Iron-based bulk metallic glasses

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1 International Materials Reviews ISSN: (Print) (Online) Journal homepage: Iron-based bulk metallic glasses C Suryanarayana & A Inoue To cite this article: C Suryanarayana & A Inoue (2013) Iron-based bulk metallic glasses, International Materials Reviews, 58:3, , DOI: / Y To link to this article: Published online: 12 Nov Submit your article to this journal Article views: 4322 View related articles Citing articles: 126 View citing articles Full Terms & Conditions of access and use can be found at

2 Iron-based bulk metallic glasses C. Suryanarayana* 1 and A. Inoue 2 The current status of research and development in Fe-based bulk metallic glasses (BMGs) is reviewed. Bulk metallic glasses are relatively new materials possessing a glassy structure and large section thickness. These materials have an exciting combination of properties such as high mechanical strength, good thermal stability, large supercooled liquid region and potential for easy forming. Ever since the first synthesis of an Fe-based BMG in an Fe Al Ga P C B system in 1995, there has been intense activity on the synthesis and characterisation of. These BMGs exhibit some unique characteristics which have not been obtained in conventional Fe-based crystalline alloys. This uniqueness has led to practical uses of these bulk glassy alloys as soft magnetic and structural materials. This review presents the recent results on the glassforming ability, structure, thermal stability, mechanical properties, corrosion behaviour, soft magnetic properties and applications of Fe-based bulk glassy alloys developed during the last 15 years. This review also highlights the advanced analysis of their properties which has contributed significantly to the progress in understanding and developing of the. The future prospects of have also been presented. Keywords: Fe-based bulk metallic glasses, Soft magnetic materials, Structural materials, Glass-forming ability, Applications Introduction Several advanced materials have been developed during the past few decades to meet the ever-growing requirements and demands of the industry. 1 Novel and advanced materials have been synthesised which are stronger, harder, stiffer, lighter, highly corrosion resistant, or for use at elevated temperatures. All these have been possible by processing materials through different non-equilibrium processing techniques such as rapid solidification processing, mechanical alloying, plasma processing, spray forming, laser processing and different types of vapour deposition methods. 2 Several new and advanced materials have also been developed during the last few decades, which include metallic glasses in thin ribbon form, 3 5 bulk metallic glasses, 6,7 quasicrystals, 8,9 high-temperature superconductors, superhard carbonitrides, thin-film diamond materials and nanostructured materials Metallic glasses are one such class of materials that display interesting combination of properties and have found applications in a variety of industries. 3 7 The first synthesis of a metallic glass was achieved in 1960 by rapidly solidifying a metallic liquid of Au 25 at-%si composition at cooling rates of y10 6 Ks Since then a very large number of alloy compositions in different alloy systems were synthesised in the glassy state. 3 5,7,15 19 However, these metallic glasses 1 Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL , USA 2 Tohoku University, Katahira, Aoba-ku, Sendai , Japan *Corresponding author, Challapalli.Suryanarayana@ucf.edu were produced in the form of thin ribbons (typically mm in thickness), powders or wires, since the critical cooling rate required to form the glassy phases in these alloy compositions was of the order of Ks 21. And these rates could be achieved only when at least one of the dimensions of the solidified material (section thickness) was a few micrometres in thickness. This limitation on the section thickness has prevented widespread use of these glassy materials in different application fields. For example, glassy alloys in sheet or wire form in the Fe Si B and Co Fe Si B systems have been used as soft magnetic materials, 20 but there have not been many reports regarding applications of these novel materials in other fields such as structural materials or as functional materials. Intensive developments during the last 20 years have resulted in the synthesis of metallic glasses with large section thickness or diameter, known as bulk metallic glasses (BMGs). 6,7 The observation of the stable phenomenon of supercooled liquid in metallic alloys made it possible to produce glassy alloys in a bulk form with diameters up to y72 mm. 6,7,21 24 Pioneering work has been done at the Tohoku University in Sendai, Japan and at the California Institute of Technology in Pasadena, California, USA, in this field of study. Research investigations in this field are now being taken up in many other laboratories around the world, notably in China, Korea and India. Bulk metallic glasses have been produced in a large number of alloy systems, for example, those based on Co, 25 Cu, 26 Fe, 27 lanthanum (Ln), 28 Mg, 29 Ni, 30 Pd, 31 Pt, 32 Ti 33 and Zr, 34,35 with critical rod diameters exceeding 10 mm. 36 Bulk metallic glasses have also been synthesised in other alloy systems based on ß 2013 Institute of Materials, Minerals and Mining and ASM International Published by Maney for the Institute and ASM International DOI / Y International Materials Reviews 2013 VOL 58 NO 3 131

3 Au, 37 Ca, 38 Ce, 39 Hf, 40 Nd, 41 Pr, 42 Sm 43 and Y. 44 Subsequent efforts by different groups also resulted in the development of BMGs with critical diameters of.20 mm in alloy systems based on Cu, 45 Ln, 46 Mg, 47 Ni, 48 Pd, 49 Pt 50 and Zr. 51,52 The general features of BMGs have been described in some reviews and books. 6,7,22 24 Among the BMG alloys, Fe-based bulk glassy alloys are very attractive because of their excellent soft magnetic properties with rather high saturation magnetisation, high electrical resistivity, high mechanical strength and low materials cost, among others. 53 Ever since the first Fe-based bulk glassy alloy was synthesised in the Fe Al Ga P C B system in 1995 by a coppermould casting technique, 27 a very large number of Febased BMGs including Fe-, Fe Co-, Fe Ni- and Fe Co Ni-based alloy systems have been developed to date. This article aims to summarise the details of the developments on the structure, synthesis, properties and applications of Fe-, Fe Co- and Fe Co Ni-based BMGs developed during the last 15 years or so. In spite of the industrial importance of Fe-based alloys, and the possibility of achieving novel attributes for these alloys in the glassy state, there has been only one major review on this topic so far, 54 which was written at the early stage of development of. The present article plans to fill that gap by presenting a comprehensive review of the literature in this developing field. Characteristics of The first synthesis of an Fe-based BMG with a diameter of 1 mm was reported in 1995 in the Fe Al Ga P C B alloy system by Inoue and co-workers. 27,55 Since then a large number of Fe-based alloy compositions have been synthesised, with the maximum thickness reaching y16 mm. 56 Table 1 presents the details of the composition, synthesis method and the thermal properties of the that have been synthesised till recently. 27,55, The alloy systems have been arranged in such a way that the metalloid content increases on going down the table. Metallic glasses produced by the rapid solidification processing (RSP) method were traditionally classified into the metal metalloid and metal metal types. In the metal metalloid type, the metal content is typically y80 at-% and the metalloid content y20 at-%. The metal can be either just one element or a combination of more than one element and similarly, the metalloid element also can be either just one or a combination of more than one metalloid element. Typical examples of metal metalloid type Fe-based metallic glasses produced by the RSP method are: Fe 80 B 20, Fe 78 B 12 Si 10, Fe 80 P 13 C 7, Fe 40 Ni 40 P 14 B 6, Fe 71?3 Cr 10 Mo 9 P 8 C 1?7 and Fe 73?4 Cu 1 Nb 3?1 B 9?1 Si 13?4. On the other hand, in the metal metal glasses, there is no such strict compositional requirement. The metal elements can be present in any proportion, and some typical compositions studied include Fe 90 Zr 10 and Fe 60 Zr 40. In principle, BMGs can also be produced in both the metal metalloid and metal metal varieties. But it is interesting to note that all the synthesised so far are of the metal metalloid type. Here again, the metallic component constitutes y80 at-% and the metalloid component y20 at-%. Furthermore, like in the case of rapidly solidified alloys, the metal component can be either only Fe or a mixture of different elements. Majority of the alloying elements are typically from the Fe-group elements, i.e. Fe, Co and/or Ni. Occasionally, other metallic elements such as Cr, Mn, Al, Ga, Mo, Zr, Nb and Ta are also added, with their concentrations ranging from as small as a few per cent to as large as nearly 15 20% total. Rare-earth elements such as Y, Er, Gd and Tm are also occasionally added with beneficial effects of increasing the glass formability. The metalloid elements added are typically B, C, P and Si, with their total content amounting to nearly 20 at-%. Bulk metallic glasses have also been produced in Nd Fe Al 123 and Pr Fe Al 124 even when the Fe content was.50 at-%. But alloy compositions with useful properties were found to be very lean in Fe content. Accordingly, these will not be discussed further in this review. Table 2 summarises the alloy compositions of typical with critical diameters of.1 mm together with the calendar years of their first synthesis. From this table, it becomes clear that the different alloy compositions and the critical diameters listed can be basically classified into two groups based on the solute additions. The first group contains Fe and metalloid elements along with Al and Ga as the main metallic solute elements. These alloys, sometimes, also contain the early transition metals (ETMs). 55,104,125,126 On the other hand, the second group is composed of Fe and metalloid elements, and sometimes they also contain ETM elements or lanthanum (Ln). 91,98,99,103,105, The second group can also be described as the Fe ETM/Ln metalloid type by putting the ETM and Ln together as ETM/Ln due to their similarity in the location in the periodic table. The second group can be more specifically divided into two series by paying attention to the amount of Fe present in the alloy: alloys with Fe >50 at-% (nos. F5 F9) and the others with Fe,50 at-% (nos. F10 F13). The addition of rare-earth elements (e.g. Y and Tm) seems to increase the glass-forming ability (GFA) of these alloys as evidenced by the synthesis of large diameter BMGs reaching up to 16 mm. 56,191 It may be emphasised that all the Fe-based BMGs contain metalloid elements such as P, C, B and Si and belong to the metal metalloid type system. The significant difference in alloy compositions between BMG alloys and the previously reported thin metallic glass ribbons is attributed to the addition of special alloying elements such as Ln, Ga, Zr, Nb and Mo having a significant atomic size difference and large negative heat of mixing between Fe and the metalloid (P, C, B and/or Si) elements. Glass-forming ability To produce metallic glasses in a reasonable and reliable way, and also to produce them in large quantities, it is essential that we understand the basic reasons regarding ease of glass formation. The ability of a crystalline metallic alloy to transform into the glassy state is defined as the GFA. There has been reasonable success in predicting the compositions and alloy systems in which thin ribbon glasses could be synthesised by RSP methods during the 1970s and 1980s. But with the discovery of BMGs, the activity in this area has been resumed in recent times. Both the alloy systems and their compositions that are likely to be transformed into the glassy condition have been predicted and experimentally verified in some cases. However, as we will see, it has been noted that these predictions seem to have been 132 International Materials Reviews 2013 VOL 58 NO 3

4 Table 1 Thermal properties of * Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference Fe 61 Co 7 Zr 8 Mo 7 B 15 Y 2 Drop casting Fe 61 Co 6 Zr 8 Al 1 Mo 7 B 15 Y 2 Drop casting Fe 61 Co 5 Zr 8 Cr 2 Mo7B 15 Y 2 Drop casting Fe 60 Co 8 Zr 10 Mo 5 W 2 B 15 Cu-mould casting Fe 61 Co 7 Zr 10 Mo 5 W 2 B 15 Suction casting Fe 59 Co 9 Zr 10 Mo 5 W 2 B 15 Suction casting Fe 57 Co 11 Zr 10 Mo 5 W 2 B 15 Suction casting Fe 66 Co 4 Hf 5 Mo 7 B 15 Y 3 Suction casting Fe 62 Co 8 Hf 5 Mo 7 B 15 Y 3 Suction casting Fe 60 Co 10 Hf 5 Mo 7 B 15 Y 3 Suction casting Fe 58 Co 12 Hf 5 Mo 7 B 15 Y 3 Suction casting Fe 61 Co 7 Zr 9?5 Mo 5 W 2 B 15?5 Cu-mould casting Fe 84 Si 4 B 12 Melt spinning Fe 83?3 Si 4 B 12 Cu 0?7 Melt spinning Fe 83?3 Si 4 B 8 P 4 Cu 0?7 Melt spinning Fe 83?3 Si 4 B 8 P 4 Cu 0?7 Melt spinning Fe 80 P 12 B 4 Si 4 Cu-mould casting Fe 79 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 76 Si 9 B 10 P 5 Cu-mould casting Fe 80 P 12 C 4 B 4 66 Fe 79 Ga 1 P 12 C 4 B Fe 76 Ga 4 P 12 C 4 B 4 66 Fe 75 Ga 5 P 12 C 4 B 4 Cu-mould casting Fe 78 Ga 2 P 12 C 4 B 4 Cu-mould casting , 67 Fe 78 Ga 2 P 10?5 C 4 B 4 Si 1?5 Cu-mould casting 40 67, 68 Fe 78 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting , 68 Fe 78 Ga 2 P 9?5 C 4 B 4 Si 2?5 Cu-mould casting , 68 Fe 78 Ga 2 P 9 C 4 B 4 Si 3 Cu-mould casting 40 67, 68 Fe 78 Ga 2 P 8?5 C 4 B 4 Si 3?5 Cu-mould casting Fe 77 Ga 3 P 12 C 4 B 4 Cu-mould casting 40 66, 67 Fe 77 Ga 3 P 10?5 C 4 B 4 Si 1?5 Cu-mould casting Fe 77 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 77 Ga 3 P 9?5 C 4 B 4 Si 2?5 Cu-mould casting Fe 77 Ga 3 P 9 C 4 B 4 Si 3 Cu-mould casting Fe 77 Ga 3 P 8?5 C 4 B 4 Si 3?5 Cu-mould casting Fe 76 Al 4 P 12 B 4 Si 4 Cu-mould casting Fe 74 Al 4 Ga 2 P 12 B 4 Si 4 Cu-mould casting Fe 76 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 76 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Injection moulding Fe 74 Cr 2 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Injection moulding Fe 72 Cr 4 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Injection moulding Fe 70 Cr 6 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Injection moulding Fe 70 Co 5 Ga 5 P 12 C 4 B 4 Cu-mould casting 71 Fe 67?5 Co 7?5 Ga 5 P 12 C 4 B 4 Cu-mould casting 71 Fe 65 Co 10 Ga 5 P 12 C 4 B 4 Cu-mould casting Fe 65 Co 10 Ga 5 P 12 C 4 B 4 Atomisationzspark plasma sintering Fe 62?5 Co 12?5 Ga 5 P 12 C 4 B 4 Cu-mould casting Fe 60 Co 5 Ga 15 P 12 C 4 B 4 Cu-mould casting Fe 57?5 Co 17?5 Ga 5 P 12 C 4 B 4 Cu-mould casting 71 Fe 55 Co 20 Ga 5 P 12 C 4 B 4 Cu-mould casting 71 Fe 74 Al 4 Ga 2 P 11 C 5 B 4 Injection casting Fe 73 Al 5 Ga 2 P 11 C 5 B 4 Injection casting Fe 72 Al 5 Ga 2 P 11 C 6 B 4 Melt spinning Fe 72 Al 5 Ga 2 P 11 C 6 B 4 Melt spinning Fe 78 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting 68 International Materials Reviews 2013 VOL 58 NO 3 133

5 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference Fe 77 Mo 1 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 76 Mo 2 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 74 Mo 4 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 72 Mo 6 Ga 2 P 10 C 4 B 4 Si 2 Cu-mould casting 68 Fe 77 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting 68 Fe 76 Mo 1 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 75 Mo 2 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 73 Mo 4 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 71 Mo 6 Ga 3 P 10 C 4 B 4 Si 2 Cu-mould casting Fe 65 Co 10 Ga 5 P 12 C 4 B Fe 74 Mo 6 P 10 C 7?5 B 2?5 Cu-mould casting Fe 70?3 Ni 3?7 Mo 6 P 10 C 7?5 B 2?5 Cu-mould casting Fe 69 Ni 5 Mo 6 P 10 C 7?5 B 2?5 Cu-mould casting Fe 66?6 Ni 7?4 Mo 6 P 10 C 7?5 B 2?5 Cu-mould casting Fe 62?9 Ni 11?1 Mo 6 P 10 C 7?5 B 2?5 Cu-mould casting Fe 72 Al 5 Ga 2 P 10 C 6 B 4 Si 1 Injection casting Fe 72 Al 5 Ga 2 P 10 C 6 B 4 Si 1 Rapid quenching Fe 70 Al 5 Ga 2 P 9?65 C 5?75 B 4?6 Si 3 Cu-mould casting (ring-shape) Fe 70 Al 5 Ga 2 P 9?65 C 5?75 B 4?6 Si 3 Melt spinning Fe 70 Al 5 Ga 2 P 9?65 C 5?75 B 4?6 Si 3 Cu-mould casting Fe 67 Co 7 Zr 6 B 20 Cu-mould casting Fe 64 Co 7 Zr 6 Nd 3 B 20 Cu-mould casting Fe 67 Co 9?5 Nd 3 Dy 0?5 B 20 Melt spinning Fe 67 Co 9?5 Nd 3 Dy 0?5 B 20 Injection casting Fe 57 Co 9?5 Nd 3 Dy 0?5 B 30 Injection casting Fe 63 Co 7 Zr 10 B 20 Melt spinning Fe 63 Co 7 Nb 2 Zr 8 B 20 Melt spinning 80 Fe 63 Co 7 Nb 4 Zr 6 B 20 Melt spinning Fe 63 Co 7 Nb 6 Zr 4 B 20 Melt spinning 80 Fe 64 Co 7 Zr 6 B 20 Nd 3 Cu-mould casting Fe 56 Co 7 Ni 7 Zr 10 B 20 Melt spinning Fe 56 Co 7 Ni 7 Zr 8 Nb 2 B 20 Melt spinning Fe 56 Co 7 Ni 7 Zr 6 Nb 4 B 20 Melt spinning Fe 56 Co 7 Ni 7 Zr 7?5 Nb 2?5 B 20 Suction casting Fe 56 Co 7 Ni 7 Zr 6 Nb 2?5 Ta 1?5 B 20 Suction casting Fe 56 Co 7 Ni 7 Zr 6 Nb 2?5 Ti 1?5 B 20 Suction casting Fe 56 Co 7 Ni 7 Zr 6 Nb 2?5 Mo 1?5 B 20 Suction casting Fe 61 Co 13?5 Zr 1 Pr 4?5 B 20 Fe 61 Co 13?5 Zr 1 Pr 3?5 Dy 1 B 20 Suction casting into Suction casting into, Fe 67 Co 9?5 Nd 3 Dy 0?5 B 20 Cu-mould casting Fe 61 Co 10 Zr 5 W 4 B Fe 72 Si 4 B 20 Nb 4 Cu-mould casting Fe 72 Si 4 B 20 Nb 4 Cu-mould casting Fe 70 Si 4 B 20 Nb 6 Cu-mould casting Fe 72 Si 9?6 Nb 4 B 14?4 Melt spinning Fe 30 Co 30 Ni 15 Si 8 B (Fe 0?5 Co 0?5 ) 72 B 20 Si 4 Nb 4 Cu-mould casting Fe 72 B 22 Y 6 Injection into Fe 72 B 22 Y 4 Ti 2 Injection into International Materials Reviews 2013 VOL 58 NO 3

6 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference Fe 72 B 22 Y 4 Ta 2 Injection into Fe 72 B 22 Y 4 Nb 2 Injection into Fe 72 B 22 Y 4 Hf 2 Injection into (Fe 72 B 22 Y 6 ) 98 Ti 2 Injection into (Fe 72 B 22 Y 6 ) 98 Ta 2 Injection into (Fe 72 B 22 Y 6 ) 98 Nb 2 Injection into Fe 68 Zr 10 B 22 Melt spinning (Fe 0?75 Mn 0?25 ) 70 Zr 9 B 21 Melt spinning (Fe 0?70 Mn 0?25 Cr 0?05 ) 68 Zr 7 Nb 3 B 22 Melt spinning (Fe 0?69 Mn 0?26 Cr 0?05 ) 68 Zr 10 B 19 C 3 Melt spinning (Fe 0?69 Mn 0?26 Cr 0?05 ) 70 Zr 4 Nb 4 B 22 Melt spinning (Fe 0?69 Mn 0?26 Cr 0?05 ) 68 Zr 6 Nb 2 B 24 Melt spinning (Fe 0?69 Mn 0?26 Cr 0?05 ) 68 Zr 4 Nb 4 B 24 Melt spinning (Fe 0?70 Mn 0?30 ) 65 Zr 4 Nb 4 Mo 3 B 24 Melt spinning Fe 72 (Nb 0?6 Zr 0?4 ) 6 B 22 Cu-mould casting Fe 71 Nb 6 B 23 Cu-mould casting Fe 71 (Nb 0?8 Zr 0?2 ) 6 B 23 Cu-mould casting Fe 71 B 23 Nb 6 Cu-mould casting 1 93 Fe 76 Si 9?6 B 8?4 P 6 Cu-mould casting (Fe 0?76 Si 0?096 B 0?084 P 0?06 ) 99?9 Cu 0?1 Cu-mould casting Fe 65 Mo 14 C 15 B 6 Fe 64?75 Mo 14 C 15 B 6 Er 0?25 Fe 64?5 Mo 14 C 15 B 6 Er 0?5 Fe 64 Mo 14 C 15 B 6 Er 1 Fe 63?5 Mo 14 C 15 B 6 Er 1?5 Fe 63 Mo 14 C 15 B 6 Er 2 Fe 64?5 Mo 14 C 15 B 6 Dy 0?5 Fe 64 Mo 14 C 15 B 6 Dy 1 Fe 63 Mo 14 C 15 B 6 Dy 2 (Fe 0?9 Co 0?1 ) 58?5 Cr 6 Mo 14 C 15 B 6 Er 0?5 (Fe 0?9 Co 0?1 ) 58?5 Cr 6 Mo 14 C 18 B 3 Er 0?5 (Fe 0?9 Co 0?1 ) 58?5 Cr 6 Mo 14 C 19 B 2 Er 0? Fe 50 Cr 15 Mo 14 C 15 B 6 Cu-mould casting, Fe 50 Cr 15 Mo 14 C 15 B 6 Fe 49 Cr 15 Mo 14 C 13 B 8 Er 1 Fe 49 Cr 15 Mo 14 C 15 B 6 Er 1 Fe 49 Cr 15 Mo 14 C 17 B 4 Er 1 Fe 49 Cr 15 Mo 14 C 18 B 3 Er 1 Fe 49 Cr 15 Mo 14 C 19 B 2 Er 1 Injection casting into International Materials Reviews 2013 VOL 58 NO 3 135

7 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference Fe 49 Cr 15?3 Mo 15 C 15 Y 2 B 3?4 N 0?3 Cu-mould casting Fe 48 Cr 15 Mo 14 C 15 B 6 Tm 2 Cu-mould casting (Fe 0?8 Co 0?2 ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 Cu-mould casting (Fe 0?6 Co 0?4 ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 Cu-mould casting Fe 78 Mo 1 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 77 Mo 2 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 76 Mo 3 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 76 Mo 2 Ga 2 P 10 C 4 B 4 Si Fe 75 Mo 4 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 74 Mo 5 P 10 C 4 B 4 Si 3 Cu-mould casting Fe 73 Mo 6 P 10 C 4 B 4 Si 3 Cu-mould casting 64 Fe 62?8 Co 10 B 13?5 Si 10 Nb 3 Cu 0?7 Cu-mould casting Fe 52 Cr 15 Mo 9 C 15 B 6 Er 3 Fe 48 Cr 15 Mo 14 C 15 B 6 Er 2 Fe 48 Cr 19 Mo 10 C 15 B 6 Er 2 Fe 48 Cr 10 Mo 19 C 15 B 6 Er 2 Fe 48 Cr 15 Mo 14 C 15 B 6 Dy 2 Injection casting into Injection casting into Injection casting into Injection casting into Injection casting into Fe 50 Cr 15 Mo 14 C 15 B 6 Suction casting Fe 49?5 Cr 15 Mo 14 C 15 B 6 Y 0?5 Suction casting Fe 49 Cr 15 Mo 14 C 15 B 6 Y 1?0 Suction casting Fe 48?5 Cr 15 Mo 14 C 15 B 6 Y 1?5 Suction casting Fe 48 Cr 15 Mo 14 C 15 B 6 Y 2 Suction casting Fe 48 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 48 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 45 Co 3 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 43 Co 5 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 41 Co 7 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 39 Co 9 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 41 Co 7 Cr 15 Mo 14 C 15 B 6 Y 2 Fe 61 Cr 4 Mo 14 C 15 B 6 Fe 59 Cr 6 Mo 14 C 15 B 6 Fe 50 Cr 15 Mo 14 C 15 B 6 Fe 60?5 Cr 4 Mo 14 C 15 B 6 Er 0?5 Fe 63 Mo 14 C 15 B 6 Er 2 Fe 55 Cr 8 Mo 14 C 15 B 6 Er 2 Fe 48 Cr 15 Mo 14 C 15 B 6 Er 2 Drop casting in Injection casting into Drop casting in Drop casting in Drop casting in Drop casting in Drop casting in Fe 64 Cr 10 Mo 5 C 8 P 13 Melt spinning International Materials Reviews 2013 VOL 58 NO 3

8 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference (Fe 0?9 Co 0?1 ) 64?875 Mo 14 C 15 B 6 Er 0?125 (Fe 0?9 Co 0?1 ) 64?75 Mo 14 C 15 B 6 Er 0?25 (Fe 0?9 Co 0?1 ) 64?5 Mo 14 C 15 B 6 Er 0?5 (Fe 0?9 Co 0?1 ) 64?25 Mo 14 C 15 B 6 Er 0?75 (Fe 0?9 Co 0?1 ) 64 Mo 14 C 15 B 6 Er 1 (Fe 0?9 Co 0?1 ) 63 Mo 14 C 15 B 6 Er 2 (Fe 0?8 Co 0?2 ) 65 Mo 14 C 15 B 6 (Fe 0?8 Co 0?2 ) 64?75 Mo 14 C 15 B 6 Er 0?25 (Fe 0?8 Co 0?2 ) 64?5 Mo 14 C 15 B 6 Er 0?5 (Fe 0?8 Co 0?2 ) 64 Mo 14 C 15 B 6 Er 1 (Fe 0?7 Co 0?3 ) 65 Mo 14 C 15 B 6 (Fe 0?7 Co 0?3 ) 64?75 Mo 14 C 15 B 6 Er 0?25 (Fe 0?7 Co 0?3 ) 64?5 Mo 14 C 15 B 6 Er 0?5 (Fe 0?7 Co 0?3 ) 64 Mo 14 C 15 B 6 Er Fe 71 Cr 4 Mo 4 P 11 B 5 C 5 Water quenching Fe 67 Cr 4 Mo 4 Ga 4 P 11 B 5 C 5 Water quenching Fe 62 Co 5 Cr 4 Mo 4 Ga 4 P 11 B 5 C 5 Water quenching Fe 65 Sb 2 Cr 4 Mo 4 Ga 4 P 11 B 5 C 5 Water quenching (Fe 66 Cr 4 Mo 4 Ga 4 P 12 C 5 ) 94?5 B 5?5 Water quenching Fe 64 Mo 14 C 15 B 7 Fe 54 Mn 10 Mo 14 C 15 B 7 Fe 50 Mn 10 Mo 14 Cr 4 C 15 B 7 Fe 50 Mn 10 Mo 14 Cr 4 C 16 B 6 Fe 49 Mn 10 Mo 14 Cr 4 W 1 C 16 B 6 Fe 51 Mn 10 Mo 14 Cr 4 C 15 B 6 Fe 48 Mn 10 Mo 16 Cr 4 C 15 B 7 Fe 49 Mn 10 Mo 14 Cr 4 W 1 C 15 B 7 Fe 48 Mn 10 Mo 13 Cr 4 W 3 C 15 B 7 Fe 49 Mn 10 Mo 13 Cr 3 W 3 C 15 B 7 Fe 46 Mn 10 Mo 16 Cr 4 Ga 2 C 15 B 7 Fe 49 Mn 10 Mo 14 Cr 4 V 1 C 15 B 7 Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Injecting into a Fe 49 Mn 10 Mo 13 Cr 3 W 3 C 15 B 7 Cu-mould casting Fe 48 Mn 10 Mo 16 Cr 4 C 15 B 7 Melt spinning Fe 45 Cr 16 Mo 16 C 18 B 5 Melt spinning Fe 45 Cr 16 Mo 14 Nb 2 C 18 B 5 Melt spinning Fe 45 Cr 16 Mo 14 TA 2 C 18 B 5 Melt spinning International Materials Reviews 2013 VOL 58 NO 3 137

9 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference Fe 46 Cr 16 Mo 16 C 18 B 4 Cu-mould casting Fe 44 Cr 16 Mo 16 C 18 B 6 Cu-mould casting Fe 42 Cr 16 Mo 16 C 18 B 8 Cu-mould casting Fe 43 Cr 16 Mo 16 C 15 B 10 Cu-mould casting Fe 43 Cr 16 Mo 16 C 17?5 B 7?5 Cu-mould casting Fe 43 Cr 16 Mo 16 C 15 B 5 P 5 Cu-mould casting Fe 43 Cr 16 Mo 16 C 20 B 5 Cu-mould casting Fe 43 Cr 16 Mo 16 C 10 B 15 Cu-mould casting (Fe 0?775 B 0?125 Si 0?10 ) 98 Nb 2 Cu-mould casting (Fe 0?75 B 0?2 Si 0?05 ) 99 Nb 1 Cu-mould casting (Fe 0?75 B 0?2 Si 0?05 ) 98 Nb 2 Cu-mould casting (Fe 0?75 B 0?2 Si 0?05 ) 96 Nb 4 Cu-mould casting (Fe 0?75 Si 0?1 B 0?15 ) 96 Nb 4 Cu-mould casting Fe 71 Mo 5 P 12 C 10 B 2 Cu-mould casting Fe 69 Cr 2 Mo 5 P 12 C 10 B 2 Cu-mould casting Fe 69 Mo 7 P 12 C 10 B 2 Cu-mould casting Fe 66 Mo 10 P 12 C 10 B 2 Cu-mould casting Fe 64 Cr 3 Mo 10 P 10 C 10 B 3 Cu-mould casting Fe 63 Cr 3 Mo 10 P 12 C 10 B 2 Cu-mould casting Fe 63 Cr 3 Mo 12 P 10 C 7 B 5 Cu-mould casting Fe 65 Cr 2 Mo 9 P 10 C 8 B 6 Cu-mould casting Fe 71?2 B 24 Y 4?8 Cu-mould casting, (Fe 71?2 B 24 Y 4?8 ) 98 Nb 2 Cu-mould casting (Fe 71?2 B 24 Y 4?8 ) 96 Nb 4 Cu-mould casting (Fe 71?2 B 24 Y 4?8 ) 94 Nb 6 Cu-mould casting (Fe 71?2 B 24 Y 4?8 ) 92 Nb 8 Cu-mould casting (Fe 0?75 B 0?2 Si 0?05 ) 96 Nb 4 Cu-mould casting [(Fe 0?9 Co 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?8 Co 0?2 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?7 Co 0?3 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?5 Co 0?5 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?9 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Melt spinning [(Fe 0?6 Ni 0?4 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Melt spinning [(Fe 0?5 Ni 0?5 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Melt spinning (Fe 0?75 B 0?2 Si 0?05 ) 96 Nb 4 Cu-mould casting [(Fe 0?8 Ni 0?2 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Ni 0?4 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?8 Co 0?1 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?8 Co 0?1 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Co 0?1 Ni 0?3 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Co 0?2 Ni 0?2 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Co 0?3 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 Cu-mould casting [(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Nb 0?04 Cu-mould casting {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Cu-mould casting Nb 0?04 } 99 Cr 1 {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Cu-mould casting Nb 0?04 } 98 Cr 2 {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Cu-mould casting Nb 0?04 } 97 Cr 3 {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Nb 0?04 } 96 Cr 4 Cu-mould casting [(Fe 0?5 Co 0?5 ) 0?75 B 0?2 Si 0?05 ] 0?96 Nb 0?04 Cu-mould casting (Fe 0?75 B 0?15 Si 0?10 ) 99 Zr 1 Cu-mould casting (Fe 0?75 B 0?15 Si 0?10 ) 99 Nb 1 Cu-mould casting (Fe 0?75 B 0?15 Si 0?10 ) 98 Nb 2 Cu-mould casting (Fe 0?75 B 0?15 Si 0?10 ) 96 Nb 4 Cu-mould casting International Materials Reviews 2013 VOL 58 NO 3

10 Table 1 Continued Alloy composition Synthesis method t max / mm T g /K T x /K T l /K DT x 5 T x 2 T g /K Heating rate/ K min 21 Reference (Fe 44?3 Cr 5 Co 5 Mo 12?8 Mn 11?2 C 15?8 Drop casting B 5?9 ) 98?5 Y 1?5 (Fe 44?3 Cr 10 CMo 13?8 Mn 11?2 C 15?8 B 5?9 ) 98?5 Y 1?5 Drop casting Fe 43 Cr 16 Mo 16 C 15 B 10 Melt spinning Fe 60 Mo 15 C 15 B 10 Cu-mould casting Fe 60 Mo 15 C 15 B 10 Cu-mould casting Fe 52?5 Cr 15 Mo 7?5 C 15 B 10 Cu-mould casting Fe 52?5 Cr 7?5 Mo 15 C 15 B 10 Cu-mould casting Fe 52?5 Cr 7?5 Mo 15 C 15 B 10 Cu-mould casting Fe 45 Cr 15 Mo 15 C 15 B 10 Cu-mould casting Fe 45 Cr 15 Mo 15 C 15 B 10 Cu-mould casting Fe 45 Cr 7?5 Mo 22?5 C 15 B 10 Cu-mould casting Fe 43 Cr 16 Mo 16 C 15 B 10 Cu-mould casting Fe 37?5 Cr 22?5 Mo 15 C 15 B 10 Cu-mould casting Fe 37?5 Cr 22?5 Mo 15 C 15 B 10 Cu-mould casting Fe 37?5 Cr 15 Mo 22?5 C 15 B 10 Cu-mould casting Fe 30 Cr 30 Mo 15 C 15 B 10 Cu-mould casting Fe 30 Cr 30 Mo 15 C 15 B 10 Cu-mould casting Fe 60 Mo 15 C 15 B 10 Cu-mould casting Fe 52?5 Cr 7?5 Mo 15 C 15 B 10 Cu-mould casting Fe 45 Cr 15 Mo 15 C 15 B 10 Cu-mould casting Fe 37?5 Cr 22?5 Mo 15 C 15 B 10 Cu-mould casting Fe 30 Cr 30 Mo 15 C 15 B 10 Cu-mould casting Fe 71?2 C 7?0 Si 3?3 B 5?5 P 8?7 Cr 2?3 Al 2?0 Fe 68?7 C 7?0 Si 3?3 B 5?5 P 8?7 Cr 2?3 Al 2?0 Mo 2?5 Fe 66?7 C 7?0 Si 3?3 B 5?5 P 8?7 Cr 2?3 Al 2?0 Mo 4?5 Fe 64?7 C 7?0 Si 3?3 B 5?5 P 8?7 Cr 2?3 Al 2?0 Mo 6? *t max : maximum thickness or diameter; T g : glass transition temperature; T x : first crystallisation temperature; T l : liquidus temperature; DT x 5T x 2T g : width of the supercooled liquid region. The subscripts represent the atomic percentage of the component elements. Table 2 Typical Fe-based BMG alloy systems reported together with the calendar years when the first synthesis was reported Number Group Series Alloy Year of report Critical Key alloying of first synthesis diameter/mm additions Remarks Reference F1 1 1 Fe (Al,Ga) (P,C,B) Al, Ga First synthesis 55 F2 1 1 Fe (Cr,Mo) (Al,Ga) 1996 ETMzAl, Ga Extension of F1 125 (P,C,B) F3 1 1 Fe (Co,Cr,Mo) (Ga,Sb) (P,C,B) ETMzGa Largest in series F4 1 1 Fe Ga (P,C,B,Si) ETM- and Al-free F1 for applications 126 F5 2 2 Fe (Zr,Hf,Nb,Ta) B ETM Origin of series F6 2 2 Fe Ln B 1999 Ln Extension of F5 128 F7 2 2 Fe (Cr,Mo) (C,B) 2001 CrzMo Improved corrosion 105 resistance F8 2 2 Fe (B,Si) Nb Nb from ETM Bs.1. 5 T 129 F9 2 2 Fe (Si,B,P) ETM-free Js T 130 F Fe Mn Cr Mo C B Mn, Mo, C Origin of series 3 91 F Fe Cr (Ln,Y) Mo C B Y and Ln 99 F Fe-(Co,Cr,Mo)-(C,B)-Y ETMzY Largest in series F Fe (Cr,Mo) (C,B)-Tm Tm Improved corrosion 98 resistance International Materials Reviews 2013 VOL 58 NO 3 139

11 made based on a limited amount of data from a few alloy compositions and therefore, these do not seem to be universally applicable. The topic of GFA has been covered in detail earlier 6,7 and therefore we will not go into many details. But for the sake of completeness and later use, we will briefly summarise the different criteria, especially with reference to BMGs. The GFA of alloys is determined both by structural and kinetic parameters. 18 Here, structural criteria deal with the geometrical arrangement of atoms, bonding and atomic size effects to predict glass formation. On the other hand, the kinetic criterion considers the rate of cooling relative to the kinetics of crystallisation. While both the structural (thermodynamic) and kinetic factors assume importance, it appears that the basic GFA is mostly determined by thermodynamic parameters. This is where structural parameters such as the atomic size and chemical interactions between atoms are important. (It is not, however, implied here that BMGs, and much less the thin-ribbon metallic glasses, are thermodynamically stable). Once these criteria are satisfied, then the actual formation of the BMGs is determined by the kinetic parameters. In the early years of research on metallic glasses produced by RSP methods, a large number of empirical criteria were proposed to explain the GFA of alloys. Majority of these also appear to be applicable to glass formation in BMG alloy compositions. A very generic and simple criterion proposed was that a glassy phase is obtained only when the liquid is undercooled (e.g. by quickly solidifying above a critical solidification rate, which is dependent on the alloy system and its composition), to a temperature below the glass transition temperature T g. However, since this is an experimentally determined value (and is time-consuming and difficult to measure) and its estimation is also an involved process, other simple and empirical predictive criteria have been proposed to explain glass formation in alloy systems and these are briefly described below. (i) Turnbull 131 had suggested, purely on the basis of kinetics of crystal nucleation and the viscosity of melts, that the reduced glass transition temperature T rg, defined as the ratio of T g to the liquidus temperature of the alloy T l, should be a good indicator of the GFA of the alloy. The higher the T rg value, the higher the viscosity of the melt and therefore the alloy melt could be easily solidified into the glassy state (ii) it was also suggested that deep eutectic alloy compositions (a deep eutectic is one, wherein the alloy has a eutectic temperature that is much lower than the melting points of the individual metals) are good glass formers. This is because the value of T g changes slowly and the value of T l decreases very rapidly with solute content as one approaches the eutectic composition. Accordingly, the T rg value is very high at the deep eutectic composition (iii) Egami and Waseda 132 had suggested that one of the possible ways by which a crystalline metallic material can become glassy is by the introduction of lattice strain. The lattice strain introduced disturbs the crystal lattice and once a critical strain is exceeded, the crystal becomes destabilised and glassy. In fact, Egami 133 took pains to state that In general, alloying makes glass formation easier, not because alloying stabilises a glass, but because it destabilises a crystal. Using the atomic scale elasticity theory, these authors 132 calculated the atomic level stresses in the solid solution (the solute atoms are assumed to occupy substitutional lattice sites in the solid solution) and the glassy phase. They observed that in a glass, neither the local stress fluctuations nor the total strain energy vary much with solute concentration, when normalised with respect to the elastic moduli. But in a solid solution, the strain energy was observed to increase continuously and linearly with solute content. Thus, beyond a critical solute concentration, the glassy alloy becomes energetically more favourable than the corresponding crystalline lattice. That is, a minimum solute concentration was necessary in a binary alloy system to obtain the stable glassy phase. And, this minimum solute concentration CB min was found to be inversely correlated with the atomic volume mismatch (V A 2V B )/V A, where V A and V B are the atomic volumes of the solvent and solute respectively. The minimum solute concentration required can be obtained from the relation V A {V B Cmin B ~l~0: 1 (1) V A Yan et al. 134 and Ueno and Waseda 135 had used this concept to calculate the l value in the case of multi-component alloy systems. The main contribution of Yan et al. 134 was that they had used the mathematical description of regular polytopes and the dense random packing of hard spheres, a cluster model to describe the structure of metallic glasses, as the basis to reach this value. They suggested that the glass structure would have an optimum defect concentration when l<0?18. The model based on the atomic size mismatch was subsequently developed by Miracle and co-workers The major difference in the models of Egami and Miracle was that while the Egami model assumed that all the solute atoms, irrespective of their size, occupy the substitutional positions in the host lattice, the Miracle model considered that the smaller atoms occupy the interstitial positions while atoms with sizes larger or comparable with those of the solvent atoms occupy the substitutional positions. Furthermore, Miracle proposed that all the solute atoms could be classified into essentially three groups of atoms, which occupy the sites designated as a, b and c by Miracle. 136 Such a designation and grouping of atoms was also found to be helpful in identifying the chemical compositions of alloy systems that led to easy glass formation. Inoue criteria Based on the extensive data generated on the synthesis of BMGs for over a decade, Inoue had formulated three basic empirical rules to predict formation of BMGs. These may be stated as: (i) the alloy must contain at least three components. The formation of glass becomes easier with increasing number of components in the alloy system (ii) a significant atomic size difference should exist among the constituent elements in the alloy. It is suggested that the atomic size differences should 140 International Materials Reviews 2013 VOL 58 NO 3

12 be above y12% among the main constituent elements (iii) there should be negative heat of mixing among the (major) constituent elements in the alloy system. These rules have been of immense value in identifying alloy compositions for the synthesis of BMGs, even though some apparent exceptions have been found for these empirical rules. For example, BMGs have been reported to form in binary alloy systems such as Ca Al, 38 Cu Hf, 143 Cu Zr, 144 Ni Nb 145 and Pd Si. 146 But the section thickness of these glasses is usually small, i.e. typically a maximum of only about 1 2 mm. It may, however, be noted that with the addition of more alloying elements, the GFA of these alloys can improve significantly, resulting in the formation of glasses with larger section thicknesses. With increasing number of alloy compositions that have been made glassy, it was noted that some of the alloy systems did not follow the above-mentioned Inoue criteria. Therefore, in addition to the above general criteria that have been widely accepted by researchers and mostly followed in a large number of alloy systems, a number of new criteria have been developed since 2003, and these were based on the thermal properties of the alloys and physical characteristics of the component atoms These include the so-called a, b, c, c m, d, w, etc. parameters. Table 3 summarises the different criteria developed to rationalise and predict the GFA of alloys. In spite of this large number of parameters, the predictabilities of glass formation have not significantly improved 158 and it has been difficult to exactly specify which alloy compositions would produce the BMG alloy phases. Furthermore, the applicability of these new criteria appears to be limited in most cases to a few typical alloy systems. That is, majority of these new criteria are followed in a limited number of systems, and not more universally applicable, with the caveat that the c parameter proposed by Lu and Liu 149,150 seems to be Table 3 Summary of the quantitative criteria proposed to evaluate the GFA of liquid alloys Criterion/parameter Equation Reference Reduced glass transition T temperature rg ~ T g 131 T l DT x parameter DT x 5T x 2T g 21 a parameter a~ T x 147 T l b parameter New b parameter c parameter c m parameter d parameter K gl parameter w parameter T rx parameter v parameter b~1z T x T l ~1za 147 b~ T x T g 148 ðt l {T x Þ 2 c~ T x 149, 150 T g zt l c m ~ 2T x{t g T l 151 d~ T x T l {T g 152 K gl ~ T x{t g 153 T m {T x DT 0:143 x 154 w~t rg T g T rx ~ T x 155 T s v~ T g { 2T g 156, 157 T x T g zt l applicable in a majority of the alloy systems. The interested reader is advised to refer to a recent critical analysis of the results by Suryanarayana et al. 158 and to the book on BMGs 7 for full details. The three empirical rules proposed by Inoue belong to the structural criteria, and the kinetic criterion involving nucleation and growth controls will be discussed later. In developing and analysing BMG alloys, experimentalists have been using mostly the structural criteria (including the three above empirical rules) rather than the kinetic ones on account of their simplicity and wide applicability to actual alloy systems. Also note that some advanced investigations have been conducted to evaluate the GFA and stability of glassy phases based on statistical- and numerical approaches based on structural criteria In these approaches, the second and third factors of the Inoue empirical rules mentioned above or other relevant ones are qualified with corresponding thermodynamic and/or physical quantities such as mixing enthalpy and electronegativity for the second factor. Specifically, the effects of atomic bonding nature and atomic size mismatch were evaluated using factors of mixing enthalpy and mismatch entropy for ternary amorphous alloys, 159 mixing enthalpy and atomic size differences, 160 composition dependence of mixing enthalpy as a sub-regular solution model, 161 and electronegativity and atomic size for a number of bulk glassy alloy systems. 162 The results suggest the usefulness and importance of such approaches for further development of BMG alloys in general, and Fe-based bulk glassy alloys in particular, in the near future. Glass formation by solid-state processing methods All the criteria described above have been developed to describe the GFA of metallic glasses (and BMGs) processed through the solidification route. Mechanical alloying (MA) is another important non-equilibrium processing method to produce metastable phases in general and amorphous alloys, in particular. Mechanical alloying is a completely solid-state powder processing method that involves repeated cold welding, fracturing and rewelding of powder particles in a high-energy ball mill. The process involves placing the powders and a grinding medium (usually stainless steel balls) in a mill and agitating the whole mixture at a high speed. The repeated mechanical impacts of the grinding medium on the powder particles flatten and cold weld them. Because of this, the powder particles get strain hardened leading to their fracture, on continued milling, resulting in the creation of fresh surfaces. These fresh and active surfaces will again get cold welded when brought together by mechanical impacts. Owing to these repeated processes occurring, the powder particles form a layered structure and the interlamellar distance decreases with milling time. Furthermore, the introduction of crystal defects (dislocations, grain boundaries, stacking faults, etc.) enhances diffusion, further aided by a slight rise in the powder temperature. All these factors lead to alloying between the powder particles and formation of both equilibrium and metastable phases A very large number of metallic glasses have been produced in a variety of alloy systems and in different compositions using MA. 163,164 However, there have not been many systematic investigations conducted to study International Materials Reviews 2013 VOL 58 NO 3 141

13 1 a XRD patterns of blended elemental powder mix of Fe 42 Al 28 Zr 10 B 20 as a function of milling time. Note that the amorphous phase has started to form on milling for y10 h and that the amorphous phase was stable up to y40 h. 168 b XRD patterns of Fe 42 Co 28 Zr 10 B 20 powder mix as a function of milling time. Note that an amorphous phase had not formed in this case; instead only a solid solution phase was obtained on milling for 30 h (Ref. 168) the conditions under which amorphous phases are formed by this method. It would be useful and instructive to see if the criteria applicable to the solidification methods would also be applicable to the solid-state processed amorphous alloys or other criteria need to be formulated to predict glass formation in alloy systems processed by MA. A systematic and comprehensive investigation has recently been reported on the glass formation behaviour and stability of several Febased glassy alloys processed by MA. 168 Based on a systematic and comprehensive investigation on Fe-based alloys of the generic composition of Fe 42 X 28 Zr 10 B 20 (where the subscripts represent the composition of the alloy in atomic percentage and X5Al, Co, Ge, Mn, Ni or Sn), it was noted that amorphisation had occurred only in some alloy systems and not in all. For example, amorphisation occurred only in alloy systems with X5Al, Ge or Ni, as evidenced by the presence of a broad diffuse peak centred at the (110) Fe position in the X-ray diffraction (XRD) patterns (Fig. 1a). Amorphisation did not occur in alloy systems with X5Co, Mn or Sn (Fig. 1b). The time required for amorphisation, which was considered a measure of GFA was also different for the different powder blends. Table 4 presents the results obtained, including the equilibrium number of intermetallic phases present between X and the constituent elements (Zr, Fe or B) in the powder bend. 168,169 A close examination of Table 4 clearly reveals that the ease of amorphisation (i.e. GFA) increased with the number of intermetallics present in the constituent Zr X binary phase diagrams. This is apparent from the powder blends containing Al or Ni, which amorphised in 10 or 20 h respectively. While the quaternary Fe Zr Al B contains eight intermetallic phases in the binary system between Zr and Al, the Fe Zr Ni B system contains seven intermetallic phases in the binary system between Zr and Ni. Similarly, the Ge-containing system, which also amorphised in 10 h, contains five intermetallic phases between Zr and Ge. The Zr Co, Zr Mn and Zr Sn systems which did not show amorphisation on milling contained five, one and three intermetallic phases respectively. However, when the total number of intermetallic phases was considered, it can be clearly seen that the systems which amorphised on milling contained >10 intermetallic phases in all the constituent binary phase diagrams. That is, if the alloy system contained,10 intermetallic phases, then amorphisation was not observed. Note, however, that when the total number of intermetallics was only 10 (with the Ge-containing alloy), the time required for amorphisation was only 10 h. But this is a special case because Ge is a semi-metal with strong directional bonds. Thus, it becomes easier to amorphise alloys containing Ge (or other semi metals). From a critical analysis of the constituent binary phase diagrams, it also becomes clear that when the phase diagrams contain extensive solid solutions, it will be very difficult to amorphise them. 169 Table 4 Summary of the results of amorphisation in the Fe 42 X 28 Zr 10 B 20 (where X5Al, Co, Ge, Mn, Ni or Sn) systems* X Number of intermetallics between X and other elements X and Zr X and Fe X and B Total number of intermetallics Milling time required for amorphisation/h Al Co 5 Nil 3 8 No amorphisation Ge 5 5 Nil Mn 1 Nil 5 6 No amorphisation Ni Sn 3 2 Nil 5 No amorphisation *The number of intermetallics in the constituent binary alloy systems with X (i.e. Zr X, Fe X and B X) is also listed International Materials Reviews 2013 VOL 58 NO 3

14 2 a Schematic illustration of local atomic structure models for Fe-based BMG alloys in Fe RE B (RE5rare earth metals) and Fe TM B (TM5Zr, Nb or Mo) systems determined by advanced structural analytical methods. b A preliminary structural model giving the pseudo-tenfold diffraction pattern in which the three types of tiles as structural motifs of Fe 23 B 6 are linked to each other 182 Amorphisation by MA was reported to occur when the free energy of the crystalline phase G C is higher than that of the hypothetical amorphous phase G A,i.e.G C.G A. 170 A crystalline phase normally has a lower free energy than the amorphous phase. But its free energy can be increased by introducing a variety of crystal defects. If an intermetallic has formed, then additional energy can be introduced by disordering the crystal lattice. By this approach, it is then possible to obtain an amorphous phase when G C zg D wg A (2) where G C represents the free energy of the crystalline phase, G D represents the increase in free energy due to introduction of defects and G A represents the free energy of the amorphous phase. The magnitude of energy increase is different for different types of defects. As an example, increasing the dislocation density to m 22 increases the free energy by y1 kj mol 21, while decreasing the grain size down to 1 nm increases the free energy by y10 kj mol The only way a solid solution could contribute to an increase in the energy of the system is by grain refinement. But this increase in energy is not sufficiently high to amorphise the system. On the other hand, the presence of intermetallics in an alloy system can contribute to an increase in the energy to favour amorphisation. This is due to two important effects. First, disordering of intermetallics contributes an energy of y15 kj mol 21 of atoms to the system. For example, in strongly ordered intermetallics such as NiAl and c-tial that continue to be in the ordered state till melting, the disordering energy has been estimated to be y17?5 kj mol Second, a slight change in the stoichiometry of the intermetallic phase increases the free energy of the system drastically. Additionally, grain size reduction contributes an energy of y5 kj mol 21. Furthermore, disordering of intermetallics has also been shown to be possible by heavy deformation. 173 Since MA reduces the grain size to nanometre levels and also disorders the usually ordered intermetallics, the energy of the milled powders is significantly raised. In fact, in most cases, it is raised to a level above that of the hypothetical amorphous phase, leading to the preferential formation of the amorphous phase over the crystalline phase. Phase diagram features have been used to predict glass formation by RSP and other methods as well. As mentioned above, alloys in the vicinity of deep eutectics exhibit high T rg values, and therefore they exhibit high GFA. 131 Furthermore, elemental solids exhibiting a large number of polymorphic phases have been shown to exhibit higher GFA than elements that do not have a large number of polymorphs. 174 Difficulty in amorphisation has been observed with RSP in alloys having melting maxima, phase diagrams featuring too many peritectic reactions, high-temperature eutectics and also alloys having positive heats of mixing. However, amorphisation has been observed in most of these cases by MA. 163,164 Thus, even though phase diagrams are useful guidelines in choosing alloy compositions for easy glass formation by both the methods, the features to look for appear to be quite different for the RSP and MA methods. It is just fortuitous that some alloy compositions can be amorphised by both the methods. Structure of the glassy phase It was mentioned earlier that the Fe-based BMG alloys were composed of multi-component alloy systems which satisfy the three-component rules. The requirement for the three-component rule has been interpreted to originate from the formation, in the supercooled liquid and glassy structure, of unique distorted trigonal prisms and anti-archimedean prisms consisting mainly of Fe and B or C. These prisms are connected to each other in edge- and face-shared configuration modes through glue atoms of M (M5Ln, Zr, Nb, Mo and Ga) elements. 175,176 As an example, Fig. 2a shows a schematic illustration of the local atomic configurations in Fe RE B (RE5rareearth metals) and Fe TM B (TM5transition metals such as Zr, Nb or Mo) glassy alloys derived from the experimental data using the anomalous XRD, pulsed neutron diffraction, high-resolution TEM and reverse Monte Carlo computer simulation techniques. Although the maximum size of the structure model was limited to 2 4 nm, due to the lack of periodic atomic configurations, this size corresponds to the upper limit which can be determined by the high density of synchrotron X-ray and the pulsed neutron beam sources. The resulting networklike long-range atomic configurations with attractive bonding nature can effectively suppress the long-range atomic rearrangements of the constituent elements that are necessary for the progress of crystallisation. This can consequently lead to the formation of bulk glassy alloys through stabilisation of the supercooled liquid. International Materials Reviews 2013 VOL 58 NO 3 143

15 3 X-ray diffraction patterns of [(Fe 0?8 Co 0?1 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 bulk glassy alloys subjected to annealing treatments leading to precipitation of the primary crystalline phase 112 It has been believed for quite some time now that the structure of the liquid, from which the BMG phases form, determines the structure of the glass. Since the liquid contains icosahedral-type clusters, it is also believed that the structure of the glass also contains icosahedral -type units. 177 In fact, several authors have reported that the first phase to precipitate out of the Zrbased BMG alloys during crystallisation is a quasicrystalline phase Recent high-resolution TEM investigations have clearly indicated the presence of medium-range order (MRO) in Fe-based glassy alloys. 181,182 These authors have reported pseudo-tenfold diffraction patterns during the course of crystallisation of Fe-based BMG alloys. For example, during the course of crystallisation of an Fe 48 Cr 15 Mo 14 C 15 B 6 Tm 2 BMG, the authors found a nanoscale metastable state with a x- FeCrMo-like structure, which changed into the [113] zone-axis pattern of the x-fecrmo structure on annealing at higher temperatures. 181 Similarly, nanoscale quasicrystal-like structural states exhibiting pseudo-tenfold nanobeam electron-diffraction patterns were observed in the course of crystallisation process of an (Fe 0?5 Co 0?5 ) 72 B 20 Si 4 Nb 4 BMG alloy. 182 In this case, the metastable structure transformed to the Fe 23 B 6 -like intermediate structure and eventually to the Fe 23 B 6 structure. Thus, these metastable states can be considered as the pre-crystallisation stage, similar to the preprecipitation stage in precipitation-hardenable alloys. 183 Existence of the intermediate states between the quasicrystal-like and Fe 23 B 6 structures indicates that the quasicrystal-like structure is an approximant to the Fe 23 B 6 structure. The pseudo-tenfold electron-diffraction patterns are understood to arise from a combination of the three types of tiles found in the Fe 23 B 6 structure. These three types of tiles can produce decagonal units which do not have any icosahedral atomic arrangement. Thus, the presence of icosahedral atomic arrangement is not a prerequisite to the formation of quasicrystal or quasicrystal-like phases. Figure 2b shows a preliminary structure model giving the pseudo-tenfold diffraction pattern, in which the three types of tiles as structural motifs of Fe 23 B 6 are linked to each other. 182 The larger and smaller circles in the model denote the Fe and B atoms respectively. The arrows in the model mean orientations of the structural motifs with five different directions, which contribute to give the isotropic diffraction pattern with 10 strong spots. This report provides the first diffraction evidence for the presence of icosahedral-like structures in an Fe-based bulk glassy alloy. The diffraction and spectroscopic methods generally provide only the average structural information. Even though a number of different atomic models have been proposed to explain the structure of metallic glasses, direct observation of the local atomic structure in disordered materials has not been achieved till recently. Hirata et al. 184 have recently reported direct observation of local atomic configurations in a binary Zr 66?7 Ni 33?3 metallic glass obtained by melt spinning. By employing a state-of-the-art nanobeam electron diffraction technique combined with ab initio molecular dynamics simulation, the authors reported observation of subnanoscale ordered regions. Distinctly symmetric nanobeam electron diffraction patterns were found to originate from individual and interconnected atomic polyhedra. These observations offer compelling evidence of the local atomic order in the disordered metallic glass, which is consistent with the recent cluster models and previous predictions that metallic glasses possess short-range order and MRO as opposed to the long-range periodicity of a crystalline solid. 136,137,185 Furthermore, it is important to point out that the spontaneous formation of such unique long-range 144 International Materials Reviews 2013 VOL 58 NO 3

16 network-like atomic configurations in special multicomponent alloy systems significantly affects the nature of the primary precipitate phase forming from the supercooled liquid. Figure 3 shows the XRD patterns of the [(Fe 0?8 Co 0?1 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb 4 BMG alloy subjected to annealing at different temperatures ranging from 853 to 1133 K. 70,112 The primary precipitate phase can be identified as (Fe,Co) 23 B 6 with a complex fcc structure and a lattice parameter of y1?12 nm. The (Fe,Co) 23 B 6 phase is in a metastable state and changes to a mixture of equilibrium crystalline phases with further increase in the annealing temperature. Thus, the primary precipitate phase in all Fe-based bulk glassy alloys reported to date consists of a complex fcc M 23 (B,C) 6 compound with a large unit cell volume containing a large number of atoms. The local atomic configuration of the M 23 (B,C) 6 phase includes an anti-archimedean atomic configuration which is similar to that of the Fe Nb B glassy alloys. 175,176 It is because of the similarity between the local atomic configurations of the supercooled liquid and the primary precipitate phase that this is the first phase to form during crystallisation of the BMG alloys. It is thus interpreted that the spontaneous formation of the long-range network-like atomic configurations in the supercooled liquid in multi-component alloy systems with at least three components is the origin for the high stability of the supercooled liquid against crystallisation in Fe-based BMG alloys. In addition to the above-described experimental observations, computational approaches were also undertaken to obtain local atomic configurations in Febased BMG alloys. These approaches used the concepts of clusters and their packing leading to MRO in clusterpacked structures. For instance, Kazimirov et al. 186 carried out ab initio calculations on Fe TM RE Me BMG alloys (where TM5transition metals such as Mn, Cr and Mo, RE5rare-earth elements such as Er and Y and Me5metalloid atoms such as C and B). Some of the alloys studied included Fe 64 Mo 14 C 15 B 7, Fe 50 Cr 15 Mo 14 C 15 B 6 and Fe 51 Cr 14 Mo 12 C 15 B 6 Y 2. Since atomic diffusion determines the GFA and stability of the glass (slow atomic diffusion may lead to better GFA by reducing the critical cooling rate required to form the glass), the authors calculated the atomic mean-square displacements nr(t) 2 m. They showed that the nr(t) 2 m for C, B and Fe are different and that these atoms diffuse much faster in the Fe 79 C 15 B 6 alloy than in the Fe 49 Cr 15 Mo 14 C 15 B 6 Er 1 alloy, explaining why the latter alloy is a better glass former. It may be noted that the Fe 49 Cr 15 Mo 14 C 15 B 6 Er 1 alloy can be cast into 6 mm diameter rods 97 while Fe 79 C 15 B 6 can only be made in ribbon form and is not entirely glassy. The role of rareearth elements in improving the GFA of alloys has also been explained using these results. The slow diffusion of RE atoms has been explained as not just due to the large size of the atom but mostly because RE atoms form structurally complex clusters that slow diffusion down for other chemical species. Kazimirov et al. 186 had also reported the effects of short- and medium-range atomic clustering on the diffusion behaviour and reported that the short-range clusters persisted well into the liquid state. Using these structural models, the authors 186 calculated the bulk moduli as a function of the Er content in these glasses and reported that the bulk modulus decreased with increasing Er content. The bulk modulus values and their variation with composition were consistent with the experimentally determined values. 96 Alloy development studies A large number of Fe-based alloy compositions have been quenched into the glassy state in recent years. All the developments in Fe-based metallic glasses, including the BMGs, can be traced to the early years of research on the binary Fe B metallic glasses, formed in the form of a thin ribbon by RSP methods. Later developments with respect to alloy development and improvement of mechanical and magnetic properties were also based on the Fe B system. As mentioned earlier, all the Fe-based BMGs can be classified into three groups, and these can be typically represented by the systems: Fe M (P,C,B,Si) (M5Al, Ga, Mo); Fe B Si Nb based; and Fe Cr Mo C B Ln. Let us now look at these systems in some detail. Fe M (P,C,B,Si) (M5Al, Ga, Mo) BMGs from Fe (Al,Ga) metalloid type As mentioned earlier, the first synthesis of Fe-based BMGs by the copper-mould casting process was made in the Fe Al Ga P C B system in 1995, 27 even though the maximum diameter of the glassy alloy was,2 mm. The success of producing was important for significant extension of application fields of BMG alloys as structural and functional materials. This incentive played a trigger effect on the subsequent development of new BMG alloys in the alloy series: Fe (Cr,Mo) Al Ga P C B, Fe Mo Ga P C B, Fe Ga P C B Si, Fe Mo P C B Si, Fe Cr Mo P C B Si, Fe Co Ga P C B Si and Fe Co Mo P C B Si systems. 68,104,187 The maximum rod diameter was reported to be 2?5 mmfor the Fe Ga P C B Si system, 126 4mmfortheFe Mo P C B Si system 68 and 6 mm for Fe Co Mo P C B Si system. 187 Fe B Si Nb based bulk glassy alloys from the Fe ETM/Ln metalloid type Another type of Fe-based soft magnetic BMG alloy has been developed in the Fe Co B Si Nb system for the past few years. It was reported in 2002 that the addition of small amounts (2 4 at-%) of Nb to Fe Co B Si amorphous alloys caused a drastic change to the glassy type leading to the appearance of distinct glass transition, followed by a large supercooled liquid region before crystallisation. 107 As a result, the glass transition phenomenon was observed over the whole composition range in the [(Fe,Co,Ni) 0?75 B 0?20 Si 0?05 ] 96 Nb 4 system. 188 The highest T rg of 0?60 and the largest supercooled liquid region DT x of y65 K were obtained at the Fe Co-rich composition around [(Fe 0?6 Co 0?4 ) 0?75 B 0?20 Si 0?05 ] 96 Nb 4. The alloy with the highest T rg and the largest DT x had the highest GFA which enabled the production of bulk glassy alloy rods with diameters up to 5 mm, as shown in Fig It has also been shown that the use of the B 2 O 3 flux melting treatment causes a further increase in the maximum diameter to y9 mm, 189,190 indicating that elimination of heterogeneous nucleation sites through refinement of the alloy melt is effective to increase the GFA even for Fe-based BMG systems. International Materials Reviews 2013 VOL 58 NO 3 145

17 6 The critical diameter for the formation of a glassy phase as a function of Co content for (Fe 12x Co x ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 BMG alloys Compositional dependence of the maximum diameter obtained in [(Fe 12x2y Co x Ni y ) 0?75 B 0?20 Si 0?05 ] 96 Nb 4 BMG alloys produced by the copper-mould casting method 115 Fe Cr Mo C B Ln BMGs from nonferromagnetic Fe ETM/Ln metalloid type By altering the Fe content in Fe-based BMG alloys, the magnetic nature of the glassy phase changes from ferromagnetic to non-ferromagnetic type at room temperature. It is also noted that in these cases, the maximum diameter of the BMG alloy rod increases to high values reaching as much as 16 mm. 99,103,119,191 For instance, as a forerunner for the non-magnetic type Fe-based BMG alloys, it was reported that Fe 48 Cr 15 Mo 14 C 15 B 6 Er 2 is formable as a BMG alloy with a critical diameter of 12 mm, 99 exceeding 1 cm. Subsequently, the largest diameter BMG rod of 16 mm was produced in the Fe 41 Co 7 Cr 15 Mo 14 C 15 B 6 Y 2 alloy. 103 The fact that partial replacement of Fe by Co increases significantly the critical diameter has also been confirmed in the Fe 48 Cr 15 Mo 14 C 15 B 6 Tm 2 glassy alloy. 98 Figure 5 shows the XRD 5 X-ray diffraction patterns of cast (Fe 0?8 Co 0?2 ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 BMG alloy rods with diameters of 6 18 mm patterns of the cast (Fe 0?8 Co 0?2 ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 alloy rods with different diameters ranging from 6 to 18 mm. No distinct diffraction peak due to a crystalline phase is seen for the rod specimens with diameters,16 mm, indicating that the critical diameter of the glassy alloy was as large as 16 mm. Figure 6 shows the change in the critical diameter for the formation of a glassy phase in the cast (Fe 12x Co x ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 alloys as a function of the Co content. 191 The critical diameter is 12 mm for the Co-free alloy, increases to 16 mm for the alloys with x50?2 and 0?4 and then decreases to 10 mm with further increase in the Co content. It is noticed that the Co Cr Mo C B Tm alloy also has a large critical diameter of 10 mm. 191 The effectiveness in increasing the GFA of alloys by partial replacement of Fe with Co seems to result from the necessity of the long-range atomic rearrangement of the constituent elements required for the simultaneous precipitation of Fe 23 (C,B) 6 and Co 23 (C,B) 6 phases from the supercooled liquid for the Fe Co C B Tm alloys. It has also been found that the critical diameter increases further to 18 mm by more fine-tuning of the composition of the alloy to (Fe 0?8 Co 0?2 ) 47 Cr 15 Mo 14 C 15 B 6 Tm Considering that the critical diameter of Fe 50 Cr 15 Mo 14 C 15 B 6 is y2 mm, addition of 2 3 at-%tm is concluded to be very effective in increasing the GFA. Furthermore, the addition of 2%Tm does not cause distinct change in the melting temperature. Therefore, the significant effectiveness for glass formation is presumably due to the increase in the ease of formation of the network-like long-range atomic configurations due to the coexistence of two kinds of glue atoms (Mo and Tm). This may be in addition to the atomic size effect since Tm has the largest atomic size as compared with all other constituent elements. Role of alloying elements on GFA The synthesis of the Fe 80 B 20 glass (commercially referred to as METGLAS alloy no. 2605), with a high strength of y3630 MPa (370 kgf mm 22 ), triggered the further development in metal metalloid type of Fe-based metallic glasses by maintaining a standard composition of y80 at-% of the transition metal(s) and y20 at-% of the metalloid element(s). All subsequent investigations were concerned with alloy modifications in terms of changing the nature and relative proportions of the 146 International Materials Reviews 2013 VOL 58 NO 3

18 transition metal and/or metalloid elements. The goal has always been to increase the GFA (to synthesise larger diameter BMG alloy rods) and improve the mechanical (simultaneous increase in strength and ductility) and magnetic properties (higher saturation magnetisation, high permeability and low magnetostriction). Recent results on increasing the GFA will be briefly described below. In a series of developments of Fe-based BMG alloys, it is widely accepted that addition of small amounts of some specific alloying elements greatly affects the GFA. For instance, it was reported 193 that 6 at-% addition of Y as well as Sc, Dy, Ho and Er in ternary Fe 72 M 6 B 22 alloys made it possible to produce BMG alloys with 1 mm in diameter. On the other hand, other elements such as Zr, Hf, Nb, Ta and lanthanide elements ranging from La to Tb did not contribute to the formation of BMG alloys. Here, special attention should be paid to the recent report, 92 investigating the role of Nb, Zr and NbzZr on GFA of Fe Nb Zr B alloy system. The authors found that the optimal glass formers were located at Fe 71 Nb 6 B 23 and Fe 77 Zr 4 B 19 in the ternary systems, and Fe 71 Nb 4?8 Zr 1?2 B 23 in the quaternary system. Accordingly, the authors were able to produce 1?5 mm diameter rod in the Fe Nb B system, 1 mm diameter rod in the Fe Zr B system and 2 mm diameter rod in the Fe Nb Zr B system. This report suggests that Nb and Y are effective alloying elements to improve GFA, supporting the recent trends to add Y and Nb together in Fe-based BMG alloys. 194 A similar early study discussing the effect of Y on the enhancement of GFA of Fe-based alloys was reported by Lu et al. 57,150 in which they reported that addition of Y to Fe Cr Co Mn Mo C B alloy enhances the GFA. It was mentioned that while the critical diameter for glass formation was only,7 mm for the Y-free alloy, it increased to >12 mm for the alloy containing 1?5 at-%y. 150 Similarly, while the Tm-free Fe 50 Cr 15 Mo 14 C 15 B 6 alloy could be cast into rods of,2 mm in diameter, the alloy containing 2 at-%tm could be cast into rods of 12 mm diameter. 98 Replacement of Fe with Co increased the critical diameter to 16 mm. 191 These reports support that the effects of Y and Tm on enhancing the GFA are, in part, due to the large atomic size differences between the principal elements among Fe, Y (or Tm) and B. In a framework for the development of Fe-based BMG alloys containing Y, it was also reported 194 that replacement of Fe with Zr or Co significantly affected the GFA and the soft magnetic properties. The effects of addition of minor amounts of solute elements in Fe-based and other glassy alloys were more widely and thoroughly reviewed. 195,196 The principal aspects of the formation and properties of Fe-based BMG alloys have been highlighted in these reviews as follows: (i) by the addition of some metals with high melting temperature, such as Zr, Nb, Ta, W and Mo, with 5 mm diameter can be obtained by Cu-mould casting 197 (ii) with the proper rare-earth element additions, the Fe- as well as Mg-based BMGs can be successfully fabricated by a conventional Cumould casting method even in air atmosphere (iii) a minor Ni addition can significantly enhance the soft magnetic properties of Fe-based glassforming alloys without deteriorating their high GFA (iv) the addition of metalloid elements usually makes the more brittle (v) enhancement of GFA by the additions of C, B and Y, each at levels of,5 at-%, in Fe-based BMG alloys, has been shown to occur in several instances. 195 Phase stability and evolution of glassy alloys associated with crystal nucleation and growth of undercooled alloys have been reviewed by Perepezko, 198 even though this was with special reference to Al-based glassy alloys. This approach involving nucleation and growth controls belongs to the kinetic criteria for formation of glassy alloys towards the structure criteria mentioned earlier. Besides, in a framework of phase evolution of Fe-based BMG alloys, Nouri et al. 199 investigated the effects of changes in test temperature on the microhardness/ strength of Fe 48 Mo 14 Cr 15 Y 2 C 15 B 6 BMG alloys by paying attention to structure evolution under a variety of different test conditions over the temperature range from room temperature to 620uC. Although a very high microhardness value (e.g..12 GPa) was exhibited at room temperature, significant hardness reductions were exhibited near T g. In addition, the effect of exposure time (up to 5 h) at elevated temperatures on the evolution of microhardness/strength was also evaluated for selected temperatures. Such results are of great value when using Fe-based BMG alloys as structural materials, in particular, at elevated temperatures ranging from room to glass transition temperatures. Mechanical properties Metallic glasses, and more specifically the BMGs, have been known to exhibit very high strength and hardness and wear resistance in comparison with their crystalline counterparts. Therefore, in view of their low cost, Febased BMGs can be exploited for structural applications. Associated with the high strength is the drawback that the Fe-based metallic glasses have little or no plasticity. However, materials used for structural applications are required to have some plasticity and therefore recent research efforts have been directed to develop BMG alloys that possess simultaneously both high strength and some amount of plasticity. The plasticity of the BMGs can be evaluated by measuring the amount of plastic deformation that the material undergoes beyond the elastic limit. The mechanical properties of BMGs have been recently reviewed As mentioned earlier, majority of the developments in Fe-based BMG alloys have been based on the early years of research on binary Fe B metallic glasses, formed in the form of thin ribbons by RSP methods. A special attribute of the Fe 80 B 20 glass is its high strength, and it was reported 203 that among the metallic glasses, Fe 80 B 20 glass exhibits the highest strength of y3630 MPa (370 kgf mm 22 ), which is superior to another type of Fe-based metallic glass (Fe 80 P 16 C 3 B 1 METGLAS alloy no. 2615) with a yield strength of 2440 MPa (249 kgf mm 22 ) and to other Ni-based metallic glasses commercialised as METGLAS series. The Fe 80 B 20 glass was known as the strongest metallic glass found in the middle of 1970s, and this situation continued till the beginning of the twenty-first century until the development of a Co 43 Fe 20 Ta 5?5 B 31?5 BMG alloy in 2004 with a yield strength of.5000 MPa. 204 Thus, the history of the development of International Materials Reviews 2013 VOL 58 NO 3 147

19 has been very closely linked with the high mechanical strength of Fe-based metallic glasses. Table 5 lists the mechanical properties of some typical Fe-based BMG alloys. It may be noted from this table that the yield strength of the is very high. The research on mechanical properties of Fe-based BMGs has been a major topic of research during the last decade or so. An important aspect of this research has been on improving the plasticity of these alloys, since Fe-based glassy alloys that belong to the metal metalloid type are intrinsically brittle in nature due to the covalent-like bonding between metal metalloid atomic pairs. These efforts to improve the plasticity of Fe-based BMG alloys, and also BMG alloys in general, are aided by alloy design principles through modification of alloy compositions to alter their fracture behaviour. Some of these studies are presented below with the aim of identifying their essential features. The studies can be grouped under three categories, namely alloying additions, compositional modifications to modify the shear and elastic moduli and the Poisson s ratio, and development of composites. However, it should be realised that the first two aspects are intimately related to each other. It has been reported 206 that the chemistry of the alloy basically determines its ductility and fracture toughness. The Fe Mn Mo Cr C B system was shown to exhibit very low toughness, approaching values typical of the inherently brittle oxide glasses (0?06 kj m 22 ). But by replacing Mn with Er and increasing the Cr content from 4 to 15 at-%, the fracture toughness was increased significantly (from y5 to 26 MPa m 1/2 ). Similarly, in the Fe Mo C B system, the addition of Cr and Er increased the shear and elastic modulus and reduced the compressive plasticity. 95 However, replacing Er with P improved the compressive plasticity. Similarly, substitution of carbon for boron in Fe 49 Cr 15 Mo 14 C 13 B 8 Er 1 and (Fe 0?9 Co 0?1 ) 58?5 Cr 6 Mo 14 C 15 B 6 Er 0?5 steels resulted in a decrease in the shear modulus, while the bulk modulus remained essentially constant. These modified alloys exhibited high fracture strengths and some compressive plasticity. 97 The relationship between ductility and alloy design was provided by Gu et al., 109 who reported that the ductility of Fe-based BMG alloys can be improved by partial replacement of the elements that enhance the nature of ionic and covalent bonds. The analysis based on first-principles electronic structure calculations was also performed for ductile Fe Cr Mo P C B amorphous steels. Furthermore, it was pointed out that the enhanced ductility of amorphous steels was attributed to the decrease in the shear modulus in the ductile region, where the Poisson s ratio is 0?33 0?34. The relationship between the mechanical properties and relevant physical quantities, such as shear modulus and Poisson s ratio, has also been clarified empirically by Gu et al., 96 suggesting that the onset of plasticity in the Fe 65 Mo 14 C 15 B 6 BMG alloy doped with lanthanides was associated with an increase in the Poisson s ratio. In fact, Chen et al. 207 had earlier pointed out way back in 1975 that it is the high Poisson s ratio (n) which is responsible for the ductile behaviour of many metallic glasses. The decreasing n with falling temperature, together with a relatively lower n (,0?40) results in a rapid increase in the fracture strength and the brittle behaviour of Fe-based glasses. Recent work 208 has further indicated that the levels of compressive plasticity may be affected by various test details (e.g. alignment, stress concentrations, etc.) and that fracture toughness/energy may be a more discriminating test 205,209 as shown below. It has been suggested that the ratio of the elastic shear modulus to the bulk modulus G/K can be utilised to predict the ductile or brittle behaviour of solids. For instance, the fracture energy (toughness) of Fe-based BMG alloys was correlated with changes in the G/K ratio as well as the Poisson s ratio by Lewandowski et al. 205 Through a broad survey of the literature, it was shown 209,210 that the critical value of G/K differentiating a brittle BMG from a ductile BMG appears to be in the range of 0?41 0?43. Bulk metallic glass alloys with a G/K value of over 0?41 0?43 are brittle. For example, the Fe 50 Mn 10 Mo 14 Cr 4 C 16 B 6 BMG alloy with the G/K ratio of 0?423 has been shown to be brittle, when compared with the Zr-based BMG alloys. A recent mechanics-based Table 5 Mechanical properties of some typical Fe-based BMG alloys* Alloy composition s y /GPa s f /GPa e pl. /% G/GPa K/GPa E/GPa n K c /MPa m 1/2 Reference Fe 65 Mo 14 C 15 B , 205 (Fe 0?9 Co 0?1 ) 64?5 Mo 14 C 15 B 6 Er 0? , 205 Fe 59 Cr 6 Mo 14 C 15 B , 205 Fe 49 Cr 15 Mo 14 C 13 B 8 Er Fe 49 Cr 15 Mo 14 C 15 B 6 Er , 205 Fe 49 Cr 15 Mo 14 C 18 B 3 Er Fe 49 Cr 15 Mo 14 C 19 B 2 Er (Fe 0?9 Co 0?1 ) 58?5 Cr 6 Mo 14 C 15 B 6 Er 0? , 205 Fe 71 Mo 5 P 12 C 10 B Fe 66 Mo 10 P 12 C 10 B Fe 65 Cr 2 Mo 9 P 10 C 8 B Fe 63 Cr 3 Mo 10 P 12 C 10 B Fe 61 Mn 10 Cr 4 Mo 6 C 15 B 6 Er Fe 53 Cr 15 Mo 14 C 15 B 6 Er (Fe 0?75 B 0?2 Si 0?05 ) 96 Nb [(Fe 0?5 Co 0?5 ) 0?75 B 0?2 Si 0?05 ] 96 Nb [(Fe 0?8 Ni 0?2 ) 0?75 B 0?2 Si 0?05 ] 96 Nb [(Fe 0?5 Ni 0?5 ) 0?75 B 0?2 Si 0?05 ] 96 Nb [(Fe 0?6 Co 0?3 Ni 0?1 ) 0?75 B 0?2 Si 0?05 ] 96 Nb *s y : yield strength; s f : fracture strength; e pl. : plastic strain; G: shear modulus; K: bulk modulus; E: Young s modulus; n: Poisson s ratio; K c : notch toughness. 148 International Materials Reviews 2013 VOL 58 NO 3

20 model also predicts the importance of elastic constant ratio on the toughness of metallic glasses. 211 The ductility of conventional crystalline metals has been shown to increase by the introduction of thin metallic glassy ribbons into them. 212 Therefore, researchers have utilised this concept to increase the ductility of BMG alloys and have achieved some amount of success But majority of these investigations have been concerned mostly with Zr-based BMG alloys and others. There have been very few investigations on Fe-based BMG alloys. For example, Shen et al. 117 produced a composite alloy by adding 0?25 at-%cu to an (Fe 0?5 Co 0?5 ) 75 B 20 Si 5 BMG alloy. The microstructure consisted of y13?6 vol.-% of precipitates of a-(fe,co) and (Fe,Co) 23 B 6 in a glassy matrix. While the fully glassy alloy (without any Cu in it) had a yield strength of 3700 MPa with no ductility, the composite alloy (with 0?25 at-%cu in it) exhibited a fracture strength of 4500 MPa and also a plastic strain of 0?6%. Thus, it appears worthwhile to further explore this approach to increase the ductility of Fe-based BMGs. Similarly, the cast glassy rods in the alloy series of (Fe 12x Co x ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 exhibit high fracture strength of.4000 MPa over the whole Co content range and the strength level of MPa, which is almost independent of the Co content. 191 In addition, all the BMG alloys with a diameter of 2 mm did not fracture within the elastic elongation limit and exhibited nearly a constant fracture strength as well as the same elastic elongation limit of 2%. This strength value and the fracture behaviour indicate that these Fe-based BMG alloys also possess a relatively good ductile nature. The important term in the literature on mechanical properties, and more specifically with reference to ductility, is intrinsic, and this term, along with the other term extrinsic, suggests new ways of approaching the field of BMG alloys. For instance, Yavari et al. 216 pointed out that the current intense interest in the mechanical response of glassy alloys is due to intrinsic and extrinsic factors, and that these factors explain remarkably well the extensive plastic deformation during compression or bending, serrations in the stress strain curves, shear softening, sharp temperature rise around shear bands and resultant growth of nanocrystals that block the propagation of shear bands. Besides, Weibull modulus of Fe-based BMG alloys was investigated by paying attention to the intrinsic and extrinsic effects. 206 It was reported that extrinsic factors such as the presence of processing defects including inclusions or porosity are responsible for the large scatter in the toughness of some samples. In addition to toughness, a number of research papers have also dealt with the basic concepts of brittleness and plasticity. For example, Xi et al. 217 established a clear correlation between the fracture toughness and the length scale of the plastic process zone for various brittle and tough BMGs. They suggested that the fracture surfaces in brittle BMGs (e.g. those based on Fe and Mg) also revealed dimple structures, suggesting that the existence of damage microvoids does not depend on the chemical composition of the BMGs and that the nucleation of microvoids is related to the glassy structure, which contains free volume and inherent atomic density fluctuations at the nanometre level. 23 But the dimple structures in the relatively brittle BMGs is on the nanometre level, indicative of the activation of plastic flow processes, possibly as a result of the local softening mechanism. In fact, Xi et al. 217 have demonstrated a linear relationship between the square of the ratio of fracture toughness to yield strength and the measured plastic zone size; the more ductile BMGs showed a larger measured plastic zone. Varying degrees of compressive plasticity were also recently observed in Ti-based BMGs by changing the sample sizes and stress states. 218 The relationship between mechanical properties and the glass transition temperature (and GFA) was reported in some publications. For instance, the relationship between ductility and glass transition of metallic glasses was analysed by first-principle calculations for Fe Cr Mo P C B BMG alloys containing up to 27 at-% metalloids, leading to atomic bonding and connectivity in the amorphous network. 219 The elastic moduli of Fe-based BMG alloys were studied in FeCrMoCBErMe (Me5Al, Be, In, Nb, Ni and Pb) amorphous steels with high Fe content (.58 at-%). 220 The authors showed that the elastic moduli of the alloys are much larger than those of Zr- and Cu-based BMGs, and can be simply described approximately by a sum of elastic constants and the atomic percentage of components. These alloys also show a low Poisson s ratio similar to that of Mg-based BMGs, indicating that they belong to the brittle BMG family. Figure 7 shows the relationship between the fracture strength and Young s modulus for Fe-based BMG alloys, together with the data for other BMG alloys and conventional crystalline alloys. 23 The Fe-based BMG alloys have very high fracture strength of y3300 MPa for the Fe TM (P,C,B,Si) system and y4200 MPa for the Fe Co B Si Nb system. It is also noticed that there is a linear relationship between the fracture strength and the Young s modulus. As seen in the figure, the slope of the linear relation corresponding to the elastic strain limit is about three times larger than that for crystalline alloys. Thus, it is suggested that Febased BMG alloys have a much higher fracture strength in conjunction with much larger elastic strain as compared with conventional crystalline alloys. In addition to the static mechanical properties, the fatigue behaviour of the (Fe 0?5 Co 0?5 ) 72 B 20 Si 4 Nb 4 BMG alloy was examined as one of the dynamic mechanical properties. The fatigue test was performed under a tension tension stress condition with a stress ratio of 0?1 and a frequency of 10 Hz for the rod specimen with a maximum diameter of 1?85 mm. The fatigue strength limit defined as the ratio of stress amplitude to fracture strength after 10 7 cycles was 0?55 for the Fe Co-based BMG alloy which is slightly higher than that (0?49 0?54) for conventional alloy steels (Cr Mo steels, SCM435 and alloy tool steel, SKD61). 192 Corresponding to the very high fracture strength of the Fe Co-based bulk glassy alloy, the stress amplitude after 10 7 cycles shows very high value of 2310 MPa for the Fe Co-based alloy which is much higher than those ( ) for SCM435 and SKD61 steels, as shown in Fig. 8. The fatigue crack always initiated at a defect site located on the outer surface of the specimen and then propagated into the inner region, accompanying the distinct striation pattern. This behaviour was independent of the type International Materials Reviews 2013 VOL 58 NO 3 149

21 7 Relationship between fracture strength and Young s modulus for Fe-based BMG alloys. The data of other BMG alloys and conventional crystalline alloys are also included for comparison 23 of the BMG. Similar behaviour was noted in Fe-, Co-, Ti- and Cu-based BMG alloys. 221 A fatigue ratio of 0?16 0?17 was reported in an Fe 48 Cr 15 Mo 14 Er 2 C 15 B 6 glassy steel using a four-point bend testing method. 222 The fatigue fracture behaviour of BMGs suggests that elimination of defect sites on the outer surface could lead to an increase in the fatigue strength. It should be, however, mentioned that results on the fatigue behaviour of BMGs have not been conclusive; very widely differing results have been reported. The reader is advised to consult Ref. 7 for full details of the present situation. Corrosion behaviour For an effective use of BMGs, it is necessary to fully characterise them for their chemical behaviour also. The corrosion behaviour of BMGs becomes important when these materials need to be used in aggressive and hostile environments (high temperatures, oxidising atmospheres and corrosive media). Knowledge of the corrosion behaviour becomes critical when the BMGs are considered for biomedical applications and for decorative applications, or when surface appearance becomes important. The corrosion behaviour of metallic glassy 8 Relationship between fatigue stress amplitude and cycles to failure (S N curves) for Fe-based BMG alloys. The data for other BMG alloys and some conventional crystalline alloys are also shown for comparison International Materials Reviews 2013 VOL 58 NO 3

22 9 a Decreasing corrosion rate with increasing B content in an Fe 502x Cr 16 Mo 16 C 18 B x BMG alloy 7 and b Decreasing corrosion rate with increasing Cr content in 0?5 N NaCl solution at 298 K open to air for 168 h for the {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Nb 0?04 } 1002x Cr x alloy 7 alloy ribbons (about mm in thickness) produced by RSP methods was evaluated starting from ,224 It was reported that Cr-containing Fe-based Fe 802x Cr x P 13 C 7 glassy ribbons exhibited much higher corrosion resistance than the crystalline Fe Cr alloys. While the crystalline Fe Cr alloys corroded at a rate of about 0?5 1 mm/year, the glassy Fe Cr P C alloy did not show any measurable corrosion rate under identical conditions of exposure in 1 N NaCl solution at 30uC. Another important observation made was that the minimum amount of Cr required to achieve this corrosion resistance was only 8 at-%, much less in the glassy state than that required (.12 at-%) in the crystalline state. Furthermore, the glassy alloy did not exhibit any measurable weight change with the concentration of HCl (from 0?01 to 1 N) on exposure for 1 week at 30uC. On the other hand, the corrosion rate of the crystalline 18-8 austenitic stainless steel [(Fe 18Cr 8Ni (wt-%)] increased from mm/year in 0?01 N HCl to over 10 mm/year in 1 N HCl solution; severe pitting corrosion occurred in the range of 0?5 1 N HCl in the crystalline alloy. A number of investigations were also carried out on as well and an overview of the corrosion behaviour of BMG alloys is presented in Ref Fe-based BMG alloys investigated for their corrosion behaviour generally contained Cr, Mo, C and B in varying proportions. Cr has been shown to be essential in forming a passive layer and this is further facilitated by Mo addition. Both carbon and boron (together to the extent of y20 at-%) were found to be necessary to achieve glass formation. A typical Fe-based glassy alloy composition appears to be Fe 45 Cr 16 Mo 16 C 18 B 5. In some of the investigations, either additional metalloid elements, especially P, have been added or Mo has been partially replaced with Ta or Nb. All the glassy alloys investigated exhibited good corrosion resistance in concentrated HCl with the measured corrosion rates of as low as 1 10 mm/year. The corrosion rate of the Fe-based glassy Fe 502x Cr 16 Mo 16 C 18 B x alloys was found to decrease with increasing B content in the alloy. 106,226 A similar decrease in corrosion rate with increasing Cr content was also noted. 116,227 Figure 9a and b shows the decreasing corrosion rate in the Fe 502x Cr 16 Mo 16 C 18 B x glassy alloys with B content and in the {[(Fe 0?6 Co 0?4 ) 0?75 B 0?2 Si 0?05 ] 0?96 Nb 0?04 } 1002x Cr x glassy alloy with Cr content respectively. It may be noted that the corrosion rate decreased from 700 to 1?6 mm/year as the Cr content increased from 0 to 4 at-%. Addition of Mo to Fe-based alloys has also been reported to improve their corrosion resistance in HCl solution, since it prevents dissolution of Cr during passivation. 228 However, it has been noted that when the dissolution rate of Fe-based alloys is very high as in the active region, Mo selectively remains in the alloy because the dissolution rate of Mo is slower than that of other constituents. Furthermore, Mo has not been able to form its own passive film in the passive region of the alloys. Mo is also known to dissolve even at lower potentials in the passive region of the alloys, indicating a lower stability of the passive film of Mo in comparison with passive hydrated chromium or iron oxyhydroxide film. Consequently, excessive amounts of Mo addition to replace Fe have been reported to be detrimental for the corrosion resistance of Fe-based glassy alloys. 229 The corrosion resistance of the BMG alloys has been shown to be higher due to the presence of P in the alloy. Figure 10a shows the potentiodynamic polarisation curves of Fe 43 Cr 16 Mo 16 C 15 B 10 and Fe 43 Cr 16 Mo 16 C 10 B 5 P 10 bulk glassy alloys in 1 N HCl at 298 K. 230 From these curves, it may be noted that both the alloys passivate spontaneously. However, the passive current density is y10 21 Am 22 for the Fe 43 Cr 16 Mo 16 C 15 B 10 alloy, whereas it is approximately half ( Am 22 ) for the Fe 43 Cr 16 Mo 16 C 10 B 5 P 10 alloy. The lower passive current density in the P-containing alloy clearly demonstrates that this alloy has a better corrosion resistance. Figure 10b shows that a similar improvement in the corrosion resistance of Fe 45 Cr 16 Mo 16 C 18 B 5 and Fe 45 Cr 16 Mo 14 M 2 C 18 B 5 (where M5Nb or Ta) BMG alloys in 6 N HCl solution has also been obtained by substituting Mo with Nb or Ta 105 and also by addition of Nb to Cu-based BMG alloys. 231 The effect of test environment and structural changes on the corrosion behaviour of was also evaluated. 106 It was shown that the corrosion rate increased with increasing concentration of the corrosive medium, irrespective of the B content in Fe 502x Cr 16 Mo 16 C 18 B x BMG alloys, even though the corrosion resistance increased with increasing B content for a International Materials Reviews 2013 VOL 58 NO 3 151

23 10 a Potentiodynamic polarisation curves of Fe 43 Cr 16 Mo 16 C 15 B 10 and Fe 43 Cr 16 Mo 16 C 10 B 5 P 10 bulk glassy alloys in 1 N HCl solution open to air at 298 K showing that the passive current density for the P-containing alloy is lower 230 and b Anodic polarisation curves of Fe 45 Cr 16 Mo 16 C 18 B 5 and Fe 45 Cr 16 Mo 14 M 2 C 18 B 5 (where M5Nb or Ta) BMG alloys in 6 N HCl solution open to air at 298 K. Note the lower passive current density when Nb or Ta is present in the alloy 105 given concentration of the corrosive medium (Fig. 11). It was also shown that pitting occurred on the surface of the alloy after immersion for 1 week in 12 N HCl at room temperature, especially when the B content was only y4 at-%; pitting did not occur at higher B levels. Such a phenomenon of pitting did not occur in 1 and 6 N HCl solutions; instead they passivated spontaneously. The influence of structurally relaxing or crystallising the BMG alloys on the structure and corrosion behaviour of Fe-based BMG alloys has also been investigated. Pardo et al. 232,233 studied the effect of Cr content on the corrosion behaviour of Fe 73?5 Si 13?5 B 9Nb 3 Cu 1 BMG alloys in different concentrations of H 2 SO 4 (1, 3 and 5 N). They investigated the corrosion behaviour of this alloy in three different conditions: in the as-solidified fully glassy condition, by annealing it for 1 h at 813 K to obtain a nanocrystalline structure (10 15 nm grain size) and by fully crystallising the samples (to achieve a grain size of 0?1 1 mm) through annealing at 973 K for 1 h. The corrosion resistance was higher with increasing Cr content in the range studied (0 8 at-%). However, a minimum Cr concentration of 8 at-% was found necessary to generate a stable passive layer. Among all the conditions studied, the glassy structure showed the best corrosion resistance, followed by that in the nanocrystalline state. The fully crystallised alloy showed the least corrosion resistance. Magnetic properties Magnetic properties of materials are of fundamental importance for several applications in the electrical and electronic industries. A very large number of studies have also been conducted on Fe-based melt-spun ribbons starting from the pioneering investigation of Duwez and Lin on the Fe C P system in Since the most important application to which the melt-spun glassy ribbons have been put to is in transformer core laminations based on the excellent soft magnetic properties of these alloys, a significant amount of effort has also been spent on investigating the magnetic properties of Febased BMG alloys. However, an important difference between the investigations on melt-spun ribbons and BMGs is that while both metal metalloid and metal metal type alloys have been investigated in the thin film category, only metal metalloid type of alloys have been studied in the BMG group. Studies on the magnetic properties of BMG alloys in the metal metal type category are conspicuous by their absence. The nature of magnetic investigations in BMG alloys has followed trends very similar to what were done in the case of melt-spun glassy ribbons. And, in fact, even for BMG compositions, several researchers have been studying the magnetic behaviour using melt-spun 11 Increasing corrosion rate with increasing concentration of HCl in an Fe 502x Cr 16 Mo 16 C 18 B x BMG alloy International Materials Reviews 2013 VOL 58 NO 3

24 12 a Hysteresis loops for the glassy a Fe 65 Co 10 Ga 5 P 12 C 4 B 4 alloy in the melt-spun ribbon and gas-atomised powder conditions, 235 and b Fe 62?8 Co 10 B 13?5 Si 10 Nb 3 Cu 0?7 alloy in the melt-spun ribbon of 20 mm thickness and bulk rod of 1?5 mm diameter 101 ribbons. Some minor differences were noted in the magnetic properties of melt-spun ribbons and bulk rods, especially in those properties that are affected by structural relaxation, e.g. magnetostriction and coercivity. For example, it has been shown that the saturation magnetisation of the alloys is not any different whether measured on the melt-spun ribbon form or powder 235 or between the melt-spun ribbon and bulk rods of different diameters. 101 These are illustrated in the hysteresis loops presented in Fig. 12 for Fe 65 Co 10 Ga 5 P 12 C 4 B 4 and Fe 62?8 Co 10 B 13?5 Si 10 Nb 3 Cu 0?7 alloys respectively. The saturation magnetisation of the Fe M (P,C,B,Si) (where M5Ga or Mo) BMG alloys with optimum alloy components in each alloy system was in the range of 1?53 1?10 T, depending on the Fe content and the type of metalloid element. The coercivity was,3 Am 21 and the effective permeability at 1 khz was in the range of about Besides, rather good high-frequency permeability with the level of at 100 khz was achieved for Fe Ga P C B Si, Fe Mo P C B Si, Fe Co Ga P C B Si and Fe Co Mo P C B Si BMG alloys. The soft magnetic BMG alloys in the Fe (Cr,Mo) (P,C,B,Si) system have been used in various application fields as discussed later. Figures 13 and 14 show the compositional dependence of saturation magnetisation I s and coercivity H c respectively, for [(Fe 12x2y Co x Ni y ) 0?75 B 0?20 Si 0?05 ] 96 Nb 4 glassy alloys. 7 These soft magnetic data were obtained from the melt-spun ribbon samples with a cross-section of 0?02610 mm and the samples were also subjected to annealing treatment for 5 min at a temperature which is 50 K lower than the glass transition temperature. As shown in Fig. 13, I s shows high values of.1?3 T for Ferich Fe Co B Si Nb alloys and decreases gradually with increasing Ni and Co contents. On the other hand, the H c shows low values of,2?5 Am 21 over the whole composition range and decreases gradually with increasing Co content (Fig. 14). It is noticed that the Co-rich Co Fe B Si Nb glassy alloys exhibit very low H c values (,1 Am 21 ), presumably because of very low saturation magnetostriction of Reflecting the very low H c as well as the nearly zero saturated magnetostriction, 13 Compositional dependence of saturation magnetisation in the [(Fe 12x2y Co x Ni y ) 0?75 B 0?20 Si 0?05 ] 96 Nb 4 BMG alloys 7 International Materials Reviews 2013 VOL 58 NO 3 153

25 14 Effect of composition on the coercivity in [(Fe 12x2y Co x Ni y ) 0?75 B 0?20 Si 0?05 ] 96 Nb 4 BMG alloys 7 the [(Co 0?9 Fe 0?1 ) 0?75 B 0?020 Si 0?05 ] 96 Nb 4 glassy alloy exhibits excellent high-frequency permeability of at 100 khz. The important features of the soft magnetic properties for Fe- and Co-based glassy alloys in the Fe TM (P,C,B,Si) and (Fe,Co) B Si Nb systems are summarised below. Figure 15 shows the relationship between coercivity and electrical resistivity for Fe- and Co-based glassy alloys. 7 The data for conventional amorphous and nanocrystalline soft magnetic alloys are also presented for comparison. It is clear that the Fe- and Co-based BMG alloys have unique combination of very low coercivity and high electrical resistivity, a combination that cannot be obtained for any other kind of soft magnetic metallic material including amorphous and nanocrystalline alloys. Based on detailed studies on glassy alloys produced by RSP techniques, it was suggested earlier 237 that a clear relationship exists between coercivity and internal stresses; a low coercivity is obtained when the internal stresses are low. Such a situation of low internal stresses can be achieved in BMG alloys produced at relatively low cooling rates due to the formation of a more homogenised disordered structure consisting of unique network-like atomic configurations. Consequently, BMG alloys exhibit a low coercivity. Results obtained to date show that in order to achieve a high saturation magnetisation, it is necessary to have as high an Fe content as possible and decrease the other metal and metalloid contents to the minimum value that is necessary to obtain the glassy phase. Investigating on these lines, Makino et al. 238 selected the Fe 942x Nb 6 B x alloy system to arrive at the optimum composition to achieve the best soft magnetic properties. They noted that a fully glassy phase was obtained only when the B content was a minimum of 10?5 at-%. The saturation magnetisation of.1?6 T was obtained in an alloy with x(9, where the alloys have a mixed glassyza-fe composite structure. The highest permeability m e of and the smallest grain size of 10?5 nm were obtained at x511. Thus, a single composition was not going to offer the best soft 15 Relationship between coercivity and electrical resistivity for Fe- and Co-based glassy alloys. The data for nanocrystalline alloys and conventional amorphous alloys are also included for comparison International Materials Reviews 2013 VOL 58 NO 3

26 magnetic properties. Therefore, these investigators have limited the Nb content in the alloy to 6 at-% and small amounts of Cu and P were added to this alloy to optimise the composition at Fe 84?9 Nb 6 B 8 P 1 Cu 0?1. This alloy, fabricated by melt spinning in air, contained the a-fe phase with an average grain size of 10 nm dispersed in a glassy matrix and showed excellent soft magnetic properties with a saturation magnetisation of 1?61 T, coercivity of 4?7 Am 21 and permeability of The core losses for this alloy were very low at 0?11 W kg 21. The best soft magnetic properties are achieved by crystallising the glassy phase to produce a uniform nanostructure, 239 referred to as FINEMET alloys. And in FINEMET and other alloys, the volume fraction of the nanocrystalline phase is very large, reaching a value of almost 90%. Since a glassy precursor is a pre-requisite to achieve this microstructure, one needs to have a minimum amount of different metals and/or metalloids to first produce the glassy phase. Furthermore, to achieve the large volume fraction of the nanocrystalline phase, the presence of metals like Cu, Nb, Mo, W and Ta is required. However, the presence of these non-magnetic metals can significantly reduce the saturation magnetisation. Furthermore, these metallic elements are quite expensive. To overcome these difficulties and the requirement of metallic elements and also to achieve a very high saturation magnetisation, Makino et al have developed novel Fe-based alloys that contain only metalloids. The generic composition of their alloys is Fe 83?3 84?3 Si 4 B 8 P 3 4 Cu 0?7 and is based on the Fe 82 Si 9 B 9 alloy that proved most promising among the Fe-based melt-spun magnetic alloys. In the modified composition, P substitutes for B and Cu for Fe. (The presence of a small amount of Cu is necessary to achieve nanocrystallisation). The large concentration of Fe provides a high saturation magnetisation and the absence of the metallic elements (other than Fe) ensures that the alloys are not expensive. Applications Bulk metallic glass alloys exhibit very high strength (both in tension and compression), large elastic elongation limit, very high hardness, excellent corrosion resistance and a good combination of soft magnetic properties. Recent rapid progress in the development of BMG alloys has made it possible to exploit these novel materials in a variety of application fields. For instance, Zr-, Ti-, Fe-, Co-, Ni- and Cu-based BMG alloys are already in use for magnetic sensing, chemical and structural applications. Specifically, it was suggested 7,243 that typical industrial applications of BMG alloys include: magnetic applications such as linear actuators, magnetic cores, choke coils and high-frequency magnetic-shielding sheets; chemical applications such as fuelcell separators; and structural applications such as sporting goods, precision optical parts, precision gears for micromotors, diaphragms for pressure sensors, tubes for Coriolis mass flowmeters, aircraft parts, automobile valve springs, etc. On the basis of these current trends of applications of BMG alloys, this section focuses on the applications of Fe-based BMG alloys. This will be based on the two important attributes of the Fe-based BMG alloys, namely good soft magnetic properties and high strength for use as soft magnetic and structural materials respectively. Soft magnetic materials Soft magnetic glassy alloys in the Fe TM P C B Si system have been commercialised under the trade name of Liqualloy, 244,245 whose magnetic properties are summarised in Table 6, 246 together with those of Mn Zn ferrite and Fe Al Si alloy (Sendust) for comparison. The application fields of the soft magnetic powder cores Liqualloy cores have been in choke coils of AC DC converters, DC DC converters, noise suppression sheets, etc. The Liqualloy magnetic powder cores exhibit good soft magnetic properties, e.g. nearly constant relative permeability in a wide frequency range up to several MHz; good linear relation with small degradation slope in the relation between permeability and DC bias field, i.e. good DC-superposed characteristics; and much lower core losses as compared with Ni Fe Mo permalloy and Sendust. The excellent core loss characteristics are due to the reduction in eddy current loss resulting from much higher electrical resistivity (168 mv cm) of Liqualloy than that of Sendust (82 mv cm), as shown in Table 6. In addition, as shown in Fig. 16, the Liqualloy powder core shows much better DC-superposed characteristics than those of Mn Zn ferrite, in addition to nearly the same low core losses between Liqualloy and Mn Zn ferrite. 246 Furthermore, in comparison with Mn Zn ferrite, the Liqualloy power inductor shows better DC-superposed characteristics in the elevated temperature range up to 393 K. Through these advantages, the Liqualloy powder core has been used as the power inductor in lap-top type personal computers because of the advantages of higher efficiency and much smaller heat generation than those for commercial power inductors, as shown in Fig. 17. It may be noted that the shape of the Liqualloy powder produced by water atomisation can be changed from the original spherical shape to flake shape with a small thicknesses of about 1 3 mm by bead milling treatment. The resulting Liqualloy sheet consisting of the flaky powder and resin was found to exhibit good electromagnetic noise suppression characteristics because of high conversion ability from electromagnetic noise to heat. 246 The use of Liqualloy noise suppression sheet has resulted in significant reduction in noise level and hence the sheet has been used in digital still camera as exemplified in Fig. 18. Furthermore, the Liqualloy electro-magnetic sheet was found to exhibit good radio frequency identification characteristics. As shown in Fig. 19, the Liqualloy sheet shows better function as Table 6 Magnetic properties of Liqualloy* 246 Commercial name Alloy B s /T H c /A m 21 m max Resistivity/mV cm Liqualloy Fe TM P C B Si Mn Zn ferrite Sendust Fe Al Si *B s : saturation magnetic flux density; H c : coercive field. International Materials Reviews 2013 VOL 58 NO 3 155

27 16 Direct current-superposed characteristics and core losses of Liqualloy powder core. The data for Mn Zn ferrite are also shown for comparison 246 magnetic deflection yokes. One can obtain much longer transmission distance as compared with the case of loop antenna/metal parts. These better characteristics are attributed to the fact that high real and low imaginary parts of permeability can be achieved at a high carrier frequency of 13?56 MHz. As a result, the Liqualloy electromagnetic sheet for radio frequency identification system has been used in higher functional cell phones. In addition, another type of soft magnetic powder core has been produced in the Fe Nb B Si and Fe Nb Cr P B Si systems 247 with a high saturation magnetisation of 1?3 T. The cores were also produced by a procedure similar to that for Liqualloy, i.e. mass production of spherical glassy alloy powder by water atomisation, followed by cold consolidation of the mixture of glassy alloy powder and epoxy resin and then annealing. The new magnetic powder cores have been named as SENNTIX-I and SENNTIX-II for the Fe Nb B Si and Fe Nb Cr P B Si systems respectively, and can be characterised as exhibiting the lowest core losses among all kinds of magnetic powder cores developed to date. Besides, the higher saturation magnetisation has enabled use of the cores in a higher current range. The fundamental characteristics of SENNTIX are summarised in Table 7, 247 together with other Fe Si B M (M5alloying additions) systems named as FINEMET 248 and other Fe-based alloys. In addition to power inductor and electromagnetic sheet applications, soft magnetic BMG alloys have been tested for applications to position sensor, antennas for radio-controlled watch, solenoid valve, magnetic sensor, etc. 249 For instance, by using the sensing method with saturable coil, a Co Fe B Si Nb magnetic sensor has been identified to show sharper transient signal, higher output voltage and higher mechanical strength than those for conventional magnetic sensors. Such a high performance is expected to cause improved sensitivity, high S/N ratio, miniaturisation and easy handling and operation. Besides, an Fe Co B Si Nb position sensor was confirmed to detect a higher impedance change by movement of core output voltage which resulted in improved sensitivity and miniaturisation. Applications to structural materials The high yield (or fracture) strength, low Young s modulus, large elastic strain limit and easy formability in the supercooled liquid region are the main attributes of BMGs that make them attractive for structural applications. The large supercooled liquid region in BMG alloys offers an excellent opportunity to form complex shapes easily. This is mainly because the plastic flow of the material in this temperature regime is Newtonian in nature (i.e. the strain rate is proportional to the applied stress). This attribute of BMG alloys has been extensively exploited to produce different types of parts with complex shapes such as gears, coiled springs and other complex parts. The parts using BMG alloys have complex shapes and the sizes of these parts are much smaller than what have been achieved using conventional crystalline alloys. But in the context of this review, it should be made clear that the applications that have been described in the literature are not just based on Fe-based BMG alloys, but for many different types of BMG alloys. However, these have been included for the sake of completion and also to bring awareness to the readers of the potential structural applications of BMG alloys. Sporting goods Bulk metallic glasses have first found widespread application in sporting goods due to their desirable mechanical properties, namely high strength and large elastic elongation limit. The excellent mechanical properties of Zr-based BMGs were exploited commercially in golf clubs, followed by tennis rackets, baseball and softball bats, skis and snowboards, bicycle parts, scuba gear, fishing equipment and marine applications. 250 The modulus of resilience U calculated as the area under the elastic portion of the stress strain curve works out to be U~ 1 2 s ye y ~ 1 2 Ee2 y (3) 17 Application example of Liqualloy powder core to power inductor in lap-top type personal computer 246 where s y and e y are the yield stress and elastic strain limit respectively, and E is the Young s modulus. Since the e y 156 International Materials Reviews 2013 VOL 58 NO 3

28 18 Characteristics of Liqualloy noise suppression sheet and its application example to digital still camera value for BMGs is at least twice that in a crystalline material, the modulus of resilience is at least four times that of a crystalline material. It is suggested that 99% of the impact energy from a BMG golf club head is transferred to the ball. This value should be compared with 70% for a titanium head. Liquidmetal Technologies calls this high energy transfer as pure energy transfer. This company has started marketing baseball bats, tennis rackets and other sporting goods as well. The HEAD Radical Liquidmetal tennis rackets seem to offer large sweetspots, plenty of control and impressive feel, with very little vibration. Precision gears Complicated structures can be easily formed from BMGs in their supercooled liquid region and this offers many advantages including: formation of highly homogeneous features, even on a nanoscale, due to the absence of crystalline features such as grains in the glassy material; very little solidification shrinkage during superplastic forming of BMGs in the supercooled liquid region, due to the absence of first-order phase transformations which introduce significant shrinkage into the casting, reduced degree of porosity in the formed part; and superior surface imprintability and net-shape forming capability, which have been found to be very attractive in the field of micro-machines. 251 Because of the excellent filling characteristics of BMG alloys, it has been possible to produce extremely small parts of complex design using BMGs. A sun-carrier made with ani 53 Nb 20 Zr 8 Ti 10 Co 6 Cu 3 BMG alloy consisting of a 19 Functional characteristics of Liqualloy magnetic sheet for radio frequency identification system. The data of loop antenna and loop antenna/metal parts modes are also shown for comparison International Materials Reviews 2013 VOL 58 NO 3 157

29 micro-gear, carrier plate and three pins has been fabricated. The outer diameter of the gear is 0?65 mm, it has 14 teeth with a module of 0?04 and the micro-gear is seated on a carrier plate 1?7 mm in diameter. Three pins with a diameter of 0?30 mm and a length of 0?45 mm are located at the bottom of the carrier plate for rotating planetary gears. 252 Motors Micro-geared motors with high torque have been used in different engineering fields. With advancing technology and the need for miniaturisation, the size of the motors has been constantly decreasing. Bulk metallic glasses appear to be particularly suited for this application. The wear resistance behaviour of a 2?4 mm diameter gear was evaluated using sliding-wear and rolling-wear tests. 253 It was reported that the wear loss (volume) of the Ni-based BMG was larger than that of the carbon steel under sliding-wear conditions, but smaller under rolling-wear conditions. It was noted that the gear teeth of the carbon steel are worn off and heavily damaged even after just 8 h of use, while the teeth of the BMG alloy are in very good condition even after 2500 h of use. The wear life of the gear was 1?6 times higher than an all-steel gear when one of the gears was replaced by a BMG alloy gear. However, the wear life increased by seven times when more gears were replaced with BMG alloys. But in the case of all BMG alloy gears, the wear life was 313 times higher than all-steel gears. 254 Encouraged by these results, Inoue et al. 255 have fabricated the world s smallest size micro-geared motor (1?5 mm in diameter and 9?4 mm in length) using a highstrength Ni-based BMG alloy. The components of this geared motor cannot be made by any mechanical machining methods. It consists of a sun-carrier, an output shaft and six pieces of planetary gear, all made out of the Ni-based BMG with the composition Ni 53 Nb 20 Zr 8 Ti 10 Co 6 Cu 3. It was confirmed that this micro-geared motor had high rotating torques of 0?1 mn m at two stages stacked gear-ratio reduction system and 0?6 mn m at three stages which were 6 20 times higher than the vibration force for conventional geared motor with a diameter of 4?5 mm in mobile telephones. These micro-geared motors are expected to be used in advanced medical equipment such as endoscopes, micropumps, rotablator and catheter for thrombus removal, precision optics, micro-industries, micro-factory, etc. Shot peening balls By utilising the good GFA, high mechanical strength, large elastic elongation limit, high corrosion resistance and good surface smoothness of Fe-based BMG alloys, commercialisation of Fe-based BMG alloy powders with particle sizes of 0?05 1 mm has been accomplished as shot peening particles. 256 The high GFA has made it possible to produce glassy alloy powders over a wide particle size range by conventional water atomisation methods, leading to very good cost advantage. The much higher yield strength and large elastic elongation limit also permit to generate much deeper compressive residual stress region as well as much higher level of compressive residual stress. In addition, Fe-based glassy alloy shot balls have much longer endurance times which are about 10 times longer than those for cast alloy steel and high speed steel shot balls. Similarly, BMG alloys have also been used as automobile valve springs, 257 diaphragms for pressure sensors, 249 pipes for a Coriolis mass flowmeter, 249,258 optical mirror devices 259 and structural parts for aircraft. 260 Reasons for the choice of these materials and the details of applications may be found in Ref. 7. Coating applications Here, special attention should be paid to a new high velocity oxygen fuel (HVOF) spray coating technique because of the advantages of very high moving velocity of atomised powders, relatively low spray flame temperature, very compact equipment size and rather easy operation. This technique was applied to the formation of spray-coated glassy alloy layer of Fe-based glassy alloy powders on various alloy substances. 261 Figure 20 shows the cross-section of spray-coated glassy alloy layer of Fe 50 Cr 15 Mo 15 C 14 B 6 alloy on stainless steel substrate, together with the XRD patterns of the original alloy powder and the coated surface layer. The spray-coated layer with a thickness of y0?2 mm has a rather high relative density of over 99% and retains the glassy structure without any crystalline phase being present and thus the glassy structure of the original powder is retained. Here, it is important to point out that a similar spray-coated layer was formed on various Table 7 Specific temperatures, GFA and magnetic properties of Liqualloy* 247,248 Commercial name Alloy T c /K T g /K T x /K DT x /K t max /mm B s /T) H c /A m 21 m/khz W/kW m 23 Phase Reference SENNTIX Fe 77 P 7 B 13 Nb 2 Cr Glass 247 Fe 77 P 9 B 11 Nb 2 Cr Fe 77 P 11 B 9 Nb 2 Cr Fe 77 P 13 B 7 Nb 2 Cr Fe 76 P 10 B 11 Nb 2 Cr FINEMET FINEMET Ultrafine grain 248 FINEMET structure FT-1H FT-1M FT-1L Fe Si B M Amorphous Co Fe Si B M Fe 75 Si 10 B 12 Cr Amorphous 247 *T c : Curie temperature; T g : glass transition temperature; T x : crystallisation temperature; DT x : supercooled liquid range; t max : maximum thickness for glass formation; B s : saturation magnetic flux density; H c : coercive field; m: permeability; W: core loss. 158 International Materials Reviews 2013 VOL 58 NO 3

30 20 Optical micrograph and XRD patterns taken from the cross-section of Fe 50 Cr 15 Mo 15 C 14 B 6 glassy alloy coated layer on SUS 304 steel substrate produced by an HVOF spraying technique 261 metallic substrates such as carbon steel, aluminium alloy and magnesium alloy and the deposition tendency of the sprayed powder was independent of the kind of substrate material. In addition, it has been confirmed that the maximum thickness for the formation of the Fe Cr Mo C B glassy alloy layer reaches at least 0?2 mm in the case of the stainless steel substrate. However, when the crystalline alloy powder, obtained by crystallisation of the atomised Fe Cr Mo C B glassy alloy powder, was used as the powder for the HVOF spray technique, the coated layer consisted of mostly crystalline phases including a small amount of glassy phase. This result indicates two important points: the heating temperature of the powder during HVOF spray coating is not high enough to obtain a fully melted state and the incomplete melted state prevents formation of the glassy alloy coated layer; and the heating temperature reaches the (solidzliquid) twophase region and that the formation of a small amount of glassy phase is the result of rapid solidification of the partially melted liquid region. This result also suggests that the cooling rate of the HVOF spray technique is high enough to result in the formation of a glassy phase through suppression of precipitation of a crystalline phase for individual powder. Figure 21 shows the corrosion resistance and Vickers hardness of spray-coated Fe Cr Mo C B glassy alloy layer shown in Fig. 20, in comparison with the data of SUS 304 steel and hard Cr-coated plate. 262 The spraycoated glassy alloy layer exhibited much better corrosion resistance as revealed by the much lower corrosion current density and much higher anodic potential than those for SUS 304 steel in the anodic polarisation curve. These tests were conducted in 1 N H 2 SO 4 at 298 K. In addition, the Vickers hardness of the spray-coated glassy alloy layer is about which is considerably higher than that ( ) for the hard-coated chromium plate. Furthermore, the wear resistance of the coated glassy alloy surface layer is demonstrated in Fig. 22, which shows SEM images of the coated Fe Cr Mo C B glassy alloy layer subjected to the ring-on-disc friction wear test, 263 together with the data for cast iron (FC) and alloy tool steel (SKD) materials tested under similar conditions. The SEM images reveal that the surface of the glassy alloy layer keeps much better smoothness as compared with those for the conventional steel materials. In addition, the wear resistance is better for the coated glassy alloy layer made of the finer alloy powder rather than the coarse alloy powder presumably because of the achievement of higher relative density for the coated surface layer and the higher hardness of the finer alloy powder. By exploiting the above-described advantages, an appropriate application field for practical use is now under consideration. A coating was applied to the inside 21 Corrosion behaviour and Vickers hardness values of spray-coated Fe 50 Cr 15 Mo 15 C 14 B 6 glassy alloy coated layer in comparison with the data for crystalline stainless steel and hard Cr-coated plate 262 International Materials Reviews 2013 VOL 58 NO 3 159

31 22 Images (SEM) of spray-coated Fe 50 Cr 15 Mo 15 C 14 B 6 glassy alloy layer after the ring-on-disc friction wear test. The data of the FC and SKD materials tested under the same conditions were also shown for comparison 263 surface of a lead-free soldering vessel with a diameter of 50 cm and height of cm. The coating was found to be effective even after continuous use of the vessel for melting. As shown in Fig. 23, one cannot recognise any detectable damage caused by solder melt attack even after use for 6 months. This is in stark contrast to serious erosive damage of SUS 304 stainless steel vessel used only for 1 month. 263 Concluding remarks Fe-based BMG alloys are relatively new materials with interesting combination of properties such as high strength and good soft magnetic properties. Starting with the synthesis of Fe 80 B 20 high strength glassy alloys produced in the form of thin ribbons in 1970s, there has been continuous progress in the synthesis of a number of alloy compositions with increasing section thickness. A maximum section thickness of y16 mm is now achieved in (Fe 0?8 Co 0?2 ) 48 Cr 15 Mo 14 C 15 B 6 Tm 2 and Fe 41 Co 7 Cr 15 Mo 14 C 15 B 6 Y 2 BMG alloys, through additions of rareearth elements such as Y and Tm. Several reports are available discussing the recent developments in the mechanical and magnetic properties and corrosion resistance of these BMG alloys. In contrast to the developments in thin film ribbons, non-ferromagnetic Fe-based BMG alloys with an Fe content,50 at-% have also been developed. Both the ferromagnetic and non-ferromagnetic types of Fe-based BMG alloys have been developed by taking advantage of their unique characteristics such as high mechanical strength and good corrosion resistance. The achievements of the ferromagnetic Fe-based BMG alloys led to practical applications utilising the magnetic properties, while the superior GFA of the non-ferromagnetic Fe-based BMG alloys makes it possible to produce them with maximum diameter of Fe-based BMG alloys reaching.1 cm. Thus, the two types of Fe-based BMG alloys have been developed separately by utilising their unique characteristics. A number of applications have been developed for these magnetic Fe-based BMG alloys. Acknowledgement The authors are highly thankful to Dr Akira Takeuchi of Institute of Materials Research, Tohoku University, 23 Surface appearance of spray-coated Fe 50 Cr 15 Mo 15 C 14 B 6 glassy alloy surface layer inside the lead-free soldering vessel of 50 cm in diameter and 40 cm in height which was used for 6 months International Materials Reviews 2013 VOL 58 NO 3

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