechanical properties ofporous Zr–Al–Ni–Cu BGA, including spherical pores pro-duced by the unique process. The porosity and pore size inthe porous Zr-based BGA can be controlled in a widerange from �5% to 70% and from 20 to 150 lm, respec-tively, by annealing treatment in the supercooled liquidregion between 690 and 762 K. The resulting porousBGA exhibited the same X-ray diffraction pattern and dif-ferential scanning calorimetry curves as those for the corre-sponding non-porous BGA. In addition, the porous Zr-based BGA with porosities of 0–68% exhibit Young’s mod-ulus, yield strength and plastic strain ranging from 106 to10 GPa, 1850 to 160 MPa and 0 to 0.4, respectively, underuniaxial compression. The success in forming porous Zr-based BGA with a wide range of mechanical properties isattractive for the future extension of fields of applicationfor BGA.

15. Commercialized Fe-based glassy alloys

Recently, there has been a strong demand for sustainabledevelopments to create low-carbon, resource-circulating andstable societies. One way of making a contribution tosustainable developments in the present academic field is todevelop Fe-based BGA with high glass-forming ability andhighly functional properties and to achieve their industriali-zation. Since the first synthesis of Fe-based BGA in Fe–Al–Ga–P–C–B system in 1995 [102], a variety of Fe-based BGAhave been developed [103]. Table 4 summarizes alloy systemsof Fe-based BGA developed up to date. Their alloy systemscan be classified as ferromagnetic and non-ferromagnetictypes at room temperature. As shown in Table 4, some Fe-based BGA have already been commercialized with the com-mercial names “Liqualloy” [104] and “SENNTIX” [105] forthe former type and “AMO-beads” [106] for the latter type.The “Liqualloy” and “SENNTIX” soft magnetic alloys havea unique combination of lower coercivity, higher electricalresistivity and lower melting temperature, which have notbeen obtained for conventional Fe-based amorphous alloycores and nanocrystalline Fe-based soft magnetic alloycores. Such a low coercivity characteristic has been inter-preted as originating from the low internal stress and lowmagnetic anisotropy due to the formation of more homoge-neous atomic configurations in the glassy state. These

Fig. 18. DSC curve (top) and corresponding porosity curve (bottom) of porous Zr48Cu36Al8Ag8 glassy alloy containing pressurized helium gas bubblesunder 400 K/min isochronal annealing. The inset picture shows the appearance of the as-prepared and annealed samples up to 795 K. The opticalmicroscope images (a–d) show cross-section structure of porous sample annealed up to four different temperatures within the supercooled liquid range of715, 755, 775 and 795 K (denoted by (a–d) in the porosity curve). Reproduced with permission from Ref. [139].

Table 4Alloy systems of Fe-based bulk glassyalloys developed up to date. The regis-tered trademark names in practical usesare also presented.

Ferromagnetic type

1. Fe–(Al,Ga)–(P,C,B)2. Fe–Ga–(P,C,B,Si)3. Fe–(Cr,Mo)–(P,C,B,Si) –

“Liqualloy”

4. Fe–(Zr,Hf,Nb)–B5. Fe–Co–Ln–B6. Fe–(Nb,Cr)–(B,Si) – “SENNTIX-

1”

7. Fe–(Nb,Cr)–(B,P) – “SENNTIX-2”

Non-ferromagnetic type

1. Fe–(Cr,Mo)–(C,B) – GA-coat2. (Fe,Ni)–(Cr,Mo)–(B,Si) – “AMO-

beads”

3. Fe–(Cr,Mo)–(C,B)–Ln

Fig. 19. Relationship between fracture strength and Young’s modulus forBGA. The data of crystalline metallic alloys are also shown forcomparison. Reproduced with permission from Ref. [140].


Fe-based BGA exhibit a high fracture strength of 3000–4000 MPa and large elastic strain of�0.02, which are signif-icantly different from those for conventional Fe-basedcrystalline alloys, as shown in Fig. 19, in addition to high cor-rosion resistance. The “SENNTIX”-type BGA also exhibitsa high fatigue stress amplitude of �2300 MPa after 107

cycles, which is much superior to those for conventionalFe-based crystalline alloys [107]. It is further noted that theseFe-based BGA exhibit much lower melting temperatures

than those for conventional Fe-based alloys, in spite of muchhigher mechanical strength. The advantage of producing Fe-based BGA in a much lower temperature range is greatlyfavorable for sustainable development.

16. Supercooled liquid forming

When bulk glassy-type alloys are heated at an ordinaryheating rate, they always exhibit glass transition, followed

(a)

(c)

(d)

(e)

(b)

50% (0.96)

1% (0.02)

30% (0.57)

13% (0.25)

52% (1)

Fig. 20. (a) Optical microphotograph of micro-forged surface of Pt-based BMG; reflectance at k = 533 nm. (b–d) SEM micrographs of die-forgedperiodically nanostructured surface of Pt-based BMG with periodic intervals of (b–d) 800 nm and (e) 400 nm.


by a significant supercooled liquid region with a tempera-ture interval of 40–130 K and then crystallization. This isin contrast to the direct transformation to the crystallinephase for amorphous-type alloys produced by rapid solid-ification. In the supercooled liquid region, one can recog-nize the Newtonian viscosity, as evidenced by a nearlyconstant viscosity in a wide strain rate range [108]. In

Fig. 21. Imprinted glassy alloy nano-pillars with diameter 200 nm and lengthsupercooled liquid region.

addition, the supercooled liquid with the Newtonian flowproperty exhibits a good linear relation between the truestress (r) and true strain (_e) rate, and this relation can beexpressed as r = k_em, where k is a constant. The slope (mvalue) of the linear relation is called the strain-rate sensitiv-ity exponent and can be evaluated as �1.0, indicating theachievement of an ideal superplasticity. Therefore, large

�5 lm obtained by pressing the glassy alloy into porous alumina in the


tensile elongations exceeding 104% have been obtained forglassy alloys subjected to pulling treatment in the super-cooled liquid region [109].

Recently, the present authors also reported that vari-ous nano-imprinted patterns can be produced usingNewtonian flow. For example, Fig. 20 shows imprintedpatterns with intervals of 800 nm and 400 nm for Pt–Pd–Cu–P glassy alloys produced by die-forge pressingin the supercooled liquid region against focus ion beam(FIB)-machined Zr–Al–Ni–Cu glassy alloy dies [110].One can recognize regular arrays of imprinted patternswhich precisely reflect the shape of the die patterns.The nanoscale imprinting technology can create a glassyalloy surface with functional characteristics. As shown inFig. 20, the reflectance ratio of light on the glassy alloysurface changes in a wide range from 1% to >50%, andhence this technique is useful for the future developmentof novel optical materials. In addition, Fig. 21 showsimprinted glassy alloy nano-pillars with a diameter of200 nm and length �5 lm, obtained by pressing theglassy alloy into porous alumina in the supercooledliquid region [111]. The nano-pillars were tested for com-mercialization as anti-reflection material (see Fig. 21),cell culture medium for bio-chips, and electrode material.

Front view

Wit

hA

r p

re-e

tchi

ngW

itho

utA

r p

re-e

tchi

ng

Fig. 22. Preparation of 18 nm (2 Tbit in.�2) pitch nano-mold by deposited Pt mprevented by applying a Pt mask and subsequently removing contamination bpressure 0.1 Pa, HF power 50 W and duration 5 min. The resulting dots have

By the conventional nano-fabrication process using aphoto-resist coating method, various shapes of Ni alloyforging dies are fabricated, and diffraction grating patternsmade of glassy alloys are then produced by die-forgingglassy alloys to the Ni alloy dies in the supercooled liquidregion. The resulting diffraction grating patterns are justthe same as the original resist structure of grating. Whenthe interval of the grating patterns as well as the incidentand diffracted angles of light beam are accurately con-trolled, one can generate three different light colors ofred, green and blue, depending on the interval of the grat-ing patterns [112]. Using the different light color spectros-copy system, the use of BGA has been tried in hologramtechnology.

It was also confirmed that the minimum size achieved bythe imprinting technology using glassy alloys reaches12 nm for one bit and 25 nm for pitch size correspondingto the high density of 1 Tbit in.�2 [113]. In addition, theZr-based BGA imprinting die produced by the FIB tech-nique has fine precision size of 9 nm for one bit and18 nm for pitch size corresponding to the higher densityof 2 Tbit in.�2, as shown in Fig. 22. These developmentsof the nano-imprinting technique indicate the possibilityof applying this type of nanoscale imprinted glassy alloy

Tilted view (30 deg.)

ask and dry-etching for high-density data storage. Connection of dots wasy dry-etching. The etching conditions contain gas type CHF3 at 10 sccm,a conical shape.

Circuit part

Swagelok®

Metallic glass pipe CasingVibration sensor

Fig. 24. An outer appearance of the self-made Coriolis flow meter usingthe Ti–Zr–Cu–Ni–Sn glassy alloy pipe.


pattern to the next generation ultra-high density of infor-mation data storage material. Development of ways ofusing nanoscale imprinting technology is also supportedby the well-recognized road map of data-storage technol-ogy and recording density. It has generally been thoughtthat the patterned media have more advantages comparedwith the perpendicular recording system, which cannotavoid the recording density limit owing to heat fluctuation.

The formation of supercooled liquid is significantlyrelated to ductility, superplasticity and the highly anticor-rosive characteristic of BGA. These fundamental studieswere also performed by Schroers et al. as pioneering work[114–116], leading to the above recent applications usingsupercooled liquid formation of glassy alloys.

17. Applications

It was previously reported that Zr-based BGA in Zr–Al–Cu–Ni and Zr–Be–Ti–Ni–Cu systems were commercializedas sporting goods, i.e., golf clubs (driver, iron and putter),watch parts, electro-magnetic device casing, optical mir-rors, connecting parts for optical fibers, etc. In subsequentdevelopments, one can recognize some significant progressin application development studies as well as practicalapplications for BGA in Zr-, Ti- and Fe-based alloysystems.

17.1. Zr- and Ti-based alloys

The Zr–Al–Ni–Cu glassy alloy diaphragms with outerdiameters of 2.2–5 mm and heights of 4 mm produced bynet-shape casting have been applied to pressure sensorsfor automobiles and ordinary industries. The sensors haveseveral unique features of smaller size, higher sensitivityand higher pressure endurance which cannot be achievedfor conventional stainless-steel diaphragms produced by

(c)(c)

Zr-Zr-

Fig. 23. World’s smallest geared-motors with d = 0.9 mm successfully constru(b) constructed gear-reduction heads, where the top is the latest head with d

assemblies of geared-motors with d = 0.9 mm made of Zr-based metallic glassThe triangles in blue, yellow and red are set up to demonstrate the workabilmeasured rotation 20,000 rpm, no-load current 70 mA and starting torque 2 lNreader is referred to the web version of this article.)

cold working [117]. Micro-geared motor parts made ofZr–Al–Ni–Cu glassy alloy have also been produced bythe net-shape casting technique. Using these parts, theworld’s smallest geared-motor with diameter 1.5 mm andlength 9.9 mm has been produced and sold [118]. Thethree-stage micro-geared motor has a high torque �20times stronger than that for conventional vibration motorswith diameter 4 mm used in ordinary cell phones. Here, it isnoticed that the diameter of the micro-geared motor hasbeen further reduced to as small as 0.9 mm, as shown inFig. 23 [119]. In addition, various types of connectionadapters of curved sections and circular piping which canbe connected to micro-geared motors have been producedby the net-shape casting technique. Consequently, micro-geared motors in conjunction with various types of connec-tion adapters have been tested for commercialization inadvanced medical equipment, including endoscopes andmicro-pumps, precision optics and micro-machines.

Zr-based monolithicZr-based monolithic

based MG compositebased MG compositeMGMG

cted with metallic glass gears: (a) parts for geared-motor with d = 0.9 mm;= 0.9 nm, and the bottom is the conventional one with d = 1.5 mm; (c)

composites (left and center) and Zr-based monolithic metallic glass (right).ity of the motors. Specifications of geared-motors: applied voltage 0.9 V,m. (For interpretation of the references to colour in this figure legend, the

Fig. 26. Higher real part of permeability and higher imaginary part ofpermeability of “Liqualloy” sheet in the higher frequency range>500 MHz compared with Fe–Si–Al “Sendust” sheet [124].


Ti-based glassy alloy pipes 2 mm in outer diameter and0.2 mm thick produced by the suction casting techniqueexhibit a high tensile strength of 2100 MPa, large elasticelongation of �2% and high corrosion resistance [120].These properties are suitable for application as a sensingelement in the Coriolis flow meter to measure the Coriolisforce of liquid or gas flowing inside the pipe subjected toreinforced oscillation. Fig. 24 shows an outer appearanceof the self-made Coriolis flow meter using the Ti–Zr–Cu–Ni–Sn glassy alloy pipe. The sensitivity of the newly devel-oped Coriolis flow meter was reported to be 28–53 timeshigher than that for a conventional Coriolis flow meterusing SUS316 pipe [121]. The significant improvement insensitivity indicates that the new type of Coriolis flowmeter can be used in various industries, such as fossil-fuel,chemical, environmental, semiconductor and medical sci-ence fields.

17.2. Fe-based glassy alloys

Soft magnetic Fe-based glassy alloys in the Fe–Cr–P–C–B–Si system have been commercialized under the name“Liqualloy” [104,122], leading to the foundation of anew, relatively large company in which the application of“Liqualloy” in the sustainable development field is focusedon [123]. As typical soft magnetic properties, the Fe–Cr–P–C–B–Si alloy ribbon exhibits a saturation magnetization of1.25 T, a coercive force of 1–3 A m�1, effective permeabilityof 60 at 1 kHz and a high electrical resistivity of 150 lX cm,in conjunction with high glass-forming ability, good corro-sion resistance and low melting temperature. “Liqualloy”

magnetic cores have been produced on a mass productionscale by cold consolidation of a mixture of water-atomizedglassy alloy powders and epoxy resin and subsequentannealing at elevated temperatures. The application ofthe conventional water atomization process is attributedto the high glass-forming ability of bulk glassy type “Liqu-alloy”. The resulting “Liqualloy” cores exhibit nearly con-stant relative permeability in a high-frequency range up to

Appearance Structure

Edgewise coil

Metal pTerminal

4mm

Fig. 25. “Liqualloy” powder cores which have been used as power inductorsgeneration compared with “Sendust” core [122,124].

several megahertz, a good linear relation between perme-ability and DC bias field, much smaller reduction in perme-ability in a wide DC range, and much lower core lossesthan those for the Ni–Fe–Mo Permalloy core and theFe–Si–Al “Sendust” core. These excellent core loss charac-teristics are attributed partly to the reduction in eddy-cur-rent loss resulting from much higher electrical resistivity(307 kX cm) than those for Permalloy (0.5 X cm) and “Sen-dust” (1.7 kX cm) cores. When the “Liqualloy” core char-acteristics are compared with a Mn–Zn ferrite core withmuch higher electrical resistivity, the ferrite core exhibitsnearly the same low core losses as that for “Liqualloy”,but the ferrite core shows a significant reduction in perme-ability in a low current range and does not have as good alinear relation of permeability as a function of DC biasfield in a wide current range. This result demonstrates thatthe Mn–Zn ferrite cannot be applied to advanced soft mag-netic cores requiring high functional characteristics in high-current, low-voltage and high-frequency conditions as wellas small core size. Thus, one can notice that the Fe-based

Application

LiqualloyTM Core

owder

in laptop computers because of higher efficiency and much smaller heat


glassy alloy (“Liqualloy”) cores have a great advantage assoft magnetic cores compared with other conventional softmagnetic cores. At present, the “Liqualloy” powder coresexemplified in Fig. 25 are used as a power inductor in lap-top computers because of the higher efficiency and muchsmaller heat generation compared with the “Sendust” core[124].

The “Liqualloy” spherical powder produced by wateratomization can be deformed into flaky shape with thick-ness 2–3 lm and large aspect ratios of 10–30 by a colddeformation method. The “Liqualloy” sheet consisting ofthe flaky powder embedded in epoxy resin possesses a high

Fig. 27. Good antenna sensitivity results from the much higher-qualityfactor defined by the ratio of real part of permeability to imaginary part ofpermeability compared with “Sendust” sheet [124].

a b

c25002500

20002000

15001500

10001000

500500

0

Fig. 28. (a) SEM micrograph of SENNTIX-II” powder. (b) Relationship expreloss (Ph). (c) Changes in core loss in terms of hysteresis loss (Ph) and eddy-curSENNTIX I and II in glassy metal states [124].

conversion ratio from electromagnetic wave noise to heat,leading to a highly efficient suppression effect of the noise.The high efficiency is because the sheet exhibits a higherreal part of permeability and higher imaginary part ofpermeability in the high-frequency range >500 MHzcompared with Fe–Si–Al “Sendust” sheet, as shown inFig. 26 [122,124]. Therefore, the “Liqualloy” sheet has beencommercialized as an electromagnetic wave noise suppres-sion sheet in various electromagnetic instruments such asdigital still cameras.

Recently, the “Liqualloy” sheet has found anotherimportant application field as the essential componentmaterial in a radio-frequency identification system. Theuse of “Liqualloy” sheet between electro-magnetic devicesand loop antenna induces the magnetic field line into theantenna, resulting in a significant increase in the transmis-sion distance of the magnetic field line. As a result, one cansignificantly increase its antenna sensitivity at a commercialhigh carrier frequency of 13.56 MHz. The increase in sensi-tivity for the “Liqualloy” sheet compared with the Fe–Si–Al “Sendust” sheet reaches �20%. The good antennasensitivity results from the much higher-quality factordefined by the ratio of the real part of permeability tothe imaginary part of permeability at the carrier frequencycompared with the “Sendust” sheet, as shown in Fig. 27[122–124]. The radio-frequency identification system usingthe “Liqualloy” sheet has been used in NTT DoCoMo cellphones and can contribute to the development of electronicmoney systems.

Thus, the “Liqualloy” soft magnetic alloys possess manyremarkably efficient characteristics, leading to the develop-

K1

ssing the contribution of crystalline magnetic anisotropy (K1) to hysteresisrent loss (Pe) for pure iron powder in a crystal, Fe–Si–B in an amorphous,


ment of a sustainable society through the following advan-tages: low-cost raw materials, high recycling ability, highlyefficient productivity of spherical glassy alloy powder by acheap water atomization method, suitability for low-costmass production, high cold consolidation ability of the mix-ture of powder and epoxy resin to core form, high electricalresistivity of consolidated cores, low hysteresis loss of coreswith low coercive force and high permeability, and loweddy-current loss of cores with high electrical resistivity.These features have attracted special attention in respectof applications, and a new company “Green Device” wasfounded in May 2010 with funding support on the basisof recognition of contributing to “green innovation” byIndustry Innovation Corporation. The new companyintends to extend the “Liqualloy” soft magnetic materialas highly efficient high-frequency permeability and lowenergy loss magnetic materials to various fields such as vehi-cles, personal computers and smart electric power meters,and to gain wide recognition for “Liqualloy” as a new typeof international standard advanced magnetic core materi-als. The production scale of “Liqualloy” cores can be esti-mated to reach about one hundred million pieces per year.

Another type of soft magnetic powder core with a highersaturation magnetization of 1.3 T in Fe–Nb–B–Si and Fe–Nb–Cr–P–B–Si systems has also been developed

Gear

Shot + Air

Nozzle

Shot + Comp

Fig. 29. Alloy steel gears can increase fatigue strength by 50–80% compared wi

commercially in collaboration with another magnetic mate-rial company [105]. The cores were also produced by a sim-ilar procedure to that for “Liqualloy”, i.e., massproduction of spherical glassy alloy powders by wateratomization, followed by cold consolidation of the mixtureof glassy alloy powder and epoxy resin and then annealing.The new magnetic powder cores have been named “SEN-NTIX-I” and “SENNTIX-II”, respectively, and can becharacterized to exhibit the lowest core losses among alltypes of magnetic powder cores developed to date. In addi-tion, the higher saturation magnetization has enabled theuse of cores in a higher current range. The use of “SEN-NTIX-II” powder cores in the Fe–Nb–Cr–P–B–Si systemshown in Fig. 28 [124] can reduce core loss from the exist-ing metal powder cores by more than 50% and save ther-mal emissions from personal computers by �10 �C,resulting in the extension of the battery lifetime of note-book computers by >10%. Therefore, the “SENNTIX-II”

consolidated core has been supported as “Next Generation2009 and 2010 Reference Designs” by several major inter-national power corporation suppliers. The “SENNTIX-II”

consolidated cores have also been produced on a large scaleof several million pieces per month.

Fe–Ni–Cr–Mo–B–Si glassy alloy powders produced bywater atomization have also been commercialized with

ressed Air

Conven-tional

ProcessNo. 1

Conven-tional

ProcessNo. 2

NewProcess

DASP

th the steel gear subjected to peening shot with high-speed steel balls [124].

(a) (b)

Bending Area

Fig. 30. Results of bending test for (a) hard chromium plating and (b) metallic glass coating. No abrasion or swelling is found in either specimen,indicating that metallic glass coating has good adhesion with substrate.

Fig. 31. Outer appearance of Al-based alloy pipes with spray-coatedmetallic glasses in as-sprayed state and after being polished.


the commercial name “AMO-beads” [106]. The powderdiameter extends in a wide range from �0.05 to 0.6 mmbecause of the high glass-forming ability for the developedFe-based alloy. The application fields are extended to peen-ing shot balls and a fine precise polishing medium. The“AMO-beads” have the advantage of much longer endur-ance times compared with those for cast steel shot andhigh-speed steel shot. By use of good mechanical character-istics such as high Vickers hardness of �900, high fracturestrength of �3000 MPa and large elastic strain of �0.02 inconjunction with high corrosion resistance and a smoothouter surface, the peening shot treatment using “AMO-beads” can generate a higher level of residual compressivestress on the surface of high class alloy steel vehicle gearswith high Vickers hardness of 750 achieved by carburiza-tion treatment. As a result, the alloy steel gears can increasefatigue strength by 50–80% compared with the steel gearsubjected to peening shot using high-speed steel balls, assummarized in Fig. 29 [124]. This improvement causes asignificant reduction in the weight of alloy steel vehicle gearby �45%.

A highly dense Fe-based glassy-alloy-coated layer in Fe–Cr–Mo–C–B system has been produced on various metallicalloy substrates using the high-velocity powder-spray-coated layer technique [125]. The glassy-alloy-coated layerin Fe–Cr–Mo–C–B system exhibits better corrosion resis-tance than that for SUS304 plate, higher Vickers hardnessthan that for hard chromium plating plate and better wearresistance than those for hard chromium plating plate andCrN coating plate as well as SKD tool steel and FC cast iron[126]. In addition, it is confirmed in Fig. 30 that neither abra-sion nor swelling is observed in the Fe-based glassy alloy-coated layer even after repetition bending tests under theJapanese Industrial Standards (JIS) test condition, indicat-ing that the interface between the glass-coated layer andSUS304 steel substrate is in a good cohesion state. As exem-plified for Al-based alloy pipe in Fig. 31, the Fe-based glassyalloy-coated layer has a highly dense surface structure withdifferent coated layer thickness and changeable surfacesmoothness. In addition, the coated surface layer to metallicplate can be formed in a wide area up to�400 cm2. Owing toits advantages, the glassy-alloy-coating technique has beenapplied to the modification of surface properties for variousmetallic materials such as plain carbon steel, stainless steeland aluminum alloy.

18. Summary

As described in Section 2, studies of BGA originatedfrom the fundamental studies on structural relaxation ofglassy alloys that have been carried out since 1982. Itappears that the contents of these studies are independentof the BGA, but in reality the aim of obtaining more reli-able and quantitative data on structural relaxation led tothe discovery of a number of glass-type alloys with distinctglass transition and a large supercooled liquid regionbefore crystallization. Eventually, a number of BGA wereobtained in a variety of alloy systems, including systemsbased on La, Zr, Ti, Fe, Mg, Pd, Pt, Au, Cu, Ni and Ca.


Recent progress in BGA has enabled more sophisticatedand precise studies on features of glassy structure to be per-formed, such as medium-range order of glassy alloys. Fur-thermore, applications, engineering and standardizationare available for the first time since the first discovery ofthe amorphous alloy in 1960 as a result of the discoveryof BGA. Studies on the mechanical properties of glassyalloys proceed from success in obtaining glassy alloys inbulk form, leading to further detailed understanding ofmechanical properties based on shear bands. In addition,the low temperature properties of Ni- and Zr-based glassyalloys and porous Zr-based BGA have become a newresearch field in glassy alloys, which is also due to the rel-atively easy formation of BGA. Application fields of glassyalloys have spread widely, using features of the supercooledliquid formation of glassy alloys, such as nanoscaleimprinted patterns. It is expected that fields of applicationwill be significantly extended in the near future on the basisof the useful and unique engineering properties of BGAresulting from the simultaneous achievements of novelatomic configurations, unique multi-component alloy com-positions, various bulk forms, slow solidification process,Newtonian flow deformability and net-shape castingformability.

Acknowledgement

This work was partially supported by NEDO (New En-ergy and Industrial technology Development Organization)under a project “Technological Development of InnovativeComponents Based on Enhanced Functionality MetallicGlass”.

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http://project2005.msl.titech.ac.jp/symp2008/Abstract/Abstract-Saotome.pdf

http://project2005.msl.titech.ac.jp/symp2008/Abstract/Abstract-Saotome.pdf

http://www.nedo.go.jp/kankobutsu/pamphlets/kouhou/2008gaiyo_e/41_54.pdf

http://www.nedo.go.jp/kankobutsu/pamphlets/kouhou/2008gaiyo_e/41_54.pdf

http://www.alps.com/products/e/npv_product/090311_HMFTW/HMFTW-E.HTML

http://www.alps.com/products/e/npv_product/090311_HMFTW/HMFTW-E.HTML


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