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Preface to the viewpoint set on mechanical behavior of metallic glasses
Scripta Materialia 54 (2006) 317–319 www.actamat-journals.com
Editorial
Preface to the viewpoint set on mechanical behavior of metallic glasses
The lack of long-range atomic order in metallic glasses makes their mechanical properties and behavior considerably different from those of crystalline alloys. Duwez and coworkers announced the formation of the first metallic glass, a gold–silicon alloy, in 1960 [1], but study of mechanical properties required the development of techniques and alloys capable of producing larger and more stable amorphous specimens. By the 1970s, however, rapid solidification techniques for producing amorphous ribbons and wires had been perfected and many of the main features of mechanical behavior, including high strength, large elastic limit, homogeneous and inhomogeneous modes of deformation, and the novel ‘‘molten’’ appearance of fracture surfaces had been established [2,3]. The early metallic glasses had poor glass-forming ability, with critical cooling rates on the order of 105–106 K s 1 which limited the section thicknesses to 50 lm and made load-bearing applications beyond the reach of these alloys. But it was recognized that larger specimens could be produced if formation of crystalline phases could be suppressed. An important success in this area was the development of fluxing techniques for Pd-based glasses [4] which, by eliminating heterogeneous nucleation sites, inhibited crystallization and allowed the production of ‘‘bulk’’ metallic glasses (loosely defined as alloys capable of forming amorphous structures with a minimum section thickness of 1 mm). In 1993, Peker and Johnson announced the discovery of Zr41.2Ti13.8Cu12.5Ni10Be22.5 [5]. This discovery, together with new alloys developed by the Inoue group in Japan [6], spurred world-wide interest in bulk metallic glasses. Among the results of the renewed efforts have been the development of a variety of new bulk glass-forming alloys, including alloys based on Ti [7], Cu [8], Fe [9–11], and Mg [12] (among others) and an improved understanding of alloy design for glass formation. From the standpoint of mechanical properties and behavior, the development of bulk metallic glasses is of interest for two reasons. First, it creates the possibility of using metallic glasses in load-bearing applications, which has long been thought desirable due to the high yield
strength of amorphous alloys. Second, the availability of bulk specimens has enabled a variety of studies of mechanical properties that were difficult or impossible with the ribbon and wire specimens used in the early studies. (It should be pointed out, however, that instrumented indentation techniques have allowed fundamental advances to be made by probing mechanical behavior on the nanoscale, without the need for bulk specimens [13,14].) Thus, now seems to be an appropriate time to take stock of what we have learned, and to lay out visions for the future of metallic glasses, in terms of both practical applications and new research directions. These are the goals of this viewpoint set, and the papers have been selected with them in mind. The general enthusiasm for any new class of materials sometimes leads to grandiose claims for the potential of the materials in question. By now, however, enough data has been gathered on a variety of amorphous alloys that we can attempt to replace the hyperbole with more reasoned speculation. Greer and Ashby provide this in the first article [15] by considering both the strengths and limitations of metallic glasses relative to other materials. They propose that the size of the process zone for plastic deformation (such as ahead of a crack tip) relative to the size of the component in question is an important criterion. The high strength of metallic glasses makes the process zone small (1 mm), suggesting that macroscopically brittle fracture is to be expected in components that exceed this size. However, it may be possible to overcome this limitation, either by increasing the process zone size (by introducing particles [16] or pores [17] that promote shear band initiation) or by effectively decreasing the component dimensions. An example of the latter is a metallic glass foam, in which the ligaments are smaller than the process zone size [18]. Economic considerations ultimately dictate whether new materials see wide-spread use. Some glass-forming alloys involve the use of elements that are either expensive or difficult to handle safely (such as beryllium), or high purity raw materials. However, the development of alloys that are
1359-6462/$ - see front matter Ó 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2005.10.004
318
Editorial / Scripta Materialia 54 (2006) 317–319
based on less expensive elements and more tolerant of impurities make it seem likely that the cost of metallic glasses will continue to decline relative to other highperformance alloys. If so, the processing flexibility afforded by superplastic forming in the supercooled liquid region above the glass transition temperature [19] should allow metallic glasses to compete favorably for use in many more applications. At temperatures well below the glass transition, metallic glasses deform inhomogeneously, with deformation localized into narrow shear bands. Shear localization limits stable plastic flow, and the concentration of slip on shear bands promotes fracture initiation. Thus, controlling shear localization will be critical if metallic glasses are to be widely used in load-bearing applications. Understanding how the atomic-scale mechanisms of deformation lead to shear localization, besides being of scientific interest, may also contribute to design of new alloys with reduced tendency for shear localization. For instance, some amorphous alloys show a tendency for the strain to be distributed over a larger number of shear bands, allowing these alloys to support larger plastic strains in compression [20]. This behavior is not well understood, and in any event the effect is not large enough to avoid abrupt fracture in tension. But it might be possible to develop composite materials in which the second phase promotes shear band initiation and the matrix has a tendency for less severe shear localization. The combination of these effects could lead to materials with significantly improved properties. Given its importance, it is appropriate that several of the articles in this viewpoint set focus on shear localization. Flores [21] and Jiang and Atzmon [22] both consider structural changes associated with inhomogeneous deformation. The positron annihilation measurements discussed by Flores clearly show that inhomogeneous deformation is accompanied by an increase in the number of free volume sites, and calorimetry data on deformed and annealed specimens can also be interpreted in terms of free volume generation and annihilation. Jiang and Atzmon show that least some of this free volume can be found as nanometer-scale voids inside shear bands, although the formation of the nanovoids is not ubiquitous, being sensitive to both the loading condition and the alloy content. In some cases, nanocrystals form inside shear bands instead. The difference is likely important, as nanocrystal formation may stabilize shear deformation and even result in strain hardening, while nanovoids are more likely to contribute to the initiation of cracks. In many load-bearing applications, fracture and fatigue behavior are important. In their paper, Lewandowski and coworkers [23] address the toughness of metallic glasses and metallic-glass-matrix composites (MGMC). They discuss the recent observation that good toughness is associated with a low resistance to shear deformation (quantified by the shear modulus, l) coupled with high resistance to dilatation near a crack tip (quantified by the bulk modulus, B) [24–26]. Thus, intrinsically tough glasses
should have low l/B, a prediction that is borne out by data on a wide range of glasses. A low value of l/B contributes to the distribution of shear strain on multiple shear bands, giving these glasses the ability to support large plastic strains. Lewandowksi also addresses toughness in metallic-glass-matrix composites, and points out that design of the toughening phase (to resist brittle cleavage in body-centered cubic metals, for instance) is likely to be as important as design of the composite microstructure itself. In MGMC materials a crystalline phase is deliberately introduced into an amorphous matrix, with the twin goals of both promoting shear band initiation (thereby distributing the plastic strain more broadly) and hindering shear band propagation. The crystalline particles (or dendrites) are usually both weaker and more ductile than the matrix—just the opposite of conventional metal–matrix ¨ stu¨ndag and coauthors [27] describe the use composites. U of in situ neutron scattering together with self-consistent modeling to study the micromechanics of deformation. A key question of ongoing research is the mechanism by which the particles initiate shear bands in the matrix. Under cyclic loading, metallic glasses tend to show low crack-growth thresholds and low endurance limits compared to other high strength alloys. As Dauskardt and coworkers [28,29] discuss, the poor fatigue properties of single-phase glasses are also the result of the tendency for shear localization, which, under cyclic loading, leads to rapid initiation of fatigue damage. It is possible that alloys with a greater tendency to distribute plastic strain may reduce this effect somewhat, but it seems more likely that substantial improvement of fatigue properties will require introduction of second-phase particles that can help delay the onset of fatigue damage and inhibit fatigue crack growth. Turning to mechanisms of deformation, in his article Spaepen [30] reviews the free volume model for deformation, in which transition state theory is used to describe liquid flow, creep deformation, and structural relaxation. An understanding of these phenomena is important even for low temperature deformation, because the underlying atomic-scale mechanisms are likely the same. Khonik and coworkers [31] address stress relaxation of amorphous alloys in their article. Although SpaepenÕs seminal paper [32] presented the free volume model in terms of single-atom hops, measurements of the activation volume for creep suggest that the actual atomic rearrangements involve something more on the order of 10 atoms. That deformation should require collective rearrangements within ‘‘shear transformation zones’’ (STZ) was proposed by Argon [33], and since that time the STZ theory has seen significant extension, particularly by Langer and coworkers. Langer [34] provides an overview in his paper, including some discussion of shear localization. He presents simulations showing that nascent shear bands can initiate at surface flaws; interestingly, the recent molecular dynamics simulations of Li and Li [35] indicate that surface imperfections are essential to the
Editorial / Scripta Materialia 54 (2006) 317–319
development of shear localization and that without them relatively stable necking can develop during tensile deformation. In their paper, Shi and Falk [36] present the results of some very interesting molecular dynamics simulations showing that shear localization is associated with the disruption of short-range order in the glass. It has long been speculated that the extensive deformation inside shear bands might disrupt short-range order; for instance, this idea was proposed to explain why shear bands are etched preferentially by chemical attack and why existing shear bands reactivate upon removal and reapplication of a load [37]. The difficulty associated with characterization of the atomic-scale structure of the deformed material inside a shear band has, however, precluded any direct confirmation to date. Shi and Falk also suggest that percolation of short-range order over longer length scales affects deformation, with more ordered structures apparently being more resistant to shear localization. This may explain why some alloys develop more distributed shear zones [20,25], as discussed above. Some support for the importance of structure on deformation is provided by the observation that in amorphous Zr–Ta–Cu–Ni–Al enhanced structural order on length scales of 1–1.5 nm is correlated with shear band multiplication and enhanced plastic strain [38,39]. Whether such effects extend to other alloys is an interesting question for future research. Progress in this area is likely to hinge on continued developments in structural characterization of amorphous alloys. Particularly promising in this regard is fluctuation electron microscopy, which has the potential to reveal entirely new details about the nanometer-scale structure of metallic glasses [40,41]. At elevated temperatures (above about 0.7 Tg) metallic glasses deform homogeneously. Nieh and Wadsworth [42] review homogeneous deformation, with an emphasis on the importance of understanding how structural changes that may occur—particularly nanocrystallization—affect the flow behavior of the glass. Above the glass transition temperature extremely large deformations can be achieved at low stresses, making this regime of great interest for processing of metallic glasses into complex shapes. Superplastic forming above Tg is not covered in this set of articles, having been the subject of a recent review by Schroers [19]. In summary, I believe that this viewpoint set provides a useful snapshot of the current state of knowledge of mechanical behavior of metallic glasses. I fully expect that the future will see continued advances both in our understanding of fundamental aspects of mechanical behavior, as well as new applications of metallic glasses and metallic-glass-matrix composites. References [1] Klement W, Willens RH, Duwez P. Nature 1960;187:869. [2] Pampillo CA, Polk DE. Acta Metall 1974;22:741. [3] Masumoto T, Maddin R. Mater Sci Eng 1975;19:1.
[4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
[32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42]
319
Kui HW, Greer AL, Turnbull D. Appl Phys Lett 1984;45:615. Peker A, Johnson W. Appl Phys Lett 1993;63:2342. Inoue A. Acta Mater 2000;48:279. Zhang T, Inoue A. Mater Trans JIM 1998;39:1001. Inoue A, Zhang W, Zhang T, Kurosaka K. Acta Mater 2001;49:2645. Shen T, Schwarz R. Appl Phys Lett 1999;75:49. Ponnambalam V, Poon SJ, Shiflet GJ, Keppens VM, Taylor R, Petculescu G. Appl Phys Lett 2003;83:1131. Lu ZP, Liu CT, Thompson JR, Porter WD. Phys Rev Lett 2004;92:245503. Inoue A, Nakamura T, Nishiyama N, Masumoto T. Mater Trans JIM 1992;33:937. Schuh CA, Nieh TG. J Mater Res 2004;19:46. Schuh CA, Lund AC, Nieh TG. Acta Mater 2004;52:5879. Ashby MF, Greer AL. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.051. Hays CC, Kim CP, Johnson WL. Phys Rev Lett 2000;84:2901. Wada T, Inoue A, Greer A. Appl Phys Lett 2005;86:251907. Brothers AH, Dunand DC. Acta Mater 2005;53:4427. Schroers J. JOM 2005;57:35. Xing LQ, Li Y, Ramesh KT, Li J, Hufnagel TC. Phys Rev B 2001;64:180201(R). Flores KM. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.04.049. Jiang WH, Atzmon M. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.052. Lewandowski JJ, Shazly M, Nouri AS. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.010. Chen HS, Krause JT, Coleman E. J Non-Cryst Solids 1975;18:157. Schroers J, Johnson W. Phys Rev Lett 2004;93. Lewandowski J, Wang W, Greer A. Philos Mag Lett 2005;85:77. ¨ stu¨ndag E, Kim CP, Brown DW, Bourke Clausen B, Lee SY, U MAM. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.009. Hess PA, Cameron KK, Menzel BC, Dauskardt RH. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.006. Cameron KK, Hess PA, Menzel BC, Dauskardt RH. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.007. Spaepen F. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.09.046. Bobrov OP, Csach K, Khonik VA, Kitagawa K, Laptev SN, Yazvitsky MY. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.09. 056. Spaepen F. Acta Metall 1977;25:407. Argon A. Acta Metall 1979;27:47. Langer JS. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.005. Li Q, Li M. Appl Phys Lett 2005;87. Shi Y, Falk ML. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.053. Pampillo CA. J Mater Sci 1975;10:1194. Hufnagel TC, Brennan S. Phys Rev B 2003;67:014203. Li J, Gu X, Hufnagel TC. Microsc Microanal 2003;9:509. Hufnagel TC. Nature Mater 2004;3:666. Stratton WG, Hamann J, Perepezko JH, Voyles PM, Mao X, Khare SV. Appl Phys Lett 2005;86:141819. Nieh TG, Wadsworth J. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.04.052.
T.C. Hufnagel Department of Materials Science and Engineering Johns Hopkins University 102 Maryland Hall 3400 North Charles Street Baltimore, MD 21218-2681 United States E-mail address: [email protected] Available online 2 November 2005
Editorial
Preface to the viewpoint set on mechanical behavior of metallic glasses
The lack of long-range atomic order in metallic glasses makes their mechanical properties and behavior considerably different from those of crystalline alloys. Duwez and coworkers announced the formation of the first metallic glass, a gold–silicon alloy, in 1960 [1], but study of mechanical properties required the development of techniques and alloys capable of producing larger and more stable amorphous specimens. By the 1970s, however, rapid solidification techniques for producing amorphous ribbons and wires had been perfected and many of the main features of mechanical behavior, including high strength, large elastic limit, homogeneous and inhomogeneous modes of deformation, and the novel ‘‘molten’’ appearance of fracture surfaces had been established [2,3]. The early metallic glasses had poor glass-forming ability, with critical cooling rates on the order of 105–106 K s 1 which limited the section thicknesses to 50 lm and made load-bearing applications beyond the reach of these alloys. But it was recognized that larger specimens could be produced if formation of crystalline phases could be suppressed. An important success in this area was the development of fluxing techniques for Pd-based glasses [4] which, by eliminating heterogeneous nucleation sites, inhibited crystallization and allowed the production of ‘‘bulk’’ metallic glasses (loosely defined as alloys capable of forming amorphous structures with a minimum section thickness of 1 mm). In 1993, Peker and Johnson announced the discovery of Zr41.2Ti13.8Cu12.5Ni10Be22.5 [5]. This discovery, together with new alloys developed by the Inoue group in Japan [6], spurred world-wide interest in bulk metallic glasses. Among the results of the renewed efforts have been the development of a variety of new bulk glass-forming alloys, including alloys based on Ti [7], Cu [8], Fe [9–11], and Mg [12] (among others) and an improved understanding of alloy design for glass formation. From the standpoint of mechanical properties and behavior, the development of bulk metallic glasses is of interest for two reasons. First, it creates the possibility of using metallic glasses in load-bearing applications, which has long been thought desirable due to the high yield
strength of amorphous alloys. Second, the availability of bulk specimens has enabled a variety of studies of mechanical properties that were difficult or impossible with the ribbon and wire specimens used in the early studies. (It should be pointed out, however, that instrumented indentation techniques have allowed fundamental advances to be made by probing mechanical behavior on the nanoscale, without the need for bulk specimens [13,14].) Thus, now seems to be an appropriate time to take stock of what we have learned, and to lay out visions for the future of metallic glasses, in terms of both practical applications and new research directions. These are the goals of this viewpoint set, and the papers have been selected with them in mind. The general enthusiasm for any new class of materials sometimes leads to grandiose claims for the potential of the materials in question. By now, however, enough data has been gathered on a variety of amorphous alloys that we can attempt to replace the hyperbole with more reasoned speculation. Greer and Ashby provide this in the first article [15] by considering both the strengths and limitations of metallic glasses relative to other materials. They propose that the size of the process zone for plastic deformation (such as ahead of a crack tip) relative to the size of the component in question is an important criterion. The high strength of metallic glasses makes the process zone small (1 mm), suggesting that macroscopically brittle fracture is to be expected in components that exceed this size. However, it may be possible to overcome this limitation, either by increasing the process zone size (by introducing particles [16] or pores [17] that promote shear band initiation) or by effectively decreasing the component dimensions. An example of the latter is a metallic glass foam, in which the ligaments are smaller than the process zone size [18]. Economic considerations ultimately dictate whether new materials see wide-spread use. Some glass-forming alloys involve the use of elements that are either expensive or difficult to handle safely (such as beryllium), or high purity raw materials. However, the development of alloys that are
1359-6462/$ - see front matter Ó 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2005.10.004
318
Editorial / Scripta Materialia 54 (2006) 317–319
based on less expensive elements and more tolerant of impurities make it seem likely that the cost of metallic glasses will continue to decline relative to other highperformance alloys. If so, the processing flexibility afforded by superplastic forming in the supercooled liquid region above the glass transition temperature [19] should allow metallic glasses to compete favorably for use in many more applications. At temperatures well below the glass transition, metallic glasses deform inhomogeneously, with deformation localized into narrow shear bands. Shear localization limits stable plastic flow, and the concentration of slip on shear bands promotes fracture initiation. Thus, controlling shear localization will be critical if metallic glasses are to be widely used in load-bearing applications. Understanding how the atomic-scale mechanisms of deformation lead to shear localization, besides being of scientific interest, may also contribute to design of new alloys with reduced tendency for shear localization. For instance, some amorphous alloys show a tendency for the strain to be distributed over a larger number of shear bands, allowing these alloys to support larger plastic strains in compression [20]. This behavior is not well understood, and in any event the effect is not large enough to avoid abrupt fracture in tension. But it might be possible to develop composite materials in which the second phase promotes shear band initiation and the matrix has a tendency for less severe shear localization. The combination of these effects could lead to materials with significantly improved properties. Given its importance, it is appropriate that several of the articles in this viewpoint set focus on shear localization. Flores [21] and Jiang and Atzmon [22] both consider structural changes associated with inhomogeneous deformation. The positron annihilation measurements discussed by Flores clearly show that inhomogeneous deformation is accompanied by an increase in the number of free volume sites, and calorimetry data on deformed and annealed specimens can also be interpreted in terms of free volume generation and annihilation. Jiang and Atzmon show that least some of this free volume can be found as nanometer-scale voids inside shear bands, although the formation of the nanovoids is not ubiquitous, being sensitive to both the loading condition and the alloy content. In some cases, nanocrystals form inside shear bands instead. The difference is likely important, as nanocrystal formation may stabilize shear deformation and even result in strain hardening, while nanovoids are more likely to contribute to the initiation of cracks. In many load-bearing applications, fracture and fatigue behavior are important. In their paper, Lewandowski and coworkers [23] address the toughness of metallic glasses and metallic-glass-matrix composites (MGMC). They discuss the recent observation that good toughness is associated with a low resistance to shear deformation (quantified by the shear modulus, l) coupled with high resistance to dilatation near a crack tip (quantified by the bulk modulus, B) [24–26]. Thus, intrinsically tough glasses
should have low l/B, a prediction that is borne out by data on a wide range of glasses. A low value of l/B contributes to the distribution of shear strain on multiple shear bands, giving these glasses the ability to support large plastic strains. Lewandowksi also addresses toughness in metallic-glass-matrix composites, and points out that design of the toughening phase (to resist brittle cleavage in body-centered cubic metals, for instance) is likely to be as important as design of the composite microstructure itself. In MGMC materials a crystalline phase is deliberately introduced into an amorphous matrix, with the twin goals of both promoting shear band initiation (thereby distributing the plastic strain more broadly) and hindering shear band propagation. The crystalline particles (or dendrites) are usually both weaker and more ductile than the matrix—just the opposite of conventional metal–matrix ¨ stu¨ndag and coauthors [27] describe the use composites. U of in situ neutron scattering together with self-consistent modeling to study the micromechanics of deformation. A key question of ongoing research is the mechanism by which the particles initiate shear bands in the matrix. Under cyclic loading, metallic glasses tend to show low crack-growth thresholds and low endurance limits compared to other high strength alloys. As Dauskardt and coworkers [28,29] discuss, the poor fatigue properties of single-phase glasses are also the result of the tendency for shear localization, which, under cyclic loading, leads to rapid initiation of fatigue damage. It is possible that alloys with a greater tendency to distribute plastic strain may reduce this effect somewhat, but it seems more likely that substantial improvement of fatigue properties will require introduction of second-phase particles that can help delay the onset of fatigue damage and inhibit fatigue crack growth. Turning to mechanisms of deformation, in his article Spaepen [30] reviews the free volume model for deformation, in which transition state theory is used to describe liquid flow, creep deformation, and structural relaxation. An understanding of these phenomena is important even for low temperature deformation, because the underlying atomic-scale mechanisms are likely the same. Khonik and coworkers [31] address stress relaxation of amorphous alloys in their article. Although SpaepenÕs seminal paper [32] presented the free volume model in terms of single-atom hops, measurements of the activation volume for creep suggest that the actual atomic rearrangements involve something more on the order of 10 atoms. That deformation should require collective rearrangements within ‘‘shear transformation zones’’ (STZ) was proposed by Argon [33], and since that time the STZ theory has seen significant extension, particularly by Langer and coworkers. Langer [34] provides an overview in his paper, including some discussion of shear localization. He presents simulations showing that nascent shear bands can initiate at surface flaws; interestingly, the recent molecular dynamics simulations of Li and Li [35] indicate that surface imperfections are essential to the
Editorial / Scripta Materialia 54 (2006) 317–319
development of shear localization and that without them relatively stable necking can develop during tensile deformation. In their paper, Shi and Falk [36] present the results of some very interesting molecular dynamics simulations showing that shear localization is associated with the disruption of short-range order in the glass. It has long been speculated that the extensive deformation inside shear bands might disrupt short-range order; for instance, this idea was proposed to explain why shear bands are etched preferentially by chemical attack and why existing shear bands reactivate upon removal and reapplication of a load [37]. The difficulty associated with characterization of the atomic-scale structure of the deformed material inside a shear band has, however, precluded any direct confirmation to date. Shi and Falk also suggest that percolation of short-range order over longer length scales affects deformation, with more ordered structures apparently being more resistant to shear localization. This may explain why some alloys develop more distributed shear zones [20,25], as discussed above. Some support for the importance of structure on deformation is provided by the observation that in amorphous Zr–Ta–Cu–Ni–Al enhanced structural order on length scales of 1–1.5 nm is correlated with shear band multiplication and enhanced plastic strain [38,39]. Whether such effects extend to other alloys is an interesting question for future research. Progress in this area is likely to hinge on continued developments in structural characterization of amorphous alloys. Particularly promising in this regard is fluctuation electron microscopy, which has the potential to reveal entirely new details about the nanometer-scale structure of metallic glasses [40,41]. At elevated temperatures (above about 0.7 Tg) metallic glasses deform homogeneously. Nieh and Wadsworth [42] review homogeneous deformation, with an emphasis on the importance of understanding how structural changes that may occur—particularly nanocrystallization—affect the flow behavior of the glass. Above the glass transition temperature extremely large deformations can be achieved at low stresses, making this regime of great interest for processing of metallic glasses into complex shapes. Superplastic forming above Tg is not covered in this set of articles, having been the subject of a recent review by Schroers [19]. In summary, I believe that this viewpoint set provides a useful snapshot of the current state of knowledge of mechanical behavior of metallic glasses. I fully expect that the future will see continued advances both in our understanding of fundamental aspects of mechanical behavior, as well as new applications of metallic glasses and metallic-glass-matrix composites. References [1] Klement W, Willens RH, Duwez P. Nature 1960;187:869. [2] Pampillo CA, Polk DE. Acta Metall 1974;22:741. [3] Masumoto T, Maddin R. Mater Sci Eng 1975;19:1.
[4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
[32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42]
319
Kui HW, Greer AL, Turnbull D. Appl Phys Lett 1984;45:615. Peker A, Johnson W. Appl Phys Lett 1993;63:2342. Inoue A. Acta Mater 2000;48:279. Zhang T, Inoue A. Mater Trans JIM 1998;39:1001. Inoue A, Zhang W, Zhang T, Kurosaka K. Acta Mater 2001;49:2645. Shen T, Schwarz R. Appl Phys Lett 1999;75:49. Ponnambalam V, Poon SJ, Shiflet GJ, Keppens VM, Taylor R, Petculescu G. Appl Phys Lett 2003;83:1131. Lu ZP, Liu CT, Thompson JR, Porter WD. Phys Rev Lett 2004;92:245503. Inoue A, Nakamura T, Nishiyama N, Masumoto T. Mater Trans JIM 1992;33:937. Schuh CA, Nieh TG. J Mater Res 2004;19:46. Schuh CA, Lund AC, Nieh TG. Acta Mater 2004;52:5879. Ashby MF, Greer AL. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.051. Hays CC, Kim CP, Johnson WL. Phys Rev Lett 2000;84:2901. Wada T, Inoue A, Greer A. Appl Phys Lett 2005;86:251907. Brothers AH, Dunand DC. Acta Mater 2005;53:4427. Schroers J. JOM 2005;57:35. Xing LQ, Li Y, Ramesh KT, Li J, Hufnagel TC. Phys Rev B 2001;64:180201(R). Flores KM. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.04.049. Jiang WH, Atzmon M. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.052. Lewandowski JJ, Shazly M, Nouri AS. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.010. Chen HS, Krause JT, Coleman E. J Non-Cryst Solids 1975;18:157. Schroers J, Johnson W. Phys Rev Lett 2004;93. Lewandowski J, Wang W, Greer A. Philos Mag Lett 2005;85:77. ¨ stu¨ndag E, Kim CP, Brown DW, Bourke Clausen B, Lee SY, U MAM. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.009. Hess PA, Cameron KK, Menzel BC, Dauskardt RH. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.006. Cameron KK, Hess PA, Menzel BC, Dauskardt RH. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.007. Spaepen F. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.09.046. Bobrov OP, Csach K, Khonik VA, Kitagawa K, Laptev SN, Yazvitsky MY. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.09. 056. Spaepen F. Acta Metall 1977;25:407. Argon A. Acta Metall 1979;27:47. Langer JS. Scripta Mater 2006;54, doi:10.1016/j.scripta.2005.10.005. Li Q, Li M. Appl Phys Lett 2005;87. Shi Y, Falk ML. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.09.053. Pampillo CA. J Mater Sci 1975;10:1194. Hufnagel TC, Brennan S. Phys Rev B 2003;67:014203. Li J, Gu X, Hufnagel TC. Microsc Microanal 2003;9:509. Hufnagel TC. Nature Mater 2004;3:666. Stratton WG, Hamann J, Perepezko JH, Voyles PM, Mao X, Khare SV. Appl Phys Lett 2005;86:141819. Nieh TG, Wadsworth J. Scripta Mater 2006;54, doi:10.1016/ j.scripta.2005.04.052.
T.C. Hufnagel Department of Materials Science and Engineering Johns Hopkins University 102 Maryland Hall 3400 North Charles Street Baltimore, MD 21218-2681 United States E-mail address: [email protected] Available online 2 November 2005