Bulk Metallic Glasses

Bulk Metallic Glasses

C H A P T E R T H R E E Bulk Metallic Glasses: Formation, Structure, Properties, and Applications Dmitri V. Louzguine-Luzgin* and Akihisa Inoue Cont...

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C H A P T E R

T H R E E

Bulk Metallic Glasses: Formation, Structure, Properties, and Applications Dmitri V. Louzguine-Luzgin* and Akihisa Inoue Contents 131 132 136 141 146 152 152 155 156 158 163 165

1. 2. 3. 4. 5. 6.

Introduction Glass-Forming Ability and Formation Mechanisms Features of Glassy Structures Thermal Stability Mechanical Properties Magnetic Properties 6.1. Soft-magnetic alloys 6.2. Hard-magnetic alloys 7. Corrosion Resistance 8. Applications 9. Future Prospects References

1. Introduction Casting of commercial metallic alloys even into a slotted mold with a thin cavity (with a cavity thickness of about 1 mm) produces a crystalline structure. At the same time, glasses can be formed in various classes of materials: oxides, ionic compounds, polymers, etc. Metallic glassy alloys were produced by Klement et al. (1960) who showed the formation of the first Au–Si sample with an amorphous structure in 1960. Casting of metallic liquids at a very high cooling rate of 106 K s1 became possible by using a rapid solidification technique, which forced molten Au–Si and Pd–Si alloys to undergo vitrification on cooling. Pd–Cu–Si and Pd–Ni–P system glassy *Corresponding author: e-mail: [email protected], Tel: +81-22-217-5957 WPI Advanced Institute for Materials Research, Tohoku University, Aoba-Ku, Sendai, Japan Handbook of Magnetic Materials, Volume 21 ISSN 1567-2719, http://dx.doi.org/10.1016/B978-0-444-59593-5.00003-9

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2013 Elsevier B.V. All rights reserved.

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alloys were the first metallic glasses to be produced in the shape of 1–2-mm diameter rods (Chen, 1974). Larger samples can be produced after flux treatment, which helps to suppress heterogeneous nucleation (Kui et al., 1982). However, these materials did not attract much attention among the materials research community until the breakthroughs achieved at the end of the 1980s and the beginning of the 1990s. At the time, large-scale bulk metallic glassy (BMG) alloys (also called BMGs with the same abbreviation) arbitrarily defined as three-dimensional massive glassy (amorphous) objects with a size of not less than 1 mm in every dimension (10 mm by other definition) were produced (Inoue et al., 1989; Zhang et al., 1991), and today, they are attracting much attention among the scientific community. The high glass-forming ability (GFA) of some alloys allowed formation of BMGs upto about 100 mm in size by using various casting and water cooling processes. BMG alloys were obtained in a variety of alloy systems, including La-/Ln-, Zr-, Ti-, Fe-, Mg-, Pd-, Pt-, Au-, Cu-, Ni-, and Cabased alloy systems (Inoue 1995, 2000; Johnson, 1999) (Ln denotes lanthanides). Ferrous metal-based Fe- (Ponnambalam et al., 2004) and Co-based (Amiya and Inoue, 2008) BMG samples exceeding 1 cm in diameter were obtained recently. Large samples of hard magnetic Nd-based alloys were also produced (Inoue and Zhang, 1997).

2. Glass-Forming Ability and Formation Mechanisms Metallic glassy (amorphous) alloys with low GFA can be prepared in an amorphous state by condensation from a vapor phase (Gleiter, 1981) or by a solid state reaction using mechanical attrition (Fecht, 1995), for example, ball milling (Weeber and Bakker, 1988). However, the materials prepared by milling often contain residual nanoscale particles. Electrodeposition is another method for producing glassy alloys from a solution for glassy coatings (Ahmad et al., 2003; Yamasaki et al., 1998). However, homogeneity of their structure remains a question. Solidification of a liquid phase (Liebermann and Graham, 1976) by melt spinning was used to produce thin ribbons. Cu-mold gravity or die casting, liquid forging, and other techniques allowed production of bulk glassy alloys (Egami, 2010; Guo et al., 2003; Inoue et al., 1989, 1996a; Loffler, 2003; Peker and Johnson, 1993) through glass transition at Tg (Fig. 3.1). The bulk glassy alloys possess three common features summarized by Inoue (2000): (1) these alloys belong to multicomponent systems, (2) the constituent elements have significant atomic size ratios exceeding 1.12, and (3) most of the alloying elements in such alloys have a large and negative mixing

133

500

Tbg

550

600

650

700

Exothermic

Heat flow (a.u.)

Bulk Metallic Glasses

Tx

750

800

850

900

950 1000

Temperature (K)

Figure 3.1 DSC trace of Cu36Zr48Al8Ag8 BMG alloy on heating scanned at 83 mK s1. One can see the formation of a supercooled liquid from the glassy phase at the beginning of glass-transition temperature on heating Tbg which is often taken as Tg, large supercooled liquid region, exothermic peak starting at the crystallization temperature Tx and several subsequent exothermic peaks related to the phase transformations leading to the formation of equilibrium crystalline phases. The data were obtained by LouzguineLuzgin et al. (2009c).

enthalpy with each other. However, it was found that the higher GFA of the Ge–Ni–Nd alloy compared to the Si–Ni–Nd alloy cannot be explained on the basis of the widely used parameters, geometrical and chemical factors, and viscosity and diffusion data (Louzguine et al., 2002) and that the electronic structure characteristics (Hasegawa et al., 2009), as, for example, the electronegativity difference (Louzguine and Inoue, 2001a), should be taken into consideration. Although binary BMG alloys exist (Inoue et al., 2001; Wang et al., 2004a; Xu et al., 2004), their GFA is low and the critical thickness of glassy samples does not exceed 2 mm (except for the flux-treated Pd–Si alloys). Pd–Si binary BMG alloys with diameters ranging from 7 to 8 mm were prepared using a combination of fluxing and water cooling (Yao and Chen, 2008). However, some binary BMG alloys were reported to contain nanoparticles and they are formed in the narrow composition ranges. On the other hand, the addition of a certain third element drastically enhances their GFA. The role of minor additions in the formation of metallic glasses is discussed by Wang (2007). Recent investigations showed that large BMG samples have similar structures and possess the thermal properties close to those of the corresponding glassy alloy samples produced by melt spinning (Louzguine-Luzgin et al., 2009c). Glassy or amorphous powder samples produced by ball milling or gas atomization techniques (Inoue, 1998) can be consolidated into bulk form

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Exothermic

DT (a.u.)

using hot pressing (Inoue and Kimura, 2001a), spark plasma sintering (SPS) (Lee et al., 2007; Xie et al., 2006a), or microwave furnace sintering (Louzguine-Luzgin et al., 2009b) processes. For example, Ni52.5Nb10Zr15Ti15Pt7.5 BMG samples produced by Xie et al. (2007) with a size of 20 mm and nearly 100% relative density were fabricated by SPS of gas-atomized glassy powders at the sintering temperature of 773 K under a loading pressure of 600 MPa. The consolidation was achieved using a relatively low sintering temperature and short holding time. Al-based bulk glassy samples of high relative density were obtained by warm extrusion of the atomized amorphous powders (Inoue and Kimura, 2001b) though they have low GFA (Allen et al., 1998), which imposes a limit on the critical thickness of glassy samples below 1 mm. Microwave heating induces fast volumetric heating of a powdered sample within a few minutes (Buchelnikov et al., 2008) and rapid sintering (Xie et al., 2009). Different criteria used to correlate with the GFA include the reduced glass transition temperature, Trg ¼ Tg/Tl by Turnbull and Cohen (1961) where Tg is the glass-transition temperature and Tl is the liquidus temperature (Figs. 3.1 and 3.2) (though the overall validity of this criterion has been questioned recently, e.g., by Senkov, 2007; Louzguine-Luzgin et al., 2008a). Another well-known criterion is the width of the supercooled liquid region (DTx ¼ Tx  Tg) (Inoue et al., 1993) where Tx is the onset crystallization temperature; g ¼ Tx/(Tg þ Tl) is a parameter introduced by Lu and Liu (2002), which combines both DTx and Tg/Tl criteria into a single parameter, d is presented by Louzguine-Luzgin and Inoue, 2007a. Many other criteria have been summarized by Suryanarayana and Inoue (2010). Louzguine-Luzgin and Inoue (2007a) predicted an “ideal” isochoric

Tbc Tfc

1070

1075

1080

1085

1090

1095

1100

1105

Temperature (K)

Figure 3.2 DTA trace of Cu36Zr48Al8Ag8 BMG alloy on cooling scanned at 17 mK s1. Beginning of crystallization Tbc and finish of crystallization Tfc temperatures are marked. Because of liquid supercooling Tl is obtained by extrapolating Tbc obtained at different cooling rates to infinitely slow cooling.

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135

Tg (temperature at which the specific volume of the liquid and solid phases becomes equal) of about 1100 K for pure Ni and 1200 K for pure Fe. Quantum mechanics (Louzguine-Luzgin et al., 2008a) as well as classical (Evteev et al., 2004) molecular dynamics simulations of Ni and Fe confirmed vitrification of the liquid phase at temperatures very close to these temperatures. Moreover, as has been shown, purely extrinsic factors also have a significant influence on the GFA (Louzguine-Luzgin et al., 2008c). The factors influencing GFA are categorized as intrinsic (belonging to a glass itself) or extrinsic (depending upon external conditions) factors. Intrinsic factors assume that homogeneous nucleation competes with glass formation and include a number of fundamental and derived thermal parameters (Tg, Tx, Tl, Trg); physical properties such as heat capacity, thermal conductivity and diffusivity, and thermal expansion coefficient; compositional proximity to a deep eutectic reaction; and a topological contribution from efficient atomic packing in the atomic structure. Extrinsic factors are usually operative when heterogeneous nucleation intervenes during solidification. Important extrinsic factors include inclusions or dissolved impurities in the melt, poor mold surface finish or cleanliness, turbulence during solidification, and the degree of liquid metal superheat. GFA of the glasses with high intrinsic GFA can be significantly limited by extrinsic factors, which should be taken into consideration to predict an actual GFA of the alloy. Because casting conditions of bulk glassy samples are far from the equilibrium by Louzguine-Luzgin et al. (2010a), the best glass-forming alloy compositions do not correspond to the equilibrium eutectic point, but shift usually toward a more refractory eutectic component (Tan et al., 2003) owing to the shift of the eutectic point with supercooling/undercooling at high enough cooling rates. The importance of efficient atomic packing for the formation of metallic glasses was shown recently by Miracle et al., 2003; Miracle, 2004, 2006 and Sheng et al. (2006), as the specific radius ratios are preferred in the compositions of metallic glasses. These features are also connected to the l criterion for good GFA (Egami and Waseda, 1984). Although the glass-transition phenomenon in metallic glasses has been studied extensively there is still no common conclusion on its nature (Angell, 2001; Cohen and Grest, 1979; Van Den Beukel and Sietsma, 1990). In some works, the glassy phase is treated as a frozen liquid, and thus, glass transition is a kinetic phenomenon and no thermodynamic phase transformation takes place. On the other hand, glass transition may be a second-order transformation as follows from the shape of the curves for the thermodynamic parameters, which exhibit a continuity at the glass-transition temperature while their derivatives such as thermal expansion coefficient or heat capacity exhibit a discontinuity (in a certain approximation) at the glass-transition temperature. Glass transition may

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also be a first-order transformation and a thermodynamic aspect of glass transition is known as the Kauzmann paradox (Kauzmann, 1948). Ferrous metal-based BMGs comprise an important group of such materials. Although atomic size differences between Fe and metalloids such as B, C, Si are significant, the differences in these values between Fe and the alloying elements that caused the formation of BMGs (e.g., Nb, Mo, or Al) are not so large compared to other well-known BMGs such as La–Cu–Al, Mg–Y–Cu, Zr–Ni–Cu–Al, or Pd–Ni–Cu–P alloys. Mo, Cr, and Mn are known as elements that enhance the hardenability of steels upon quenching by hampering the eutectoid decomposition of austenite. Nb likely produces the same effect. Thus, one can suggest that both Mo and Nb may retard crystallization of the liquid phase in BMGs as well. On the other hand, soft magnetic Fe–B–C (Wang et al., 2011) and Fe– Si–B–P–C (Makino et al., 2009) BMGs without any glass-forming metallic alloying elements were developed recently though their diameters are limited to less than 1.5 and 3 mm, respectively. B-rich Fe66Nb4B30 ferromagnetic bulk glassy alloy was produced by fluxing and casting (Stoica et al., 2009). Typical ferrous and ferromagnetic metal-based bulk glassy alloy systems reported up to date together with the calendar year when the first scientific report, paper, or patent on each alloy system was published are listed as follows: Fe–(Al,Ga)–(P,C,B,Si,Ge) (1995); Fe–(Nb,Mo)–(Al, Ga)–(P,B,Si) (1995); Co–(Al,Ga)–(P,B,Si) (1996); Fe–(Zr,Hf,Nb)–B (1996); Co–(Zr,Hf,Nb)–B (1996); Fe–Co–Ln–B (1998); Fe–Ga–(Cr, Mo)–(P,C,B) (1999); Fe–(Cr,Mo)–(C,B) (1999); Ni–(Nb,Cr,Mo)–(P,B) (1999); Co–Ta–B (1999); Fe–Ga–(P,B) (2000); Ni–Zr–Ti–Sn–Si (2001); Ni–(Nb,Ta)–Zr–Ti (2002); Fe–Si–B–Nb (2002); Co–Fe–Si–B–Nb (2002); Ni–Nb–Sn (2003); Co–Fe–Ta–B–Si (2003); Ni–Pd–P (2004); Fe–(Cr, Mo)–(C,B)–Ln (Ln ¼ Y, Er, Tm) (2004); Co–(Cr,Mo)–(C,B)–Ln (Ln ¼ Y, Tm) (2005); Ni–(Nb,Ta)–Ti–Zr–Pd (2006); Ni–Pd–P–B (2009); Fe–(Nb,Cr)–(P,B,Si) (2010).

3. Features of Glassy Structures The structure of various glassy alloys has been studied by diffraction experiments and computer modeling (Bernal, 1960; Haruyama et al., 2007; Hirata et al., 2006a; Kramer et al., 2007; Matsubara and Waseda, 1995; Miracle et al., 2007; Suzuki, 1982; Waseda and Chen, 1978a; Yavari, 2006). For example, the atomic structure of amorphous Fe–B and Fe–Nb–B alloys was studied by means of electron diffraction and high-resolution electron microscopy by Waseda and Chen (1978b) and Hirata et al., (2004, 2006b). The obtained experimental results and structure models created for Fe80B20

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137

and Fe70Nb10B20 indicated that the prismatical clusters observed in some Fe-boride compounds form around B atoms. The atomic arrangements of deformed bcc clusters with a large coordination number were frequently observed around Fe atoms in Fe80B20. Also, the higher GFA of Fe70Nb10B20 alloy compared to Fe–B is ascribed to the high atomic packing density around Fe atoms. The structure of Fe–B alloys was also investigated by EXAFS studies on the iron K-edge (Defrain et al., 1984; Gurman, 1982; Waseda and Chen, 1978b). The iron atom was found to have 2.2 neighbor boron atoms at 0.206 and 8.2 neighbor iron atoms at 0.255 nm. The Fe–B distance was found to be close to the sum of the covalent radii, while the Fe–Fe distance is close to that found in metallic iron. Partial pair distribution functions (PDFs) for amorphous Ni81B19 were also obtained by Lamparter et al. (1982) and compared with the calculated functions. Important steps toward a better understanding of the atomic structure in the glassy state have been taken recently. In the case of weakly interacting Cu–Zr atoms, an “ideal solid solution” behavior (weak chemical ordering) was suggested and recently proven by Georgarakis et al. (2009). The structure of Zr–Cu binary metallic glasses except for Zr67Cu33 one can be modeled by statistics of Cu–Zr, Cu–Cu, and Zr–Zr nearest neighbor (NN) populations in an “ideal solid solution” (Fig. 3.3). As an example of the effect of the third component in the alloy on the structure of the glassy phase, one can consider Cu–Zr–Al and Cu–Zr–Ag alloys. The addition of Al strongly modifies the atomic structure of Cu–Zr BMG alloys because of strong attractive interactions between Al and Zr leading to the formation Zr–Al NN populations far in excess of the random atomic distribution of an “ideal solid solution” (Fig. 3.4). With Al atomic size intermediate between those of Cu and Zr, formation of a large number of Zr–Al NN pairs results in a broader dispersion of first NN interatomic distances apparently leading to a higher packing efficiency. Cu45Zr45Ag10 alloys exhibited one of the highest GFAs among Cu–Zr– Ag glassy alloys (Zhang and Inoue, 2006). Ag also drastically improves the GFA of a Cu–Zr–Ti alloy (Dai et al., 2006) and the Cu–Zr–Al alloys (Sung et al., 2004; Zhang et al., 2007). Figure 3.5 shows PDFs of Cu50Zr50, Cu45Zr45Ag10, and Cu35Zr45Ag20 glassy alloys studied by synchrotronradiation XRD, as an example, to illustrate the effect of the third element on the structural features (Louzguine-Luzgin et al., 2009a). The Gaussian function fitting of the first PDF maximum from 0.20 to 0.38 nm (baseline corrected) using two Gaussian peaks gave a reasonable correspondence to the original PDF data of the alloys studied and the results are shown in Table 3.1. As one can see, the increase in Ag content causes a shift in two peaks’ maxima in the first coordination shell toward larger distances owing to the larger atomic size of Ag compared to Cu. At the same time, the shape of the first coordination shell becomes more symmetrical

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Zr40Cu60

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

Goldschmidt radii

0.35

0.30

Zr–Zr

Zr–Cu Cu–Cu

Zr–Zr

Cu–Cu

Zr–Cu

Pair distribution function

0.25

Zr50Cu50

(b)

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

Pair distribution function

(a)

Goldschmidt radii

0.25

0.30

r (nm) Zr67Cu33

(c)

Zr–Cu

Zr–Zr

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

Cu–Cu

Pair distribution function

0.35

r (nm)

Goldschmidt radii

0.25

0.30

0.35

r (nm)

Figure 3.3 Gaussian fitting of the first PDF maximum for (a) Zr40Cu60, (b) Zr50Cu50, and (c) Zr67Cu33 metallic glasses. The thick continuous black curves correspond to the experimental PDF, the thin reddish continuous curves correspond to partial PDFs derived from the deconvolution of the experimental PDF into three Gaussians corresponding to Cu–Cu, Zr–Cu, and Zr–Zr partial PDFs. The thin blue discontinuous curves represent the Gaussian partial PDFs if they are scaled to the weight factors of an ideal solution and the thick discontinuous bluish curves represent the expected total PDFs based on the weight factors. Reprinted from Georgarakis et al. (2009) with permission of the American Institute of Physics.

and the second peak (shoulder) of the first diffraction maximum becomes weaker. The integration of the main RDF maximum (both peaks from 0.22 to 0.36 nm) produces the area under the peak (N) as shown in Table 3.1. According to Fig. 3.5, a detectable degree of medium-range order (MRO) (Waseda et al., 2008) maintains until about 2–2.5 nm of the interatomic distance. Such a high degree of MRO is well consistent with recent models, which predict that metallic glasses are not a random packing of atoms (Bernal, 1960), but a dense packing of clusters (Miracle, 2006; Sheng et al., 2006); this is also deduced from the small volume difference between glassy and the corresponding crystalline phases (Yavari, 2006). An icosahedral-type MRO is expected to be a dominant local configuration in liquids and metal–metal-type metallic glasses (Kelton et al., 2003).

139

Bulk Metallic Glasses

0.30 r (nm)

(c)

Zr–Al

Cu–Al

Al–Al

0.25

0.35

Goldschmidt radii

0.30 r (nm)

0.35

Zr50Cu20Al30

0.25

Zr–Zr

Zr–Al

Al–Al

Cu–Al

Zr–Cu

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

Cu–Cu

Pair distribution function

Zr–Cu

Cu–Cu

Goldschmidt radii

Zr–Zr

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

Pair distribution function

Zr–Zr Zr–AI

Cu–AI

AI–AI

Zr–Cu

PDF Gaussian fitting Partial Gaussian fitting Partial / weight factor Total / weight factor

0.25

Zr50Cu30AI20

(b)

Zr50Cu40AI10

Cu–Cu

Pair distribution function

(a)

Goldschmidt radii

0.30

0.35

r (nm)

Figure 3.4 Gaussian fitting of the first PDF peak for (a) Zr50Cu40Al10, (b) Zr50Cu30Al20, and (c) Zr50Cu20Al30 metallic glasses. The thick continuous black curves correspond to the experimental PDFs, and the thin reddish continuous curves correspond to Gaussian partial PDFs derived from the deconvolution of the experimental PDF. The thin blue discontinuous curves represent the partial PDFs scaled to the ideal solution weight factors and the thick discontinuous bluish curves represent the expected ideal solution total PDFs. Reprinted from Georgarakis et al. (2009) with permission of the American Institute of Physics.

Table 3.1 Structural information obtained from PDF and RDF curves of the studied alloys within the first coordination shell Glassy alloy

Cu55Zr45

Cu45Zr45Ag10 Cu35Zr45Ag20

First peak of PDF (Cu–Zr): CM (nm), RA Second peak of PDF (Zr–Zr): CM (nm), RA N from RDF

0.269, 0.60 0.275, 0.68

0.284, 0.89

0.315, 0.40 0.318, 0.32

0.333, 0.11

11.0

10.7

11.3

CM, center of mass; RA, relative area; N, coordination number. Reproduced from Louzguine-Luzgin et al. (2009a) with permission of Materials Research Society.

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Pair distribution function

2.5

Pair distribution function

2.0

1.5

1.01

1.00

0.99

0.98 1.0

1.5

2.0

2.5

R (nm)

1.0 Cu55Zr45 0.5

Cu45Zr45Ag10 Cu35Zr45Ag20

0.0 0.2

0.3

0.4

0.5

0.6

0.7

R (nm)

Figure 3.5 PDFs of Cu–Zr–Ag alloys. The insert is the PDF of Cu45Zr45Ag10 alloy at relatively large R values. Reprinted from Louzguine-Luzgin et al. (2009a) with permission of Materials Research Society.

Although PDFs of good glass formers such as Zr–Cu–Al (Georgarakis et al., 2009), Zr–Cu–Ni–Al, and Zr–Cu–Ni–Al–Ti (Dmowski et al., 2007) exhibit two peaks in the first coordination shell similar to the first PDF maximum in the first coordination shell of the Cu45Zr45Ag10 and Cu–Zr– Al alloys, the split of two peaks in the first coordination shell is much less pronounced compared to Zr–Ni–Al alloys (Georgarakis et al., 2010). A similar difference is observed in the case of Cu–Zr and Ni–Zr glasses, which have significantly different GFA, and the Zr–Zr distance is much better resolved in the PDF of Ni–Zr alloys (Hirata et al., 2007). Splitting of the first maximum in Zr–Ni–Al alloy systems also correlates to lower GFA of these alloys compared to Zr–Cu–Al alloys. Coefficient of determination R2 obtained for a single Gaussian function fitting indicates a more symmetrical shape of the first PDF maximum of the Cu45Zr45Ag10 alloy (R2 ¼ 0.971) in comparison with that of Cu55Zr45 (R2 ¼ 0.919) also correlates with its higher GFA compared to Cu55Zr45 alloy. Cu55Zr45 alloy clearly demonstrates two peaks in the first coordination shell and its GFA is significantly lower. On the other hand, the Cu35Zr45Ag20 alloy, which demonstrates a nearly similar symmetrical shape of the first peak, has lower GFA compared to the Cu45Zr45Ag10 glassy alloy likely due to the phase separation observed in this alloy on heating (Louzguine-Luzgin et al., 2007b). The shape of the maximum in the first coordination shell (at least in the Cu–Zr–Ag alloy systems) may

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Bulk Metallic Glasses

2 nm

Figure 3.6 MRO zones in Ni50Pd30P20 alloy, encircled in the high-resolution TEM (HRTEM) image.

possibly be connected with the GFA of the alloy as it is defined by the type of atomic clusters found in the glassy alloy and its MRO. It may be responsible for the formation of a more dense glassy structure. One should also mention that some bulk glassy alloys, especially Pd–Si (Ohkubo and Hirotsu, 2003), Pd–Ni–P (Haruyama et al., 2007; Hirata et al., 2007; Hirotsu et al., 1998), Pd–Cu–Si (Hirotsu et al., 1986), and Ni–Pd–P (Louzguine-Luzgin et al., 2007d) contain clear MRO zones and even nanoscale particles (Jiang et al., 2003; Louzguine-Luzgin et al., 2012) in as-solidified state even though these precipitates do not produce diffraction peaks in the XRD and selected-area electron diffraction pattern owing to their small volume fraction. MRO zones observed in the Ni50Pd30P20 alloy are shown in Fig. 3.6. The thermal expansion, glass-transition Tg, and volume change upon heating have been studied by dilatometry by Schermeyer and Neuha¨user (1997) and by X-ray radiation diffraction for Fe40Ni40P14B6 by Egami (1978) (structural relaxation), for Pd40Cu30Ni10P20 by Mattern et al. (2003), for Zr55Cu30Ni5Al10 by Yavari et al. (2004a,b) and Yavari et al. (2005), and for La-based BMG by Jiang et al. (2010) as well as for Cu55Hf25Ti15Pd5 and Cu55Zr30Ti10Ni5 glassy alloys by Louzguine et al. (2005).

4. Thermal Stability On heating, metallic glasses exhibit relaxation below Tg and then many of them transform to a supercooled liquid before crystallization. The process of structural relaxation leads to densification of the glassy phase and causes variation of its properties. In some cases (Cu55Hf25Ti15Pd5

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Table 3.2 Density (r) and volume changes (DV) in the Cu55Zr30Ti10Pd5 BMG alloy (2 mm) in different structure states

Parameter

Peak 1 Relaxed I-phasea Asprepared glass (740 K, (785 K, 120 s) 300 s) glass

r (Mg m3) 7.518 DV (%) Origin

7.580 0.8

7.592 1.0

Peak 2 MIMCb (850 K, 180 s)

Peak 3 SIMCc (993 K, 300 s)

7.650 1.8

7.591 1.0

a

Icosahedral phase. Metastable intermetallic compound. Stable intermetallic compounds.

b c

and Cu55Zr30Ti10Pd5 BMG alloys), large volume changes up to 0.8% (see Table 3.2) were detected by Louzguine et al. (2005) and by LouzguineLuzgin et al. (2007c), which may indicate the beginning of clustering in these alloys below thermally manifested Tx. Bulk glassy alloys can be thermomechanically shaped or welded in the supercooled liquid region before they crystallize. This property is used in micromachine technology for creating small gears having diameters ranging from 0.1 to 1 mm using the Newtonian flow and by making submicronscale patterns (Kumar et al., 2009; Saotome et al., 2001a). Shaping of BMGs can also be done by the electromechanical shaping technology at low applied stresses due to the high electrical resistivity of glassy alloys (Yavari et al., 2004a,b). As shown by Kawamura et al. (1997), some of the glassy alloys exhibit “superplasticity” (actually good fluidity) on heating to a supercooled liquid region. Bonding of glassy alloys can be achieved by laser (Louzguine-Luzgin et al., 2008b), electron beam (Louzguine-Luzgin et al., 2007a), and friction welding (Shoji et al., 2003). Metallic glasses are metastable at room temperature and devitrify/crystallize on heating. Such a process leads to the formation of a highly homogeneous dispersion of nanoparticles in various alloys. Nanostructured materials exhibit unique and superior properties, and thus, they are subjects of great interest to scientists in various fields of physics, chemistry, and materials science (Gleiter, 1989; Greer, 1995). High nucleation rates leading to a high number density of the precipitates above 1021 m3 and lowgrowth rates of the precipitating phase are required in order to obtain a nanostructure (Perepezko and Hebert, 2002). The difference in the devitrification pathways of glassy alloys is often connected with the state of the matrix phase prior to devitrification: amorphous, glassy, or supercooled liquid. Amorphous alloys do not transform to a supercooled liquid upon conventional heating. Glassy alloys form the supercooled liquid region on heating (see Fig. 3.1) prior to crystallization

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143

and in general have a better GFA compared to amorphous alloys, which do not exhibit Tg on heating. The formation of the supercooled liquid influences the devitrification process in metallic glasses (Louzguine-Luzgin and Inoue, 2007b). It may be connected with the change of the local atomic structure in the supercooled liquid region due to higher atomic mobility compared to that in the glassy phase by Louzguine and Inoue (2004). Nanostructured alloys are readily obtained on the primary devitrification of glasses with a long-range diffusion-controlled growth. Another type of phase transformation in an amorphous solid leading to the formation of a nanostructure is spinodal decomposition (Cahn, 1961). It is also found that ultrasonic vibrations promote the crystallization of Pd40Ni40P20 BMG alloy (Ichitsubo et al., 2004). In many cases, diffusive redistribution of the alloying elements on a short scale precedes crystallization. The changes in the amorphous halo peak occur after heating Mg–Ni–Mm and Mg–Ni–Y–Mm glassy alloys (where Mm denotes Mischmetal) to the first broad DSC exothermic peak, which indicates redistribution of the alloying elements in the amorphous matrix forming Mg-enriched zones without an incubation period (Louzguine et al., 2003a). Four types of phase transformations take place in glassy alloys on heating: polymorphous (the product phase has the same composition as the glassy phase), primary (its composition is different from that of the glassy phase), eutectic/eutectoid transformation (two or more phases nucleate and grow conjointly), and spinodal or binodal decomposition involving a phase separation of the glassy phase prior to crystallization/devitrification (LouzguineLuzgin and Inoue, 2005). The devitrification kinetics of glassy alloys can be studied by Kolmogorov– Johnson–Mehl–Avrami analysis (Kolmogorov, 1937, for example) and many glassy alloys obey this mechanism though there are many exceptions from the rule. Nanoscale size intermetallic compounds can also be formed. For example, the devitrification of the Ti50Ni20Cu23Sn7 alloy begins from the primary precipitation of nanoscale equiaxed particles of cF96 (Pearson Symbol) Ti2Ni solid solution (He et al., 2003; Louzguine and Inoue, 2000). The growth rate of the cF96 phase at a constant temperature is nonlinear, which indicates the diffusion-controlled growth mechanism. An extremely low growth rate of the cF96 crystals was observed on the primary crystallization of the Hf55Co25Al20 glassy alloy (Louzguine et al., 2003b). Very small cF96 Hf2Co clusters 2–5 nm in size are formed in the sample annealed at 907 K for 0.9 ks, which causes the appearance of the broad X-ray diffraction peaks. These clusters are not visible in the bright-field TEM images, though the XRD pattern of the sample annealed for 0.9 ks at 907 K shows a broad diffraction peak from about 30–45 2y, which can be separated into five narrower peaks (Fig. 3.7) corresponding to the cF96 phase.

144

25

30

(511)

(440)

(331)

(400)

Intensity (a.u.)

(422)

Dmitri V. Louzguine-Luzgin and Akihisa Inoue

35

40

45

50

2q (⬚)

Figure 3.7 XRD pattern of the Hf55Co25Al20 glassy alloy annealed at 907 K for 0.9 ks. The location of the five strong peaks (Gaussian fitting) corresponds to that of the cF96 Hf2Co phase. Reprinted from Louzguine et al. (2003b) with permission of Elsevier.

Comparison of the long-term thermal stabilities of different metallic glasses has been carried out using continuous heating transformation (CHT) diagrams (Louzguine and Inoue, 2002) constructed by applying a corollary from the Kissinger analysis method. CHT diagrams also can be recalculated from the isothermal diagrams using a method (LouzguineLuzgin and Inoue, 2009). Nanoscale quasicrystals are formed on devitrification in various metallic glassy alloys (Louzguine-Luzgin and Inoue, 2008). An icosahedral quasicrystalline phase (a three-dimensional quasicrystal, though two- and one-dimensional quasiperiodic structures also exist) having a long-range quasiperiodic translational and an icosahedral orientational order, but with no three-dimensional translational periodicity was initially discovered in Al–Mn alloys by Shechtman et al. (1984) and later in some other binary Al–TM (TM-transition metal) alloys as well as in different ternary Albased alloys. After this, the icosahedral phase was observed in Ga-, Ti-, Mg-, and Pd-based alloys as well as in Cd-, rare earth-, and Zn-based alloys as reviewed by Kelton (2000) and Ranganathan and Inoue (2006). The formation of the nanoscale icosahedral phase was observed in the devitrified Zr–Cu–Al, Zr–Al–Ni–Cu (Ko¨ster et al., 1996) and Zr–Ti–Ni– Cu–Al (Xing et al., 1998) glassy alloys containing an oxygen impurity above about 1800 mass ppm, although no icosahedral phase is formed if the oxygen content is lower than 1700 mass ppm. The nanoscale icosahedral phase was obtained in devitrified Zr–Al–Ni–Cu–Pd (Inoue et al., 1999),

145

Bulk Metallic Glasses

Zr–Pd, Zr–Pt (Murty et al., 2001), and other alloy systems at a much lower (about 800 mass ppm) oxygen content. The nanoscale icosahedral phase has been produced in the noble metal-free Zr–Cu–Ti–Ni (Louzguine and Inoue, 2001b) and Zr–Al–Ni–Cu (Saida et al., 2002) glassy alloys with low oxygen content, below 500 mass ppm. It has been found that reduced supercooling before crystallization from the melt was found to be the smallest for quasicrystals, larger for crystal approximants (crystals having structures somewhat similar to those of certain quasicrystals), and the largest for crystalline phases (Kelton et al., 2003). The nucleation barrier scales with the supercooling, and thus, local icosahedral order is considered to exist in some supercooled liquids and glasses. A low energy barrier for nucleation of the icosahedral phase may explain the fact that growth of the preexisting icosahedral nuclei was observed in the Zr65Ni10Al7.5Cu7.5Ti10Ta10 alloy (Ouyang et al., 2003). The nanoscale icosahedral quasicrystalline phase has been also produced upon heating glassy Hf-based alloys containing Pd or Au (Li et al., 2000; Louzguine et al., 2000). The formation of the nanoscale icosahedral phase was observed in the Cu-based alloys containing Pd (Louzguine and Inoue, 2003) or Au. A clearly heterogeneous nucleation was observed during formation of the Fe nanocrystals (Yavari and Drbohlav, 1995). The devitrification of the Fe73.5Cu1Nb3Si13.5B9 alloy starts from the formation of Cu-enriched zones (Hono, 2002). Precipitation of a nanoscale a-(Fe,Co) phase was observed in the Fe40Co40Cu0.5Zr9Al2Si4B4.5 alloy (Fig. 3.8; Mitra et al., 2004). As has been shown by means of atom probe field ion microscopy as well as by high-resolution transmission electron microscopy (Hono et al., 1992), Cu atoms form nanoclusters in the Fe73.5Si13.5B9Nb3Cu1 amorphous matrix, which act as the nucleating sites for heterogeneous nucleation of

100 nm

Figure 3.8 Bright-field TEM image of Fe40Co40Cu0.5Zr9Al2Si4B4.5 alloy annealed for 15 min at 873 K showing the formation of nanoparticles. Inset shows the selected-area electron diffraction pattern.

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the bcc Fe particles on devitrification (Botta et al., 1999). The density of the clusters estimated by Ohnuma et al. (1999) using a three-dimensional atom probe is on the order of 1024 m3 at the average cluster size of about 2–3 nm. The studies for X-ray absorption fine structure also showed that Cu clusters with near-fcc structure were present from very early stages of the devitrification process (Ayers et al., 1998). EXAFS analysis of grain boundaries in the nanocrystalline Fe85Zr7B6Cu2 alloys also showed a low Fe content in the grain boundaries between the bcc Fe solid solution nanograins. Yavari and Negri (1997) discussed the nanocrystallization process of soft magnetic Fe-based amorphous alloys using the concentration gradients of the elements that are insoluble in the primary crystalline phase. One should also mention the electron-beam irradiation-induced crystallization (Nagase and Umakoshi, 2004; Xie et al., 2006b). A general observation is that it causes primary nanocrystallization even in alloys that exhibit a eutectic transformation mechanism on heating. It rather indicates a nonthermal crystallization.

5. Mechanical Properties Owing to the absence of crystalline lattice and dislocations, a unique deformation mechanism is realized in bulk glassy alloys (Argon, 1982; Liu et al., 2006), which thus exhibit high strength (2 GPa for Cu-, Ti-, and Zr-based, 3 GPa for Ni-based,  4 GPa for Fe-based, and 5 GPa for Co-based alloys), high hardness, good wear resistance (Togashi et al., 2008), and large elastic deformation (Yavari et al., 2007). For example, (Fe,Co)– Cr–Mo–C–B–Tm glassy alloys prepared by Amiya and Inoue (2008) in a cylindrical form with a diameter of 18 mm demonstrate an excellent GFA and high strength exceeding 4 GPa. Nevertheless, as localized shear deformation is a dominant mode of plastic deformation at room temperature, tensile ductility of metallic glasses is not found except in a few special cases in very small samples or/and high strain rates (Song et al., 2011; Yokoyama et al., 2009). The tensile deformation behavior of Zr-based glassy thin foils has also been studied recently in situ in TEM (Guo et al., 2007; LouzguineLuzgin et al., 2010b). Novel fracture behavior of Pd-based metallic glasses was reported recently by Chen et al. (2009). Co43Fe20Ta5.5B31.5 glassy alloys exhibited an ultrahigh fracture strength exceeding 5 GPa, a high Young’s modulus of 270 GPa, high specific strength, and a high specific Young’s modulus (Inoue et al., 2003). The strength and specific strength values exceed those reported for any bulk glassy alloys. Large elongations of 1400% were obtained in the supercooled liquid region indicating good fluidity. Strong bulk glassy alloys exhibiting a wide supercooled liquid region before crystallization were produced (Inoue et al., 1997a) in

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Bulk Metallic Glasses

Fe–(Co,Ni)–(Zr,Nb,Ta)–(Mo,W)–B systems. These alloys have a high Tg of about 870 K and the supercooled liquid region close to 90 K. The bulk glassy alloys with diameters up to 6 mm exhibit a high compressive strength of 3.8 GPa, a high Vickers hardness of HV1360, and high corrosion resistance. Although the majority of BMG alloys are brittle even under compression, some BMGs exhibit a significantly higher compressive plasticity (Greer and Ma, 2007; Schuh et al., 2007) compared to the others (Fig. 3.9). Such an improved plasticity can be related to Poisson’s ratio n (Lewandowski et al., 2005), in situ nanocrystallization (Hajlaoui et al., 2006), or glassy phase separation (Kim et al., 2006). Some glassy alloys contain preexisting nanoparticles (Inoue et al., 2005) or demonstrate nanocrystallization within shear bands on deformation. It was supposed and illustrated experimentally (Louzguine-Luzgin et al., 2007d) that a glassy alloy, which crystallizes by the nucleation and growth mechanism, may be less prone to nanocrystallization during deformation while the alloy having preexisting nuclei may be predisposed to show such a behavior. While the strain-rate sensitivity (SRS) of polycrystalline alloys is generally positive, for BMGs it was reported to be positive, negative, or zero. Hufnagel et al. (2002) reported a negative SRS for Zr57Ti5Cu20Ni8Al10 alloy at strain rates ranging from 104 to 3  103 s1. Also, a negative value was obtained by Dalla Torre et al. (2006) for a similar composition, for example, Zr52.5Ti5Cu17.9Ni14.6Al10 (Vit105) and for a Vit105-based composite, at strain rates of 3.3  104 and 3.7  103 s1. However, Lu et al. (2003) reported that Zr41.2Ti13.8Cu12.5Ni10Be22.5 is insensitive to the strain rate below 473 K. In order to illustrate SRS of BMG alloys, Zr65Cu20Fe5Al10 rod (as one of the relatively ductile BMG alloys) glassy samples were tested at room temperature at the four loading conditions. The strain rate was suddenly changed upon deformation. The plastic deformation is always above 5% and the magnitude of stress drops increases with the deformation until failure occurs (Fig. 3.9). Figure 3.9 shows the true stress–plastic strain curves calculated from engineering stress values of test 1 and test 2 and the evolution of the stress drop magnitude with the compressive strain for plastic deformations up to 3% and 5%, respectively. After an initial region of apparent strain hardening, the magnitude of the stress drops increases slightly. A large number of stress drops of very small magnitude, close to 2 MPa, are observed especially for small plastic deformations and small loading rates. The value of the SRS parameter (m) has been calculated from the true stress–strain curves according to the relationship: 

 @ lns m¼ @ ln e_ e;T

ð1Þ

1600

40

25 S0.2 = 1600 MPa

15 10 5 0.5

1.0

1.5 2.0 Plastic strain (%)

2.5

3.0

1600

0

35 5 ⫻ 10-5 s-1

5 ⫻ 10-4 s-1

30

1500 1450

25

S0.2 = 1520 MPa

20 15 10

Stress drop (MPa)

1550 5 ⫻ 10-6 s-1

Calculated true stress (MPa)

20

1550

1400 5 1350 0.0

0.5

1.0

1.5

2.0 2.5 3.0 Plastic strain (%)

(c)

3.5

4.0

4.5

0 5.0

(d) 300

300

250

250

Frequency

Frequency

35 30

1500 0.0 (b)

5 ⫻ 10-6 s-1

5 ⫻ 10-5 s-1

Stress drops (MPa)

1650

5 ⫻ 10-4 s-1

Calculated true stress (MPa)

(a)

200 150 100

150 100 50

50 0 1533

200

1536

1539

1542

1545

1548

Calculated true stress (MPa)

1551

0 1533

1536

1539

1542

1545

1548

1551

Calculated true stress (MPa)

Figure 3.9 Calculated true stress–plastic strain curves (elastic strain is subtracted) and stress drop magnitude for (a) test (1) performed at strain rates 5  104, 5  105, and 5  106 s1 and (b) test (2) performed at strain rates 5  106, 5  105, and 5  104 s1. The diamonds represent the magnitude of stress drops and the continuous line represents the calculated true stress. Minimum and maximum stress values are marked. (c) Normal (Gaussian) and (d) log-normal distributions and fittings for the calculated mean true stress in test (2) at 5  105 s1 from 1.6% to 2.1% plastic deformation. Reprinted from Gonza´lez et al. (2011) with permission of Elsevier.

Bulk Metallic Glasses

149

where s is the flow stress and e_ is the strain rate. It is important to note that the strain rate changes have been performed in the regions where the stress– strain curves are generally linear, which may correspond to the formation of multiple shear bands and the macroscopically homogeneous deformation of the sample. “Elephant–plank” curves at larger strain correspond to a highly inhomogeneous deformation through propagation of a single major shear band. The mean and maximum flow stresses and the confidence interval (probability P ¼ 0.95) have been calculated in a 0.25% plastic deformation interval using true stress values (Fig. 3.9) on both sides of each point of the strain rate change as shown in Table 3.3. For better statistics, a larger sampling interval was used in test 2 for a plastic deformation interval from 1.6% to 2.1% at 5  105 s1. The resulting frequency distribution of stress values fitted with normal (Gaussian) and log–normal distribution functions are shown in Fig. 3.9c and d, respectively. For both distributions the calculated mean stress is similar 1543.8  0.3 MPa. Another parameter, which is an average maximum stress before each stress drop (see Fig. 3.9 and Table 3.3) satisfied normality tests and produced similar results. Not only the flow stress but also the average maximum stress before each stress drop corresponding to a serration for each test is practically the same at each loading rate in the vicinity of the strain rate change point and the differences are practically within the confidence interval, which indicates that the Zr65Cu20Fe5Al10 BMG is strain rate insensitive (m  0) in the range of strain rates from 5  106 to 5  104 s1. A compression jump test was also performed at higher strain rates (i.e., 5  104, 5  103, and 5  102 s1) to analyze the SRS in a broader strain rate range (Fig. 3.10). Between 5  104 and 5  103 s1, the average flow stress is similar, 1684.4  0.5 and 1685.2  1.6 MPa, respectively, and the difference is within the confidence interval. It indicates that the alloy is not strain rate sensitive. However, when the strain rate increases from 5  103 to 5  102 s1, the average flow stress meaningfully decreases from 1684.1  1.1 to 1675  2.0 MPa, and according to Eq. (2), the SRS is negative (m ¼ 0.0026). From the statistical analysis of the mean and maximum flow stress at each loading rate, it was calculated that the alloy does not exhibit SRS within the confidence interval from 5  106 to 5  103 s1. The SRS values derived by using the confidence interval are smaller (m  0) than the value of m reported for Zr52.5Ti5Cu17.9Ni14.6Al10 at room temperature (e.g., about 0.002) by Song et al. (2008) and for Cu50Zr50 BMG by Dalla Torre et al. (2007). However, the SRS becomes meaningfully negative when the strain rate increases from 5  103 to 5  102 s1 because the strain rate is so fast that the relaxation time is not enough to build up the stress, which explains the decrease in the flow stress.

Table 3.3 Mean and maximum stress values for test 1 and test 2 calculated considering a 0.25% plastic deformation interval at both sides of the strain rate change Test 1 5  104 to 5  105 (s1)

Mean stress (MPa) 5  104 1609.4  1.2 5  105 1611.9  0.5 Mean max. stress (MPa) 5  104 1614.6  1.4 5  105 1619.7  3.0

Test 2 5  105 to 106 (s1)

5  106 to 5  105 (s1)

5  105 to 5  104 (s1)

5  105 5  106

1611.9  0.5 1612.0  0.2

5  106 5  105

1543.2  0.2 1543.5  0.4

5  105 5  104

1543.8  0.3 1542.4  0.9

5  105 5  106

1618.1  1.2 1618.1  3.1

5  106 5  105

1548.2  3.2 1549.1  2.3

5  105 5  104

1549.7  1.3 1547.4  1.1

Mean maximum stress indicates average value of maximum stress before stress drop. Confidence intervals were calculated using 0.95% probability. Reproduced from Gonza´lez et al. (2011) with permission of Elsevier.

151

Bulk Metallic Glasses

Calculated true stress (MPa)

1750

1700

5 ⫻ 10-3 s-1

5 ⫻ 10-4 s-1

5 ⫻ 10-2 s-1

1650

1600

1550

1500 0.0

S0.2 = 1600 MPa

1.0

2.0

3.0

Plastic strain (%)

Figure 3.10 Plastic region of the calculated true stress–strain curves for the jump test 3 at 5  104, 5  103, and 5  102 s1 showing serrated flow. Reprinted from Gonza´lez et al. (2011) with permission of Elsevier.

While the minimum strain rate at which this alloy exhibits negative SRS is 5  103/5  102 s1, it should be taken into account that it depends on the loading conditions. For example, while Zr52.5Ti5Cu17.9Ni14.6Al10 exhibits negative SRS in compression (Dalla Torre et al., 2006), it is strain rate insensitive under indentation (Trichy et al., 2005). The different behavior could be mainly attributed to the stress state, which is closer to three-axial compression (Sanditov et al., 2004; Srikant et al., 2006) when tested by indentation. Moreover, there is a strong size effect on deformation behavior upon indentation in which the deformed volume is significantly smaller, while local plastic deformation is larger than in compression. Metallic glasses have been proven to deform homogeneously on the nanoscale by Guo et al. (2007a,b) and Louzguine-Luzgin et al. (2010a,b). All these factors may be responsible for the different SRS obtained in case of nanoindentation. The magnitude of stress drops in test 1 and test 2 upto a few percent of engineering strain is also similar, which suggests that the internal structural relaxation at 5  104 and 5  106 s1 is similar and fast enough to keep up with the loading rates. Deformation behavior of Zr-based BMG alloys was also tested at liquid nitrogen temperature. Strength of the sample increases with temperature and no clear serrated flow typical for bulk glassy samples tested at room temperature is observed in case of the samples tested in liquid nitrogen (Louzguine-Luzgin et al., 2011). The mechanical behavior and the kinetics of shear deformation in BMGs were also investigated at room and liquid nitrogen temperature using the acoustic emission technique. It was demonstrated that the intensive acoustic emission reflecting the activity of strongly

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localized shear bands at room temperature vanishes at the transition from serrated to non-serrated plastic flow at low temperature. The disappearance of acoustic emission signals clearly suggests that the shear band propagation velocity significantly decreases at low temperature, and sliding along the principle shear band is observed at the machine-driven rate (Vinogradov et al., 2010). Fatigue properties of BMG alloys have been also studied and fatigueendurance limits of some Zr-based alloys are comparable with those of high-strength structural alloys (Wang et al., 2004a,b).

6. Magnetic Properties 6.1. Soft-magnetic alloys Ferromagnetic alloys can exhibit hard or soft magnetism depending on their coercivity. Magnetic materials having a coercivity above about 104 A m1 are considered to be hard, while soft magnetic materials have a coercivity below 103 A m1. In general, Fe- and Co-based BMG alloys exhibit good soft magnetic properties (Amiya et al., 2007; Inoue and Shen, 2004; Makino et al., 1995), while Nd-based alloys show hard magnetic properties. For a long time, soft magnetic alloys were limited to marginal glass formers. However, ever since a ferromagnetic Fe–(Al,Ga)-metalloid bulk glassy alloy was produced by Inoue et al. (1995), Fe-based bulk metallic glasses have attracted significant interest in the scientific community owing to their good soft magnetic properties and good GFA. The bulk glassy alloys having Curie temperature Tc of 580–610 K were developed in the Fe–(Al,Ga)–(P,C,B) and Fe–(Al–Ga)–(P,C,B,Si) systems (Bitoh et al., 2004; Mizushima et al., 1999). These bulk glassy alloys also exhibited good soft magnetic properties. For example, a ring-shaped sample of a Fe70Al5Ga2P9.65C5.75B4.6Si3 glassy alloy with a thickness of 1 mm, an outer diameter of 10 mm, and an inner diameter of 6 mm formed by the copper mold casting method exhibits a high saturation magnetization of 1.2 T, low coercivity Hc of 2.2 A m1, and a rather low saturated magnetostriction (ls) of 21  106. The maximum permeability (mmax) is 110,000 (Fig. 3.11). The Hc and mmax values are superior to those (3.7 A m1 and 27,000) of the melt-spun samples. The remarkable improvement in the soft magnetic properties has been demonstrated to result from the significant difference in the magnetic domain structure. The domain walls are arranged along the circumference direction for the cast-ring alloy and radial direction for the ribbon ring sheet. The difference in the domain wall structure was reported to originate from the difference in residual stress during the preparations of the ring sample and the melt-spun ribbon. The glassy Fe–(Al,Ga)–(P,C,B,Si) alloys

153

Bulk Metallic Glasses

Fe70Al5Ga2P9.65C5.75B4.6Si3

0.6

Magnetization (T)

0.4

0.2 Melt-spun ribbon Hc = 3.7 A m-1

0

(mmax = 27,000)

-0.2

Cast bulky sample Hc = 2.2 A m-1 (mmax = 110,000)

-0.4

-0.6

-16

-8

0

8

16

Applied magnetic field (A m-1)

Figure 3.11 Hysteresis I–H loops of the as-cast ring-shaped Fe70Al5Ga2P9.65C5.75B4.6Si3 glassy alloy with a thickness of 1 mm, an inner diameter of 6 mm, and an outer diameter of 10 mm. The data for the ring-shaped sheet with a thickness of 0.02 mm made from the melt-spun glassy ribbon are also shown for comparison. Reprinted from Makino et al. (2000) with permission of JIM.

exhibit low coercivity, though their magnetostriction is relatively large. The origin of the low coercivity of the Fe–(Al,Ga)–(P,C,B,Si) glassy alloys is explained based on the low density of quasidislocation dipole-type defects. In the case of the (Fe,Co,Ni)70Zr10B20 glassy alloys produced by Inoue et al. (1997b), Hc decreases gradually from 6 to 3 A m1 with increasing Fe content. The Is increases from 0.3 to 0.9 T with increasing Fe content, while ls equals zero in the Co-rich composition range and increases monotonously to 15  106 with increasing Fe content. me reaching the maximum of about 20,000 is obtained in the Fe- and Co-rich composition ranges. These glassy alloys exhibit good soft magnetic properties including Is up to 0.9 T, Hc of 3–6 A m1, ls of 12–15  106, and me of 20,000 in the Fe-rich range (Table 3.4), and Is of 0.5 T, Hc of 6 A m1, nearly zero ls, and me of 20,000 in the Co-rich range. Compared to other BMGs, (Fe–Co)-based alloys are particularly attractive for engineering applications because of their combination of ultrahigh strength (the highest reported in the literature), superior wear resistance, good GFA, and rather low cost (Inoue et al., 2003). Co43Fe20Ta5.5B31.5 glassy alloy exhibited an extremely low value of the coercive force as low as

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Table 3.4 Composition and magnetic properties (Js, Hc) of typical ferromagnetic BMGs, and their critical diameter (Dcr) where specified (Inoue et al., 2000; Makino et al., 2008) Composition

Js (T)

Hc Dcr (A m1) (mm)

References

Fe76Si9B10P5 (Fe0.75Si0.1B0.15)96Nb4 Fe74Al4Ga2P12B4Si4 Fe73Al5Ga2P11C5B4 Fe77Ga3P9.5C4B4Si2.5 Fe72Al5Ga2P10C6B4Si1 Fe76Mo2Ga2P10C4B4Si2 Fe30Co30Ni15Si8B17 Co43Fe20Ta5.5B31.5 Fe74Nb6Y3B17 Fe56Co7Ni7Zr8M2B20, see below M ¼ Ti M ¼ Hf M¼V M ¼ Nb M ¼ Ta M ¼ Cr M ¼ Mo M¼W M ¼ Ti Fe–(Co,Ni)–(Zr,Nb, Ta)–(Mo,W)–B

1.51 1.47 1.14 1.29 1.36 1.14 1.32 0.92 0.49 0.81

0.8 2.9 6.4 6.3 4.25 2.8 2.9 3.4 0.25 15

2.5 1.5 – 1 2.5 2 2 1.2 3 2

Makino et al. (2008) Makino et al. (2008) Makino et al. (2008) Makino et al. (2008) Shen and Inoue (2002) Makino et al. (2008) Makino et al. (2008) Makino et al. (2008) Inoue et al. (2003) Makino et al. (2008)

0.82 0.8 0.83 0.75 0.74 0.75 0.73 0.7 0.82 0.74–0.96

1.9 2.2 4.4 1.1 2.7 1.9 4.9 5.7 1.9 1.1–3.2

– – – – – – – – – 2–6

Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Makino et al. (2000) Inoue et al. (1997a,b)

–, Means no data.

0.25 A m1 together with the maximum permeability of 550,000. However, the saturation magnetization was only 0.49 T. Fe–(Co,Ni)–(Zr,Nb,Ta)–(Mo,W)–B system glassy alloys exhibit a large saturation magnetization of 0.74–0.96 T, low coercivity of 1.1–3.2 A m1, high permeability exceeding 1.2  104 at 1 kHz, and low magnetostriction of about 12  10–6 (Inoue et al., 1997a). Recently developed Fe-metalloid Fe80P11C9 BMG alloy showed good soft-magnetic properties and high strength of 3.2 GPa (Wang et al., 2011). The partially devitrified alloy also exhibited good soft-magnetic properties including magnetic polarization of 1.49 T and coercivity of 4 A m1. A similar Fe76Si9B10P5 alloy exhibited a saturation magnetization Js of 1.51 T, Hc of 0.8 A m1, and a critical diameter of Dcr 2.5 mm (Makino et al., 2008).

Bulk Metallic Glasses

155

Fe82(Zr,Hf,Nb)7B10Cu1 alloys exhibit good soft magnetic properties especially in the high-frequency range (Moon et al., 1998). The Fe–M–B (M ¼ Zr, Hf, or Nb) alloys also show low core losses (Suzuki et al., 1993). The soft magnetic properties of (Fe,Co)–RE–B glassy alloys with large thicknesses were studied by Zhang and Inoue (2001). Y addition was confirmed to improve GFA of the Fe72B24Nb4 alloy. The maximum critical diameter for glass formation in Fe–B–Nb–Y system alloys was 7 mm. Saturation magnetization and coercivity of the as-cast (Fe0.72B0.24Nb0.04)95.5Y4.5 glassy ring were found to be 0.8 T and 0.8 A m1, respectively, in the relaxed state after annealing at 821 K (Tg  50 K) for 300 s (Lee et al., 2008). 0.8 A m1 is an extremely low value of coercivity. A large increase in coercivity from 5 A m1 upto 4.4 kA m1 was reported for Fe75Si11B10Nb3Sn1 glassy samples after the first stage of crystallization. This outstanding change was ascribed to the generation of metastable nanocrystallites that disappear at higher temperatures (Cremaschi et al., 2000).

6.2. Hard-magnetic alloys Hard magnetic alloys (also called magnetically hard alloys) have sufficiently high coercive force as a resistance to demagnetizing fields with coercivity exceeding 10 kA m1. These alloys with high magnetic induction which is retained because of a strong resistance to demagnetization can be used as permanent magnet materials, for example, as a result of high anisotropy. The development of permanent magnets led to an increase in the coercive force, saturation magnetization, and the ability to store energy, characterized by the energy product BHmax. Ferromagnetic Nd90xFexAl10 bulk amorphous alloys with high coercive force at room temperature were obtained by a copper mold casting method. For some alloys, the maximum diameter of the cylindrical amorphous samples reaches 15 mm (Inoue and Zhang, 1997). Neither glass transition nor supercooled liquid region was observed in these alloys before crystallization, which makes them different from typical bulk glassy alloys exhibiting a wide supercooled liquid region before crystallization. The bulk amorphous Nd70Fe20Al10 alloy shows ferromagnetism with the Curie temperature (Tc) of about 600 K, which is much higher than the highest Tc (about 480 K) for the Nd–Fe binary amorphous alloy ribbons. The remanence (Br) and intrinsic coercive force (iHc) for the bulk Nd60Fe30Al10 alloy are 0.122 T and 277 kA m1, respectively, in the as-cast state and 0.128 T and 277 kA m1, respectively, in the annealed state for 600 s at 600 K. The Br and iHc decrease to 0.045 T and 265 kA m1, respectively, for the crystallized sample. The hard magnetic properties of the bulk amorphous alloys are presumably due to the development of homogeneous ferromagnetic clusters with a large random magnetic anisotropy (Inoue et al., 1996b).

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Hard magnetic alloys can be produced by crystallization of the glassy phase. For example, permanent magnetic materials consisting mainly of Fe3B with Nd2Fe14B phase were obtained by annealing Nd4.5Fe77B18.5 rapidly solidified alloys (Coehoorn et al., 1988). High remanence (Br) of 0.8 T is obtained due to the remanence enhancement effect of exchangecoupled magnetic grains. The remanent polarization (Jr) of this material is 1.2 T and maximum energy product (BH)max ¼ 97 kJ m3 while its coercivity Hc ¼ 240 kJ m1. High coercivity values exceeding 300 kA m1 were obtained in amorphous Nd5Fe72Cr5B18 crystallized into the Fe3B/ Nd2Fe14B state (Hirosawa and Kanekiyo, 1996).

7. Corrosion Resistance Many of the metallic glassy alloys have a high corrosion resistance (Pang et al., 2004; Peter et al., 2002). For example, the anodic polarization curves of the (Ti0.45Zr0.1Pd0.1Cu0.31Sn0.4)98Nb2 and (Ti0.45Zr0.1Pd0.1Cu0.31Sn0.4)98Ta2 glassy rod samples were measured in 1 mass% lactic acid, PBS (phosphate-buffered saline salts solution containing 8 g l1 NaCl, 0.2 g l1 KCl, 1.15 g l1 Na2HPO4, 0.2 g l1 KH2PO4), and HBSS (Hank’s balance salts solution containing 8 g l1 NaCl, 0.4 g l1 KCl, 0.09 g l1 Na2HPO47H2O, 0.06 g l1 KH2PO4, 0.35 g l1 NaHCO3, 1.0 g l1 glucose) aqueous solutions at 310 K open to air by Oak et al. (2009). The results are shown in Fig. 3.12 in comparison with those for pure Ti, Ti–6Al–4V alloy, and other Ti-based BMG alloys. All anodic polarization curves exhibit spontaneous passivation behavior upon an increase in the anodic polarization. (Ti0.45Zr0.1Pd0.1Cu0.31Sn0.4)100x(Nb/Ta)x glassy rod samples exhibit lower passive current densities than their competitors before pitting corrosion by increasing potential indicating that these glassy rod samples have spontaneously passivated regions. (Ti0.45Zr0.1Pd0.1Cu0.31Sn0.4)98Nb2 glassy alloy has exceptional corrosion resistance accompanied by lower current density. It exhibits excellent passivation with the lower passive current densities in PBS and in HBSS, which are lower than those of the competitors (pure Ti metal, Ti–6Al–4V alloy, and the basic BMG alloy). However, it is not passivated above þ540 mV (Epit) compared to Ti and Ti–Al–V alloy in PBS solution. Nevertheless, human cells and organs use bioelectricity in the range of 70–90 mV. Also, the current density of the Ti44.1Zr9.8Pd9.8Cu30.38Sn3.92Nb2 bulk glassy alloys is below 101 A m2 in lactic acid, 102 A m2 in PBS, and 101 A m2 in HBSS, and the passivated region exceeds the range of the body potential (Eisenbarth et al., 2004; Steinemann et al., 1980). Fe-based BMGs also have high corrosion resistance. The corrosion rates of the Fe50xCr16Mo16C18Bx glassy alloys were in the range of 103–102 mm year1 in 1, 6, and 12 N HCl solutions. These bulk glassy

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1 mass% lactic acid, 310 K

102

Current density, I (A m-2)

101 100

(Ti0.45Zr0.1Pd0.1Cu0.31Sn0.04)98Nb2

10-1 10-2 Ti6Al4V

10-3 10-4 10-5

Ti45Zr10Pd10Cu31Sn4 Ti

0.5 1.0 1.5 Potential, E (V) vs. Ag/AgCl

0.0

2.0

PBS(-), 310 K

103 Current density, I (A m-2)

102

(Sn4)99Nb1 (Sn4)98Nb2 (Sn4)98Ta2 Pd20Cu20Sn5

101

Sn3

100

Sn5

10-1

Sn4(the basic alloy) Ti6Al4V4 Ti

10-2 10-3 10-4 10-5

-0.2

0.0 0.2 0.4 0.6 Potential, E (V) vs. Ag/AgCl

0.8

1.0

Figure 3.12 Anodic polarization curves in 1 mass% lactic acid, PBS(), and HBSS at 310 K, as indicated. All samples were rods of 3 mm in diameter (the exposed area was 1 cm2). Reprinted from Oak et al. (2009) with permission of Elsevier. Continued

alloys are spontaneously passivated in 1 and 6 N HCl solutions and do not exhibit pitting corrosion in 12 N HCl solution up to the potential of 1.0 V (Ag/AgCl). The high corrosion resistance results from the formation of chromium-rich passive films during immersion in HCl solutions by Pang et al. (2002).

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HBSS, 310 K

103

Current density, I (A m-2)

102 101

(Ti0.45Zr0.1Pd0.1Cu0.31Sn0.04)98Nb2

100

SUS316L

10-1 10-2

Ti6Al4V

10-3 10-4

Ti45Zr10Pd10Cu31Sn4

Ti

10-5 -0.5

0.0 0.5 Potential, E (V) vs. Ag/AgCl

1.0

Figure 3.12—cont’d

8. Applications Owing to their high GFA, good casting ability, good formability in the supercooled liquid region as well as good mechanical and chemical properties, BMGs have important applications (Ashby and Greer, 2006; Inoue et al., 1989, 1995), which are predicted to expand in the near future (Nishiyama et al., 2007a,b). BMGs are used in sporting equipment, watch parts, electromagnetic wave shields, parts of optical devices, power inductors, magnetic field identification systems, microgeared motor parts, pressure sensors, Coriolis flow meters, coating materials, shot peening balls, and in medical instruments. In hydrogen energy technologies, materials selection is critical as hydrogen tends to decrease the mechanical properties of the metallic alloys, while glassy alloys exhibit high hydrogen solubility and significant embrittlement resistance. Hydrogen permeation characteristics of melt-spun Zr–Hf–Ni (Hara et al., 2003), Ni–Nb–Zr (Yamaura et al., 2003), and other amorphous alloy membranes (Jayalakshmi et al., 2012) were studied. Such materials are promising for future applications in separators for fuel cells. Ti–Zr–Pd–Cu–Sn–Nb BMGs can be applicable as biomaterials. Porous BMG samples produced by hydrogenation treatment (Wada et al., 2005), infiltration (Brothers and Dunand, 2004), mixing the melt with hydrated B2O3 (Schroers et al., 2003) or by glassy powder sintering (Xie et al., 2006a) can be used as bioimplants or damping materials owing to their reduced Young’s modulus compared to monolithic BMGs.

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The Fe–Cr–P–C–B–Si system glassy alloy was commercialized under the name LiqualloyTM (Koshiba et al., 2008; Mizushima et al., 2008). With high GFA and good corrosion resistance, the Fe–Cr–P–C–B–Si alloy exhibits a saturation magnetization of 1.25 T, a coercive force of about 2 A m1, effective permeability of 60 at 1 kHz, and a high electrical resistivity of 1.5 mO m. “LiqualloyTM” magnetic cores have been produced on a mass production scale by cold consolidation of a mixture of water-atomized glassy alloy powders with binder and subsequent annealing at elevated temperatures as shown in Fig. 3.13. The LiqualloyTM cores exhibit nearly constant relative permeability in a high-frequency range up to several megahertz and much lower core losses than those for the Ni–Fe–Mo Permalloy core and the Fe–Si–Al Sendust core. These excellent core loss characteristics are attributed partly to the reduction in eddy-current loss resulting from much higher electrical resistivity. Magnetic properties of soft magnetic cores made of LiqualloyTM surpass these properties of conventional soft magnetic cores. At present, the LiqualloyTM cores are used in laptop computers owing to their high efficiency and small heat generation; see, for example, Inoue and Takeuchi (2010) . The spherical powder of LiqualloyTM produced by water atomization can be deformed into flaky shapes with thicknesses of 2–3 mm and large

Micro structure

Elastomer

Flaky powder

SEM image

Figure 3.13 Cross-section of a LiqualloyTM sheet. Reproduced with permission from ALPS Electric Co., Ltd. all rights reserved.

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aspect ratios of 10–30 by a cold deformation method. The LiqualloyTM sheet consisting of the flaky powder embedded in resin possesses a high conversion ratio from electromagnetic wave noise to heat, leading to a highly efficient suppression of noise. The high efficiency comes from a higher imaginary part of permeability in the high-frequency range >500 MHz in comparison with Fe–Si–Al Sendust sheet. Therefore, the LiqualloyTM sheet has been commercialized as an electromagnetic wave noise suppression sheet in various electromagnetic instruments such as digital still cameras (Fig. 3.14). The use of Liqualloy sheets between electromagnetic devices and loop antenna results in a significant increase in the transmission distance. As a result, one can significantly increase the sensitivity of the antenna at a high carrier frequency of 13.56 MHz. The good antenna sensitivity results from the much higher quality factor defined by the ratio of the real part of permeability to the imaginary part. Another type of soft-magnetic powder core with a higher saturation magnetization of 1.3 T was developed in Fe–Nb–B–Si and Fe–Nb–Cr–P– B–Si systems by Matsumoto et al. (2010). The cores were also produced by a similar procedure, that is, production of spherical glassy alloy powders by water atomization, followed by cold consolidation of the mixture of glassy

Figure 3.14 Liqualloy sheets and rolls. Reproduced with permission from ALPS Electric Co., Ltd. all rights reserved.

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alloy powder and epoxy resin. These new magnetic powder cores named SENNTIX exhibit the lowest core losses among all types of magnetic powder cores developed to date. In addition, the higher saturation magnetization enabled the use of cores in a higher current range. SENNTIX-II powder cores have low core losses and reduce thermal emissions from personal computers. The SENNTIX-II consolidated cores are produced on a large scale of several million pieces per month. Fe–Ni–Cr–Mo–B–Si glassy alloy powders produced by water atomization have also been commercialized with the commercial name AMObeads. The powder diameter ranges from 0.05 to 0.6 mm owing to the high GFA for the Fe-based alloy. The application fields are extended to shot peening balls and precise polishing medium. The AMO-beads have much longer endurance times compared to those made of steel. Good mechanical properties, such as high Vickers hardness of 900, high fracture strength of 3000 MPa, and large elastic strain of 0.02 together with high corrosion resistance and a smooth outer surface make AMO-beads a good tool for the shot peening treatment. They generate a higher level of residual compressive stress on the surface of steel vehicle gears with high Vickers hardness of 750 achieved by carburization treatment. As a result, the alloy steel gears can increase fatigue strength by 50–80% in comparison with the steel gear subjected to shot peening using steel balls. This improvement causes a significant reduction in the weight of steel vehicle gear by 45%. Dense Fe-based glassy alloy-coated layers in the Fe–Cr–Mo–C–B system have been produced on various metallic alloy substrates using the highvelocity powder-spray coating layer technique (Kobayashi et al., 2008a). The Fe–Cr–Mo–C–B coating layer exhibits better corrosion resistance than that of SUS304 and higher Vickers hardness than that of hard chromium plating plate (Kobayashi et al., 2008b). BMGs have a smooth surface and soften on heating about Tg, which allows creation of micron and nanoscale patterns on their surfaces by molding (Saotome et al., 2001b, 2002, 2004). An example is shown in Fig. 3.15. After cooling, such a material becomes hard again and can be used as a stamp for other materials. A growing number of research activities are now focused on metallic glassy nanoobjects, for example, nanowires (Nakayama et al., 2010; Carmo et al., 2011). Metallic nanowires may have analytical applications ranging from interconnects to sensors (Walter et al., 2002). Micrometer-long multisegment nanowires have been used as barcodes for biological multiplexing, which can be easily visualized by an optical microscope (Keating and Natan, 2003). The micrometer-long nickel nanowires can be internalized by cells allowing the manipulation of living cells through a magnetic field (Tanase et al., 2005). Metallic glasses are suitable materials for scanner micromirrors. In order to achieve the large scanning angle without mechanical failure during actuation, the micromirror structure was fabricated using Fe-based metallic

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A

Dd=30 nm

C

Lp=50 nm

100 nm B

100 nm

Dd=23 nm

D

Lp=40 nm

100 nm

100 nm

Figure 3.15 SEM microphotographs of dies having nanodot arrays with pitches of (a) 50 nm and (b) 40 nm. (c) and (d) are the nanoimprinted surfaces of the Pt-based metallic glass using the dies of (a) and (b), respectively. Reprinted from Fukuda et al. (2011) with permission of JIM.

glass (Lee et al., 2011). High values of mechanical strength and elastic strain limit are desired for the torsion bar for providing high performance of the mirror, including large tilting angle and good stability. Millimeter-size mirror device was successfully fabricated from Fe-based MG ribbon. An extremely large optical tilting angle exceeding 100 was obtained on activation of the mirror by magnetic force when electrical current of 100 mA was applied. The large tilting angle of the mirror is due to the torsion bar, which was fabricated with Fe-based MG material that has alarge elastic strain limit and high fracture toughness. Recently developed Au-based metallic nanoglasses with a large surface area produced from a BMG forming alloy (Fig. 3.16) by Chen et al. (2011) have opened up a new application area with such a material as catalyst. This material showed a good catalytic activity for the following reaction: PhMe2 Si  H þ H2 O ½1

Aubased NGMG

!

acetone; rt; 24 h

PhMe2 Si  OH þ H2 ½2

ð2Þ

of dimethylphenylsilane (PhMe2Si–H) with water in the presence of Aubased metallic nanoglass. The reaction proceeded at room temperature for 24 h and the desired dimethylphenylsilanol was obtained. As similar granular structure was obtained in the case of Pd78Si22 nanoglass other catalysts can be prepared soon.

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200 nm

Figure 3.16 SEM image of the surface morphology of a nanoglassy Au46Ag6Pd2Cu27Si14Al5 sample. The insert demonstrates a TEM nanobeam diffraction pattern.

9. Future Prospects BMG alloys compose one of the most dynamically developing scientific fields connected to metallic materials. According to Web of Science (http:// apps.webofknowledge.com) (one of the largest sources, but still not covering all scientific publications), since the late 1980s and the beginning of the 1990s, the number of research papers related to BMGs/bulk glassy alloys published per year have increased nonlinearly, stabilizing after 2006 with several hundreds of papers per year. The Scopus database (http://www.scopus. com) also shows a large number of research papers on the subject, which indicates the intensive research activity related to such materials. Almost every year, vital papers containing scientific reports related to BMG alloys are published reporting unusual general physical, mechanical, and chemical properties. Gradual success in understanding their structure, formation mechanism, mechanical behavior, and properties leads to the creation of BMG alloys with higher performance compared to early works on this subject. For example, as a result of substantial research efforts, GFA and plasticity of BMGs have been drastically improved, increasing interest in such materials among the industrial circles. Magnetic BMG alloys developed since the mid-1990s are at present being used in various applications and this field is still expanding. BMG alloys have also become important materials for micro and nanomolding. Success achieved in the field of BMGs provides new insights in materials science offering the opportunity for further studies of these novel materials and promotes their applications. Structural materials should have high strength, fracture toughness, impact fracture toughness, and stiffness. At present, polymeric composites

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have started to compete in a big way with metallic crystalline alloys with regard to specific strength (the strength/mass ratio). Here, metallic glasses have a great advantage because of their high strength and specific strength. Recent achievements related to ductile Zr–Ti–Be-based crystal-glassy composites (Hofmann et al., 2008a,b) and the invention of Pd79Ag3.5P6Si9.5Ge2 BMG alloys with high fracture toughness (Demetriou et al., 2011) have inspired another set of optimistic viewpoints on metallic glasses. Moreover, although Zr–Ti–Nb–Cu–Be bZr-BMG composites show necking owing to stress-softening of the glassy phase, it was reported that Cu–Zr–Al–Co (Hofmann, 2010; Wu et al., 2010) and Cu–Zr–Al composites (Pauly et al., 2010) show strain-hardening owing to transformation-induced plasticity upon martensitic transformation. Although most of these ductile and tough materials are based on the expensive elements such as Zr, Pt, and Pd, the success achieved allows us to suggest that BMG alloys and the composites containing a glassy phase are promising structural and functional materials of the present century. New Fe-based bulk glassy alloys exhibiting a high Js of 1.51 T reported recently require further investigation as new promising magnetic materials owing to their high Fe, more than three-fourth, content and absence of other metallic elements, except for Fe. After the successful attempts at the simultaneous achievement of high Fe content and high GFA in Fe-based BMGs, further research activities are expected in this field. Fe, C, Si, and P are the constituent elements in pig-iron produced in a blast furnace. As lowpriced Fe–Si and Fe–P ferroalloys are in mass production, there is no restriction in the availability of such materials, which offers the advantage of lower material cost for industry. It was shown that a Fe-based metallic glass with good soft-magnetic properties is a promising material for mirror actuation. Metallic glasses provide a high value of elastic strain limit and fracture toughness resulting from its amorphous structure without structural defects like dislocations and grain boundaries. These attractive functional properties are claimed for microelectromechanical systems. In addition, large-scale consolidated powder cores are made from metallic glassy powders. This technique allows overcoming limitations associated with the small GFA of some metallic glassy alloys. Recent work has shown the possibility to form magnetic glassy-ceramic composites with reduced core losses (Xie et al., 2012). The addition of the SiC particulates was effective in improving the high-frequency magnetic properties. This approach is very promising for creation net shape products the shape and dimension of which require minimum or no further adaption. Besides massive glassy samples, metallic glassy nanowires and nanoglasses are attracting increasing attention at present. Here, there is scope for further employment of magnetic materials. Arrays of magnetic nanowires are attracting considerable interest from the viewpoint of perpendicular magnetic

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recording. The nanowires are potentially capable of producing recording densities in excess of several tens of Gbits per square inch. It has also been reported (Nielsch et al., 2001) that the magnetic nanowires of Fe, Co, and Ni show a much enhanced magnetic coercivity than those of their bulk counterparts. Such an approach can be applied to metallic glassy nanowires as well.

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