What we have learned from studies on chemical properties of amorphous alloys?

Applied Surface Science 257 (2011) 8141–8150

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Review Article

What we have learned from studies on chemical properties of amorphous alloys? Koji Hashimoto ∗ Tohoku Institute of Technology, Sendai, 982-8577, Japan

a r t i c l e

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Article history: Available online 31 December 2010 Keywords: Homogeneous alloy 12 M HCl Spontaneous passivation Amorphous bulk alloy Catalyst CO2 methanation

a b s t r a c t Amorphous alloys have many attractive characteristics including extremely high corrosion resistance if the sufficient amounts of corrosion-resistant elements are added. The superiority of amorphous alloys is based on the homogeneous single phase nature without any chemical and physical heterogeneities. Although there are processing limitations to avoid the formation of heterogeneous crystalline structure in addition to no welding technology without crystallization, the application of corrosion-resistant amorphous alloys is expected particularly to the very aggressive environments, where any conventional crystalline metallic materials cannot be used. Some amorphous bulk alloys showed zero corrosion mass loss due to spontaneous passivation even in 12 M HCl. Production of amorphous bulk alloys became possible for selected compositions. The homogeneous single phase nature is also effective to form useful catalysts with unique composition and structure. An example of catalysts is for carbon dioxide methanation useful for supply of renewable energy in the form of methane. © 2011 Elsevier B.V. All rights reserved.

Contents 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8141 High corrosion resistance due to homogeneous nature of amorphous alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8142 Enhanced passive film formation due to the metastable nature of amorphous alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8142 Difficulty in producing and processing amorphous alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8143 Corrosion-resistant amorphous alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8143 Production of amorphous bulk alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8145 Amorphous bulk alloys resistant to 12 M HCl . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8145 Hot corrosion resistance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8148 Catalysts . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8148 Reactivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8149 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8149 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8149

1. Introduction Masumoto and Maddin in succeeding preparation of amorphous Pd–20 at%Si alloy by rapid quenching from the liquid state first discovered extraordinarily high mechanical strength with toughness in 1970 [1,2]. (Alloy formulae for amorphous alloys in this article are expressed in at%.) After that, the preparation of amorphous, high strength Fe-metalloid alloys became feasible. However, amorphous Fe-metalloid alloys were easily corroded forming reddish iron rust within one day after preparation in a laboratory environment. Our

∗ Corresponding author. Tel.: +81 904049 1121; fax: +81 22 372 5224. E-mail address: [email protected] 0169-4332/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2010.12.142

attempt to enhance the corrosion resistance of the amorphous Fe–13P–7C alloy by adding chromium led to astonishing results; there was no mass loss of an amorphous Fe–8Cr–13P–7C alloy immersed for 168 h in 10% FeCl3 at 60 ◦ C, whereas type 316 stainless steel disappeared leaving the oxide film [3–5]. As shown in Fig. 1, polarization curves of the amorphous Fe–Cr–13P–7C alloys measured in 2 M HCl showed no active region because of spontaneous passivation and no pitting corrosion up to the potential region of oxygen evolution. If we wish for ferritic stainless steels to avoid pitting corrosion up to the potential region of oxygen evolution in 1 M HCl we need to add 28 wt% chromium, although the Fe–28wt%Cr alloy shows a wide active region [6]. Along with the findings of extraordinarily high mechanical strength with toughness in 1970 [1,2] and extremely high corrosion resistance in 1974 [3,4], the finding of soft magnetic property

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Fig. 1. Potentiodynamic anodic polarization curves of amorphous Fe–Cr–13P–7C alloys in 2 M HCl.

for amorphous Fe–P–C alloys in 1974 [7] led to the worldwide expansion of the study of amorphous alloys. This was the dawn of materials science of novel metallic materials. 2. High corrosion resistance due to homogeneous nature of amorphous alloys Ordinary metals have been found in the form of oxides and/or salts in nature. Thus, in nature where oxygen and water exist, the oxidized state is chemically more stable than the metallic state, and the metallic materials produced by reduction are readily oxidized. However, the metallic materials are generally separated from oxygen and water in nature due to coverage by an air-formed oxide film, and further oxidation is significantly slower because oxidation is controlled by outward diffusion of cations and inward diffusion of oxygen through the film. This is called spontaneous passivity. If the air-formed film is broken by the attack of given environments, corrosive dissolution and/or rusting of materials occur. Thus, metallic materials other than precious metals should be and can be practically used only in the spontaneously passive state. However, the uniform passive film with uniform protectiveness cannot be formed on heterogeneous metallic materials. In crystalline materials, heterogeneities in the form of second phases, precipitates, segregates, grain boundaries, dislocations and stacking faults exist. The corrosion resistance expected from homogeneous solid solution alloys, where the solute atoms are homogeneously distributed, has never been realized in heterogeneous crystalline metallic materials. Corrosive degradation of crystalline materials generally occurs based on the heterogeneity. In contrast, amorphous alloys are composed of a single phase of homogenous solid solution without any physical and chemical heterogeneities. Thus, amorphous alloys with sufficient concentrations of corrosion-resistant elements show superior corrosion resistance that has never been found in any crystalline metallic alloys. Almost all superior chemical properties as well as physical properties of amorphous alloys are based on its homogenous single phase nature. The material of such a composition showing the synergistic effects of all constituents cannot be found even in the form of a single crystal. Consequently, the homogenous single phase nature of amorphous alloys guarantees their high corrosion resistance due to the formation of a uniform passive film without defects and their high resistance to depassivation. No corrosion of the amorphous Fe–8Cr–13P–7C alloy by immersion for 168 h in 10% FeCl3 at 60 ◦ C is a typical example of the high corrosion resistance of homogeneous single phase alloy [3,4]. On the other hand, the fact that the oxide film remained after dissolu-

Fig. 2. Mass fraction of chromium in films formed on Fe–Cr alloys polarized for 1 h in 0.5 M H2 SO4 , Fe–30Cr and Fe–30Cr–2Mo alloys polarized for 1 h in 1 M HCl and amorphous Fe–10Cr–13P–7C alloys exposed to air and immersed for 168 h in 1 M HCl.

tion of type 316 stainless steel by immersion for 168 h in 10% FeCl3 at 60 ◦ C suggests that type 316 stainless steel will not be corroded in this solution, if its oxide film were able to cover uniformly the steel without having any heterogeneities. 3. Enhanced passive film formation due to the metastable nature of amorphous alloys Passivation of alloys containing corrosion-resistant elements occurs as a result of formation of a passive film in which these corrosion-resistant elements are highly concentrated. When the chromium content of crystalline binary Fe–Cr alloys is increased above 13%, these alloys are classified as stainless steels with wellknown high corrosion resistance. As can be seen in Fig. 2, stainless steels are characterized by the significant enrichment of chromium in the passive film, whereas the cationic composition of air-formed films on binary Fe–Cr alloys is almost the same as the alloy composition [8,9]. In other words, anodic polarization in the passive region of the Fe–Cr alloys the higher chemical stability of the chromiumenriched film. In this connection, the chromium content in the air-formed film on an amorphous Fe–10Cr–13P–7C alloy is almost the same as that in the passive film on the crystalline Fe–20wt%Cr alloy [10]. Thus, the amorphous Fe–10Cr–13P–7C alloy is spontaneously passive without suffering active dissolution and pitting corrosion up to the potential region of oxygen evolution even in 2 M HCl as shown in Fig. 1. Furthermore, immersion in 1 M HCl for 168 h gives rise to the enrichment of chromium of up to 97% of the cations as shown in Fig. 2. In general, the amorphous alloys are characterized by the fact that corrosion-resistant solute elements with a high affinity to oxygen and a high chemical stability in the oxidized state tend to be concentrated in the film during exposure to air and/or corrosive solutions. This has been interpreted in terms of the higher chemical reactivity due to the metastable nature of amorphous alloys. As mentioned above, the high activity led to rapid rusting of the amorphous Fe-metalloid alloys without chromium within one day after preparation. In contrast, if a solute element is able to form a stable oxidized solid, the high activity of the alloy enhances the enrichment of the oxidized element in the film, and

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if other solute or matrix elements are unstable in their oxide or oxyhydroxide form the high activity of the alloy accelerates the dissolution for those unstable elements into aqueous environments. Therefore, those amorphous alloys containing sufficient concentrations of corrosion-resistant solute elements are characterized by a rapid formation of uniform corrosion-resistant passive film where those corrosion-resistant elements are highly concentrated. Furthermore, whenever that film is mechanically broken, repassivation in the breached area rapidly occurs unless that alloy itself is massively fractured by a shear stress or hydrogen embrittlement failure mechanism.

4. Difficulty in producing and processing amorphous alloys In spite of superior characteristics, the preparation of amorphous alloys is restricted. For the basic studies of amorphous alloys, specimen preparation is commonly carried out by rapid quenching from the liquid state to form solid amorphous alloy ribbons, fibers or powders. Since corrosive degradation is the surface phenomenon, the protective surface coating of conventional metallic materials with corrosion-resistant amorphous alloys has been attempted. Laser and electron beam processing was one of expected methods for the preparation of amorphous surface alloys [11,12]. It has, however, been found that laser and electron beam processing of amorphous surface alloys was the most difficult technology to form amorphous alloys. A linear molten alloy formed by a linear single scan of a beam on the surface was easily quenched to form the amorphous solid by thermal conduction to the substrate metal to extract the heat from the solidifying linear surface alloy. However, this procedure should be repeated along the previously formed linear amorphous alloy to cover the entire surface by the amorphous surface alloy. Thus, the previously formed amorphous alloy is inevitably heated by the linear scan of a beam on the neighbor of the alloy surface, and this procedure often induces the crystallization of the previously formed amorphous alloy. Consequently, the formation of a uniform amorphous surface alloy layer requires alloy selection from a very limited number of alloy systems that have a high glass-forming ability. In addition, the rapidly quenched surface alloy layer that is bonded to the underlying solid substrate is subject to very high tensile residual stresses. These stresses are substantially high to bend the specimen and often induce cracking of the amorphous surface alloy thus formed. Therefore, laser and electron beam processing for the formation of amorphous alloy coating seriously limits its application. Sputter deposition was widely used method for the preparation of amorphous alloys and once was thought to be suitable for the production of amorphous surface alloys. However, from a corrosion point of view, the main problem of all kinds of surface coating is always the presence of physical defects. Sputter deposition is not an exception. Even if highly corrosion-resistant materials were coated, perfect coating without defects was impossible. Thus, the corrosion rate has generally been determined by the density of physical defects in the coated materials regardless of high corrosion resistance of coated materials themselves. Consequently, the idea of the corrosion-resistant coating of amorphous alloys was abandoned, and further effort was directed at producing corrosion-resistant amorphous alloys in a bulk form. However, even if amorphous bulk alloys are prepared, welding of amorphous alloys is not possible because heating during welding induces the metastable amorphous phase to transform into multiple crystalline phases with a consequent loss of almost all kinds of superiority of the uniform single phase. Because of the sophisticated nature for the production and processing of amorphous bulk alloys and because welding cannot be done, most attempts to

Fig. 3. Corrosion rates of sputter-deposited Cr–Ti and Cr–Zr alloys in 6 M HCl and Cr–Nb and Cr–Ta alloys in 12 M HCl.

directly replace crystalline metallic alloys with comparable amorphous alloys are not suitable; thus, the practical application of amorphous alloys is restricted to only using them when the specific beneficial properties are not realized in any crystalline materials. The typical example of such environments for corrosion-resistant amorphous alloys is concentrated hydrochloric acids. 5. Corrosion-resistant amorphous alloys For the purpose of choosing the solute elements effective in enhancing the corrosion resistance of homogeneous alloys in concentrated hydrochloric acids, sputter deposition was used for the preparation of binary single-phase alloys consisting only of corrosion-resistant elements such as amorphous Cr–Ti [13], Cr–Zr [14], Cr–Nb [15], Cr–Ta [15] and Mo–Zr [16] alloys and bcc single phase Mo–Ti [17], Mo–Nb [18], Mo–Ta [19] and Mo–Cr [20] alloys. The corrosion rates of crystalline and amorphous Cr-valve metal alloys in concentrated hydrochloric acids are shown in Fig. 3. The corrosion rates of those amorphous binary chromium-valve metal alloys are clearly lower than those of individual component elements of the alloys, and the corrosion rate decreases as the chromium content of the alloy increases, in spite of the fact that the corrosion rate of pure chromium metal is far higher than those for these valve metals. Among all metallic materials the amorphous Cr–Ta alloys are the only ones that do not experience corrosion in 12 M HCl, since these alloys are all spontaneously passive in concentrated hydrochloric acids. The surface characterization was performed using X-ray photoelectron spectroscopy for Cr–Ti [21], Cr–Zr [22], Cr–Nb [23] and Cr–Ta [24,25] alloys. Because chromium and valve metals have similar high affinity to oxygen, the spontaneously formed passive films were composed of both chromium and valve metal cations. Similarly, the anodically formed films on sputter-deposited amorphous Al–Ta and Al–Nb alloys in a borate buffer solution have known to consist of both aluminum and valve metal cations [26]. In general, the composition of the spontaneously formed passive film was not largely different from that of the air-formed film; this indicates that spontaneous passivation occurs because the air-formed film has a high stability in concentrated hydrochloric acids. Anions in the

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Fig. 4. The binding energies of Cr3+ 2p3/2 and Ta5+ 4f7/2 electrons in the surface films formed on the binary Cr–Ta alloys.

films on Cr-valve metal alloys were O2− and OH− , where the relative amount changed with cationic composition of the films from CrO1−x (OH)1+2x to ZrO1.4 (OH)1.2 [22] or TaO2 (OH) [25], although the O2− content was higher in the inner part of the film near the film/alloy interface whereas the OH− content was higher in the outer part of the film near the film/environment interface. As can be seen in Fig. 4, the binding energies of Cr3+ 2p3/2 and Ta5+ 4f7/2 show the parallel change with cationic composition of the film [25]. This fact suggests the charge transfer from Cr3+ ions to Ta5+ ions in the films. This electronic interaction between Cr3+ ions and Ta5+ ions in the film indicates that two kinds of cations located very closely in the film. Thus, the film is not a heterogeneous mixture of two kinds of oxyhydroxides but rather consists of a homogeneous double oxyhydroxides, such as Cr1−x Tax Oy (OH)3+2x−2y . Because the stability of Cr1−x Tax Oy (OH)3+2x−2y is significantly higher than that of TaO2 (OH), those amorphous Cr–Ta alloys did not suffer corrosion when exposed to concentrated hydrochloric acids in contrast to pure tantalum metal that dissolved at the rate of about 2.8 × 10−4 mm y−1 in 12 M HCl. In this manner, the synergistic stability of multiple oxides and oxyhydroxides was often higher than the stability of oxides and oxyhydroxides of the individual component elements. Thus, we expect enhanced corrosion resistance for such newly designed homogeneous alloys, even though the stability constants of nonstoichiometric multiple oxides and oxyhydroxides have not been determined. No concentration gradient of cations was detected in the airformed film. However, with increasing time of immersion in 12 M HCl, after a week there was a slight decrease in the chromium content near the exterior of the film formed on Cr–Nb alloys [23] and Cr–Ta alloys [25]. In contrast, a slight decrease in the zirconium content near the exterior of the film formed on Cr–Zr alloys in 6 M HCl was observed [22]. Fig. 5 shows the corrosion rates of a series of Mo-corrosionresistant element alloys in 12 M HCl for Mo–Ti [17], Mo–Zr [16], Mo–Nb [18], Mo–Ta [19] and Mo–Cr alloys [20]. In contrast to the Cr-valve metal alloys, the formation of Mo-valve metal alloys resulted in only slight decrease in their corrosion rates in 12 M HCl. The affinity of valve metals to oxygen was significantly higher than that of molybdenum. Thus, preferential oxidation of valve metals by air exposure led to the formation of the film enriched with valve metals. Similarly, the anodic oxidation of sputter-deposited amorphous Al–Mo and Al–W alloys led to preferential oxidation of aluminum in a borate buffer solution [27]. Nevertheless, the molybdenum addition to valve metals resulted in a slight decrease in the corrosion rate. Molybdenum is a unique element. Fig. 6 shows potentiostatic polarization and dissolution curves of molybdenum in 6 M HCl [28]. Molybdenum is passive forming a Mo4+ oxide or

Fig. 5. Corrosion rates of sputter-deposited Mo–Ti, Mo–Zr, Mo–Nb, Mo–Ta and Mo–Cr alloys in 12 M HCl.

oxyhydroxide film at potentials higher than the hydrogen equilibrium potential, but at the potentials higher than about 0.15 V vs. SCE, molybdenum experiences transpassive dissolution forming molybdate ions. The open circuit potential of Mo-valve metal alloys was at about 0.1–0.2 V vs. SCE. Thus, molybdenum near the surface of the film exposed to 12 M HCl was dissolved as a molybdate, but the Mo4+ was protected by the overlaying valve metal oxyhydroxide film and thus, found in the inner part of the film. The XPS analysis revealed no electronic interaction between Mo4+ ions and valve metal cations, and higher concentrations of molybdenum particularly Mo4+ in the inner part of the film. Therefore, the spontaneously formed passive film on the Mo-valve metal alloys consisted of bi-layer structure of the outer valve metal oxyhydroxide and the inner Mo4+ oxide. The presence of the inner Mo4+ oxide layer adjacent to the alloys is responsible for their higher corrosion resistance in comparison with the corrosion resistance of the

Fig. 6. Potentiostatic polarization and dissolution curves of molybdenum metal in 6 M HCl.

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component valve metals. The role of molybdenum in enhancing the corrosion resistance will be further discussed in the section of the amorphous bulk alloys. It was found from these studies that chromium, tantalum and molybdenum were effective solute elements in preventing corrosion of amorphous alloys in concentrated hydrochloric acids.

6. Production of amorphous bulk alloys The most of amorphous alloy specimens were prepared by rapid quenching from the liquid state or sputter deposition. However, since about 1994 consolidation of amorphous alloy powders of a specific composition enabled to form amorphous “bulk” alloys [29,30]. When the temperature is increased by heating, the most of traditional amorphous alloys result in crystallization, but some alloys of selected compositions transform to the supercooled liquid state just below the crystallization temperature. If the supercooled liquid is stable enough to process the alloys, consolidation of amorphous alloy powders in the supercooled liquid state enables to form the amorphous alloys in any desirable bulk shapes. The presence of the supercooled liquid state can be detected by thermal analysis. When an amorphous alloy specimen having the supercooled liquid state is heated at a constant heating rate of 20 K min−1 , the endothermic behavior corresponding to the formation of the supercooled liquid is observed just below the crystallization temperature up to the appearance of an exothermic peak for crystallization. The lower temperature limit of the endothermic behavior is called the glass transition temperature, Tg, because cooling of the supercooled liquid under the Tg results in transformation to the amorphous (glass) solid. The lower bottom of the exothermic peak corresponds to the crystallization temperature, Tx. In the temperature range between Tg and Tx this alloy is in the supercooled liquid state. If the temperature interval T = Tx − Tg is approximately 50 K or higher, that is, 2.5 min or longer heating time after reaching Tg at the heating rate of 20 K min−1 , it is possible to form an amorphous alloy plate by consolidation of the amorphous alloy powders using a heated twin-roller type rolling mill after quick pre-heating these amorphous alloy powders into the supercooled liquid state. The first attempt for consolidation of corrosion-resistant amorphous alloy powders was performed for Ni–10Cr–5Nb–16P–4B alloy, for the HCl dew point corrosion test in waste incineration environments with high concentrations of steam and HCl gas [31]. The thermal analysis of the melt-spun ribbon-shaped Ni–10Cr–5Nb–16P–4B alloy at a heating rate of 20 K min−1 revealed that there was a supercooled liquid state between Tg of 676 K to Tx of 740 K. The gas-atomized Ni–10Cr–5Nb–16P–4B alloy powder was consolidated as follows: about 50 g of powder was sealed in a stainless steel tube of 15 mm inner diameter in vacuum; the stainless steel tube specimen was cold-rolled to the thickness of 7 mm; the cold-rolled specimen was quickly heated to 708 K in the supercooled liquid state of this amorphous alloy in a furnace; and immediately warm-rolled at 708 K to the total thickness of 4 mm. After removing the sheath, the alloy specimen of about 2 mm in thickness was obtained. The amorphous structure of the bulk alloy specimen was confirmed by X-ray diffraction. A high-resolution transmission electron microscopy image of ultramicrotomed section of the specimen revealed maze patterns typical of the amorphous structure. The corrosion rate in 6 M HCl at 303 K for this sheath-rolled specimen was 1.7 × 10−2 mm y−1 , whereas that of an as-cast amorphous alloy ribbon specimen was less than 1 × 10−3 mm y−1 . The lower corrosion resistance of that sheathrolled specimen was attributed to the presence of crystallized powder phase in the amorphous matrix, because, in general, the early product of gas atomization was not well amorphized, although the early product was removed as much as possible.

Fig. 7. Corrosion rates of copper-mold cast amorphous Ni–Cr–Mo–xTa–(36–x)Nb–4P alloy rods in 6 M and 12 M HCl and that of copper-mold cast Ni–3Cr–2Mo–22Ta–4Nb–4P alloy rod consisting of crystalline phases with the amorphous matrix in 12 M HCl.

The performance of the sheath-rolled amorphous Ni–10Cr–5Nb–16P–4B alloy specimen was examined by flue gas exposure in a waste incinerator [32]. The amorphous alloy sheet showed no detectable mass loss corrosion after exposure for 20 days at 393 and 433 K; in contrast, the mass losses of Type 316L stainless steel and Alloy B samples after exposure at 433 K were more than 80 and 360 mg cm−2 year−1 , respectively.

7. Amorphous bulk alloys resistant to 12 M HCl Production of amorphous alloy powders and consolidation in the supercooled liquid state are suitable for the production of engineering materials of amorphous bulk alloys but not suitable for alloy design because of time-consuming, sophisticated and expensive procedure. On the other hand, production of amorphous bulk alloy by consolidation in the supercooled liquid state is empirically possible, if the glass-forming ability is high enough to form an amorphous alloy rod of 1 mm diameter by copper-mold casting and if the temperature interval T is 50 K or higher at the heating rate of 20 K min−1 . Thus, conventional examination of amorphous bulk alloy formation was performed by copper-mold casting and thermal analysis of melt-spun amorphous alloy ribbons. The objective was to identify metallic materials resistant to 12 M HCl that can be used widely in HCl environments to which no conventional crystalline materials are resistant. Sputter-deposited amorphous binary Cr–Ta alloys suffered no corrosion in 12 M HCl. However, binary Cr–Ta alloys cannot be solidified as amorphous bulk alloys. Because a Ni–40Nb alloy had relatively high glass-forming ability, enhancement of corrosion resistance was attempted by substituting some niobium with chromium, molybdenum and tantalum. Some addition of phosphorus was made to further enhance the glass-forming ability. Preparation of amorphous Ni–Cr–Mo–(36–x)Ta–xNb–(4–7)P alloy rods was attempted by copper-mold casting. All these alloys were amorphous after melt spinning and showed a wide supercooled liquid region at relatively high temperatures, because of the presence of tantalum. For instance, using thermal analysis at the heating rate of 20 K min−1 , the supercooled liquid region of Ni–4Cr–1Mo–22Ta–14Nb–7P alloy was about 100 K (between Tg = 840 K and Tx = 940 K). Thus, if the amorphous alloy rods of 1 mm or thicker diameter are prepared by copper-mold casting, it would be possible to form desirable bulk shapes of the amorphous alloys by consolidation of those amorphous alloy powders in the supercooled liquid state.

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Fig. 8. Potentiodynamic anodic polarization curves of copper-mold cast amorphous Ni–Cr–Mo–Ta–Nb–4P alloy rods and Ni–3Cr–2Mo–22Ta–4Nb–4P alloy rod consisting of crystalline phases with the amorphous matrix in 6 M HCl.

Fig. 7 shows the corrosion rates in 6 and 12 M HCl for Ni–Cr–Mo–xTa–(36–x)Nb–4P alloy rods of 1 mm diameter made by copper-mold casting [33]. Even in amorphous single phase alloys the addition of sufficient chromium, molybdenum and tantalum was necessary to assure no detectable corrosion in 12 M HCl. A smaller increase in the chromium content was more effective than an equivalent increase in molybdenum or tantalum content. For instance, Ni–2Cr–2Mo–22Ta–14Nb–4P alloy was immune to corrosion in 12 M HCl, but if the chromium content was dropped to 1 at%, the tantalum content had to be increased to 36 at% for the same level of immunity in 12 M HCl, that is, Ni–1Cr–2Mo–36Ta–4P alloy. Although the increase in the chromium content was effective in enhancing the corrosion resistance, “excess” additions of chromium promoted the formation of crystalline phases within the amorphous matrix with a consequent increase in the corrosion rate as can be seen by the higher corrosion rate of the Ni–3Cr–2Mo–22Ta–14Nb–4P alloy. Even if zero corrosion mass loss was observed in concentrated hydrochloric acids, potentiodynamic polarization curves measured in 6 M HCl clearly differentiated the beneficial effects of increases in chromium and molybdenum contents as shown in Fig. 8. These alloys are spontaneously passive without showing active region and the open circuit potentials are higher than 0 V vs. Ag/AgCl. In Ni–Cr–Mo–22Ta–14Nb–4P alloys, an increase in the chromium content from 1Cr–2Mo to 2Cr–2Mo clearly decreases the anodic current, and an increase in molybdenum content from 2Cr–2Mo to 2Cr–3Mo further decreases the anodic current. However, even if chromium is enhancing the corrosion resistance, the “excess” addition of chromium leads to the formation of crystalline phases within the amorphous matrix that also leads to an increase in the anodic current; an example is the Ni–3Cr–2Mo–22Ta–14Nb–4P alloy. Fig. 9 shows the analytical results obtained by X-ray photoelectron spectroscopy, XPS, of the roles of alloying elements on corrosion performance. The Ni–4Cr–1Mo–24Ta–12Nb–7P alloy was immune to corrosion in both 6 and 12 M HCl. The upper right portion shows the cationic fraction in the surface films and the lower right portion the atomic fraction in the underlying alloy surface. Because no nickel was found in the surface film, the alloy composition without nickel is shown in the upper left portion and the nominal alloy composition is shown in the lower left portion. Air exposure before immersion results in preferential oxidation of chromium, tantalum and niobium where the chromium enrichment in the air-formed film is remarkably high.

Fig. 9. Cationic fractions and thickness of surface films and atomic fractions in underlying alloy surfaces for copper-mold cast amorphous Ni–4Cr–1Mo–24Ta–12Nb–7P alloy rods before and after immersion for 168 h in 6 and 12 M HCl.

After immersion in hydrochloric acids, the surface film composition is not largely changed from the air-formed film. This indicates that the air-formed film itself is stable and protective in concentrated hydrochloric acids. Thus, spontaneous passivation occurs. This alloy is able to thicken the protective film depending upon the aggressiveness of the environments as shown in the top of Fig. 9. In this connection, Fig. 10 shows the impact of film thickening on the “no corrosion” phenomena. Because the chromium and molybdenum contents in the Ni–2Cr–1Mo–24Ta–12Nb–4P alloy are not sufficiently high, the alloy is corroded in 12 M HCl although no mass loss is observed in 6 M HCl. The alloy can increase the thickness of the protective film to resist against corrosion in 6 M HCl. However, the corrosion resistance of the alloy is not sufficient in 12 M HCl where it cannot increase the thickness of the protective film. As shown in Fig. 3 zero corrosion mass loss is found only when the double oxyhydroxide film of Cr3+ and Ta5+ is formed, whereas the double oxyhydroxide film of Cr3+ and Nb5+ is not highly stable and less protective particularly in 12 M HCl. Consequently, although Nb5+ is found in the film, the contribution of Nb5+ to the high corrosion resistance in the concentrated hydrochloric acids is low and the formation of triple oxyhydroxide particularly containing Cr3+ and Ta5+ is responsible for no dissolution in 12 M HCl. Another important fact is that the enhancement of the protectiveness of the film depending upon the increase in the aggressiveness of hydrochloric acids is associated with the increasing the amount of oxidized molybdenum in the film. As can be seen in Fig. 9, the fraction of molybdenum in the air-formed film on the Ni–4Cr–1Mo–24Ta–12Nb–7P alloy was 0.034 but after immersion in 6 and 12 M HCl, the fraction of molybdenum in the film increased to 0.039 and 0.043, respectively. In Fig. 10 the mass of molybdenum

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Fig. 10. Masses of cations in the surface films on copper-mold cast amorphous Ni–2Cr–1Mo–22Ta–14Nb–4P alloy rod before and after immersion for 168 h in 6 and 12 M HCl.

in the air-formed film on the Ni–2Cr–1Mo–24Ta–12Nb–7P alloy was 17.5 ng cm−2 but after immersion in 6 M HCl the mass of molybdenum in the spontaneous passive film increased to 25.2 ng cm−2 . However, this alloy which was not sufficiently corrosion-resistant could not increase sufficiently the mass of molybdenum in the less protective film in 12 M HCl and was 20.5 ng cm−2 . Angle-resolved XPS gives the information on the distribution of elements as a function of depth below the surface. When the electron take off angle in XPS is low, the signal from the inner part of the film must travel a longer distance before leaving the specimen and hence be weaker. Consequently, at the lower the take off angles of electron, the strongest intensity signal is from the outer part of the surface film. Fig. 11 shows the analytical results of Xray photoelectron spectra measured at two different electron take off angles taken from the same specimen. The tantalum content is higher for the results analyzed using the data measured at the 30◦ electron take off angle, whereas molybdenum content is higher in the analytical results obtained using the data measured at the 90◦ electron take off angle. These results indicate that tantalum is rich in the outer portion of the film whereas molybdenum is rich in the inner portion of the film. The chromium content shows a similar trend as that for molybdenum. The slight chromium deficiency in the exterior region of the film is similar to that found for the Cr1−x Tax Oy (OH)3+2x−2y film formed on the binary Cr–Ta alloys in 12 M HCl [25]. The affinity of molybdenum for oxygen is significantly lower than those of chromium, tantalum and niobium. After preferential oxidation of chromium, tantalum and niobium, the concentration of molybdenum under the film at the underlying alloy surface becomes high enough so that oxidation of molybdenum can occur forming molybdenum oxide mostly located under the chromium, tantalum and niobium-enriched outer portion of the film. The analytical results of angle-resolved XPS for molybdenum is shown in Fig. 12. Molybdenum, particularly Mo4+ is richer in the inner portion of the film. Oxidation of molybdenum under the overlaying triple oxyhydroxide film of Cr3+ , Ta5+ and Nb5+ will form MoO2 near the underlying alloy surface. If molybdenum is oxidized to the high valence Mo6+ on the outer film surface, the Mo6+ would be dissolved in the form of molybdate.

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Fig. 11. Analytical results of X-ray photoelectron spectra measured at electron take off angles of 30◦ and 90◦ for cationic fractions in the surface film on copper-mold cast amorphous Ni–2Cr–8Mo–22Ta–14Nb–7P alloy rod after immersion for 336 h in 12 M HCl.

Consequently, the high chemical stability of the triple oxyhydroxide of Cr3+ , Ta5+ and Nb5+ is protective without dissolving in 12 M HCl and is responsible for zero corrosion mass loss. The enhanced corrosion resistance caused by molybdenum is due to the formation of a MoO2 film under the triple oxyhydroxide film of Cr3+ , Ta5+ and Nb5+ where the MoO2 film acts as the diffusion barrier suppressing the outward diffusion of cations and inward diffusion of oxygen. In this manner, the passive film formed on the alloys with chromium, molybdenum, tantalum and niobium consists of the bi-layer structure where the outer triple (of Cr3+ , Ta5+ and Nb5+ ) oxyhydroxide portion of the film is stable even in 12 M HCl and the inner MoO2 portion of the film acts as the diffusion barrier that sup-

Fig. 12. Analytical results of X-ray photoelectron spectra measured at electron take off angles of 30◦ and 90◦ for fractions of molybdenum ions in the surface film on copper-mold cast amorphous Ni–2Cr–8Mo–22Ta–14Nb–7P alloy rod after immersion for 336 h in 12 M HCl.

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Fig. 13. Parabolic rate constants of molybdenum metal and Al–Mo alloys for sulfidation and Al–Mo and Al–Mo–Si alloys for oxidation.

presses the corrosive dissolution and oxidation of the metal alloy substrate. 8. Hot corrosion resistance Some sputter-deposited homogeneous alloys containing necessary elements have high resistance to both high temperature sulfidation and oxidation. All oxidation-resistant conventional alloys undergo very rapid degradation in sulfur containing environments due to poor protective properties of the sulfide scales. It has been known that pure molybdenum has a high sulfidation resistance, because MoS2 formed acts as the effective barrier layer to sulfidation [34]. Although molybdenum is oxidized very rapidly, Al–Mo alloys were expected to be resistant to sulfidizing–oxidizing atmospheres encountered particularly in coal conversion and other energy systems. Although the boiling point of aluminum is lower than the melting point of molybdenum, sputter deposition does not require melting for alloy formation and various amorphous Al–Mo alloys were prepared by sputter deposition. Fig. 13 shows sulfidation and oxidation rate constants of molybdenum and Al–Mo [35] and Al–Mo–Si [36] alloys. The sulfidation rate constants of the Al–Mo alloys were many orders of magnitude lower than those of oxidation-resistant alloys, and were more than an order of magnitude lower than those of pure molybdenum. The scale was heterogeneous and composed of a major MoS2 and a minor Al0.5 Mo2 S4 phases. The lower sulfidation rate of these alloys than that of pure molybdenum was explained in terms of doping effect [37,38]. The MoS2 phase constituting the main part of the inner barrier layer of the scale may be substitutionally doped with trivalent aluminum ions. The doping of aluminum ion reduces the concentration of interstitial sulfide ions which are the predominant point defects in this system. Thus, the defect concentration in an Al2 S3 –MoS2 solid solution should be lower than that in pure MoS2 , and consequently, the sulfide scale on Al–Mo alloys were more protective than that on pure molybdenum. The oxidation rate of the Al–Mo alloy was lower than oxidation rates of chromia forming materials, but still slightly higher than those of alumina formers. In order to suppress oxidation, silicon addition to Al–Mo alloys was performed [36]. The oxidation rates of ternary Al–Mo–Si alloys were much lower than those of binary Al–Mo alloys.

Fig. 14. The CO2 conversion in CO2 + 4H2 on the catalyst formed by calcination of the mixture of 50Ni(NO3 )2 + 41.67ZrO2 sol + 8.33Sm(NO3 )3 at 800 ◦ C.

In this manner, amorphous Al–Mo–Si alloys are candidates for sulfidation–oxidation resistant materials. However, they crystallize to form intermetallic compounds by exposure to high temperature environments. Although in laboratory experiments grains formed were very fine and uniform, for practical application, the durability test is required in addition to development of a defect free coating method of these alloys. 9. Catalysts The metastable nature and homogeneous structure of amorphous alloys are attractive as catalysts [39]. One of the useful results has been obtained when amorphous Ni–Zr alloys were used as the catalyst precursor for methane formation by the reaction of carbon dioxide with hydrogen. The catalysts showed the very fast reaction rate and the almost 100% selectivity for methane formation [40]. This was quite different from usual catalysts on which the conversion was very slow and the main product was carbon monoxide. The amorphous Ni–Zr alloys were converted to metallic nickel supported on ZrO2 type oxide during catalytic reaction. The stable ZrO2 is monoclinic but most of ZrO2 formed from the amorphous Ni–Zr alloys were tetragonal. During oxidation of amorphous Ni–Zr alloys in which nickel and zirconium atoms are adjacent to each other, some Ni2+ ions substitute the Zr4+ lattice points in the ZrO2 lattice with a consequent stabilization of the tetragonal ZrO2 . The catalyst consisting of nickel supported on the tetragonal ZrO2 was extraordinarily active for carbon dioxide methanation [41–44]. Amorphous alloys were effective as the catalyst precursors but not suitable for mass production of catalysts for carbon dioxide methanation. However, the study of amorphous alloys as the catalyst precursor revealed that the effective catalyst is metallic nickel supported on tetragonal ZrO2 . The tetragonal ZrO2 can be stabilized by inclusion of rare earth element, Ca2+ and Mg2+ ions. We created the catalysts in the form of powder. Aqueous zirconia sol was used as the zirconia source in which rare earth element, calcium or magnesium salt together with nickel salt were dissolved. Calcination and subsequent reduction in hydrogen gave rise to the formation of the powder catalysts consisting of nickel supported on tetragonal ZrO2 stabilized by rare earth element, Ca2+ or Mg2+ ions. An example of the performance of powder catalyst of Ni supported on the tetragonal ZrO2 stabilized by Sm3+ ions is shown in Fig. 14 [45].

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The performance of the catalyst powder was almost the same as that of the catalyst obtained from the amorphous Ni–Zr-rare earth element alloy precursors. Using these catalysts, the research and development of the technology for the worldwide supply of methane through hydrogen from intermittent, fluctuating and remote electricity generated from renewable energy are in progress [46–51].

10. Reactivity This paper presented at 10th International Symposium on Electrochemical/Chemical Reactivity of Metastable Materials. When international exchange had become free by dissolution of the Soviet Union system, Prof. Maria Janik-Czachor, Polish Academy of Sciences, who had been taking an interest in high reactivity of amorphous and other metastable materials, decided to organize the International Symposium to discuss this subject and to stimulate international and interdisciplinary collaboration. A series of symposia were held in Eastern Europe, in Warsaw in 1993, 1996, 2003, 2007 and 2010, in Szeged in 1995 and 2005, and in Dresden in 1997, with the exception of Sendai in 1998 and Mt. Trembland, Quebec in 2001. She attempted the further enhancement of the reactivity of metastable materials by electrochemical treatment, hydrogen charging, devitrification, mechanical milling, etc. [52–58], for which she made joint researches with Prof. Árpad Molnár, University of Szeged, and others. The metastable nature was responsible for enrichment of corrosion-resistant elements in the passive films on the amorphous alloys as mentioned in Section 3. However, in general, the reactivity is dependent upon not only the metastable nature but also the treatment for enhancement of the reactivity utilizing the homogeneous structure and unique composition of metastable materials. The relation between their characteristics and catalytic properties has been reviewed by Prof. Árpad Molnár [39]. These studies and a series of the International Symposia contributed greatly to a better understanding of chemical properties of metastable materials and their application. The author expresses sincere thanks to Prof. Maria Janik-Czachor for her farsighted efforts.

11. Summary Amorphous alloys have many attractive characteristics, including extremely high corrosion resistance if the sufficient amounts of corrosion-resistant elements are added. The superiority of amorphous alloys is based on the homogeneous single phase nature without any chemical and physical heterogeneities. However, for the formation of the amorphous structure, the movement of constituent atoms to coordinate themselves into crystalline structure during solidification must be prevented. Heating during welding will destroy the superior characteristics by introducing the heterogeneities due to crystallization. Nevertheless, the application of corrosion-resistant amorphous alloys will be expected to the very aggressive environments where any conventional crystalline metallic materials cannot be used. Concentrated hydrochloric acids are examples of such environments. Some amorphous bulk alloys showed zero corrosion mass loss due to spontaneous passivation even in 12 M HCl, although the alloy compositions suitable for forming corrosion-resistant amorphous bulk alloys are limited. The technology for producing amorphous bulk alloys in desirable shapes was also interpreted. The homogeneous single phase nature is also useful as the precursor of catalysts. An example for carbon dioxide methanation is described, which will be useful for supply of renewable energy in the form of methane.

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