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Metallic Glasses Formation and Properties Edited by Behrooz Movahedi

Metallic Glasses: Formation and Properties Edited by Behrooz Movahedi

Published by ExLi4EvA Copyright © 2016 All chapters are Open Access distributed under the Creative Commons Attribution 3.0 license, which allows users to download, copy and build upon published articles even for commercial purposes, as long as the author and publisher are properly credited, which ensures maximum dissemination and a wider impact of our publications. After this work has been published, authors have the right to republish it, in whole or part, in any publication of which they are the author, and to make other personal use of the work. Any republication, referencing or personal use of the work must explicitly identify the original source. As for readers, this license allows users to download, copy and build upon published chapters even for commercial purposes, as long as the author and publisher are properly credited, which ensures maximum dissemination and a wider impact of our publications. Notice Statements and opinions expressed in the chapters are these of the individual contributors and not necessarily those of the editors or publisher. No responsibility is accepted for the accuracy of information contained in the published chapters. The publisher assumes no responsibility for any damage or injury to persons or property arising out of the use of any materials, instructions, methods or ideas contained in the book.

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ISBN-10: 953-51-2512-5 ISBN-13: 978-953-51-2512-9 Print ISBN-10: 953-51-2511-7 ISBN-13: 978-953-51-2511-2

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Contents

Preface Chapter 1 Metallic Glasses from the Bottom-up by Aras Kartouzian and Jerzy Antonowicz

Chapter 2 Structural and Dynamical Properties of Metallic Glassy

Films by Hui Li, Weikang Wu and Kun Zhang

Chapter 3 Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level by HaiJun Chang Chapter 4 Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses by Shanlin Wang Chapter 5 Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification by Celal Kursun and Musa Gogebakan Chapter 6 Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites by Devinder Singh, R.K. Mandal, R.S. Tiwari and O.N. Srivastava Chapter 7 On the Prospects of Using Metallic Glasses for Invessel Mirrors for Plasma Diagnostics in ITER by Vladimir S. Voitsenya, Alexandra F. Bardamid, Martin Balden, Flaviu Gostin, Sergey V. Khovrich, Vladimir G. Konovalov, Konstantin V. Kovtun, Petro M. Lytvyn, Sergey V. Ketov, Dmitri V. Luzguine-Luzgin, Sergei I. Solodovchenko, Anatoly N. Shapoval, Anatoly F. Shtan’, Vladislav N. Bondarenko, Ivan V. Ruzhkov, Ol’ga O. Skoryk and Andrei A. Vasil’ev

Preface Metallic glasses and amorphous materials have attracted much more attention in the last two decades. A noncrystalline solid produced by continuous cooling from the liquid state is known as a glass. From the other point of view, a noncrystalline material, obtained by any other process, for example, vapor deposition or solid-state processing methods such as mechanical alloying, but not directly from the liquid state, is referred to as an amorphous material. At this moment, bulk metallic glasses (BMG) are appearing as a new class of metallic materials with unique physical and mechanical properties for structural and functional usage. Extreme values of strength, fracture toughness, magnetic properties, corrosion resistance, and other properties have been registered in BMG materials.

Chapter 1

Metallic Glasses from the Bottom-up Aras Kartouzian and Jerzy Antonowicz Additional information is available at the end of the chapter http://dx.doi.org/10.5772/63514

Abstract The main challenμe in understandinμ the relation between the structure and properties oλ metallic μlasses is describinμ their structure at the atomic level. Currently, their structures are considered simply disordered and indeed our understandinμ oλ their structure is as undeλined as this term. Followinμ the most advanced structural models oλ metallic μlasses that are based on metal clusters, a bottom-up approach to λabrication oλ metallic μlasses usinμ cluster beam technoloμy is introduced. Usinμ metal clusters to λabricate metallic μlasses λrom the bottom-up, that is, λormation oλ cluster-assembled metallic μlasses, provides us with the possibility oλ varyinμ their structure at the atomic level while keepinμ their composition unchanμed. “ unique λeature workinμ with clusterassembled metallic μlasses is the independent control oλ their structure and composition. The advantaμes oλ this approach are presented, and its potential toward the resolution oλ structure–property puzzle in metallic μlasses is demonstrated alonμ with the main challenμes. Keywords: Cluster-assembled metallic μlasses, metal clusters, structure–property re‐ lation, cluster deposition, cluster-based structural models

. Introduction First discovered in by rapidly quenchinμ > Ks− an alloy oλ “u Si at.% [ ], metallic μlasses MGs are amonμ the most studied metallic materials. The non-periodic character oλ MGs underlies their unique properties which are oλten superior to conventional crystalline materials. Due to their reduced eddy current losses, as compared to the crystalline alloys oλ identical compositions, λor instance, λerromaμnetic MGs are commonly used as maμnetic core materials [ ]. “lso, the corrosion resistance oλ iron-based metallic μlasses was shown to be

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Metallic Glasses: Formation and Properties

much hiμher than that oλ crystalline stainless steel [ ]. “s another example, the Ti Cu Pd Zr metallic μlass is a biocompatible material that is about three times stronμer than titanium, has an elastic modulus that matches that oλ bone, and does not produce abrasion powder [ ]. Further, the combination oλ hiμh touμhness and hiμh strenμth in Pd-based MGs puts them amonμ the stronμest and most damaμe-tolerant materials ever known [ ]. “s oλ now, the main challenμe in investiμatinμ these materials is to describe their structure at an atomic level. In absence oλ an atomic description, no systematic desiμn oλ MGs has been possible, and the proμress in the λield is merely based on the costly and ineλλicient procedure oλ trial and error. The pioneerinμ work by ”ernal [ , ] on the structure oλ metallic liquids, who suμμested dense random packinμ oλ hard spheres as the structural model, was the λirst step in this direction. Further research in the λield has led to the discovery oλ many other metallic alloys that could be solidiλied into the amorphous state with moderate coolinμ rates – Ks − such as Pd Ni P μlass λor which bulk sections oλ mm across were produced at a coolinμ rate oλ Ks− [ ]. ”ased on the observations across the compositions oλ the MGs, Inoue put λorward a set oλ empirical criteria λor their λormation and stability [ ]. This hiμhly valuable classiλication accelerated the discovery oλ new μlass λorminμ alloys. “s a result, very soon a revision oλ the criteria was required [ ]. Such criteria that have been proven very helpλul in desiμninμ new μlasses, however, naturally suλλer λrom numerous exceptions. For instance, based on the binary phase diaμram oλ “u–“l alloy and the relative atomic radii oλ μold and alumi‐ num, it is expected to be possible to produce an “lx“u –xMG. However, to date, no one has succeeded in the production oλ MGs in this alloy system reμardless oλ the employed techni‐ ques [ ]. Despite the intensive research in the λield oλ MGs, the understandinμ oλ the λundamental link between their structure and properties is still missinμ [ – ]. Theoretical computations have made a larμe contribution toward our understandinμ oλ the structure oλ MGs but are damned to be inaccurate due to their restricted timescale, which imposes coolinμ rates that are many orders oλ maμnitude hiμher than what is experimentally achievable beside the ultraλast liquid quenchinμ reported by Mao et al. [ ] with a coolinμ rate oλ ∼ Ks− . Currently, there exists no perspective to solvinμ this issue without external help λrom other disciplines. Currently, the structure oλ MGs is addressed as disordered and our understandinμ oλ it is in λact as diλλuse and undeλined as this term. Oλ course, considerable proμress has been made in the past toward describinμ the structure oλ MGs at the atomic level, but we are still λar λrom havinμ a coherent and consistent model. “ssuminμ we have the correct model to describe the structure oλ MGs, the next step would be to λiμure out the structure–property relation. Obviously in the absence oλ periodicity, it will still be much more complicated to develop this relation in MGs compared to oxide μlasses considerinμ that bond anμles and lenμths are much more λlexible and distortable in MGs. In this chapter, we introduce some oλ the most advanced experimental approaches to tackle these issues usinμ cluster beam technoloμy. “ccordinμly, this chapter aims to put λorward an interdisciplinary approach and λamiliarize the material scientists workinμ on MGs with cluster beam technoloμy and how it can be used.

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

. Cluster-based structural models for MGs “lthouμh amorphous alloys lack lonμ ranμe order, they possess well deλined nearest neiμhbor shells. The resultinμ short and medium ranμe order is experimentally observed in MGs. The latest structural models λor MGs [ – ] take this observation into account and use atomic clusters to describe the structure oλ amorphous alloys. “ structure model based on eλλiciently packed solute-centered atomic clusters was introduced by Miracle. In this model, atomic clusters are idealized as spherical particles, which similar to atoms, λill the space in λacecentered cubic λcc or hexaμonal close-packed hcp arranμements [ ]. However, unlike atoms, atomic clusters can overlap and share atoms with each other. This model has μained some credit because λirstly it is consistent with a broad ranμe oλ previously established μuidelines λor metallic μlasses, and secondly, it has a predictive capability λor the experimen‐ tally observed medium-ranμe order in MGs. Shortly aλter the presentation oλ Miracle’s model, a revised version was suμμested by Fan et al. [ ] where the buildinμ blocks are aμain atomic clusters, but are arranμed randomly instead. Later, the same μroup reλined their model based on reverse Monte Carlo simulations and introduced the tiμht-bond cluster model , which includes the clusters, the λree volume between the clusters, and the interconnectinμ zones amonμ clusters [ ]. In a closely related approach, Donμ et al. [ ] introduced the clusterplus-μlue-atom model , where the structure oλ the MG is described by speciλic metal clusters that are μlued toμether by additional μlue atoms. The recent review by Liu and Zhanμ [ ] provides a concise summary oλ structural models λor MGs. Figure presents a μraphical summary oλ the cluster-based models.

Figure . Planar representations oλ cluster-based structure models λor metallic μlasses. Leλt “ Zr Cu μlass consist‐ inμ oλ Zr Cu clusters arranμed in λcc structure [ ]. Middle “ Zr Cu μlass consistinμ oλ randomly arranμed tiμht‐ ly bonded Zr Cu clusters. Riμht “ Zr Cu “l μlass consistinμ oλ randomly arranμed Zr Cu clusters μlued toμether by “l atoms.

In this context atomic clusters are small particles consistinμ oλ up to

atoms.

Notation oλ clusters “μm indicates a clusters consistinμ oλ m “μ atoms “μ indicates a cluster made out oλ “μ atoms. In the same way, Zr Cu indicates a cluster made out oλ Zr and Cu atoms.

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Metallic Glasses: Formation and Properties

“ll these cluster-based models allow short- and medium-ranμe order up to distances oλ λew cluster sizes , while the disordered nature oλ MGs on the lonμ-ranμe is retained due to local atomic stresses and topoloμical λrustration. Cluster-based models have been employed by various μroups λor the interpretation oλ their experimental and theoretical results. Probably, the best example was delivered by Hirata et al. [ ] throuμh nanobeam electron diλλraction experiments perλormed on rapidly quenched Zr . Ni . MGs in combination with ab initio molecular dynamics simulations. They have shown that sub-nanoscale-ordered reμions can produce distinctly symmetric electron diλλrac‐ tion patterns that oriμinate λrom individual and interconnected atomic clusters as buildinμ blocks oλ MGs. Cluster-based structural models have improved our understandinμ oλ MGs to a μreat extent the positive observations are exclusively limited to searchinμ λor and λindinμ oλ cluster units in MGs and occasionally relatinμ the overall composition oλ MGs to the composition oλ observed clusters. One deλinite knowledμe that has emerμed as the result oλ cluster-based structural models is, however, that MGs indeed belonμ to the cateμory oλ cluster-assembled materials C“Ms . “s such, it should be possible to λabricate metallic μlasses by puttinμ appropriate clusters toμether. This approach, which has been neμlected till quiet recently λor practical reasons as it will be outlined below, is the subject matter oλ this chapter.

. The bottom-up approach to MGs In order to veriλy the appositeness oλ cluster-based structural models λor MGs, which suμμest that metal clusters are the buildinμ blocks oλ MGs, their λabrication by deposition oλ select‐ ed metal clusters to λorm cluster-assembled metallic μlasses C“MGs was recently proposed [ ]. In the λollowinμ sections, we will μo into some details about what metal clusters are, and how are they synthesized. Despite the λact that C“MGs are still at a very early staμe oλ their development, they make up the core oλ this chapter, because they are expected to contribute larμely to our understandinμ oλ amorphous structure oλ MGs at the atomic level and also help to decipher the structure–property eniμma. Generation, selection, and deposition oλ metal clusters are all amonμ the most advanced disciplines oλ material science. The current state-oλ-the-art only allows λor the λabrication oλ C“MG samples in λorm oλ thin λilms. This temporary technical limitation, which will probably accompany us λor another decade, brinμs C“MGs very close to thin-λilm metallic μlasses TFMGs that also have attracted interest [ ]. TFMGs are also λabricated in a bottom-up approach and thus are included in this chapter. However, they will not be at the spot liμht here, mainly due to the λollowinμ two reasons Firstly, in TFMGs, the buildinμ blocks cannot be actively altered and controlled as they are always atoms or an undeλined distribution oλ clusters. Consequently, it is not possible to C“Ms are materials that are λabricated by assemblinμ atomic clusters i.e. have atomic clusters as their buildinμ blocks. In this context and throuμhout the text, selection reλers to mass-selection separation oλ clusters based on their mass.

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

actively inλluence the structure oλ TFMGs at the atomic level. Secondly, the composition and the buildinμ blocks oλ the λilms cannot be modiλied independently, so that a correlation between the buildinμ blocks and the λilm properties cannot be established. Nanoμlasses are another class oλ metallic μlasses that are closely related to C“MGs. Introduced by Gleiter et al. [ , – ], nanoμlasses are μenerated by sputterinμ or evaporatinμ the material oλ choice, and subsequently consolidatinμ the λormed μlassy droplets into a pellet-shaped sample. Here, only a very vaμue control on the structure and composition oλ the droplets may be achieved. There are number oλ published works on nanoμlasses, which deal with them in appropriate details [ , , ]. . . Cluster-assembled metallic glasses, CAMGs ”uildinμ blocks oλ C“MGs are metal clusters. In this section, we address μeneration, selection, and deposition oλ metal clusters to λorm C“MGs. Metallic clusters can be μenerated in metal cluster sources, which will be described brieλly in Section . . . The output oλ a cluster source is a distribution oλ neutral and charμed clusters, and thus, a selection step Section . . is required to pick out the desired clusters beλore deposition. Finally, the selected clusters should be deposited on to a support material in order to λabricate C“MG, as will be explained in Section . . . The three main steps oλ C“MG λabrication are schematically summarized in Figure .

Figure . ”ottom-up approach to nanoλabrication oλ metallic μlasses. Leλt mixed metal clusters are μenerated by laser vaporization oλ a metal alloy tarμet. Middle usinμ mass selection, a speciλic cluster is picked out oλ the cluster beam. Riμht mass-selected clusters are deposited on a support material to λorm a metallic λilm.

. . . Generation of metal clusters The development oλ cluster sources and subsequently the investiμation oλ clusters, started back in the s with the idea oλ utilizinμ the non-equilibrium conditions oλ an adiabatically expandinμ vapor, λor example, by supersonic expansion oλ a μas into vacuum [ , ]. Cluster Supersonic expansion is achieved when a μas expands into vacuum with a Mach number larμer than unity.

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Metallic Glasses: Formation and Properties

λormation is believed to be due to the supersaturation oλ rapidly cooled vapor which stimulates homoμenous nucleation in the beam [ ]. There are various possibilities to produce atomic clusters λrom bulk materials. Common to almost all these methods is that atoms are λirstly ejected λrom the bulk material and then are brouμht toμether to λorm clusters in the μas phase. “ review oλ all types oλ cluster sources is beyond the scope oλ this chapter. Only the laser vaporization cluster source will be introduced here in more detail as this is the only source that has ever been employed λor λabrication oλ C“MGs [ ]. The λirst μeneration oλ a laser vaporization cluster source was reported in the early s, at Rice University by Smalley et al. [ ]. Many variants oλ this cluster source have emerμed since then. The use oλ lasers λor ablation oλ material is a very important λeature oλ this kind oλ cluster source, since it allows λor the production oλ metal vapors oλ even the most reλractory metals such as W and Mo without overheatinμ any part oλ the apparatus. The supersonic expansion oλ the cluster beam is the other important λeature oλ this source also common amonμ other cluster sources . “ schematic view oλ a laser vaporization cluster source is shown in Figure .

Figure . “ schematic view oλ the laser vaporization cluster source is illustrated. “ laser beam is λocused on to a metal tarμet either pure metal, or an alloy in the presence oλ hiμh pressure oλ a buλλer inert μas. The plume is mixed with the μas and underμoes multiple collisions prior to expansion into vacuum.

The laser vaporization cluster sources produce clusters in the size ranμe λrom two to several hundreds oλ atoms per clusters. The vaporized material is cooled by collisions with inert μas atoms which μreatly outnumber the ablated metal atoms. The λormation oλ clusters requires three-body collisions between two metal atoms and a rare μas atom in the case oλ dimer λormation , such that the rare μas atom can take the collision enerμy away in the λorm oλ its kinetic enerμy and thus make it possible λor the metal atoms to stick toμether without violatinμ the conservation oλ momentum. This process then needs to be repeated many times to λorm

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

larμer clusters and that is why hiμher pressures oλ the rare μas are required λor the μeneration oλ larμer clusters. The λormation oλ larμer clusters can then proceed by either the addition oλ sinμle atoms to smaller clusters or by merμinμ smaller clusters toμether. In most cluster beam λacilities, clusters traverse a skimmer aλter leavinμ the nozzle. The main λunction oλ a skimmer is to collimate the expanded μas mixture that contains the clusters, into a directed cluster beam [ – ]. The cluster beam is then μuided λurther to the mass selection unit beλore it is deposited. . .2. Cluster selection The cluster beam that leaves the source contains neutral clusters as well as neμatively and positively charμed ones. The ion optics used to μuide the cluster beam is set to μuide either the positive or the neμative ions, but it cannot inλluence the neutral particles. For instance, iλ neμatively charμed ions are excluded λrom the cluster beam throuμh the ion optics, the cluster beam will consist oλ positively charμe ions that are actively μuided and neutral clusters that λly in the same direction. In order to exclude the neutrals λrom the beam, it is common practice to include an electrostatic bender to deλlect the charμed cluster beam while the neutral beam will not be aλλected. This separation step is a crucial prerequisite λor cluster selection as electrical and maμnetic mass λilters cannot interact with neutral particles and thus are not able to distinμuish amonμ diλλerent neutral clusters. “λter exitinμ the deλlector, the beam oλ positively charμed clusters is λurther μuided to a mass λilter, commonly a quadruple mass spectrometer, which combines DC and radio λrequency “C voltaμes to select a speciλic cluster mass λrom the cluster beam. This is the selection step where a sinμle cluster mass or a collection oλ masses are selected λor deposition. The selection criterion oλ a quadruple mass spectrometer is the mass-to-charμe ratio oλ the clusters the voltaμes can be set to make the trajectories oλ clusters that are heavier or liμhter than a set mass instable and thus exclude them λrom the cluster beam. This selects only those clusters that have a mass within the set mass window while discardinμ all the others. The width oλ the mass window can be controlled, and thus, the mass resolution oλ the device can be adjusted. In μeneral, the mass resolution is set to be just hiμh enouμh to separate adjacent masses since the transmission oλ mass λilters decreases with increased mass resolution. The intensity oλ a cluster beam is commonly evaluated by measurinμ the current that is caused by charμed clusters in the beam. To this end, a Faraday cup or alternatively a metal plate is used to collect the clusters and the current λlowinμ λrom the collector to μround upon arrival oλ clusters is measured over time. The perλormance oλ a cluster source can be determined λrom its cluster distribution by recordinμ a mass spectrum. Figure presents mass spectra λor pure “μ and CuZr alloy clusters μenerated by a laser vaporization cluster source. The number oλ clusters in the cluster beam can be deduced λrom the measured current, dividinμ the cluster beam current by the elementary charμe will μive the number oλ sinμly charμed clusters that have been detected over s. For instance, a cluster current oλ p“ translates into ∼ clusters Named aλter Michael Faraday, a Faraday cup is an electrically conductive cup-shaped plate , which is used to collect charμed particles under vacuum conditions.

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Metallic Glasses: Formation and Properties

in s. For a laser vaporization source with a repetition rate oλ Hz laser pulses in a second , this means million clusters in each sinμle laser pulse. “lthouμh these numbers may sound larμe, much hiμher cluster currents are required λor deposition purposes, as will be explained in the next section.

Figure . Mass spectra oλ pure silver clusters top and mixed CuZr cluster bottom μenerated by a laser vaporization cluster source are shown. The mass resolution oλ the quadrupole mass λilter is identical λor both cases. While in case oλ silver clusters, the ion peaks are clearly separated, λor CuZr clusters peaks cannot be resolved due to the overlap be‐ tween the masses oλ mixed clusters plus the λact that Zr with λour and Cu with two naturally stable isotopes λurther broaden the spectrum. In the top, “μm+ clusters with m = – are observed. The hiμher intensity oλ clusters with odd number oλ atoms is due to their hiμher stability based on their electronic structure. Such odd-even stability oscillation is common to s metals. In the bottom, mixed ZrnCum+ clusters oλ various compositions are observed while the spec‐ trum is dominated by pure Zr clusters. Two series oλ ZrnCum+, m = – clusters are assiμned. The Zr + cluster peak is clearly more intense than other clusters in the mass spectrum. Here, the μeometry oλ the icosahedral clusters with atoms is the stabilizinμ λactor.

Notation oλ clusters “μm indicates a clusters consistinμ oλ m “μ atoms “μ indicates a cluster made out oλ “μ atoms. In the same way, Zr Cu indicates a cluster made out oλ Zr and Cu atoms.

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

. . . Cluster deposition “λter cluster selection, the selected clusters should be deposited onto a support material. Collision with surλace may lead to λraμmentation oλ the clusters and render the selection step obsolete. “ccordinμly, special care should be taken to achieve soft-landing conditions when depositinμ clusters. The material used as support is, thereλore, very important. Iλ a conductive support material is used, soλt-landinμ can be easily achieved by applyinμ a voltaμe to the support to slow down the arrivinμ ions to kinetic enerμies lower than . eV per atom lower than eV λor a cluster consistinμ oλ atoms or lower than eV λor a cluster consistinμ oλ atoms and so on . Iλ on the other hand a non-conductive material is used, no adjustinμ voltaμe can be applied and the enerμy oλ the ion beam should be adjusted by tweakinμ the ion optics so that the kinetic enerμies oλ the ions are low enouμh to μuarantee soλt-landinμ. “nother issue reμardinμ the conductivity oλ the support material is related to the charμe oλ the clusters. “s discussed in Section . . , clusters are selected based on their mass to charμe ratio and thus only charμed clusters are suited λor mass selection. While the ions will lose their charμe once deposited onto a conductive surλace, they will keep or only partially lose their charμe while in contact with an insulator surλace. Consequently, a neutralization mechanism is required to avoid interruption in deposition due to electrostatic repulsion. Due to this eλλect, almost exclusively positively charμed clusters are used λor deposition purposes, since they can be neutralized by an electron beam, whereas a proton beam would be required λor neutrali‐ zation oλ neμatively charμed clusters. The other aspect reμardinμ the choice oλ the support material is its atomic structure, that is, whether it is crystalline or amorphous. “morphous supports are preλerred because they will not provide periodic nucleation cites and thus will not promote the rearranμement oλ clusters and crystallization oλ the λilm. However, as it will be shown in Section . ., havinμ an amor‐ phous substrate complicates the structural characterization oλ C“MGs. The Support material may also be cooled down in order to suppress diλλusion oλ clusters and to enhance the μlass λorminμ probability by stoppinμ the clusters λrom underμoinμ larμe μeometrical deλormations. “s already mentioned in the previous section, hiμh cluster currents are required λor deposition oλ C“MGs. To λurther illuminate this issue, we may use the λollowinμ example “ cluster beam can be λocused down to a round spot with a diameter oλ . cm. Usinμ such a beam to deposit clusters on to a support will result in a coated area oλ ∼ . cm , the so-called cluster spot. “n icosahedral cluster oλ atoms one oλ the larμest clusters relevant λor MGs will have a diameter oλ ∼ nm and thus cover an area oλ ∼ . nm . Fillinμ a sinμle layer -nm-thick λilm oλ the cluster spot with the total area oλ . cm , with such clusters will demand ∼ × clusters λor this estimation, the λree space between touchinμ spherical clusters was neμlected which leads to ∼ % overestimation oλ the number oλ clusters . In order to μet a -μm-thick λilm, at least times more clusters should be deposited. Now assuminμ a cluster beam current oλ p“, about h oλ deposition will be required. Obviously, an enhancement oλ at least an order oλ maμnitude in cluster beam current is necessary to have a λeasible deposition time. Such hiμh cluster beam currents are above what can currently be achieved usinμ stateoλ-the-art laser vaporization cluster sources, and consequently, the only reported C“MGs to date have used a relatively broad collection oλ clusters and not an absolutely selected beam oλ

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a sinμle cluster mass. Figure components.

depicts a cluster deposition λacility incorporatinμ all its

Figure . “ schematic view oλ a cluster deposition λacility is depicted. The λacility is composed oλ a laser vaporization cluster source μrey shade , set oλ ion optics beλore and aλter mass selection μreen shade , mass selector yellow shade , and a deposition chamber unshaded . The deposition chamber is λurther equipped with a sputter μun λor cleaninμ the surλace oλ the support material prior to deposition, a rest μas analyzer RG“ λor monitorinμ the quality oλ the vacuum in this chamber, and a transλer chamber λor sample handlinμ purposes such as removinμ the sample λrom this λacility λor transport to other λacilities λor analysis and characterization. The dashed oranμe line shows the path oλ the cluster beam λrom the cluster source throuμh the ion optics and the mass λilter down to the deposition chamber where clusters are soλt-landed onto the support material.

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

. . Thin-film metallic glasses TFMGs Driven to examine the relation between the heat oλ hole λormation and crystallization temper‐ ature in amorphous alloys, Nastasi et al. [ ] λabricated probably the λirst thin-λilm metallic μlass in binary systems oλ Cu–W and Cu–Ta just couple oλ years aλter the very λirst TFMG in La–“u system achieved by solid-state amorphization, where it was also showed that deep eutectics are not a necessary criterion λor μlass λormation in metallic alloys [ ]. Currently, TFMGs are commonly λabricated by co-deposition oλ multiple metals either λrom an alloy tarμet or multiple tarμets where the λlux oλ each component can be controlled separately. Zrbased alloys are the most studied systems to date. Interest in TFMGs is λueled by their distinct properties even compared to counterpart MGs, such as broader μlass λorminμ ranμe and hiμher strenμth [ , – ]. Their broad and contin‐ uous μlass λorminμ ranμe leads to tunability oλ their properties by simply adjustinμ their composition [ ]. Their potential application as bio-coatinμs λurther increases their relevance. Recently, the corrosion resistance oλ binary Zr–Ni and Zr–Co TFMGs was investiμated [ ]. “lthouμh TFMG enjoys a μreat technical siμniλicance, they will not play a momentous role in unravelinμ the structure–property puzzle in MGs. In the case oλ TFMGs, the μas phase entities used λor λabrication oλ metallic λilms are mainly atoms and not atomic clusters. Even iλ some clusters are available in the deposition beam, up to now no control on the structure and composition oλ the buildinμ blocks could be achieved. Moreover, beside the case oλ atomic deposition, no inλormation on the properties oλ the structural units that build up the λilms has been accessible. In contrast, in the case oλ C“MGs, atoms are very oλten deliberately excluded λrom the deposition beam and only clusters are used to build up a metallic λilm. The main advantaμe oλ C“MGs over TFMGs is that the buildinμ blocks can be altered while keepinμ the composition oλ the resultant metallic λilm unchanμed.

. Structure–property relation in CAMGs In previous sections, C“MGs were introduced and their λabrication usinμ cluster beam technoloμy was described. Here, we will have a closer look on how the study oλ C“MGs serves to explore the structure–property relation in MGs. Usinμ the example oλ Zr Cu alloy, we will demonstrate this capability. Employinμ the aλorementioned cluster beam technoloμy, a Zr Cu λilm can be λabricated usinμ many diλλerent combinations oλ various clusters as buildinμ blocks. The possibility oλ havinμ the same composition composed oλ diλλerent buildinμ blocks is a unique λeature oλ C“MGs. Investiμatinμ the properties oλ these λilms in comparison with a rapidly quenched MG oλ the same composition will be a key step in understandinμ the structure oλ amorphous alloys and thus the structure–property relation in amorphous metals. On the other hand, comparinμ the properties oλ C“MGs with the proper‐ ties oλ metal clusters used λor their λabrication will reveal the nature oλ interactions amonμ metal clusters when λorminμ an extended material. Furthermore, lookinμ at the local structure

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Metallic Glasses: Formation and Properties

oλ C“MGs will unravel the extent oλ structural deλormation that metal cluster underμo as buildinμ blocks oλ C“MGs. Figure summarizes the diλλerent aspects oλ this approach.

Figure . The scientiλic approach oλ utilizinμ C“MGs is presented. “λter λabrication oλ C“MGs λrom speciλic metal clusters as buildinμ blocks, their structure and properties should be determined. On the one hand comparinμ the prop‐ erties oλ C“MGs with the properties oλ rapidly quenched MGs with identical composition will provide us with inλor‐ mation on the structure oλ rapidly quenched MGs in relation to the metal clusters the closer the properties oλ C“MGs and MGs, the closer their structures are! On the other hand, comparinμ the properties oλ C“MGs with the properties oλ their constituent metal clusters will reveal the nature oλ inter-cluster interactions. Further, by comparinμ the local atomic structure oλ C“MGs with the structure oλ their constituent metal clusters, the deμree oλ deλormation and stabili‐ ty oλ metal clusters while perλorminμ as buildinμ blocks oλ MGs can be deduced.

It will not be practicable to consider all the possible combinations. Instead, let us choose some representative combinations to provide a more clear idea about the beneλits oλ C“MGs. Consider the λollowinμ λour scenarios .

the λilm is λabricated by deposition oλ equal number oλ Zr and Cu clusters.

.

the λilm is λabricated by deposition oλ equal number oλ Zr Cu and Zr Cu clusters.

.

the λilm is λabricated by deposition oλ a set oλ clusters in the mass ranμe between and amu, includinμ pure Zr and Cu clusters with more than and atoms, respectively,

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

and all the mixed clusters with a mass within that ranμe, so that the overall stoichiometry oλ the λilm remains is not violated “toms and smaller clusters are deliberately excluded . .

the λilm is λabricated by deposition oλ Zr Cu clusters.

For all oλ the above-mentioned hypothetical λilms, the composition is the same however, they will possess diλλerent atomic structures unless the buildinμ blocks are stronμly deλormed and are hiμhly λlexible in sharinμ atoms amonμ each other. Considerinμ the λour diλλerent scenarios introduced earlier, only the third scenario usinμ a set oλ cluster within a mass ranμe has been realized experimentally [ ] and will be presented in the next section as a prooλ oλ principle. . . Zr–Cu CAMGs: the first steps The λirst attempt to apply cluster beam technoloμy λor λabrication oλ metallic μlasses was undertaken recently usinμ binary Zr–Cu alloys [ ]. The justiλication λor this choice is threeλold. Firstly, Zr–Cu binary system shows hiμh μlass λorminμ abilities GF“ in a wide ranμe oλ compositions [ – ]. Secondly, a larμe body oλ literature on Zr–Cu MGs and TFMGs exists, which proved to be essential in interpretation oλ the experimental observations [ – ]. “nd thirdly, cluster μeneration and cluster selection oλ mixed metal clusters become more diλλicult with increasinμ the number oλ elements in the cluster, and thus, a binary system is the loμical startinμ point. Metallic Zr Cu λilms were λabricated by deposition oλ a set oλ clusters havinμ masses between and amu scenario Nr. on silicate μlass substrates, under ultrahiμh vacuum, and soλt-landinμ conditions. Films oλ various thicknesses ranμinμ λrom to nm were pro‐ duced. “t this staμe, the structure and properties oλ the clusters used to assemble the metallic λilm are not available. However, all properties oλ metal clusters can be obtained in state-oλ-theart cluster laboratories includinμ their optical, maμnetic, chemical, catalytic, electronic, and structural properties. The λirst question to be answered while λabricatinμ C“MGs λor the λirst time is, however, whether the synthesized λilm is in an amorphous state at all. Surλace X-ray diλλraction at the European Synchrotron Radiation Facility ESRF was used to answer this question. The details oλ the experiments and sample preparation are available elsewhere [ ]. Here, the most important λindinμs underpinninμ this bottom-up approach are brieλly sum‐ marized. Figure depicts the diλλraction pattern oλ the λirst cluster-assembled Zr–Cu MG. Despite the interλerence caused by the broad diλλraction band oλ the silicate μlass support, a clear halo peak, which corresponds to a scatterinμ vector oλ . nm− , is observed in excellent aμreement with literature values λor Zr Cu MGs determined λrom hiμh enerμy XRD, neutron diλλraction, and extended X-ray absorption λine structure EX“FS spectroscopy [ ]. “ccordinμly, it could be unambiμuously concluded that the λabricated cluster-assembled metallic λilm is in a μlassy state. This observation alone is suλλicient to prove the practical λeasibility oλ employinμ cluster beam technoloμy to λorm C“MGs. The μood aμreement between the observed scatterinμ vector and the literature value may suμμest that the structure oλ the Zr–Cu C“MG is very close to that oλ rapidly quenched samples. However, it should be noted that the position oλ the λirst XRD halo is not very sensitive

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Metallic Glasses: Formation and Properties

to the atomic structure oλ amorphous metals [ ]. In λact, there is almost no diλλerence in peak position λor amorphous solids and their correspondinμ liquids.

Figure . “ Diλλraction patterns oλ borosilicate μlass solid μrey line , cluster-assembled λilm at room temperature dashed line , and cluster-assembled λilm aλter annealinμ at K solid black line are shown. The arrows indicate the position oλ the peaks that emerμe as a result oλ annealinμ. ” ”y subtractinμ the μlass siμnal λrom the siμnal recorded λor cluster-assembled λilm, a halo peak is observed that can be λitted by a Gaussian dashed line . The position oλ the peak in scatterinμ vector is in excellent aμreement with literature values λor Zr–Cu MGs oλ the same composition.

Crystallization temperature, Tx, is a more sensitive probe λor the structure. “s shown in Figure , annealinμ the Zr–Cu C“MG at K λor s leads to emerμence oλ sharp crystalline peaks in the diλλraction pattern that were not there prior to heat treatment. These diλλraction peaks belonμ to λcc Cu and bcc CuZr phases [ ]. The crystallization temperature oλ Zr–Cu MGs is much hiμher than K about – K hiμher , which points to the structural

Metallic Glasses from the Bottom-up http://dx.doi.org/10.5772/63514

diλλerence between rapidly quenched samples and the μlassy λilm λabricated λrom metal clusters. “ reasonable explanation λor the lower Tx oλ Zr–Cu C“MG under discussion could be based on an increased deμree oλ structural disorder caused by the diversity oλ metal clusters used to λabricate the λilm. In Fact, lowerinμ the deμree oλ short-ranμe order is known to lead to the decrease oλ Tx in μlasses [ ]. “lthouμh the existence oλ an amorphous Zr Cu phase in the λabricated C“MG could be conλirmed, the broad ranμe oλ metal clusters used to produce the amorphous λilm detains any detailed analysis oλ the relation between the structure oλ the λilm and its constituent buildinμ blocks. The example provides a very promisinμ λirst step in a rather lonμ journey oλ mostly unλoreseeable challenμes. Some oλ the upcominμ challenμes, however, can be expected, and a number oλ research μroups are workinμ on solvinμ them. The most immediate next steps that have to be are beinμ taken in this road are brieλly listed in the next section. . . The next steps Havinμ demonstrated the λeasibility oλ C“MG λabrication usinμ a diverse set oλ clusters in the Zr–Cu binary system, and the next steps can be taken in three directions oλ diλλerent nature. The λirst and probably the most important technical issue is the improvement oλ cluster sources. “s described in Section . . , much hiμher cluster currents are needed iλ speciλic sinμle clusters should be picked out λor the deposition oλ desiμned C“MGs. Since the cluster science com‐ munity is continuously enμaμed in enhancinμ the perλormance oλ the cluster sources, we stronμly anticipate that this and other technical issues will be resolved in the near λuture. The second communicational issue is the lack oλ inλormation on clusters relevant λor MGs. Scientists workinμ on MGs and metal clusters have not been in any close contact and the inλormation λlow between these two λields has been suλλerinμ. Yet there is more than enouμh motivation λrom both λronts to come toμether and hopeλully put an immediate end to this disconnection. On the one hand, cluster science community is hiμhly interested in under‐ standinμ cluster–cluster interactions and expandinμ the borders oλ the cluster science to more complex clusters. Further, departure λrom mainly purely λundamental science and movinμ toward real application by developinμ cluster-assembled materials C“Ms has been a lonμterm μoal oλ cluster scientists. On the other hand, the vision oλ μettinμ to an atomic structural model which can be veriλied throuμh C“MGs, and eventually solvinμ the lonμ standinμ structure–property puzzle has already triμμered enthusiastic activities amonμ material scientists. “lthouμh not many individuals have been active across the borders oλ the two λields, an enthusiastic collective interest that has been missinμ in the past is currently emerμinμ. The third issue is related to the handlinμ oλ C“MGs and their characterization. Currently, none oλ the cluster deposition λacilities around the world and the equipment λor thin λilm and MG characterizations are in the vicinity oλ each other. This requires complicated sample handlinμ endeavors, which are not always compatible with the metastable state oλ C“MGs. For instance, in the case oλ the C“MG treated in Section . , the sample transλer λrom the cluster deposition λacility to the apparatus where its structure was studied has been a challenμe, which could have been λully avoided by havinμ a cluster deposition λacility present at ESRF. We consider this issue less critical because we believe as soon as the other two issues are resolved even

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Metallic Glasses: Formation and Properties

partly , the interdisciplinary collaborations will naturally lead to emerμence oλ many such laboratories.

Author details “ras Kartouzian * and Jerzy “ntonowicz *“ddress all correspondence to [email protected] Department oλ Physical Chemistry, Technical University oλ Munich, Garchinμ, Germany Faculty oλ Physics, Warsaw University oλ Technoloμy, Warsaw, Poland

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Ni

. Zr

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bulk

metallic μlass, Journal

Chapter 2

Structural and Dynamical Properties of Metallic Glassy Films Hui Li, Weikang Wu and Kun Zhang Additional information is available at the end of the chapter http://dx.doi.org/10.5772/64107

Abstract In this chapter, a series oλ molecular dynamics simulations have been carried out to explore structural and dynamical λeatures oλ monatomic liquid metallic λilms durinμ rapid coolinμ. Results show a semi‐ordered inhomoμeneous morpholoμy containinμ crystal‐like and disordered reμions. The icosahedron contributes to nucleation throuμh the synerμy with other short‐ranμe ordered structures and participates in crystal μrowth via assimilation, but the pinninμ eλλect should be overcome. The second‐peak splittinμ in pair correlation λunctions is λound as the result oλ a statistical averaμe oλ crystal‐ like and disordered structural reμions, not just the amorphous structure. The splittinμ can be viewed as a prototype oλ crystal‐like peaks exhibitinμ distorted and vestiμial λeatures. ”esides, we use the parameter P a, , ν λor predictinμ both local structural order and motion propensity. The λraction oλ crystalline clusters λollows a neμative power‐law scalinμ with the coolinμ rate increasinμ, which is the inverse oλ P a, , ν . Keywords: molecular dynamics simulation, metallic μlass λilm, structural evolution, dynamical λeature, pair correlation λunction, local icosahedral order

. Introduction Ultrathin metallic μlassy λilms with a thickness oλ one or a λew monolayers attract much attention, since they are now available as epitaxial λilms on insulatinμ substrates and are, thereλore, the best model systems λor two‐dimensional D conduction in metal systems [ ]. Disorder is known to play an important role in the phase diaμram oλ the superconductor material at low temper‐ atures and hiμh maμnetic λields [ ]. Siμniλicant eλλort is currently beinμ invested in attemptinμ to understand theoretically the interplay between disorder and the conductivity in D sys‐

26

Metallic Glasses: Formation and Properties

tems [ – ]. Speciλically, in semiconductor λield, thinner and thinner disordered λilms are needed, since the restriction to a thickness oλ a λew monolayers may lead to novel atomic structures and modiλy the physical and chemical properties dramatically [ – ]. Huanμ et al. [ ] reported the accidental discovery oλ D amorphous silica supported on μrapheme. They λound that the imaμes oλ D amorphous silica contain both the crystalline and amorphous reμions. Lichtenstein et al. [ ] studied the interλace between a crystalline and amorphous phase oλ silica λilm supported by the Ru substrate. The atomic structure oλ the topoloμical transition λrom a crystalline to an amorphous phase in the thin silica λilm can lead to a better description oλ the crystal‐to‐ μlass and the liquid‐to‐μlass transitions. “lthouμh there has been much proμress in the understandinμ oλ the properties oλ amorphous materials in three‐dimension [ – ], some important questions on the microstructural λeature and its λorminμ mechanism oλ the metallic μlassy λilms have remained unanswered. Thereλore, λurther studies on the atomic structures oλ the D disordered systems and their physical proprieties are necessary [ – ]. This chapter is orμanized as λollows. In Section  , we describe the modelinμ and simulation methods. In Section  , we discuss the structural evolution oλ liquid metallic nano‐λilm durinμ rapid solidiλication as well as the eλλect oλ coolinμ rates [ , ]. In Section  , we consider the motion propensity distribution to predict both local structural distribution and dynamical siμnature in metallic nano‐λilms [ ]. In Section  , we study the synerμy and pinninμ eλλects oλ the local icosahedral order durinμ λreezinμ [ ]. In Section  , we clariλy the oriμin oλ the splittinμ oλ the second peak in PCFs based on a statistical explanation [ ]. Conclusions are provided in Section  .

. Models and theoretical methods Molecular dynamics MD simulations were perλormed usinμ the embedded atom method E“M potential [ ] supplied in L“MMPS [ ]. The pure copper λilm and the pure cobalt λilm are studied respectively. For the copper λilm, copper atoms were distributed in a  ×    ×  lattice unit box based on the structure oλ FCC crystal and the initial box lattice was set to be .  Å. For the cobalt λilm, the initial conλiμuration consistinμ oλ cobalt atoms was distributed in a  ×   ×  lattice unit box, in which atoms were arranμed in the liμht oλ the HCP crystal structure and the box lattice was set to .  Å. Periodic control was exerted on the x‐ and y‐directions oλ the box, and the z‐direction was nonperiodic. Specially, the lower boundary oλ the simulation box alonμ the z‐direction not reλer to atoms was λixed, while the upper boundary was λree. That was to say, a virtual wall was set at the lower edμe oλ the simulation box in the z‐direction, which is similar to the substrate in experiment. The Velocity‐Verlet alμorithm was used with an MD time step oλ  λs while the temperature was controlled by a Nosé‐Hoover thermostat [ ]. “ well‐equilibrated initial system was prepared by μradually heatinμ perλect crystals to melt at a low heatinμ rate and then relaxinμ the system at  K λor  ps. Then, the liquid system was quenched to  K at diλλerent coolinμ rates. “t each coolinμ rate, the atomic conλiμuration durinμ the quenchinμ process were recorded λor λurther analysis.

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

. Structural features in liquid metallic nano‐films during rapid cooling Figure shows the potential enerμy landscape oλ the liquid copper nano‐λilm at the tempera‐ ture oλ  K, which is composed oλ sinμle‐strinμ structures and partial rinμ structures we λind the rinμ structure is the icosahedron . We call the icosahedron in our simulations the quasi‐ two‐dimensional icosahedron Q D‐I because oλ one missinμ vertex atom due to the dimen‐ sional limit. “toms in sinμle‐strinμ structures show hiμh potential enerμy and motion propensity, while Q D‐Is have a hiμh dense packinμ [ ]. “ctually, in the melt λilm, both structures disinteμrate and reunite λrequently owinμ to the random collision and enerμy transλer oλ hiμh‐enerμy atoms. ”esides, ensembles oλ atoms in diλλerent reμions oλ this liquid λilm exhibit temporarily enhanced or diminished mobility in comparison with the averaμe. Notably, the cooperative motion oλ sinμle‐strinμ structures is dominant in the metallic liquid λilm, in accordance with that in three‐dimensional liquids [ – ].

Figure . The potential enerμy landscape oλ the copper nano‐λilm at  K. The denoted circular reμion represent the quasi‐two‐dimensional icosahedron Q D‐I and the elliptical one represents the sinμle‐strinμ structure.

“s the temperature decreases, atoms in the D liquid copper tend to μather into clusters. To clariλy the liquid‐solid transition mechanism in metallic λilms durinμ the coolinμ process in the view oλ an atomic level, the structural evolutions at diλλerent coolinμ rates are investiμated as shown in Figure . “t hiμh coolinμ rates and  K/ps , a semi‐ordered morpholoμy exhibitinμ maze‐like nano‐patterns μradually λorms at  K. To be speciλic, at  K/ps, compared with the atomic conλiμuration at  K Figure , Q D‐Is increases obviously accompanyinμ with the decrease oλ sinμle‐strinμ structures at  K. When the temperature decreases into  K, Q D‐Is show continuous increase while the sinμle‐strinμ structures beμin to arranμe side by side, λorminμ crystalline zones. The size scale oλ crystalline zones in the

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simulation is beyond –  Å, meetinμ the size requirement oλ MRO –  Å [ ] . Thus, the crystalline zone can be considered as the MRO. Notably, these MRO structures are the precursor λor nucleuses, whereas Q D‐Is would have barrier eλλects on nucleation. Durinμ the rapid coolinμ process, Q D‐Is and crystalline zones competite aμainst each other and λinally determine the solid structure at  K. For example, at  K/ps, Q D‐Is play the leadinμ role since crystalline zones are limited due to the hiμh coolinμ rate, and the system exhibits the most disorder at  K. However, when the lower coolinμ rate  K/ps is perλormed, crystalline zones tend to be dominant, and λinally the λilm system shows more crystalline MRO characteristics. “lthouμh the crystalline MRO μradually λorms and develops durinμ rapid coolinμ, the transition time is so transient that crystalline zones cannot be converted into the crystalline lonμ‐ranμe order CLRO . In contrast, in the case oλ  K/ps as shown in Fig‐ ures g–i , owinμ to the absence oλ the icosahedral λrustration eλλect, atoms are mainly arranμed in an crystalline order.

Figure . Quenchinμ processes at diλλerent temperatures with three diλλerent coolinμ rates. The denoted circular reμion represent the Q D‐I and the elliptical one represents the crystalline zone.

To λurther investiμate the inλluence oλ the coolinμ rate on the phase transition oλ metallic λilms, both the pair correlation λunctions PCFs and the proportion oλ crystalline zones at diλλerent coolinμ rates are plotted. The crystallization moments can be charaterized by the pressure drops [ ], indicatinμ a linear relation between the crystallization moment and the supercooled melt liλetime. However, the λaction oλ crystalline zones in this simulation is determined

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

throuμh the statistical averaμe oλ the recorded structural inλormation [ , ], which is diλλerent λrom Morozov et al.'s method [ ]. “s shown in Figure a, the hiμhest coolinμ rate  K/ps leads to a sliμht splittinμ second peak in the PCF curve, while at  K/ps, the second peak exhibits obvious splittinμ and a small shoulder peak appears between the λirst peak and the second peak. “lthouμh the second peak splittinμ oλ the PCF curve is usually considered as an important siμnature oλ amorphous solids [ ], we would rather reμard it as the λormation oλ crystalline zones which can be seen as a precursor to the CLRO [ ]. “t a low coolinμ rate, such as  K/ps, the PCF curve presents typical crystalline λeatures the splittinμ characteristic oλ the PCF second peak becomes more obvious the small shoulder peak chanμes into a sharper peak. Durinμ the slow coolinμ process, Q D‐Is disappear while the crystalline MRO inteμrates toμether, leadinμ to the CLRO characteristic oλ the metallic λilm. Figure b shows the crystalline λraction chanμe oλ the above three quechinμ processes. “t the hiμh coolinμ rate and   K/ps , the FCC and HCP structures exhibits a similar amount and shows little chanμes durinμ the coolinμ process. However, at  K/ps, both the FCC structure and the HCP structure μrow λast between and  K, and λinally FCC structures become the dominant with the λaction oλ more than % at  K. Such results stronμly indicate a non‐linear relation between the crystallization λraction and the λreezinμ time, quite diλλerent λrom the Morozov's result in bulk liquid metals [ ].

Figure . a Pair correlation λunctions PCFs at es.

 K d the λraction oλ crystalline zones durinμ the coolinμ process‐

. Power‐law scaling of dynamical signatures In this section, the motion propensity distribution is considered to predict the structural and dynamical λeatures oλ the liquid copper nano‐λilm. First, the Common Neiμhbor “nalysis

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Metallic Glasses: Formation and Properties

CN“ and motion propensity are introduced to reveal the relation between crystalline MRO and the subsequent crystallization. The CN“ can be used to measure the local crystalline structure around an atom [ , , ], based on the Honeycutt and “ndersen bond analysis [ ]. Generally, there are λive kinds oλ CN“ patterns that L“MMPS recoμnizes, which are deλined as λollows FCC =  , HCP =  , ”CC =  , icosahedral =  , and unknown =  . The λirst three indices are all crystalline . “lso note that the CN“ calculation in L“MMPS uses the neiμhbors oλ an owned atom to λind the nearest neiμhbors oλ a μhost atom. The local motion propensity oλ a particle [ , ] is directly associated with the probability oλ a particle underμoinμ a substantial displacement within a short time interval. The motion propensity oλ a particle p is deλined as λollows



where Δrp t,t+τ

is the displacement oλ the particle p obtained λrom the quenched conλiμu‐

ration between t and t +  and a is the lenμth scale over which the motion is probed. Here a =  , and  =   λs [ ]. Figure shows the CN“ pattern and the motion propensity distribution oλ queched λilms at diλλerent coolinμ rates. When the coolinμ rate is hiμh  K/ps , althouμh the barrier eλλect oλ Q D‐Is is prominent, several ordered crystalline clusters, such as FCC and HCP structures, still exist at  K, indicatinμ that the so‐called λull amorphous state may contain λew crystalline SRO structures, namely the crystallite. However, due to the limitation oλ experi‐ mental methods, it is hard to measure the real atomic arranμement within the SRO ranμe. “ctually, these crystallites are quite stable, which can be seen λrom their motion propensity distribution as shown in Figure a. “s the coolinμ rate reduce, it can be clearly seen that crystallites are rare and discrete initially, but appear at random positions, exhibitinμ a larμe structural heteroμeneity oλ the metallic μlassy λilm [ ]. “t  K/ps, more crystallites appear and μrow in size at  K as shown in Figure b. “ two‐step crystallization process is proposed [ ] λirst, the λluctuations oλ structure and enerμy cause the λormation oλ several SRO crystal‐ lites which can precursors λor the nucleus oλ crystalline zones Then, the crystallites expand into surroundinμ and λorm crystalline zones. Notably, the FCC or HCP structures do not emerμe alone. Instead, the crystalline zone is made up oλ alternant FCC and HCP structures as presented in the CN“ patterns. When the coolinμ rate is quite low, there would be enouμh relaxation time λor the evolution λrom HCP structures to FCC structures, similar to Wolde et al.'s viewpoint [ ]. For example, at  K/ps, the weakeninμ barrier eλλect oλ Q D‐Is would result in the ordered atomic arranμement and the λormation oλ nano‐polycrystalline structures dominated by FCC structures at  K. The appearance oλ HCP structures is actually just an intermediate state durinμ the crystallization. ”esides, it is worth notinμ that, the complete‐ disorder reμion is characterized by hiμh motion propensity, whereas the low motion propen‐ sity is related to the crystalline zones. Generally, the basic principle oλ local propensity Qt a, ,

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

is that the conλiμurational order does not determine the motion oλ a neiμhborhood directly but aλλects the probability oλ the underμoinμ motion. Thereλore, a theoretical discussion is necessary to veriλy the possible relation between the structure and dynamics.

Figure . Local atomic orderinμ leλt and motion propensity riμht oλ the quenched structures at  K with diλλerent coolinμ rates a  K/ps b  K/ps c  K/ps. In the leλt column, local structures are colored based on the CN“ method blue, FCC μreen, HCP red, disordered structures. In the riμht column, colorinμ denotes the motion propensi‐ ty.

Next, we emphatically studied the relation between the total motion propensity and the crystalline‐like structure in the copper λilm at diλλerent coolinμ rates. Qsd is the standard deviation oλ the total motion propensity obtained λrom the quenched conλiμuration at  K durinμ a short relaxation time, which may reλlect the total λluctuation oλ the D system with respect to time clearly. Figure a shows the Qsd chanμe as a λunction oλ time durinμ a short period ϕ =   λs, at diλλerent coolinμ rates. For the quenched structure under the hiμhest coolinμ rates, its Qsd shows the hiμhest dynamical λluctuation durinμ the relaxation, indicatinμ that the system is in the metastable state. In contrast, the dynamical λluctuation at a lower coolinμ rate exhibits less obvious λeatures, in μood accordance with Figure . This means that the total motion propensity is reliable to measure the system stability. For λurther exploration, a μlobal measure oλ the propensity P a, , is calculated by the correlation λunction Qt a, ,

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Metallic Glasses: Formation and Properties

Figure . a The standard deviation Qsd oλ the total motion propensity and b the relationship between the total mo‐ tion propensity leλt and the λraction oλ crystalline zones riμht at diλλerent coolinμ rates.

where  = loμV, V is the coolinμ rate. Qt a, ,

is the motion propensity oλ a sinμle particle.

Figure b shows the relationship between P a, , and the crystalline λraction at  K, with respect to the coolinμ rate. “s the coolinμ rate increases, the λraction oλ crystalline zones increases and shows an index relationship with the coolinμ rate. For a low coolinμ rate    . . “t a moderate coolinμ rate, such as  =  . , a complicated amorphous‐crystalline composite λorms, with the FCC and HCP structures distributinμ randomly and exhibitinμ the crystalline MRO characteristic. Obviously, the λraction oλ crystalline zones exhibits a neμative power‐law scalinμ with the coolinμ rate, which is the inverse oλ P a, , . This indicates a close relationship between P a, , and the crystalline structure as the coolinμ rate increases. Such phenomena have not been reported in experimental research, due to some λactors associated with multiple scatterinμ and the atomic cluster distortion. To conclude, as a new parameter, P a, , may be a direct consequence oλ the local structural orderinμ and the dynamic siμnature may also result λrom the local structural orderinμ.

. Synergy and pinning effects of local icosahedral order The above results show that, Q D‐Is appear earlier than crystalline zones durinμ the coolinμ and can be preserved under the hiμher coolinμ rate since the local icosahedral order LIO μrows with severe undercoolinμ [ ]. When crystalline zones emerμe, the quantity oλ Q D‐Is decreases, indicatinμ that Q D‐Is inλluence the λormation oλ the CMRO.

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

Figure a shows how crystalline zones nucleate throuμh the synerμy eλλect oλ Q D‐Is. It should be pointed out that, the sinμle‐strinμ structure is the basic unit oλ crystalline zones since it can extend to λorm the close‐packed plane CPP , similar to the expansion oλ a μraphene piece into a μraphite layer. Figure d reveals the similar topoloμical unit between Q D‐Is and CCPs, which is the structural basis λor the assimilation oλ Q D‐Is. Initially, in the supercooled melt, owinμ to the stronμ atomic activity, the SRO structures includinμ the sinμle‐strinμ structure and Q D‐Is disinteμrate and reunite λrequently. Durinμ the nucleation, several Q D‐Is are decomposed and inteμrate into the nucleus with the help oλ sinμle‐strinμ structures, indicatinμ that the LIO should be the raw component oλ the crystal nucleus. Figures b, c show the μrowth oλ crystalline zones with the assistance oλ Q D‐Is in the parallel and perpendicular direction to CPPs, respectively. In the parallel μrowth, when Q D‐Is are contacted by the λront oλ crystalline zones, Q D‐Is would be touched by the CPPs naturally due to their similar topoloμical unit, leadinμ to the μradual assimilation see Figure in detail . On the other hand, when the crystalline zones μrow in the perpendicular direction and contact a Q D‐I, a suitable connection to the Q D‐I by the surroundinμ sinμle‐strinμ structures or CPPs would assimilate the Q D‐I into the crystalline structures, leadinμ to the perpendicular μrowth oλ crystalline zones. Thus, Q D‐Is show the synerμy eλλect durinμ the whole crystallization process. Notably, icosahedra are usually seen as a barrier to crystallization [ ], but these results demonstrate that icosahedra can participate the crystal μrowth throuμh the synerμy eλλect in the D system, providinμ a new view oλ the correlation between the LIO and crystallization.

Figure . a “ crystal nucleus λorms throuμh the synerμy between Q D‐Is and sinμle‐strinμ structures. The crystal nu‐ cleus is characterized by white solid lines. b, c crystalline zones μrow throuμh throuμh the assimilation oλ Q D‐Is in the direction parallel and perpendicular to the close‐packed planes CPPs , respectively. The Q D‐Is are characterized by dotted circles and rounded rectanμles. The IS, TS and FS represent the initial state, transition state and λinal state, respectively. d Q D‐Is and CPPs have the similar topoloμical unit. From top to bottom the side structure oλ Q D‐Is, lonμitudinal section oλ Q D‐Is and CPP.

Next, the detailed assimilation oλ Q D‐Is by crystalline zones is investiμated to explore the mechanism oλ the synerμy eλλect. Figure a shows the assimilation process oλ a Q D‐I by

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Metallic Glasses: Formation and Properties

crystalline zones in the direction parallel to CPPs. The contaction between CPPs oλ crystalline zones and a Q D‐I would result in a halλ‐icosahedral and halλ‐crystalline μeometrical λrustra‐ tion, as illustrated in Figure b. Then, the Q D‐I is μradually assimilated as the μeometrical λrustration evolves into the crystalline order. “ccordinμ to Figure c, durinμ the assimilation, the Q D‐I experiences an enerμy λluctuation accompanyinμ with an obvious oscillation λor the averaμe potential enerμy. In order to measure the enerμy λluctuation, the root mean squared deviations RMSDs [ ] are calculated based on the mean value oλ enerμy λluctuations in the whole system RMSD=

N

N

¯ | , where N is the atomic number, ΔE is the potential ∑ | ΔEi‐ΔE i i=

¯ is the averaμe oλ ΔE . The averaμe λluctuation oλ the system enerμy λluctuation per atom, ΔE i

¯  =  .  eV and durinμ the process in Figure is .  ±  .  eV where, ΔE RMSD =  .  eV . The λluctuation oλ Q D‐I just in the ranμe indicates that enerμy λluctuations lead to the assimilation durinμ the parallel μrowth. Thus, the Q D‐I can be actually seen as a pin durinμ the assimilation, which needs to be drawn out λor the μrowth oλ crystalline zones. In another word, Q D‐Is have pinninμ eλλects on the crystal μrowth due to the enerμy λluctuation. Iλ the pinninμ eλλect can be overcome, the assimilation oλ Q D‐Is would assist the μrowth oλ crystalline zones, presentinμ the synerμy eλλect oλ the LIO.

Figure . a “ Q D‐I is assimilated by crystalline zones in the direction parallel to CPPs. White solid arrows show the crystal μrowth direction. b Detailed structural evolution durinμ the assimilation. Three adjacent CPPs naturally con‐ nect to the Q D‐I with red solid arrows showinμ the direction. c Potential enerμy chanμe durinμ the assimilation. The curve oλ the Q D‐I shows potential enerμy λluctuations durinμ the assimilation.

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

Figure shows that the crystalline zone assimilates a Q D‐I in the direction perpendicular to CPPs, which presents more complicated λeatures than that in the parallel μrowth. However, under closer observation, the assimilation process is very similar to that durinμ a parallel μrowth. When crystalline zones touch a Q D‐I, the Q D‐I would be μradually connected by the surroundinμ CPPs, revealinμ that the crystal μrowth in the direction perpendicular to CPPs oriμinate λrom the connection between CPPs and Q D‐Is in the parallel direction. Figure b presents that the Q D‐I also underμoes a potential enerμy λluctuation and the averaμe enerμy oscillates durinμ the assimilation. The potential enerμy λluctuations durinμ the perpendicular μrowth indicate the pinninμ eλλect oλ Q D‐Is. The averaμe λluctuation oλ system is ¯  =  . .  ±  .  eV where, ΔE  eV and RMSD =  .  eV , and the Q D‐I's enerμy λluctuation leads to the leap over the pinninμ eλλect durinμ the perpendicular μrowth. Once the pinninμ eλλect succeeds, Q D‐Is miμht survive as shown in Figure . In this case, the surroundinμ CPPs have to bend to cater to the Q D‐I due to the incompatibility between Q D‐Is and the crystals, exhibitinμ the μeometrical λrustration. In another word, the Q D‐I is actually a pin, causinμ the μeometrical λrustration oλ the surroundinμ crystalline order.

Figure . a “ Q D‐Is is assimilated by crystalline zones in the direction perpendicular to CPPs. White solid arrows show the crystal μrowth direction. The assimilation in a is very similar to that in the parallel μrowth. b Potential enerμy chanμe durinμ the assimilation. The curve oλ the Q D‐I shows potential enerμy λluctuations durinμ the assimi‐ lation.

Figure . Preservation oλ a Q D‐I due to the pinninμ eλλect. The solid arc line shows the λrustrated CPP caterinμ to the Q D‐I.

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Metallic Glasses: Formation and Properties

. Origin of the second peak splitting in pair correlation functions “ splittinμ oλ the second peak in the PCF curve in three‐dimensional materials is usually reμarded as a characteristic indication oλ disordered structure λorminμ [ , ]. However, owinμ to the lack oλ one dimensionality as well as the diλλerence oλ atomic arranμement, the explanation on the second peak oλ the PCF in D systems should be clariλied. In addition, previous studies are usually based on the viewpoint oλ how the clusters connect to each other to λorm a larμe supercluster with a speciλic μeometric structure. Due to the λact that the PCF is the statistical averaμe oλ the atomic conλiμuration, it seems more appropriate to use the statistical methods to interpret the nature oλ the second peak splittinμ oλ the PCF in D system. “ctually, λewer eλλorts have λocused on the relation between the splittinμ oλ the second peak and the crystalline or μlass λormation oλ D disordered λilms by statistical averaμe analysis. In order to understand these questions well, this section would provide a statistical explanation on the nature oλ the splittinμ oλ the PCF in the D copper system.

Figure . PCF evolutions oλ the D copper at diλλerent coolinμ rates a  K/ps b  K/ps c  K/ps d   K/ps. ”rown curves represent the PCF without the second peak splittinμ ”lue curves show the second peak splittinμ Green curves show typical crystal peaks.

The PCFs oλ the D copper are shown in Figure . It is worth notinμ that the main peak heiμht oλ the PCFs, which represents the nearest‐neiμhbor shell, increases siμniλicantly with the decreasinμ temperature, and the second peak beμins to split. “s shown in Figures a, b, at the coolinμ rates oλ Q  =   K/ps and Q  =   K/ps, the second peak beμins to split into two subpeaks at  K. Interestinμly, a small shoulder peak appears between the λirst and second peaks at  K in Figure b, which means the short‐ or medium‐ranμe ordered structures λorm. With the coolinμ rate decreasinμ to Q  =   K/ps, the splittinμ emerμes at  K moreover, the small shoulder peak between the λirst and second peaks arises on the leλt at  K with siμniλicant heiμht, which indicates that the lenμth oλ the ordered structure is λurther extended to a larμe scale. “t the coolinμ rate oλ Q  =   K/ps, the leλt shoulder peak arises at  K and becomes more prominent than the riμht subpeak as the temperature decreases,

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

suμμestinμ that the orientation oλ the crystalline structure becomes more consistent. It is widely known that the atomic structure oλ amorphous materials is similar to that oλ liquid metals, and the λact that the second peak oλ the PCFs splits into two subpeaks is reμarded as a characteristic indication oλ disordered structures. However, this is not the true case in D systems. “s shown in Figures c, d, the evolution oλ the PCFs clearly indicates how the second‐peak splittinμ converts into crystal peaks. For example, as shown in Figure d, at  K, the splittinμ oλ the second peak oλ the PCF appears, but at  K, the splittinμ oλ the second peak becomes three peaks, and λinally these three peaks evolve into three typical crystal peaks at  K. ”ased on this evolution trend, it can be seen that the splittinμ second‐peak has a close relationship with crystal peak, and that the splittinμ second‐peak is the rudiment oλ the crystal peaks. Our simulations do not support other hypothesis which states that the splittinμ oλ the second peak occurs as a result oλ the connection oλ some small clusters to a supercluster with a special μeometrical structure.

Figure . a Snapshot oλ the D copper at  K with the coolinμ rate oλ  K/ps. Crystal‐like orderinμ reμions are constituted by blue atoms, while other reμions show λully disordered orderinμ. b–d PCF curves oλ diλλerent reμions local crystal‐like reμions, λully disordered reμions, the μlobal reμion, respectively.

In order to λurther clariλy the oriμin oλ the splittinμ oλ the second peak, the structural conλiμ‐ uration oλ the D copper at  K at the coolinμ rate oλ Q  =   K/ps is supplied in Figure . It is seen that the overall atomic structure consists oλ two types oλ reμions the well‐orμanized reμion with crystal‐like order and the λully disordered reμion with some packinμ λrustration. Our theoretical results are in μood aμreement with the experiments by Huanμ and Lichtenstein [ , ], which prove our MD simulation result is reliable. It is also worth notinμ that in the disordered reμion there are some sinμle strinμs, arcs, and rinμs which clearly illustrate the packinμ λrustration oλ the atoms in the quick coolinμ process. In λact, Figure a shows that the structure oλ the D amorphous Cu is the mixture oλ crystal‐like and λully disordered structural reμions with a certain percentaμe. The local PCFs in these two distinct reμions diλλer λrom one another. “s shown in Figures b–d, the local PCFs oλ the crystalline reμion have some crystal‐like subpeaks, showinμ typical crystalline λeatures, while the local PCFs in the λully disordered structural reμion show no splittinμ on the second peak. However, the μlobal

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PCFs averaμed by the overall atomic structures oλ the two types oλ reμions show a sliμht splittinμ in the second peak. It is known that the PCF is the statistical averaμe oλ the structural conλiμuration, thus, the sliμht splittinμ oλ the μlobal PCF is caused by the combined averaμe results oλ the crystal‐like and λully disordered reμions. Moreover, the very similar results are also obtained λor the simulations oλ the D cobalt, which indicate the coexistence oλ crystal‐ like and λully disordered reμions. The splittinμ oλ the second peak in D systems may not be the siμnature oλ the μlass λormation, but the appearance oλ both the crystal‐like and disordered structures. The splittinμ second‐peak can be viewed as an embryonic λorm oλ the crystal peak. The above results arouse us to λurther investiμate the oriμin oλ these two subpeaks in the second‐peak splittinμ. Figure shows the respective PCF curves oλ the liquid, amorphous, and ideal crystalline solid Cu. It is known that in ideal FCC crystal there are λour nearest coordinated shells, namely, R  =  .  Å, R  =  .  Å, R  =  .  Å, and R  =  .  Å. The positions oλ the λirst peak in both the liquid and amorphous Cu correspond to the ones in the ideal FCC crystal at R  =  .  Å, while the second peaks correspond to three peaks oλ the ideal FCC crystal at the locations R  =  .  Å, R  =  .  Å, and R  =  .  Å. From the correspondence oλ the peak positions and the evolution trend, it is concluded that the two subpeaks on the second peak are due to the appearance oλ a small amount oλ the short‐ or medium‐ranμe ordered structures.

Figure . PCF curves oλ the liquid, amorphous, and ideal crystal copper. Red line liquid at  K ”lue line amor‐ phous solid at  K Q =   K/ps Green line amorphous solid at  K Q =   K/ps ”lack line ideal crystal. R , R , R , and R represent the ideal FCC peaks.

Figure shows the λraction oλ the crystal‐like reμions with the temperature at diλλerent coolinμ rates. Iλ the splittinμ occurs on the riμht without the leλt subpeak, this point is labelled with a blue arrow, while iλ both the leλt and riμht subpeaks appear, it is labelled with a red arrow. “ccordinμ to this rule, Figure may be divided into three zones. When the λraction oλ the crystal‐like reμion is less than . %, it belonμs to the non‐splittinμ reμion correspondinμ to the λully disordered structure in the liquid state. Iλ the λraction oλ crystal‐like reμion exceeds to %, the leλt shoulder subpeak appears, which indicates it almost becomes the polycrystal‐

Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

line structure. However, iλ the λraction ranμes λrom . to %, it is the mixture oλ the disordered and crystal‐like structural reμions, which results in the splittinμ second peak oλ the PCF. The disordered structure in D may be a simple mixture oλ crystal‐like and λully disor‐ dered reμion, which sheds a new liμht on the understandinμ oλ the atomic structure oλ the low‐ dimensional materials.

Figure . Relationship between the λraction oλ crystal‐like reμions and the second peak splittinμ at diλλerent coolinμ rates  K/ps blue line ,  K/ps red line ,  K/ps blue line ,  K/ps brown line ,  K/ps purple line ,  K/ps dark blue line .

. Conclusion In summary, we systematically investiμate the structural evolution and dynamical properties oλ μlassy metallic λilms durinμ rapid coolinμ. Our results have shown several aspects as λollows • “ semi‐ordered amorphous morpholoμy with maze‐like nano‐patterns emerμes as the temperature decreases at a hiμh coolinμ rate. The μrowth competition between two typical dominatinμ structures, Q D‐Is and crystalline zones, siμniλicantly aλλects the λinal solid‐state structure. The FCC and HCP strutures are alternant in the crystalline zone reμion, actinμ as the precursor oλ CLRO structures. • The disordered reμion usually has a hiμh motion propensity distribution, whereas a lower motion propensity corresponds to the crystalline‐like order reμion. ”oth the local structural distribution and the dynamical siμnature in metallic nano‐λilms can be predicted by an excellent indicator P a, , . The reμion with a lower P a, , can accommodate a larμer crystallin‐like order, and vice versa. • The LIO has synerμy and pinninμ eλλects on the λreezinμ behavior oλ a monatomic liquid λilm. crystalline zones would μradually λorm throuμh the synerμy oλ Q D‐Is with other SRO

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Metallic Glasses: Formation and Properties

structures, but the pinninμ eλλect should be overcome when crystals μrow in both the directions parallel and perpendicular to close‐packed planes, consuminμ enerμy. • The oriμin oλ the splittinμ oλ the second peak in PCFs is a statistical result oλ the disordered and crystal‐like ordered structure with a certain percentaμe rather than the λully disordered structure. The results show that the shoulder peak on the leλt side oλ the second peak is due to the appearance oλ a small amount oλ the short‐or medium‐ ranμe ordered structures. The structure in D disordered λilm may be a simple mixture oλ the crystal‐like and disordered structural reμions.

Acknowledgements We would like to acknowledμe the support by the national λunds λrom Chinese μovernment Nos. , , and No. . This work is also λunded by the National ”asic Research Proμram oλ China Grant No. C” . This work is also λunded by the μrants λrom Shandonμ Province in China Nos. , JQ , and ZR FM . This work is also supported by the Special Fundinμ in the Project oλ the Taishan Scholar Construction Enμineerinμ. We also thank the support λrom Shandonμ University No. JQ and No. and Shandonμ excellent younμ scientist science λounda‐ tion ”S CL .

Author details Hui Li*, Weikanμ Wu and Kun Zhanμ *“ddress all correspondence to [email protected] KeyLaboratory λor Liquid‐Solid Structural Evolution and Processinμ oλ Materials, Ministry oλ Education, Shandonμ University, Jinan, People's Republic oλ China

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Structural and Dynamical Properties of Metallic Glassy Films http://dx.doi.org/10.5772/64107

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] Kobayashi K., Salam M.U. Comparinμ simulated and measured values usinμ mean squared deviation and its components. “μronomy Journal. – .

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43

Chapter 3

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level HaiJun Chang Additional information is available at the end of the chapter http://dx.doi.org/10.5772/63676

Abstract Metallic μlass MGs has many unique properties such as low density, low Younμ's modulus, and so on. These unique physical and mechanical properties attract much attention on their application in manuλacturinμ production. While, structural proper‐ ties such as complete absence oλ the lonμ‐ranμe order and most MGs are consist oλ equal or more than ternary constituent which complex λactors make that the atomic level structure oλ metallic μlass still have not well known by researchers. The limited methods and data sets obtained by experiment make the acknowledμe in uncoverinμ atomic structure oλ melt states oλ alloy, and the supercooled liquid about the alloy is absent as well. These messaμes are important to improve and increase the understandinμ oλ MGs, as we know that μlasses are essentially λrozen liquid made by quenchinμ oλ their hiμh‐ temperature melts. Computer simulation method is an useλul tool to obtain structure messaμes oλ melt and the supercooled liquid. The static, dynamical properties as a λunction oλ temperature can also be investiμated by ab initio MD simulation. The atomic level rearranμement consists oλ both local topoloμical structure chanμe and chemical reorderinμ and evolution oλ electronic properties oλ the “l Ni Nd and Ca Mμ Cu alloy durinμ the μlass transition process is investiμated, and the discussion oλ the results is μiven in this chapter. Keywords: supercooled liquid, μlass transition, microstructure, electronic evolution, metallic μlass, atomic structure, electronic properties

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Metallic Glasses: Formation and Properties

. Introduction For amorphous metals, the expandinμ application in manuλacturinμ production makes metallic μlass MGs attract much attention to uncoverinμ their atomic structure. “s a consequence oλ the complete absence oλ the lonμ‐ranμe order, equal or more than ternary constituent ele‐ ments and the varyinμ chemical aλλinity between the constituent elements which complex λactors known by researchers. Topoloμical and chemical short‐ranμe order is existed in MGs and is the most pronounced structure λactor in these amorphous alloys. “toms packinμ λrom micro‐observation view and structural diλλerent toμether with chanμes by temperature decrease durinμ the μlass transition proμress are important to explore these interestinμ properties oλ MGs. The electronic structure and the bond between paired atoms in amorphous alloy are still little reported. The limited methods and data sets obtained by experiment make the acknowledμe in uncov‐ erinμ atomic structure oλ melt states, and the supercooled liquid about the alloy is absent as well. These messaμes are important to obtain useλul knowledμe such as μeneral structural models and λundamental principles oλ atom packinμ and the correlations oλ thermodynamic, kinetic and mechanical properties with the MG structures. Computer simulation method is an useλul method to obtain messaμes oλ melt and the super‐ cooled liquid. The static, dynamical properties as a λunction oλ temperature can also be investiμated by ab initio MD simulation λrom the obtain results toμether with the electronic structure, bond and electronic density oλ states evolution durinμ the μlass transition process [ , ]. Usinμ the λirst‐principle computer simulation method, the atomic level structure oλ metallic μlass was indicated, which is vital to understandinμ oλ the behaviour oλ these materials. The research was λocussed on λirst‐principle molecular dynamic simulation oλ the rapid solidiλi‐ cation process oλ Ca Mμ Cu alloy [ ] and the evolution oλ structural and electronic proper‐ ties durinμ the μlass transition process oλ “l‐Ni‐Nd alloy [ ]. Some theory means have used to propose the evolution oλ the structure oλ this alloy durinμ the μlass transition such as H“ index method and bond‐anμle distribution λunction. The result indicates that the increasinμ pentaμonal bipyramids orderinμ suppress the λormation oλ the crystal structures durinμ the rapid solidiλication processes. Evidence λor the existence oλ covalent bonds is provided by our research, the polyhedral local topoloμical structures toμether with the chemical short‐ to medium‐ranμe order SRO structures play an important role in μlass transition and also increase the μlass‐λorminμ ability oλ these alloy. The charμe density oλ the alloy was also μiven. The static and dynamical properties oλ Ca Mμ Cu alloy and “l‐Ni‐Nd alloy durinμ the rapid solidiλication process are investiμated usinμ ab initio MD simulations. The results we calcu‐ lated aμree well with the previous report oλ these MGs.

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

. Part , atomic level structure of metallic glass Since the unique physical and mechanical properties oλ the metallic μlass, there has been increasinμ interest in developinμ and understandinμ this new λamily oλ materials. “ number oλ experimental techniques, such as X‐ray diλλraction XRD and neutron diλλraction, have been carried out to understand the structure oλ the metallic μlass as well as modellinμ methods, which include Reverse Monte Carlo RMC and molecular dynamics MD . These results μiven by researchers have more attraction and more attention on their application in manuλacturinμ production. “tomic level structure oλ metallic μlass was investiμated usinμ ab initio MD simulations, and the local atomic structures oλ Ca‐Mμ‐Cu and “l‐Ni‐Nd amorphous metallic μlasses have been investiμated by the pair distribution λunction PDF analysis oλ neutron diλλraction data. The pair distribution λunction PDF is a pair correlation λunction to indicate that the probability oλ an atom is located at distance r λrom an averaμe centre atom. PDF is used to characterize the structural properties oλ liquids and amorphous solids. The ‐β partial PDF is deλined in Eq. gab (r ) =

L3 Na N b

å 4pibr 2Dr N

n (r )

i =1

where in the system oλ N atoms, N and Nβ are the number oλ atoms oλ type and type β, respectively. r is the interatomic distance between two atoms i oλ type and j oλ type β .

Figure . The total paid distribution λunctions and partial pair distribution λunctions oλ the Ca Mμ Cu alloy at diλλer‐ ent temperatures.

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Metallic Glasses: Formation and Properties

It can be seen that the λirst peak oλ total PDF oλ Ca Mμ Cu Figure μets hiμher as the temperature decreases, which indicates that atomic orderinμ in the λirst coordination shell increases as the temperature decreases. The λirst peak oλ total PDF starts splittinμ at  K. “t the lowest temperature,  K, a splittinμ oλ the λirst peak into two peaks around . “ and . “ and a splittinμ oλ the second peak are seen. For the PDF oλ Ca atom, we can λind a shoulder which appears at  K and becomes more pronounced as the temperature is decreased. This prepeak is located at . “° which is equal to the interatomic distance oλ Ca‐ Cu pair and it is believed to oriμinate λrom increasinμ interaction oλ the Ca‐Cu atomic pairs. For Mμ atom, the prepeak located at . “°, which appears at  K, indicates that aλλinity oλ Mμ and Cu atoms becomes stronμer below  K. ”esides, the shoulder oλ PDF oλ Cu atom at  K is due to the λormation oλ Cu‐Cu atomic pair.

Figure . The partial pair distribution λunctions oλ Ca‐Ca, Ca‐Mμ, Ca‐Cu, Mμ‐Mμ, Mμ‐Cu and Cu‐Cu atomic pairs oλ the alloy at diλλerent temperatures.

The partial pair distribution λunctions oλ the diλλerent atom pairs such as Ca-Ca Mμ‐Mμ, Mμ‐ Cu and Cu‐Cu are shown in Figure . “ll oλ these λunctions reveal an increase in the heiμht oλ the λirst and second peaks with decreasinμ temperature, which demonstrates the increase oλ the short to medium ranμe order durinμ the solidiλication process, while the location oλ the λirst peak chanμes little at the same time. We can distinμuish splittinμs oλ the second peak oλ all six partial pair‐correlation λunctions. However, the splittinμs occur at diλλerent tempera‐ tures λor these λunctions. For the Ca‐Mμ and Cu‐Cu pairs, the splittinμ λirst occurs at  K and is well developed at  K. While λor other pairs, the splittinμs occur at lower temperature than

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

 K the experimental Tμ =   K [ ]. It demonstrates that some diλλerent substructures in preλer atom pairs have been λormed beλore reachinμ the λinal μlassy state. Figure shows the total PDF and partial PDF “l‐“l, “l‐Ni and “l‐Nd oλ “ Ni Nd alloy. The peaks about λirst and second oλ all the PDFs μet hiμher durinμ the μlass transition process, which indicates that short to medium structure orderinμs oλ the associated atom increases as the temperature decreases. The increase oλ Ni atom PDF is obvious, which leads to the sharpness peak oλ the partial PDF Ni atom at  K eventually. It determines the composi‐ tional and μeometrical order became stronμer and λormed durinμ the μlass transition process around Ni atom. Total PDF shows a splittinμ oλ the second peak at  K, and dislocation oλ the two well‐developed new peaks is around . and .  Å at  K.

Figure . The total and partial pair distribution λunctions oλ the “l Ni Nd alloy at diλλerent temperatures. a Total, b “l‐“l, c “l‐Ni and d “l‐Nd.

The splittinμ oλ the second peak with the diλλerent shapes oλ partial PDFs may caused by the diλλerent radii oλ the atoms in the atomic pairs, which indicate the complex structures oλ the disorder systems. The partial coordination numbers oλ “l‐Ni and “l‐Nd pairs estimated λrom the DRP model are . and . [ ]. The CNs oλ the “l‐Nd pair decreases with decreasinμ temperature, which determines that “l turns to preλer to be near Ni rather than Nd atoms durinμ the μlass transition process.

. Part , the evolution of structural properties during the glass transition process Honeycutt‐“ndersen H“ bond‐type index is compute to obtain detailed inλormation [ ] oλ the local structure oλ the “ Ni Nd alloy. Three‐dimensional imaμe oλ the local conλiμurations

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Metallic Glasses: Formation and Properties

is μiven durinμ the solidiλication process. The H“ bond‐type index represents the number and properties oλ common nearest neiμhbours oλ atom pair it can used to analysis the local environment surroundinμ the atomic pair, which is under consideration. Each H‐“ bond index is classiλied by the relation amonμ their neiμhbours with λour indices oλ inteμer. The λirst inteμer is iλ the pair is bonded and the considered atomic pair is closer than cut‐oλλ distance or else . The common neiμhbour oλ the considered atomic pair is deλined as the second inteμer, the number oλ bonds amonμ common neiμhbours is the third number and the λourth is used to distinμuish the atomic pair iλ the λormer three inteμers are not suλλicient. More than types oλ the bond pairs are λound in the alloy at diλλerent temperatures with the H“ bond‐type index method, which indicate the complex structures oλ the disorder systems. The variation oλ seven typical bond pairs oλ Ca‐Mμ‐Cu alloy is shown in Figure whereas the other pairs oλ , , , and types are not shown because they are quite rare and every type oλ them is < % in the whole temperature ranμe. It can be seen that pair has a λraction oλ . % at  K and increases rapidly with temperature decreasinμ . % at  K . The does not chanμe much in the whole temperature ranμe the λraction oλ pairs at  K is . %. These two pairs that contain λive coordinated vertices represent the pentaμonal bipyramids. “s the melt beinμ quenched, a larμe quantity oλ pentaμonal bipyra‐ mids structures are λormed in undercooled and μlassy alloy. Moreover, it can be observed that the quantity oλ pair, which is the most popular type at  K, decreases as the temper‐ ature decreases and the pair increases about % in the supercooled reμion.

Figure . Evolution top seven most populated H“ bond‐type index in the Ca Mμ Cu alloy as a λunction oλ tempera‐ ture.

The population oλ them, which indicate the λourλold bond structure, is much less than the most prevalent type oλ pair. In addition, the pair, which presents the sixλold bond, increases λrom . % at  K to . % at  K. Furthermore, the pair with hiμher enerμy and larμer distortion decreases obviously in the solidiλication process. Much inλormation oλ chemical and topoloμical orderinμs oλ the hiμher order correlations can be provided by the bond‐anμle distribution λunction, which is deλined about a μroup oλ three

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

atoms. One is deλined as a central atom the other two within a cut‐oλλ distance that is deter‐ mined by the location oλ the minimum oλ PDF are denoted as the side atoms. The two side atoms toμether with the central atom deλine the bond anμle. The bond‐anμle distribution λunction can be obtained by statistically summinμ the bond anμles oλ all the μroups oλ three atoms. Two peaks located at about . ° and . ° at  K oλ total bond‐anμle distribution about Ca Mμ Cu alloy Figure , and the heiμht oλ the oλ the λirst is hiμher. Peaks become hiμher durinμ the μlass transition process. The total bond‐anμle distribution have two peaks near . °and . ° at  K. Close packinμ oλ three neiμhbour atoms is suμμested by the λirst peak at around °, and the second peak located near ° is consist with the pentaμon conλiμura‐ tions. The shoulder appeared at  K at around ° relates to threeλold coordinated atoms.

Figure . Calculated total bond‐anμle distribution λunctions oλ Ca Mμ Cu alloy at diλλerent temperatures.

The peaks oλ all the partial bond‐anμle distribution λunction oλ Ca‐Mμ‐Cu alloy Figures – become more pronounced as the temperature decreases, which exhibits that the distribution oλ the local structure types around the relevant atoms decreases at lower temperature. For the bond‐anμle distribution around Ca atoms at  K, predominant anμles correspond to °, ° and ° can be seen clearly, and the λlat peaks located at °, ° and ° are shown at   K. The N‐Mμ‐N bond‐anμle distribution λunctions in the amorphous phase T =   K have a prominent peak at ° and °, while they present two broad peaks at about ° and ° at  K. In addition, the plot oλ N‐Cu‐N shows main peaks at ° and ° at  K, and the location oλ the peaks is at around °and ° at  K. The second peak oλ N‐Cu‐N shiλts to larμe anμle durinμ the solidiλication process, which is more evident than the other peaks.

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Metallic Glasses: Formation and Properties

Figure . Calculated partial bond‐anμle distribution λunctions oλ the N‐Ca‐N N = Ca, Mμ, Cu at diλλerent tempera‐ tures.

Figure . Calculated partial bond‐anμle distribution λunctions oλ the N‐Mμ‐N N = Ca, Mμ, Cu at diλλerent tempera‐ tures.

Figure . Calculated partial bond‐anμle distribution λunctions oλ the N‐Cu‐N N = Ca, Mμ, Cu at diλλerent tempera‐ tures.

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

The detailed three‐dimensional imaμe inλormation about the evolution oλ local atomic conλiμuration oλ the “l‐Ni‐Nd alloy is obtained usinμ the method oλ Honeycutt‐“ndersen H‐ “ bond‐type index durinμ the μlass transition process. The and bond pairs are the most two popular types seen λrom Figure in the molten alloy liquids, which λraction is and %, respectively. The two pairs do not chanμe very much in whole temperature ranμe. Moreover, λraction oλ pair is % at  K and increases with decreasinμ temperature rapidly which is % at  K. The and pairs represent the pentaμonal bipyramids that contain λive coordinated vertices. Durinμ the μlass transition process, in the undercooled and μlassy alloy, larμe number oλ pentaμonal bipyramid structures are λormed. The pair increases % durinμ the rapid solidiλication process. Furthermore, and pairs decreases obviously, which consists with hiμher enerμy and larμer distortion. The bond-type index around “l‐Nd atomic pair is only, due to the larμe radius oλ Nd atom and λraction oλ this index is < % durinμ the whole temperature ranμe.

Figure . Evolution oλ the top six most popular H‐“ indices types in the “l Ni Nd alloy as a λunction oλ temperature.

Figure shows the total and partial bond‐anμle distributions oλ the “l‐based alloy at diλλerent temperatures. The partial bond‐anμle distributions have λive classes “l‐“l‐“l, Ni‐“l‐“l, Nd‐ “l‐“l, N‐Ni‐“l and N‐Nd‐“l N = “l, Ni, Nd . Total bond‐anμle distribution exhibits two peaks at  K that located at . ° and . °, but the heiμht oλ the λirst peak is hiμher. The peaks become hiμher by the temperature decreases. Two peaks located . ° and . ° at  K oλ the total bond‐anμle distribution. “ll the partial bond‐anμle distribution λunctions have more pronounced peaks by μlass transition proμress, and this shows that the distribution types oλ the local structure around the atomic pair decreases at low temperature. The hump between two peaks decreases as the temperature decrease and becomes less distinμuishable oλ the “l‐ “l‐“l, Ni‐“l‐“l, Nd‐“l‐“l and N‐Ni‐“l N = “l, Ni, Nd .

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Metallic Glasses: Formation and Properties

Figure . Calculated total and partial bond‐anμle distribution λunctions oλ the “l Ni Nd alloy at diλλerent tempera‐ tures N=“l, Ni, Nd . a “l‐“l‐“l, b Ni‐“l‐“l, c Nd‐“l‐“l, d N‐Ni‐“l, e N‐Nd‐“l and λ total.

The close packinμ oλ three neiμhbour atoms is related to λirst peak at . ° position oλ the total bond anμle. For “l‐“l‐“l, it is at around ° shows the conλiμuration oλ equilateral trianμle, which is λormed by the three neiμhbour atoms at  K. The second peak oλ the total bond anμle is located near °, it indicates the pentaμon conλiμurations, which is aμreement with the exist oλ the pentaμonal bipyramids obtained by the H“ bond‐type index. The shoulder appeared at  K at around ° relates to the threeλold coordinated atoms. We also λound that the increase in the peaks heiμht oλ the bond‐anμle associate with the Ni is more evident than the others. The siμniλicant μeometrical order around Ni is λormed by the μlass transition process. The shapes oλ the peaks λor N‐Nd‐“l do not chanμe very much, indicatinμ that the neiμhbours’ order oλ the elements Nd does not chanμe much as the temperature decreases.

. Part , the dynamical properties of the nucleation and glass‐forming process of liquids The dynamical properties are important λor describinμ the nucleation and μlass‐λorminμ process oλ liquids. The variation oλ mean‐square displacement MSD as a λunction oλ temper‐ ature is also calculated throuμh Eq.

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

r 2 (t ) =

1æ N 2ö ç å ri (t ) - ri (0) ÷ N è i =1 ø

lim r 2 = c + 6 Dt x ®¥

These quantities oλ MSD [ ] λor Ca, Mμ and Cu atoms are plotted in Figure a–c. The linear behaviour can be clearly seen and the slope decreases with temperature decreasinμ λor all the three atoms. The selλ‐diλλusion coeλλicients, as the λunctions oλ temperature, are shown in Figure d. It is λound that the selλ‐diλλusion coeλλicients oλ Ca atoms are less than the others at all temperatures, which is due to the larμer radius oλ Ca atom. The relation between the selλ‐ diλλusion coeλλicients and the temperature obeys an “rrhenius relationship.

Figure . The MSD and selλ‐diλλusion coeλλicient oλ the alloy at diλλerent temperatures a – c The MSD oλ Ca, Mμ and Cu atoms d variation oλ diλλusion coeλλicients and the ln D oλ Ca, Mμ and Cu atoms at diλλerent temperatures.

“ccordinμ to the “rrhenius relationship, there will be a linear relationship between the ln D and /T, which is also plotted in Figure d, whereas, it shows non‐“rrhenius in the super‐ cooled reμion in which the selλ‐diλλusion coeλλicients decrease more rapidly than expected. “s mentioned above, the development oλ polyhedron local structures in the undercooled liquid delays the diλλusive reμime and leads to the non?“rrhenius behaviour [ ], which also inhibits the crystallization and promotes the λormation oλ the amorphous solid. The activation enerμy Ea is . , . and .  kJ/mol λor the Ca, Mμ and Cu selλ‐diλλusions, respectively. The pre‐ exponential λactor D is .  ×  - , .  ×  - and .  ×  -  m /s λor Ca, Mμ and Cu,

55

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Metallic Glasses: Formation and Properties

respectively. The polyhedral order and CSRO, which are enhanced in the supercooled reμion, decrease the selλ‐diλλusion coeλλicients.

. Part , electronic density of states of the glass transition of the alloy In order to investiμate the electronic oriμin oλ the μlass transition oλ the “l Ni Nd alloy, we calculate the DOS λor “l Ni Nd alloys at diλλerent temperatures and plot in Figure . The snapshots oλ the conλiμuration at and  K are also shown. The local DOS oλ N‐d, Ni‐d and “l‐spd oλ Figure a–c indicates that the covalent bond is λormed due to the hybridization between the electronic states oλ “l‐p and Ni‐d. Peak oλ DOS Ni atom becomes narrower and hiμher, the resonance between the electronic states oλ Ni and “l around the -  eV is more siμniλicant durinμ the μlass transition process, shows the stronμer hybridization between the electronic orbitals are λormed and more ordered local structure around Ni λormed. The lower peak oλ Nd than Ni, and the heiμht oλ the peak λor Nd increases sliμhtly indicates the transition metal oλ Ni is more active than the rare earth oλ Nd durinμ the solidiλication process. The chemical bonds contained in the snapshots d are calculated by the same cut‐oλλ distances oλ the correspondinμ atomic pairs at and  K, which presents the interaction oλ nearest neiμhbour atomic pairs are enhanced durinμ the μlass transition and the larμer number oλ bonds are λormed at the same time. It is contributinμ to the hiμher peak oλ DOS imaμe the Ni and Nd atom at the lower temperature as well.

Figure

. DOS a – c and snapshots “l

pink, Ni

blue, Nd

μreen d

 K, e

 K oλ the “l Ni Nd alloy.

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

Figure . Contour plots oλ the valence electronic‐charμe density oλ amorphous solid Ca Mμ Cu λor diλλerent atomic conλiμurations a Ca‐Cu‐Ca, b Cu‐Cu‐Ca, c Cu‐Mμ‐Ca, d Ca‐Mμ‐Ca, e Ca‐Ca‐Ca and λ Mμ‐Cu‐Cu.

. Part , electronic charge densities of the amorphous solid “ combined neutron and x‐ray diλλraction study [ ] have been perλormed λor Ca Mμ Cu bulk metallic μlass‐λorminμ system, which indicates that the averaμe distances oλ Cu‐Mμ and Cu‐Ca are consistent with the sum oλ covalent radii, whereas all other interatomic distances are consistent with the sum oλ metallic radii. Our simulation shows an accumulation oλ charμe between Ca‐Cu and Mμ‐Cu atomic pairs, which evidence the persistence oλ Ca‐Cu and Mμ‐ Cu covalent bonds Figure . On the other hand, the covalent bonds oλ Cu‐Cu pair are also existed as presented in the plot. The Cu‐Cu interaction is not as attractive as Cu‐Ca and Cu‐ Mμ. In other words, Cu atoms preλer to cluster with Ca and Mμ atoms rather than Cu itselλ. ”ut this does not mean that Cu atoms repel each other. This is what the CSRO tells us. “s a result, Cu‐Cu covalent bonds are still λound. The electrons with an sp character oλ the Mμ atoms are transλerred to the d states oλ Ca and Cu atoms. The accumulation oλ charμe between Ca‐ Mμ pairs indicates that there is Ca‐Mμ covalent bond in this alloy as well. While, this phe‐ nomenon is not λound about Ca-Ca pair, which presents that the Ca‐Ca bond is metallic bond. Our result provides strinμ evidence λor the existence oλ Ca‐Mμ, Ca‐Cu, Mμ‐Cu and Cu‐Cu covalent bonds, which is in accord with that covalent bonds between the elements dominate in the Ca‐Mμ‐Cu alloy proposed by the λormer experiment. “n increasinμ λraction oλ covalent bondinμ is beneλit to increase the μlass-λorminμ ability [ ], and this is in λavour oλ the μlass λormation λor the Ca Mμ Cu alloy.

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Metallic Glasses: Formation and Properties

. Part , evolution of the electronical properties of the metallic glass In our simulation, we also study the μround‐state electronic charμe densities. The charμe density ρ r in a plane deλined by three neiμhbourinμ atoms at three temperatures oλ , and  K are plotted in Figure . The Ni ions tend to λorm stronμ interactions with “l, by the neutron diλλraction data report [ ], whereas the Nd bond lenμth to “l is close to the expected sum oλ atomic radii about neutral ions. The accumulation oλ charμe between Ni‐“l atomic pairs at all temperatures is shown in the imaμe it indicates persistence oλ the Ni‐“l covalent bonds. “ll diλλerent atomic conλiμurations are chosen oλ the most nearby atomic pairs at the correspondinμ temperature. The interatomic distance oλ  K is less than the others due to the broad distribution oλ interatomic distance at hiμh temperature Figure . The accumulation oλ charμe between “l‐“l atomic pair and Nd‐“l atomic pair at  K is more evident than the other temperature. While eλλect oλ temperature λor Ni‐“l atomic pair is insiμniλicant, as a result oλ durinμ the μlass transition process “l atom preλer to transλorm near Ni atom than “l atom and Nd atom. The “l‐“l and Nd‐“l atomic pairs are weakly covalent bond properties in the molten liquid states, whereas Nd‐“l is metallic bond properties at  K because oλ that the electronic charμe density around Nd atom is a spherical distribution almost. The electrons with a sp character oλ the “l atom are likely transλerred to the d states oλ Ni atoms at  K. The increasinμ λraction oλ covalent bond and metallic bond can increase the GF“ oλ the alloy, and it is in λavour oλ the μlass λormation durinμ the solidiλication process oλ the multicomponent molten alloy.

Figure . Contour plots oλ the valence electronic charμe density oλ “l Ni Nd alloy λor diλλerent atomic conλiμurations at diλλerent temperatures.

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

. Conclusion “b initio molecular dynamic simulations are used to understand the evolution oλ structural and electronic properties durinμ the μlass transition process oλ the “l Ni Nd and Ca Mμ Cu alloy based on the density λunctional theory. The pair correlation λunction, CN, diλλusion coeλλicient, mean square displacement, H“ bond‐type index and bond‐anμle distribution at diλλerent temperatures are observed by our simulation. The PDF imaμes investiμates that the short to medium structure orderinμs are enhanced with decreasinμ temperature durinμ the rapid solidiλication process. The interaction between diλλerent atomic pairs is strenμthened, and the splittinμ oλ the second peaks occurs at diλλerent temperatures λor the partial PDF, which indicates that some diλλerent substructures in preλer atom pairs have been λormed beλore reachinμ the λinal μlassy state. The CNs we calculated aμree well with the previous report oλ Ca‐Mμ‐Cu MGs [ ]. The CN oλ Ca atoms chanμes λrom . to . , it is λrom . to . λor Mμ atoms and it is λrom . to . λor Cu atoms, which are very close to the previous theoretical result. “ccordinμ to the ECP model [ , ], which allows the number oλ structural sites to be counted, each Mμ atom will be surrounded by Ca atoms in the binary Ca‐Mμ MGs, and three additional solute sites occur at the interstices oλ these Mμ‐centred clusters. The CNs around Ca atom are about while which is more than result oλ in the Ca‐Mμ‐Zn MGs [ ], which shows that the packinμ around Ca atoms is more eλλicient in the Ca‐Mμ‐Cu alloy. The interatomic distance we obtained and the partial coordination numbers oλ “l‐Ni and “l‐Nd pairs in our simulation are consistent with the previous experimental oλ the neutron diλλraction and theoretical results oλ “l‐Ni‐Nd alloy [ ]. The pentaμon conλiμurations and the three neiμhbour atoms packinμ are the most primary short‐ranμe order in both these two alloys, which local structural order is enhanced obviously by the decreases oλ temperature durinμ the μlass transition proμress concluded by the bond‐ anμle distribution methods in these two alloys, which is consist with results evaluated by the Honeycutt‐“ndersen H‐“ bond‐type index that the amount oλ pentaμon conλiμurations increases durinμ the μlass transition and become the primary short‐ranμe order SRO as temperature decreases. Our modellinμ studies oλ multicomponent oλ “l‐based MGs and Ca‐based MGs show that the covalent bonds oλ Ca‐Mμ, Ca‐Cu, Mμ‐Cu and Cu‐Cu pairs are existed in Ca‐Mμ‐Cu alloy and accumulation oλ charμe between Ni‐“l atomic pairs which evidences the persistence oλ the Ni‐ “l covalent bonds in “l‐Ni‐Nd alloy at all the temperatures. The evolution oλ structural and electronic properties durinμ the μlass transition process oλ “l‐ Ni‐Nd alloy is also simulated. “ result oλ the λact is that “l turns to preλer to be near Ni than “l and Nd durinμ the μlass transition process. In the molten liquid oλ  K, weakly covalent properties oλ “l‐“l and Nd‐“l atomic pairs are existed, while the electronic charμe density around Nd is almost a spherical distribution at  K. Glass‐λorminμ ability GF“ is also another important research issue. Inoue have provide three empirical rules λor ”MG λormation [ , ] which suμμested that multicomponent system oλ

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Metallic Glasses: Formation and Properties

more than three elements, diλλerence oλ atomic size mismatch larμe than %, and neμative heat oλ mixinμ between the components atoms tends to transλorm to μlass easily oλ this multicomponent alloy system. The λormation oλ the crystal structures is suppressed by the increasinμ polyhedral local structures orderinμ and chemical SRO oλ covalent bonds and metallic bonds the increasinμ oλ the electronical interaction durinμ the rapid solidiλication processes plays an important role durinμ the μlass transition and increases the μlass‐λorminμ ability oλ the “l‐Ni‐Nd alloy. We also λound that the electrons with a sp character oλ the “l atoms are more likely to transλer to the d states oλ the Ni atoms. The more disordered amorphous state and the optimum bondinμ state seem to achieve the conλused compositions and multicomponent MGs with μood μlass‐λorminμ ability and unique mechanical properties.

Author details HaiJun Chanμ “ddress all correspondence to chanμhaijun

@

.com

School oλ Materials Science and Enμineerinμ, ShanμHai JiaoTonμ University, Shanμhai, China

References [ ] H. J. Chanμ, L. F. Chen and X. F. Zhu, J. Phys. D “ppl. Phys.

,

[ ] H. J. Chanμ, L. F. Chen, and X. F. Zhu, J. “ppl. Phys. . / . .

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[ ] O. N. Senkov, D. ”. Miracle, and J. M. Scott, Intermetallics. [ ] “hn K, Louca D, Poon S J and Shiλlet G J

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[ ] J. D. Honeycutt and H. C. “ndersen. J. Phys. Chem.

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[ ] J. P. Hansen and I. R. McDonald, Theory oλ Simple Liquids “cademic, London, [ ] “. Pasturel, E. S. Tasci, M. H. F. Sluiter, and N. Jakse, Phys. Rev. ”

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[ ] E. R. ”arney, “. C. Hannon, O. N. Senkov, J. M. Scott, D. ”. Miracle, and R. M. Moss, Intermetallics. , – . [ ] R. W. Cahn, P. Haasen, and E. J. Kramer, In μlasses amorphous metals, edited by Weinheim VCH, NY, .

Structure of the Metallic Glass and Evolution of Electronical Properties during Glass Transition in Atomic Level http://dx.doi.org/10.5772/63676

[

] K. “hn, D. Louca, S. J. Poon and G. J. Shiλlet. J. Phys. Condens. Matter. .

[

] “hn K, Louca D, Poon S J and Shiλlet G J

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] D. ”. Miracle, D. Louzμuine‐Luzμin, L. Louzμuina‐Luzμina, and “. Inoue, Int. Mater. Rev. , .

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] D. ”. Miracle, “cta Mater.

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] O. N. Senkov, D. ”. Miracle, E. R. ”arney, “. C. Hannon, Y. Q. Chenμ, and E. Ma, Phys. Rev. ” , .

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61

Chapter 4

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses Shanlin Wang Additional information is available at the end of the chapter http://dx.doi.org/10.5772/63677

Abstract Metallic μlasses exhibit excellent corrosion resistance and electrocatalytic properties, and present extensive potential applications as anticorrosion, antiwearinμ, and catalysis materials in many industries. The eλλects oλ minor alloyinμ element, microstructure, and service environment on the corrosion resistance, pittinμ corrosion, and electrocatalytic eλλiciency oλ metallic μlasses are reviewed. Some scarcities in corrosion behaviors, pittinμ mechanism, and eletrocatalytic reactive activity λor hydroμen are discussed. It is hoped that the overview is beneλicial λor some researcher payinμ attention to metallic μlasses. Keywords: metallic μlass, corrosion resistance, pittinμ corrosion, electrocatalytic prop‐ erty

. Introduction Except λor hiμh compression strenμth, microhardness, electrical resistivity, and μood soλt maμnetic properties, most metallic μlasses exhibit excellent corrosion resistance. The excel‐ lent corrosion resistances oλ metallic μlasses are mainly attributed to the homoμeneous sinμle μlass phase, the alloy chemistry, and the presence oλ metalloids [ – ]. No μrain boundaries, dislocations, and other deλects where corrosion can occur preλerentially are expected to allow the μrowth oλ a uniλorm protective λilm. The chemical homoμeneity is believed λor rapid coolinμ rates required to produce λull amorphous structure since no enouμh time is availa‐ ble λor solid-state diλλusion, that is, it is impossible λor the λormation oλ second phases, precipitation, and seμreμations. The homoμeneity in chemical composition and microstruc‐ ture promotes amorphous oxide λormation on the surλace which retards ionic transport. The

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Metallic Glasses: Formation and Properties

improvement oλ corrosion resistance is also considered to link to the ability oλ these metasta‐ ble alloys to λorm supersaturated solid solution in one or more alloyinμ elements. The alloyinμ element available in solid solution may be incorporated into the oxide λilm to enhance its passivity. Thus, the eλλect oλ the amorphous structure, chemical and structural homoμeneity, and the possibility oλ λorminμ unique chemical composition not typical oλ near-equilibrium crystalline alloys are mostly considered as λactors that can aλλect the corrosion properties oλ metallic μlass.

Figure . “ schematic diaμram oλ potentiodynamic polarization a the theoretical anodic polarization curve, b the calculation oλ corrosion potential and corrosion current density.

In order to estimate the corrosion resistance, immersion test is one oλ method to calculate the averaμe corrosion rate in one year, while the electronic chemistry methods such as the potentiodynamic polarization are applied in most researches, where the considerable inλor‐ mation on the electrode processes can be attained, such as corrosion potential Ecorr , corrosion current density Icorr , corrosion rate, pittinμ susceptibility, passivity, and the cathodic behavior. “ schematic curve oλ the theoretical anodic polarization a and the calculation oλ corrosion potential and corrosion current density b are illustrated in Figure . “s can be seen in Figure a , the scan start λorms point and proμresses in the positive potential direction until termination at point , the open circuit potential is located at point “. “t the potential, the sum oλ the anodic and cathodic reaction rates on the electrode surλace is zero. “s a result, the measured current will be closed to zero. With the increase oλ the potential, it moves to active reμion. In this reμion, metal oxidation is the dominant reaction takinμ place. Point ” is known as the passivation potential, and as the applied potential increases above the value the current density is seen to decrease until a low, passive current density is achieved in passive reμion E. Once the potential reached a suλλiciently positive value, that is located as point C, sometimes termed the breakaway potential, the applied current rapidly increases. “round the open circuit potential, a new line is λitted accordinμ to the linear reμions oλ the polarization curve as illustrated in Figure b . The current density at that point is the corrosion current density

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

Figure . “ schematic cyclic-anodic-polarization curve [ ].

Icorr and the potential at which it λalls is the corrosion potential Ecorr . It is μenerally aμreed that the hiμher is the corrosion potential, the more diλλicult is the occurrence oλ the oxidation reaction λor the metals, moreover, the larμer is corrosion current density, the hiμher is corrosion rate, that is, the lower corrosion resistance λor metallic μlass. While λor the susceptibility oλ pittinμ corrosion, the cyclic-anodic-polarization is usually measured, and some parameters and typical characteristics with reμard to pittinμ corrosion susceptibility are deλined in the schematic polarization curves oλ Figure [ ]. “ potential scan is started below the corrosion potential, Ecorr. “t Ecorr, the current density μoes to zero, and then increases to a low and approximately constant anodic value in the passive ranμe. In this ranμe, a thin oxide/hydroxide λilm, a passive λilm, protects the material λrom hiμh corrosion rates. Iλ the current density decreases when the potential scan direction is reversed, as in path , the material is shown to be immune to pit corrosion. However, iλ on the potential up scan, the current density suddenly increases, and remains hiμh on the down scan, until λinally decreas‐ inμ to the passive-reμion value, as in path , the material is shown to underμo a λorm oλ pittinμ corrosion. The potential at which the current density suddenly increases pittinμ initiation is known as the pit potential, Epit, and the potential at which the current density returns to the passive value is known as the repassivation potential or the protection potential, Epp. ”etween Epit and Epp, pits are initiatinμ and propaμatinμ. In the case oλ path , pits will not initiate at Ecorr, the natural corrosion potential and, thereλore, the material will not underμo pittinμ corrosion under natural corrosion conditions. Iλ, on the other hand, path is exhibited, where Epp is below Ecorr, the material will underμo pittinμ corrosion at surλace λlaws or aλter incubation time periods at Ecorr. In terms oλ the overall resistance to pittinμ corrosion, two parameters oλ Epit − Ecorr and Epp − Ecorr are important. Hiμher values oλ both are desirable to reλlect hiμh values oλ Epit and Epp relative to Ecorr.

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Metallic Glasses: Formation and Properties

. Corrosion resistances of nonferrous metallic glasses Metallic μlass is comparatively newcomer to the amorphous material μroup, which is λabri‐ cated λrom a cooled liquid without crystallization under a rapid coolinμ rate. “s the λirst metallic μlass oλ “u Si was discovered in by Duwez and coworkers [ ], a series oλ metallic μlasses such as Zr-, Ti-, Pd-, Cu-, Fe-, and Mμ-based alloys are successλully λabricated by the method oλ melt quenchinμ. In order to extend the industrial application oλ metallic μlasses, the corrosion behaviors oλ metallic μlass have been oλ μreat interest. The corrosion resistance oλ nonλerrous metallic μlass oλ Cu-, Ti-, Zr-, and Mμ-based alloys will be discussed in the λollowinμ part. . . Effect of composition “monμ oλ nonλerrous metallic μlasses, Zr-based metallic μlass Zr-MG is investiμated abroad in corrosion resistances. “ddition oλ minor element such as “μ, Cu, Y, Ti, Ni, and Nb has been utilized to enhance the μlass λorminμ ability and resistance to μeneral and local corrosion. Inoue and coworkers [ – ] have investiμated the eλλect oλ Ni on the corrosion resistance oλ Zr-NiCu-“l alloy. Zr-MG without Ni shows lowest corrosion potential and no obvious passivation reμion, but pittinμ directly can be observed as potential rises. Zr-MG with Ni is spontaneously passivated with current density around − “/m beλore the occurrence oλ pittinμ corrosion in chloride solution. Since the Cu element in the Zr-MG is easily dissolved in chloride solutions, thus leads to a low corrosion resistance. The additional Ni inhibits the λormation oλ soluble Cu-Cl λilms and λacilitates λorminμ the protective surλace λilms with a hiμh concentration oλ Zr cation, leadinμ to a denser, thicker, and more pittinμ resistance ZrO passive λilm. Zhanμ and coworkers [ , ] reported that partial substitution oλ Ni and Co by “μ was eλλective in improvinμ the corrosion resistance oλ Zr-MG, as the “μ addition increases the concentration oλ Zr and decreases the concentration oλ “l in the surλace passive λilms, while Liu and coworkers [ ] λound that the addition oλ “μ could promote the λormation oλ “l O but sliμhtly suppressed the λormation oλ ZrO . The cast Zr “l Co −xNbx x = , , and at% metallic μlass are spontaneously passivated in NaCl solution with a passive current density between − and − “/cm , and the pittinμ potentials shiλt to positive direction with the increasinμ oλ Nb content [ ]. Thouμh the Nb-bearinμ alloy’s pittinμ potential and passive reμion are larμer than Nb-λree alloy, aλter pittinμ, however, the alloy with at% Nb exhibits hiμher corrosion current density than Nb-λree alloy, as shown in Figure , meaninμ that the corrosion reactions in Nbbearinμ alloy are more severe at hiμh potential [ ]. Generally, the addition oλ Nb element in Zr-MG can λacilitate the λormation oλ hiμhly protective Zr-, “l-, and Nb-enriched surλace λilm, while Cu addition will deteriorate the passivation [ ]. For Zr “l Ni Cu metallic μlass alloy, the mass loss and the averaμe corrosion rate decrease with increase oλ Ti content in M HCl solution [ ], and the Ti addition improves the stability oλ passive λilm and pittinμ resistance, while they are susceptible to pittinμ corrosion [ ]. Compared to Zr-MG, Cu-based metallic μlasses with their superiority in price and mechanical properties possess μreat potential applications in the λields, such as bipolar plate materials, biomedical instruments, and microdevices. The investiμations about corrosion resistance have

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

been carried extensively. Small addition oλ Nb, Cr, Ta, and Mo has proved to be eλλective in improvinμ the corrosion resistance [ – ]. “sami et al. [ ] investiμated the eλλect oλ small addition oλ Nb, Mo, and Ta to Cu Zr Ti at% metallic μlass in M HCl, HNO , NaOH, and . M NaCl solutions. The results demonstrate that Nb element is most eλλective in decreasinμ the corrosion rate in all oλ the solutions, moreover, the corrosion rate decreases with increasinμ the Nb content. The minor element addition can enhance the stability oλ passive λilm enriched in ZrO and TiO . Except the minor addition Mentioned above, some rare metal element additions oλ In, Y, Ce, and Ln to Cu-MG are eλλective to improve the corrosion resistance [ – ]. The results demonstrate that the dissolution oλ rare element is λavorable to λorminμ continuous Zr-, Ti-rich protective oxide λilm and alleviates the local corrosion and propaμation at the initial corrosion staμe. The Ln addition can increase the nearest neiμhbor atomic distance aλλectinμ the topoloμical instability, which is attributed to the improvement oλ corrosion resistance.

Figure . Potentiodynamic polarization curves λor the alloy in phosphate-buλλered solutions at

°C [

].

“s conventional titanium alloy, Ti-based metallic μlass with hiμh yield strenμth, low Younμ’s modulus, hiμh corrosion resistance can be applied as biomaterial [ ], and mostly possesses hiμher corrosion resistance than Ti- “l- V alloy in a simulated body λluid environment. The minor element addition oλ Zr, Nb, and Cu will chanμe the corrosion behavior oλ Ti-MG [ – ]. Nb addition can enhance the pittinμ resistance due to an improvement oλ the passive layer properties λor near-homoμenous alloys. Small addition oλ Zr promotes the corrosion potential and decreases the corrosion current density. The addition oλ Cu can shiλt the beμinninμ oλ polarization reaction to a positive voltaμe level, while provokes severe Cu-induced selective dissolution under the hiμher applied voltaμe levels, resultinμ continuous pittinμ and the depletion oλ Ti and Zr in the alloy. With increase oλ Ti + Zr /Cu ratio, the pittinμ corrosion resistance is μreatly enhanced due to the λormation oλ surλace λilm mainly composed oλ TiO and ZrO .

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68

Metallic Glasses: Formation and Properties

“s well known, maμnesium alloy is also one oλ biomaterials. Recently, the corrosion behavior oλ Mμ-based metallic μlass is investiμated. Wanμ et al. [ ] reported that the Mn addition can promote the λormation oλ a dense passive λilm, which delays the corrosion oλ the matrix. The Zn addition provokes the λormation oλ Zn HO and Mμ OH , and the evolution oλ the corrosion process oλ the MμZnCa μlass is schematically illustrated in Figure [ ]. When MμMG is immersed in body λluid, the anodic dissolution oλ maμnesium occurs and the maμne‐ sium hydroxide layer well is λormed on the surλace oλ the sample. The attack oλ Cl− occurs at the weak sites oλ the maμnesium hydroxide layer and transλorms the maμnesium hydroxide into soluble maμnesium chloride. The λresh substrate, exposed to the medium directly, suλλers λurther corrosion, and results in the releasinμ oλ Mμ + and Zn +, as shown Figure a . “s immersion prolonμed, the Zn + concentration is increasinμ due to the continuous dissolution oλ Zn. The Zn OH precipitates preλerentially, compared to that oλ Mμ OH Figure b , the Zn OH precipitations will repair the deλects in the surλace layer, and then λorms a continuous and uniλorm layer. With the corrosion proceedinμ, the corrosion product layer will be thickened and Zn OH precipitation spreads, which are evidently depicted in Figure c . Meanwhile, the undissolved Mμ OH and Zn OH precipitation can provide λavorable sites λor apatite nucleation. With Ti addition, the protective λilm oλ Mμ OH will enrich Ti, im‐ provinμ the stability [ ].

Figure . The sketch map λor the evolution oλ corrosion process oλ Mμ-Zn-Ca bulk metallic μlass immersed in S”F a initial staμe, b middle staμe, c λinal staμe [ ].

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

. . Effect of microstructure The microstructure and composition homoμeneities are destroyed with the crystallization, which is necessary to deteriorate the corrosion resistance oλ metallic μlasses. Zr “l Co metallic μlass exhibits a decrease oλ passivation potential and an increasinμ oλ penetration rate with increasinμ heatinμ temperature in Rinμer’s simulated body λluid at room temperature [ ]. The corrosion parameters oλ some metallic μlasses are summarized in the Table [ – ]. It can be attained that the corrosion resistance oλ most metallic μlasses aλter crystallized will decrease, as shown that the corrosion potential decreases and the corrosion current density increases relatively, suμμestinμ that the passive λilms λormed on the surλace oλ the μlassy alloy in the anodic process are protective and denser than those on the crystal alloys. However, another metallic μlasses exhibit more positive corrosion potential and low corrosion current density aλter crystallized, meaninμ that the crystalline alloy possesses excellent corrosion resistance, compared to metallic μlass with same composition, since the nanocrystal phase such as a-Ti, CuZr precipitation or the reduction oλ the λree volume in amorphous state that in turn reduces the averaμe atomic distance. Composition 

State Ecorr

Icorr A/

mV  

Epit

cm  

Epit−Ecorr Ipass A/

mV  

mV  

CPR

cm  

um/y

Temp

Solution 

K 

  Zr [

Cu

.

.

Fe . “l . “μ .

] 

Ti Zr Cu Pd [

Ti Zr Si Ta [

] 

] 

Zr Cu “l Fe Ti [

Zr Ni [

Cu

.

Zr

] 

.

“l [

] 

] 

“m.  −



. ×











--- 

.  



P”S 

Cry.  −



. ×











--- 

.  



P”S 

--- 

--- 

--- 

--- 

--- 

--- 

“m.  −



. ×





--- 

Cry.  −





--- 

H”SS 

. ×



--- 

--- 

“m.  −



. ×







--- 

. ×



  --- 



S”F 

Cry.  −



. ×







--- 

. ×



  --- 



S”F 

“m.  −



. ×





--- 

--- 

--- 

--- 



S”F 







--- 

--- 



S”F 

--- 

--- 



S”F 

Relx  −



. ×



Cry.  −



. ×









H”SS 

“m.  −



. ×







--- 

. ×



  --- 



. M NaCl 

Cry.  −



. ×







--- 

. ×



  --- 



. M NaCl 

“m.  −



. ×









. ×



  --- 



“SS 

Cry.  −



. ×









. ×



  --- 



“SS 

Table . Summary oλ corrosion parameters λor some metallic μlasses and its crystalline alloys λrom literature reports.

Wanμ et al. [ , ] reported that the corrosion resistance oλ Mμ-MGs was sliμhtly reduced when the in situ second phase or reinλorcement phase were induced into metallic μlass. It is believed, when corrosion is developinμ, the continuous distribution oλ μlass matrix miμht be able to prevent corrosion λrom spreadinμ λrom one a μrain to another a μrain directly across

69

70

Metallic Glasses: Formation and Properties

the μlass matrix. Then corrosion is stopped aλter the crystalline phases dissolves and a continuous μlass matrix is exposed to solution. Iλ the crystalline phase is nanoparticle and presents hiμh chemical potential, the corrosion resistance oλ metallic μlass composite will not reduce, even increase λor some metallic μlass alloy [ , ]. When bulk metallic μlass is λabricated into metallic μlass coatinμ, the corrosion resistance oλ metallic μlass coatinμ is aλλected not only by the composition, but also by the surλace and porosity oλ the metallic μlass coatinμ [ ]. The eλλect oλ porosity on corrosion resistance λor TiMG evaluated with potentiodynamic polarization is shown in Figure [ ]. The metastable current transition oλ diλλerent maμnitude can be observed λor the porous bulk metallic μlass. “lthouμh rapid increase in anodic current due to pittinμ is not observed, anodic current density sliμhtly increases, indicatinμ that some oλ metastable pittinμ occur within the pore at the same time and aλterward are stabilized by the pore wall durinμ the anodic polarization process. Undoubtedly, the existinμ oλ pores would result in crevice corrosion, where potentiodynamic polarization curve exhibits a slow increase oλ current density in the anodic polarization part. Gebert et al. [ ] reported that the state oλ surλace λinishinμ oλ Zr-based metallic μlass remark‐ ably inλluences its corrosion and passivity. It is considered that the smoothness, homoμeneity, and the modiλication oλ surλace chemistry such as Cu concentration on the surλace oλ Zr-MG are modiλied aλter polished with diλλerent polishinμ materials.

Figure . Potentiodynamic polarization curves oλ the produced porous Ti Zr Cu Pd Sn bulk metallic μlass with var‐ ious porosities in Hanks’ solution at K compared to pure Ti alloy [ ].

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

”esides chemical and physical deλects oλ μlassy alloy, the mechanically μenerated deλects can enhance the corrosion susceptibility. Gebert et al. [ ] reported that a sliμht improvement oλ spontaneous passivity but a decrease oλ resistance aμainst chloride-induce pittinμ were detected when Zr-based bulk metallic μlass was shot-peened with lonμ time, and the corrosion damaμe evolution was μoverned by the nature oλ the mechanically μenerated deλects, such as craters, cracks oλ scratches and their surroundinμ stress λields. The eλλect oλ shear bands breakinμ throuμh a sample surλace on corrosion processes in acidic environments is investi‐ μated. The preλerential sites λor corrosion initiation and propaμation are λormed alonμ the shear bands, as shown Figure [ ]. The local chemical and structural chanμes in the close vicinity oλ the shear band zone are mainly predisposinμ λactor. “n et al. [ ] λound that the Cu-MG aλter tensioned exhibited more neμative corrosion potential and larμer corrosion current density in chloride solution, which indicated the deterioration oλ corrosion resistance oλ CuMG tensioned, compared with as-cast.

Figure . SEM imaμes oλ corroded shear band reμions at surλaces oλ predeλormed Zr-based metallic μlass aλter expo‐ sure to M HCl a, b bent ribbon c, d lateral area [ ].

. . Effect of environment Thouμh the microstructure and chemical composition oλ nonλerrous aλλect the corrosion resistance, it is evident λrom Table that the environment is also a siμniλicant λactor in the corrosion properties. Table provides corrosion parameters λor some nonλerrous metallic μlass in various electrolytes. Pourμashti et al. [ ] λound that Zr-MG exhibited excellent corrosion resistance in . % NaCl solution, and showed better corrosion resistance than L in HNO and H SO solutions. It can be seen λrom Table that the tendency oλ corrosion current density Icorr is Icorr/HCl > Icorr/HNO > Icorr/NaCl > Icorr/H SO . The Zr-MG reveals to increase in passive current density and decrease oλ transpassive potential with increase in nitric acidic concen‐ tration [ ]. Gebert et al. [ , ] reported that Zr Cu “l Ni was immune to localized corrosion in alkaline solution. However, a susceptibility to pittinμ corrosion was observed

71

72

Metallic Glasses: Formation and Properties

durinμ anodic polarization experiment at chloride concentration as low as − M. Moreover, the eλλect oλ temperature on the corrosion resistance was investiμated. The Zr-MG exhibits that Epit decreases as the temperature increases, indicatinμ an increased tendency oλ pittinμ in the chloride solutions as the temperature is increasinμ.

Composition 

Ecorr

Icorr A/

mV  

Epit

cm  

Epit−Ecorr Ipass

mV  

mV  

CPR

A/cm   um/y

Temp

Solution 

K 

  Zr

.

Ti

.

Ni Cu

.

Zr Cu “l Ni [

”e

.

[

] 

Zr Ti Cu “l Ni [

]  −



. ×





--- 

--- 





. % NaCl 





. ×



  --- 

--- 

--- 





M HNO  





. ×



  --- 

--- 

--- 





. M H SO  





. ×



  --- 

--- 

--- 





M HCl 

--- 

--- 

--- 

--- 

--- 

--- 

] 



. ×





--- 

--- 

--- 



.

M NaCl 

--- 

--- 

--- 



.

M NaCl 



--- 

--- 

--- 



.

M NaCl 



--- 

. ×

--- 



M HNO  

--- 



M HNO  

--- 



  − −







. ×







--- 

. × −



. ×







--- 

] 

.   −

Ti Cu [

.

Zr

.

Co Sn Si “μ

] 

Mμ Zn Ca [

] 

.  



. × −

Cu Zr T [



. M HNO  



. ×



  --- 

--- 

--- 

.  



.

M HCl 

. ×



  --- 

--- 

--- 

.  



.

M HCl 



.  

. ×



  --- 

--- 

--- 





. M HCl 



.  

. ×



  --- 

--- 

--- 





M HCl 

.  

. ×



  --- 

--- 

--- 

.  



.

M NaCl 



.  

. ×



  --- 

--- 

--- 

.  



.

M NaCl 



.  

. ×



  --- 

--- 

--- 

.  



. M NaCl 



.  

. ×



  --- 

--- 

--- 

.  



M NaCl 



.  

. ×



  --- 

--- 

--- 

--- 





.  

. ×



  --- 

--- 

--- 

--- 



. wt% NaCl 



.  

. ×



  --- 

--- 

--- 

--- 



M HCl 



.  

. ×



  --- 

--- 

--- 

--- 



M NaOH 





--- 

--- 

--- 



S”F 





--- 

--- 

--- 



P”S 



   

   

P”S 

Table . Summary oλ corrosion parameters λor some metallic μlasses in diλλerent corrosive environment.

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

“s similar to Zr-MG, Cu Zr T metallic μlass durinμ the potentiodynamic polarization exhibits the active dissolution state in the whole anodic reμion in diλλerent solutions [ ]. The current density increases to a very hiμh value aλter the Ecorr is reached indicatinμ hiμh rate oλ metal dissolution and no siμn oλ passivity is observed. The hiμher the concentration oλ chloride, the hiμher is the Icorr value which in turn indicates that the rate oλ corrosion increases with increase in concentration oλ Cl−. The Ecorr becomes more neμative with the increase in concen‐ tration oλ solution. The tendency oλ Icorr and Ecorr also can be seen in Table λor Cu-, Ti- and Mμ-based metallic μlass.

. Corrosion resistance of Fe-based metallic glass Due to its hiμh strenμth, μood soλt maμnetic properties, excellent corrosion resistance, and low producinμ cost, Fe-based metallic μlass is attended extensively to the researchers in material science and technoloμy λields around the world. ”esides the μlass λorminμ ability, strenμth, and soλt maμnetic properties, the investiμation on corrosion resistance is interestinμ λor the industrial application oλ Fe-based metallic μlasses. . . Enhance of minor element addition The eλλects oλ the addition oλ a small amount oλ metallic elements such as Cr, Mo, Nb, W, Ni, Ta, Y, “l, Co, and Mn on the corrosion resistance oλ Fe-MGs are investiμated by means oλ electrochemical polarization and weiμht loss measurements. It is well known that chromium is an eλλective element enhancinμ corrosion resistance oλ Fe-MGs. In Fe-Co-”-Si-Nb metallic μlass [ ], the corrosion rate decreases λrom . mm/year λor Cr-λree alloy to × − mm/year λor the alloy with at% Cr in . M NaCl solution at K. For the Fe . Si . ” Nb Cu metallic μlass in marine environments, the corrosion rate is times lower λor the material with at% Cr and times lower λor the material with and at% Cr as compared with the material in amorphous state without Cr [ ]. Thouμh increasinμ Cr concentration up to at% Cr tends to stabilize the passive layer, the corrosion rate remains very hiμh. The addition oλ at% is not suλλicient to the λormation oλ a stable passive layer, and the materials are dissolved or underμo severe attack in , , and M H SO [ ]. In . M H SO solution at °C, the Fe . C . Si . ” . P . Cr . “l . metallic μlass exhibits hiμh corrosion rate oλ mm/year that is λive times lower than that oλ “ISI L [ ]. The crack width on the corrosion product layer aλter potentiostatic polarization measurement decreases with increase oλ content as shown in Figure a and b , the pittinμ morpholoμy occurs on the surλace oλ metallic μlass with . at% Cr, as showed in Figure c , however, a homoμenous surλace without cracks or pittinμ is λormed λor the alloys with Cr content exceedinμ . at% as shown in Figure d and e , due to the λormation oλ Cr-enrich passive λilm on the surλace.

73

74

Metallic Glasses: Formation and Properties

Figure . SEM microμraphs on the surλaces in . M H SO solution at e x = . [ ].

K a x= . , b x= . , c x= . , d x= . ,

With minor addition oλ Y, not only μlass λorminμ ability, but also corrosion resistance is increasinμ evidently. The dependence oλ the electrochemical parameters upon the yttrium content is shown in Figure [ ]. The corrosion current density Icorr, passive current density Ipass, corrosion potential Ecorr λrom the polarization behavior and open-circuit potentials OCP oλ FeCrMoC”Y metallic μlass aλter immersion in M HCl solution λor h as a λunction oλ Y content are presented, respectively. It can be seen that the passive current density is sensitive to the yttrium content. The eλλects oλ some metal element additions are summarized in Table . It is obvious that minor element addition into Fe-MG will evidently increase the corrosion resistance. The corrosion rate oλ Fe”SiNb alloy with . at% Ni addition is about times lower than that without Ni addition in . M NaCl [ ]. Generally, the minor element additions such as Mo, Y will provoke the λormation oλ passive λilm, resultinμ in improvement oλ corrosion resistance [ , ].

Figure . The statistical analysis oλ the various electrochemical parameters obtained λrom polarization behavior oλ im‐ mersion tests oλ Fe-MG with the variation oλ yttrium content [ ].

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

Composition 

Elem. X Content Ecorr Icorr A/ at%   mV   cm  

Epit Ipass A/ mV   cm  

Fe”SiNbX [

Ni 

--- 

--- 

--- 

. × −   . × −   --- 

Fe”CuX [

] 

] 

FeCrMo”CX [

Fe”CrX [

] 

] 

FeCSi”P“lMoCoX [ ] 

Nb, Zr  Nb, Mo  Zr. Mo  Mo  Nb 







. , .   ---  . , .   ---  . , .   ---    ---    −  

. × −   . × −   . × −   . × −   . × −   . × −   . × −   . × −   . × −   ---  ---  ---  ---  --- 

.  





.  











.  





.  











.  





.  







−  

--- 





--- 

--- 

--- 

---  ---  --- 

---  ---  --- 

---  ---  ---  ---  − − − −

             

--- 

Mo,Nb  . ,  

--- 

--- 

--- 

Mo,Nb  , .  

--- 

--- 

--- 

Mp,Nb  . , .   Cr   

--- 

--- 

--- 

. × −   . × −   . × −   . × −   . × −   . × −   . × −   . × −   ---  ---  --- 

--- 

.  















   

.  

] 

--- 

. × −   . × −   . × −   . × −   . × −   ---  ---  ---  ---  . × −   . × −   . × −   . × − *   . × − *   . × − *   . × − *   . × −   . × −   . × −   . × −   . × −   . × −   --- 

--- 

.  

FeCrMoCX [

--- 

--- 



] 

--- 

Mo, Nb   

.  

FeCrNiX [

--- 



Si 







P 







” 

     

---  ---  --- 

---  ---  ---  ---  --- 

CPR um/y    

Temp Solution  K   

. M NaCl 

.   <    

. M NaOH 

     

. M H SO  

    ---  ---  ---  ---  --- 

  NaCl+NaOH pH  

  RS pH  

---  ---  --- 



. M H SO  



. M H SO  

---  ---  ---  ---  ---  ---  --- 

M HCl 

---  ---  --- 



.

M HCl 



M HCl  M HCl  M HCl 

---  ~   –   –  

Table . Summary oλ corrosion parameters aλλected by element addition λor some metallic μlasses.

75

76

Metallic Glasses: Formation and Properties

”esides the metal elements, the metalloids element addition oλ ”, Si, P, S, Ni, and C are also important to the corrosion resistance. The Fe −xCr Mo C ”x x = , , at% μlassy alloys exhibits spontaneously passivation in and M HCl solutions with wide passive reμion and low passive current density [ ]. With increase oλ boron content in alloys, the corrosion resistance oλ μlassy alloys is improved, even in HCl solution, the μlassy alloy with at% ” do not suλλer pittinμ corrosion. With P addition in the Fe C Mo C ” μlassy alloy, the kinetics oλ passivation and composition oλ passive λilm are improved in HCl solution. While with Si replacement oλ P, the corrosion resistance can be enhanced due to the λormation oλ passive λilm composed oλ chromium oxide with some amounts oλ silica [ ]. . . Effects of microstructure homogeneity Since metallic μlasses are metastable and can be transλormed into stable crystalline phase by heat treatment or mechanical workinμ, the structural chanμe can also aλλect corrosion resist‐ ance λor metallic μlasses. “ comparison oλ passive current density Ipass, corrosion/transpassi‐ vation potential Ecorr and corrosion rate CPR λor some Fe-based metallic μlasses and their crystalline alloy is summarized in Table [ – ]. It can be observed λrom Table that the corrosion/passive current density and corrosion rate increase λor the crystalline alloys com‐ pared with metallic μlass, while corrosion/transpassivation potential depends on their compositions. Composition  FeCrMoC”Y [

State  ] 

Ecorr mV  

“mor 

   

Cryst 

   

FeSi” [

] 

FeSi”NbCu [

FeZr” [

] 

FeNbZr”Cu [

FeNb” [

] 

] 

] 

Ipass A/cm  

CPR um/y  

. ×





--- 

. ×





--- 

. ×





M HCl 

--- 

Cryst 





--- 

--- 

“mor 





--- 



Cryst 





--- 



“mor 



*



. ×





--- 

Cryst 



*



. ×





--- 



. ×





--- 



. ×





--- 

. ×





--- 

. ×





--- 



*

“mor 



Cryst 



* *





M HCl 

M HCl 

--- 



M HCl 

--- 



“mor 

Solution 

--- 



Cryst 



--- 

“mor 

*

Temp K  

Table . Summary oλ corrosion parameters aλλected by microstructure λor some Fe-MGs.



. M NaCl 



. M NaCl 



. MHO 



. MHO 



. MHO 

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

Figure . ”F-TEM microμraphs oλ a λully amorphous, b devitriλied S“M ent periods oλ time [ ].

aλter immersion in

M HCl λor diλλer‐

The decrease oλ corrosion resistance in crystalline alloys obtained by the isothermal heat treatment oλ Fe-M-” M = Nb, Zr metallic μlasses is explained by the λormation oλ the α-Fe crystalline phase that has μreater corrosion susceptibility in compared to that oλ the amorphous phase [ ]. Lonμ et al. thouμht the μalvanic eλλects between adjacent phases with diλλerent composition were resulted in the deterioration oλ corrosion resistance λor Fe-Co-”-Si-Nb metallic μlass [ ]. “ comparison oλ ”F-TEM morpholoμies λor amorphous and devitriλied S“M [ ] is shown in Figure . It indicates that the lacier morpholoμies λor devitriλied S“M mean the deμradation in the corrosion resistance. However, the abrupt increase in the corrosion potential λor crystalline alloy is attributed to the decrease oλ the residual stress durinμ densiλication, and the surλace atom electrochemically active site [ ]. Since atom at a μlassy metal surλace are in nonequilibrium conλiμuration and may eλλectively sit on hiμher enerμy wells than that correspondinμ to atoms on an equilibrium conλiμuration. Moreover, the λaster miμration oλ silicon ions to the surλace in the crystalline structure promotes the SiO λilm λormation, which enhances the corrosion properties [ ]. Durinμ the λabrication processinμ oλ bulk metallic μlass and metallic μlass coatinμ, the porosity is not avoided due to rapid coolinμ. The eλλect oλ porosity on the corrosion resistance is investiμated in some literatures [ – ].The corrosion resistance oλ coatinμ with the porosity oλ . % is better than that oλ another two coatinμs with the porosity oλ . % and . %, respectively [ ]. When the porosity decreases λorm . % oλ low deposition rate to . % oλ hiμh deposition rate, the corrosion resistance oλ FeCrMoC”Y amorphous coatinμ increases evidently due to the elimination oλ throuμh pores [ ]. However, when the porosity is lower than . %, the corrosion resistance seems more sensitive to the amorphous phase content. Iλ the thickness oλ the coatinμ decreases, the number oλ throuμh-porosity in coatinμ increases, which aλλects the corrosion resistance, as shown in Figure [ ]. It indicates that throuμhporosity is much more detrimental to the corrosion resistance oλ the coated material compared with nonthrouμh porosity.

77

78

Metallic Glasses: Formation and Properties

Figure . “nodic polarization conducted on the bare substrate and coated substrate with various thickness oλ Fe-MG coatinμ in . wt% NaCl solution [ ].

In some bulk metallic μlass or metallic μlass coatinμ, the crystallize particles such as WC, TiN, SUS , NbC, TiO , and “l O are induced [ – ]. It is obvious that the crystallized particles are deteriorated the structural homoμeneity. However, little investiμations are done about the inλluence oλ crystalline particle on the corrosion resistance, which is important λor the potential application oλ Fe-based metallic μlass as anticorrosion and antiwearinμ materials. . . Effects of service environment The corrosion behaviors oλ Fe-based metallic μlasses are aλλected by environmental λactors. Intuitively, the stronμer the aμμressiveness oλ the solution is, the weaker the corrosion resistance oλ metallic μlass exhibits. The results [ ] attained in immersion experiments λor Fe Cr Mo C ” μlassy alloy in , , and M HCl solutions exhibit, as expected, that the corrosion rate increases as the increase in concentration oλ HCl solution. The alloy occurs pittinμ on the surλace aλter h oλ immersion in the M HCl solution at room temperature. FeCrMoC”P alloy is spontaneously passivated with a passive current density oλ about − “/ m and a wide passive reμion in M HCl solution, however, its passive λilm is not stable by anodic polarization, as an anodic current density increases with increasinμ potential in M HCl solution, and no passive λilm seems λormed on the surλace with rapid increasinμ oλ current

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

density in M HCl solution [ ]. Fe . Cr . Mo . Mn . W . ” . C . Si . wt% alloy can passivate spontaneously in the H SO solution, and the passive current density is chanμed λrom × − “/cm with . M to × − “/cm with . M [ ]. The corrosion penetration rates oλ Fe Cr Mo Er C ” metallic μlass [ ] are . , . , and mm/year in M HCl, M NaOH, and . M NaCl with pH solutions, respectively. The critical passivation potential Ecorr and critical passivation current density Icorr oλ Fe-Ni metallic μlass decrease with increase oλ pH value oλ solution, as shown in Figure [ ].

Figure

. The eλλect on critical passivation potential Ecr and current density Icr oλ Fe-Ni amorphous alloy at

K[

].

The corrosion rate oλ FeNi” metallic μlass is only μm/year in . wt% NaCl solution [ ], while the corrosion rate increases to , μm/year in M HNO solution, near two thousand times larμer than that in sodium chloride solution, as shown in Table . In Table , it is attained that more neμative corrosion potential is obtained with increase oλ pH value, and the corrosion current density decreases. “s the acidic ion or hydroxyl ion concentration increase, the corrosion potential becomes more neμative and the corrosion current density increases μenerally. That is, the corrosion resistance decreases with increase oλ acidic ion concentration and hydroxyl ion content. When the concentration oλ hydroμen ion is same, the existence oλ chloride ion will deteriorate the corrosion resistance. In a word, the corrosion resistance oλ Febased metallic μlass decreases as the solution aμμressiveness increases.

79

80

Metallic Glasses: Formation and Properties

Composition  FeNi” [

Ecorr mV  

] 

FeNi”“iNb [

] 

FeCSi”PCr“lMo [ FeCoCrMoC”Y [ FeCrMoC”Y [

FeCrMnMoW”CSi [ ] 

Epit mV  

Ipass A/ cm  

CPR um/y  

--- 

. ×





--- 

--- 

--- 

. ×





--- 

--- 

,

--- 

. ×



,



--- 

--- 



--- 

--- 

Temp K 

     

Solution 



M HNO  

--- 

M NaOH 

--- 

. wt% NaCl 

. ×





--- 





. ×





--- 

. × −  

<  



. M NaCl 





. ×





--- 

. × −  





. M NaOH 





. ×





--- 

. × −  



. M H SO  

]  −



. ×





--- 

--- 

. M H SO  



--- 



--- 

--- 







. ×



]  −



. ×









. ×









. ×







. × -  





. ×







. × −  





. ×







. × −  

] 

FeCoCrMoC”Y [

Icorr A/ cm  

M HCl 

--- 

--- 





--- 

--- 



Hank’s 

M HCl 



--- 

--- 



Saliva 

--- 



. wt% NaCl 

--- 



M HCl 

--- 



M H SO  

]  --- 

--- 

--- 

--- 

.





M HCl 

--- 

--- 

--- 

--- 

.





M HNO  

--- 

--- 

--- 

--- 

.





M NaOH 

--- 

--- 

--- 

--- 

.





. wt% NaCl 





--- 

--- 

--- 

--- 



.

M H SO  





--- 

--- 

--- 

--- 



.

M Na SO  



   

--- 

--- 

--- 

--- 



. M HCl 

--- 

--- 

--- 

--- 



. M NaCl 

--- 

--- 



--- 

--- 



FeCoCrMoC”Y [ ]  −



. ×









. ×









. ×





--- 

--- 

--- 



NaCl+H SO pH .  





. ×





--- 

--- 

--- 



NaCl pH .  

Fe”Nb [

] 

FeCo”SiNb [

FeCrNi” [

] 

] 

   

“cid rain  . wt% NaCl 





. ×





--- 

--- 

--- 



NaCl+NaOH pH





. ×





--- 

--- 

--- 



NaCl+H SO pH .  





. ×





--- 

--- 

--- 



NaCl pH .  





. ×





--- 

--- 

--- 



NaCl+NaOH pH





. ×





--- 

--- 

--- 



NaCl+H SO pH .  





. ×





--- 

--- 

--- 



NaCl pH .  





. ×





--- 

--- 

--- 



NaCl +NaOH pH

Table . Summary oλ corrosion parameters aλλected by environment λor some Fe-MGs.







Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

. Pitting corrosion of metallic glasses Thouμh metallic μlass exhibits excellent μeneral corrosion resistance, it is also susceptible to pittinμ corrosion in aμμressive solutions, especially containinμ Cl− ion [ ]. Since the surλace λilm is not stable durinμ the anodic polarization, many pits are observed on the surλace oλ FeCrMoC” metallic μlass with at% ” immersed in M HCl solution λor h [ ]. Durinμ potentiodynamic polarization in . M HCl solution, many peaks oλ current density occur λor FeCrM” M = Mo, Nb metallic μlass, which is attributed λrom pittinμ corrosion. Moreover, the morpholoμies oλ pits are conλirmed by SEM analysis aλter the immersion test [ ]. Fe Mn Mo Cr ” C metallic μlass is susceptible to pittinμ corrosion, althouμh presents μood corrosion resistance characterized by a low passivatinμ current in . M NaCl solution [ ]. Fe Cr Mo ” C Y metallic μlass, as known S“M , exhibits hysteresis loop durinμ cyclic potentiodynamic polarization in M NaCl solution at K, which indicates the λormation oλ localized corrosion [ ]. Pardo et al. [ , ] λound that FeSiNb”CuCr metallic μlass was immune to pittinμ corrosion in simulated industrial environments, since the current density decreased when the potential scan direction is reversed and it is identiλied that no pit λormed on surλace aλter immersion test. Thouμh no hysteresis loop is observed and Ecorr shiλted toward more anodic values λrom the cyclic polarization curve, the λormation oλ pits occurs when the anodic branch is enlarμed durinμ the λorward scan [ ]. The size oλ corrosion pit is less than μm in the P”S solution [ ], and the pits are distributed inhomoμeneous on the surλace oλ Zr-based metallic μlass in NaCl solution [ ]. The λormation oλ pits is attributed λrom the broken oλ the passive λilm or irreμular microstructure on the surλace [ ]. Jianμ and coworkers [ ] λound that almost all pits were passed throuμh by shear bands λor as-cast sample, while the pit was distributed randomly aλter annealinμ. Gostin et al. [ ] considers that no pittinμ propaμation is attributed λrom the hiμh repassiva‐ tion ability due to the hiμh content oλ the beneλicial Mo in its composition, althouμh the yttrium oxide particle provides a λavorable location λor pit λormation and local dissolution is initiated at their interλace with the μlassy matrix. However, Gostin et al. [ ] also λound that a pittinμlike process occurred λor the μlassy alloy with hiμh concentration oλ C, as the initial breakdown oλ passive λilm caused by the sudden direct exposure oλ the alloy surλace to the electrolyte subsequent to local rupturinμ oλ the C-rich layer by μrowinμ CO bubbles. Liu and coworkers [ ] reported the pittinμ was initiated since the λormation oλ a nanoscale Cr-depleted zone near the intersplat due to oxidation eλλect durinμ thermal sprayinμ, as shown in Figure . Paillier et al. [ ] reported that Cu-rich nanocrystals oλ – nm were λormed inside the corrosion pits durinμ the corrosion process, as shown in Figure . The corrosion mechanisms oλ is λeasible that elements like Zr, Ti, and “l mainly dissolve in solution whereas Cu and probably Ni is prone to λorm nanocrystals on the surλace covered by a passive oxide layer, as in Figure a , the structure vanishes with the complete removinμ oλ the surλace oxide by HF Figure b . On the bare surλace alloy without native oxide layer, the small pits develop λirst with the bow-like morpholoμy, and then, because the corrosion appears to proceed quicker in the vertical direction, and μoes alonμ with the canyon-like morpholoμy. Very deep trenches are indeed hollowed leadinμ to a canyon aspect, as shown in Figure c.

81

82

Metallic Glasses: Formation and Properties

Figure . TEM imaμes oλ the corroded morpholoμies oλ the amorphous coatinμ aλter immersion in λor h a and h b [ ].

Figure [ ].

. SEM imaμes oλ Zr Ti Cu “l Ni aλter diλλerent immersion times in HF . % a

s, b

M NaCl solution

s, and c

s

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

. Electrocatalytic properties of metallic glasses Metallic μlasses have μained considerable attention in catalysis research due to their unique structural and chemical properties [ ], such as the unique atomic structure with a shortranμe orderinμ oλ the constituents, the larμe λlexibility in their chemical composition compared to that oλ crystalline alloys, the structural and chemical homoμeneity, the hiμh reactivity due to their metastable structure, and the excellent conductivity λor electricity and heat. Some investiμations about the electrocatalytic activity oλ metallic μlasses λor the hydroμen evolution reaction or oxyμen reduction reaction, such as Co- [ , ], Zr- [ , ], Ni- [ ], Cu- [ ], Pt- [ , ], and “u-based [ ] metallic μlasses, are done in last λew decades. However, Febased metallic μlasses are the most attractive as catalytic material. Since the λirst catalytic materials oλ Fe-based metallic μlass were reported in [ ], a larμer number oλ investiμa‐ tions have be done [ – ], such as the Fe . ” . amorphous ribbon used as a catalyst λor the Fischer-Tropsch-type reaction oλ CO + H [ ], amorphous Fe-Zr precursor λor ammonia synthesis [ ], amorphous FeNiCrP” alloy as catalysts λor acetylene hydroμenation [ ] and hydroμen evolution [ , ]. The chromium eλλect on the catalytic activity λor FeNiCrP” metallic μlass is shown in Figure . It can be seen that the catalytic activity is stronμly aλλected by Cr presence, while catalytic eλλiciency is independent oλ Cr content.

Figure

. Catalytic activity vs. % Cr ○, amorphous □, crystalline [

].

The λamous composition oλ Fe Co Si ” G [ ] is λirstly reported in , exhibitinμ μood electrocatalytic activity λor hydroμen evolution reaction HER comparable with Pt. “ comparison oλ kinetics parameters between G and pure Fe, vit. , pure Pt is illustrated in Table . With increasinμ temperature the exchanμe current densities i oλ G is siμniλicant‐

83

84

Metallic Glasses: Formation and Properties

ly μreater than that oλ the polycrystalline iron. “λter this, the relationship between electroca‐ talytic activity oλ hydroμen evolution and crystallization [ ], anodic treatment [ ], and anodic dissolution [ ] oλ amorphous G have been investiμated in KOH solution. The results indicate amorphous G exhibits hiμher electrocatalytic activities compared with their crystalline alloys. This enhancement is not related to the electronic properties oλ metallic μlass.

Electrode  Fe poly  

Fe Si ”  

Fe Co Si ”  

Pt poly  

T K 

i A/cm  

b −mV  



. ×









. ×









. ×









. ×









. ×











. ×





. ×







. ×









. ×









. ×









. ×









. ×







Table . Electrocatalytic activity parameters oλ the cathodic hydroμen evolution λor G Fe Si ” , pure Pt.

   

and pure Fe, vit.

The electrocatalytic properties are aλλected not only by the composition oλ alloy, but also by the surλace composition and/or surλace area by chemical pretreatment. The catalytic studies [ ] on the hydroμenation oλ carbon monoxide by Fe-based alloy indicated that the activity is auμmented by a treatment in HNO solution. Guczi et al. [ ] thouμht that the surλace composition and valence state determined in depth were related to the catalytic activity and selectivity revealed in CO + H reaction. “n increased number oλ nickel and iron sites by removinμ oλ the prevailinμ boron oxide, iron oxide, and iron oxide layer aλter HCl treatment was responsible λor the enhanced catalytic activity, as shown in Figure , that is, the activity oλ the Fe” sample in as-received state was about twice that observed λor the FeNi” sample, on the other hand, comparinμ the HCl etched samples, the activity oλ the FeNi” alloy was about times hiμher than that oλ the Fe” ribbon. The electrocatalytic activity oλ Fe Ni P ” metallic μlass [ ] is improved λor HER by acid pretreatment with M HF or HNO λor min, since a porous structure with hiμhly rouμhed and numerous small craters resulted λrom the selectively leachability oλ phosphorous λrom the surλace oλ Fe Ni P ” metallic μlass enhanced the electrode surλace area in comparison to the as-polished surλace.

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

The mechanism oλ hydroμen evolution reaction at the metal electrode in alkaline solution is based on three-step reactions as λollowinμ

Figure . Catalytic activity, Sc + and Sc− selectivities over a Fe ” , b Fe Ni ” . Hatched as received. Full black treated with HCl [ ].

Electronation oλ water with adsorption oλ hydroμen

Volmer reaction is shown in Eq.

M +  H 2O +  e Û MH ads +  OH Electrochemical desorption oλ H

Heyrouvsky reaction is shown in Eq.

MH ads +  H 2O +  e Û M +  H 2 +  OH Chemical desorption

Taλel reaction is shown in Eq. 2MH ads Û 2M +  H 2

Due to involvinμ the transλer oλ electron λrom the electrode surλace, the density oλ electrons close to the enerμy level oλ metal surλace is an important parameter μoverninμ electrocatalytic reaction activity. However, it is diλλicult to estimate durinμ HER, so the eλλiciency is usually evaluated with overpotential η, Taλel slope b, and exchanμe current density i . The perλormance oλ cathodic electrode with respect to HER is primarily characterized by the overpotential which is μiven by the workinμ electrode potential minus the reversible potential. “ linear relationship exists between the overpotential and the cathodic current density as shown in Eq. [ ]

85

86

Metallic Glasses: Formation and Properties

h = a + b log(i ) where a is constant, b is the Taλel slope, and the exchanμe current density i = −a/b. The Taλel slope is mainly aλλected by the reaction mechanism. In the case oλ identical Taλel slope, the exchanμe current density is mainly aλλected by the eλλective surλace area. Some Fe-based metallic μlasses with hiμh electrocatalytic eλλiciency are reported, such as a low overpotential η = mV λor amorphous Fe Mo alloy [ ] at % KOH solution at K, a low Taλel slope b = mV and a larμe exchanμe current density i = μm/cm λor amor‐ phous Fe Co Si ” alloy [ ] at M KOH solution at K, and Fe . C Si . ” . P . Mo . Co . alloy with b = mV and i = μm/cm [ ]. It is evident that the Taλel slope as well as overpotential decrease with increase oλ Mo content at the same conditions [ ]. However, the Taλel slope does not chanμe with replace oλ Co λor Fe in the amorphous FeCo” alloy [ ]. The results oλ Taλel slope are considered that the eλλective transλer coeλλicients λor the dissolution oλ boron and the metals are equal.

. Summary The investiμation oλ corrosion resistance oλ metallic μlass is attractive λor allover researchers in materials science and enμineerinμ, since the unique structure and properties, extensive potential application. Most researches are λocus on the eλλects oλ element addition and nanocrystallization on the μeneral and local corrosion resistance in various environments. Certain elements are identiλied are quite eλλective in improvinμ the corrosion resistance. “nd the crystallization oλ metallic μlass is usually deleterious λor the corrosion resistance. In μeneral, the decreasinμ is the corrosion resistance oλ metallic μlasses, the increasinμ is the solution aμμressiveness, especial λor the chloride ion concentration. Unλortunately, the mechanism oλ pittinμ corrosion oλ metallic μlass is not clear, as that oλ conventional materials such as stainless steel. “s practical application oλ the metallic μlasses in industrial λield, some metallic μlasses such as Fe-based metallic μlass are used as anticorrosion or antiwearinμ materials. The coatinμ is one oλ eλλective methods. Durinμ the coatinμ processinμ, some inclusion, oxidation, crystallization, and even second particle as reinλorcement phase in the coat layer is inevitable. However, the eλλect oλ these particles on the μeneral and pittinμ corrosion resistance is seldom reported. Thereλore, λurther investiμation about the pittinμ corrosion is necessary to the industrial application oλ metallic μlass.

Acknowledgements The work was supported by the National Natural Science Foundation oλ China and the Department oλ Education Fund oλ Jianμxi GJJ . Thanks λor their understandinμ and support oλ my wiλe, Yanmei Zhanμ, and my son, ”ob.

Corrosion Resistance and Electrocatalytic Properties of Metallic Glasses http://dx.doi.org/10.5772/63677

Author details Shanlin Wanμ “ddress all correspondence to slwanμ

@nchu.edu.cn

National Deλense Key Discipline Laboratory oλ Liμht “lloy Processinμ Science and Technoloμy, Nanchanμ hanμkonμ university, Nanchanμ, China

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Chapter 5

Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification Celal Kursun and Musa Gogebakan Additional information is available at the end of the chapter http://dx.doi.org/10.5772/63513

Abstract The amorphous ribbons oλ Cu Zr Ni “l alloy were manuλactured by rapid solidiλica‐ tion. The ribbons were investiμated by X‐ray diλλraction XRD , scanninμ electron microscopy coupled with enerμy dispersive spectroscopy SEM‐EDX and diλλerential scanninμ calorimetry DSC . The activation enerμy oλ the crystallisation in amorphous alloys was determined by Kissenμer technique. The mechanical properties oλ the ribbons were characterized usinμ Vickers microhardness HV tester. “ccordinμ to the XRD and SEM results, the Cu Zr Ni “l alloys have a λully amorphous structure. The EDX analysis oλ the ribbons showed that compositional homoμeneity oλ the Cu Zr Ni “l alloy was λairly hiμh. From the DSC curves oλ the amorphous ribbons, it was deter‐ mined that μlass transition temperature Tμ is around – °C and super‐cooled liquid reμion ΔTx = Tx - Tx beλore crystallisation is around – °C. The microhard‐ ness oλ the as‐quenched ribbons was measured about  HV. However, this micro‐ hardness value decreased with increasinμ annealinμ temperature and it was calculated about  HV aλter annealinμ temperature oλ °C. Keywords: rapid solidiλication, microhardness, copper‐based alloy, crystallisation, Kissenμer plot

. Introduction “morphous alloys, with hiμh corrosion resistant, ultrahiμh strenμth and soλt λerromaμnetic and mechanical properties, have widely been the subject oλ intense investiμation [ – ]. These excellent properties stem λrom their hiμh chemical and structural homoμeneous creation. ”esides, it is possible to synthesise the amorphous alloys without restriction a wide chemical composition ranμe. “morphous alloys are used in many applications such as deλence, electri‐

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Metallic Glasses: Formation and Properties

cal, weldinμ, automobile and aircraλts industries. Cu‐based amorphous alloys are optimal materials because oλ their excellent mechanical properties and hiμh electrical and thermal conductivities λor these applications [ , ]. In addition to these applications, copper alloys are also used as the rocket nozzles, hiμh‐perλormance switches, the heat exchanμers, the condens‐ er tubes oλ ships [ , ]. Cu‐based amorphous alloys can be produced by many diλλerent techniques such as rapid solidiλication, mechanical alloyinμ, vapour depositions, plasma processinμ and solid state reactions. In the rapid solidiλication method, the amorphous alloys are manuλactured on thin ribbons λorms, which are usually ductile and briμht surλace. Many Cu‐based binary, ternary, quaternary and quinary alloys have been manuλactured by these methods [ – ]. In this work, Cu‐Zr‐Ni‐“l quaternary amorphous alloys are produced by rapid solidiλication technique at wheel surλace velocities oλ and  ms- as ribbons λorms with very λlexible. The eλλects oλ the wheel surλace velocities and diλλerent annealinμ process on mechanical and microstructural properties oλ produced ribbons are systematically investiμated. Thereλore, it has been revealed the amorphous nature oλ Cu Zr Ni “l ribbon alloys in order to contribute the continuously improvinμ Cu‐based alloys in industry.

. Methods and materials “n inμot oλ the Cu Zr Ni “l at.% alloy was prepared by arc meltinμ the mixtures oλ the pure elements, Cu . % , Zr . % , Ni . % and “l . % in a titanium‐μettered arμon atmosphere. From this alloy, ribbon materials oλ approximately  μm thickness and  mm in width were manuλactured by a sinμle‐roller Edmund ”(hler melt spinner at wheel surλace velocities oλ and  ms- . The structure oλ the ribbon samples was examined by XRD usinμ a Philips X'Pert powder diλλractometer with Cu‐Kɑ radiation μenerated at  kV and  m“. The transλormations temperatures and heat eλλects durinμ transλormations were examined by Perkin‐Elmer Sapphire DSC unit under inert μas atmosphere usinμ continuous heatinμ mode with the heatinμ rate oλ  K min- . Moreover, the DSC analysis was carried out λor the melt‐ spun ribbon at wheel speed oλ  ms- usinμ continuous heatinμ mode with the heatinμ rates oλ –  K min- . The cross section oλ the melt‐spun ribbons was studied by Zeiss Evo LS SEM and SEM‐EDX aλter conventional metalloμraphic preparation. The ribbons were annealed λor  min at diλλerent temperatures under vacuum/inert μas atmosphere. These temperature values are , , and °C. The annealed ribbons were investiμated by XRD λrom surλace, SEM λrom cross‐section with the same conditions used λor as‐quenched ribbons. The Vickers microhardness measurements oλ the as‐quenched and subsequently annealed ribbons were perλormed usinμ a Shimadzu HMV‐ by an applied load oλ .  N with a dwell time oλ  s at ten diλλerent locations.

. Results and Discussion Figure shows the X‐ray diλλraction patterns oλ the rapidly solidiλied Cu Zr Ni “l ribbons produced at wheel surλace velocities oλ and  ms- . “s shown in Figure , the XRD patterns

Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification http://dx.doi.org/10.5772/63513

exhibit the broad maxima characteristic which is λeature oλ amorphous materials without the evidence oλ any crystalline peaks. This means that the surλace velocities oλ and  ms- are optimal to synthesize Cu Zr Ni “l alloy as λully amorphous structure.

Figure . XRD pattern oλ the melt‐spun Cu Zr Ni “l ribbons prepared at wheel speeds oλ quenched.

and

 ms- as‐

DSC traces oλ amorphous Cu Zr Ni “l alloys at wheel speeds oλ and  ms- at a heatinμ rate oλ  K min display distinct and an obvious μlass transition temperature, Tg, beλore crystallisation, as shown in Figure . From the DSC curves, it is seen a wide super‐cooled liquid temperature ranμe λollowed by a pronounced exothermic reaction λor both ribbon alloys. Table summarises the characteristic temperatures which are μlass transition temperature Tg , crystallisation temperature Tx , super‐cooled liquid reμion ΔTx ΔTx = Tx - Tg , and peak temperature Tp oλ the Cu Zr Ni “l alloy. “ccordinμ to the Table , Tx, ΔTx and Tp increase while Tg decreases with increasinμ melt‐spun wheel surλace velocity.

Figure . The DSC curves oλ the Cu Zr Ni “l ribbon alloys at wheel speeds oλ at a heatinμ rate oλ  K min- .

and

 ms- obtained durinμ heatinμ

99

100

Metallic Glasses: Formation and Properties

Wheel speed/ms-

Tg/°C

Tx/°C

ΔTx/°C

Tp/°C

Table . Thermal values obtained λrom DSC curves λor melt-spun Cu Zr Ni “l ribbons at diλλerent wheel speed.

Figure exhibits the DSC curves at , , and  K min- oλ the ribbon alloy which are manuλactured at wheel speeds oλ  ms . The obtained peak temperature values, Tg, Tx, Tp and the super‐cooled liquid reμion ΔTx λrom Figure are presented Table . “s can be seen Table , Tg, Tx, Tp and ΔTx values are moved to hiμher temperatures with increasinμ heatinμ rate. It is attributed the heatinμ rate which are depended on the parameters oλ crystallisation and μlass transition durinμ continuous heatinμ [ , ]. Thereλore, this case reveals the siμniλi‐ cant oλ the kinetic aspects oλ the μlass transition λor μlassy alloys [ ].

Figure . DSC analysis results λor the melt‐spun ribbon prepared at wheel speed oλ mode with the heatinμ rates oλ –  K min- .

φ K/min

Tg/K

Tx/K

ΔTx/K

 ms- usinμ continuous heatinμ

Tp/K

Table . Thermal values obtained λrom DSC curves λor rapidly solidiλied Cu Zr Ni “l amorphous ribbons manuλactured at wheel speed oλ  ms- at diλλerent heatinμ rates.

Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification http://dx.doi.org/10.5772/63513

The activation enerμy E λor μlass transition or crystallisation is commonly estimated by the Kissinμer [ ] equation. The Eq. is μiven below. To calculate activation enerμy oλ the amorphous alloys with this equation, it is necessary to use data λrom diλλerent heatinμ rates oλ the alloy E æf ö +A ln ç 2 ÷ = RT èT ø where T is the speciλic temperature, μlass transition temperature Tg , crystallisation temper‐ ature Tx , or peak temperature Tp , oλ crystallisation, φ is the heatinμ rate, R is the μas constant .  J/mol K , E is the activation enerμy, “ is a constant. ”y plottinμ ln φ/T2 versus / RT , nearly a straiμht line is obtained. From the slope oλ this straiμht line, the activation enerμies Eg, Ex or Ep are calculated usinμ the certain peak temperatures Tg, Tx, Tp . Figure shows the Kissenμer plots oλ Cu Zr Ni “l ribbon alloy produced at wheel speed oλ  ms- . From the Kissenμer plots, the activation enerμies oλ Eg, Ex and Ep are determined . ± , . ± and . ±  kJ/mol, respectively. These values are very hiμh compared with previous studies whose activation enerμies are Ex = , Ep =  kJ/mol λor Cu Zr Ni “l alloy [ ], Eg =  , Ex = , Ep =  kJ/mol λor Cu . Zr . Ti Ni [ ] and Eg =  , Ex =  , Ep =  kJ/mol λor Cu Zr “μ “l alloy [ ]. On the other hand, it is also possible to mention that the amorphous Cu Zr Ni “l alloy has very hiμh thermodynamic stability with Ex =  .  kJ/mol value compared with previous works.

Figure . Kissinμer plots oλ the amorphous Cu Zr Ni “l alloy produced at wheel speed oλ

 ms- .

101

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Metallic Glasses: Formation and Properties

The annealinμ oλ the amorphous alloys is a siμniλicant process to characterise their crystalli‐ sation behaviour. Thus, it miμht be revealed that the amorphous structure transλorms into what kind oλ crystalline phases with increasinμ annealinμ temperature. For this purpose, the melt‐spun ribbon oλ Cu Zr Ni “l alloy synthesised at wheel speed oλ  ms- was annealed in the temperature ranμe oλ – °C λor  min. Figure shows the XRD patterns oλ Cu Zr Ni “l alloy aλter annealinμ. “ccordinμ to Figure , beλore exothermic reaction, the XRD pattern oλ Cu Zr Ni “l alloy with annealed oλ °C exhibits λully an amorphous structure. “λter the annealinμ temperature oλ °C, intermetallic phases with sharp diλλraction peaks have been obtained λrom the amorphous matrix and λully crystallisation oλ the amor‐ phous phase. This result is in μood aμreement with crystallisation peak in DSC traces which is above °C. The obtained phases in the XRD spectrum were marked by symbols and indexed as cubic‐“lCu Zr with lattice parameters, a = b = c =   Å, orthorhombic‐Cu Zr with lattice parameters, a =  b =  c =  ,  Å, tetraμonal‐Zr Cu with lattice parame‐ ters, a =  b =  c =  ,  Å and λ.c.c‐Cu with lattice parameters, a = b = c =   Å. These phases were also observed in previous works aλter a similar annealinμ process λor Cu‐based amorphous alloys [ , – ]. Number oλ the crystalline peaks which belonμs to “lCu Z, Cu Zr , Zr Cu and Cu phases was increased by increasinμ annealinμ temperature °C , as shown in Figure .

Figure . XRD pattern oλ the melt‐spun ribbon oλ Cu Zr Ni “l alloy manuλactured at a wheel speed oλ annealed in the temperature ranμe oλ – °C λor  min.

 ms- and

In addition to XRD patterns oλ annealed ribbons, typical SEM microμraphs λrom cross section oλ the amorphous Cu Zr Ni “l alloy prepared at a wheel speed oλ  ms- as well as annealinμ ribbons at , , and °C are shown in Figure . In Figure a, b, the microstructure

Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification http://dx.doi.org/10.5772/63513

with λeatureless morpholoμy oλ unannealed and annealed at °C ribbons are exhibited. This λeatureless morpholoμy is a typical characteristic oλ the amorphous materials. In previous works, similar SEM imaμes taken surλace oλ amorphous structured materials were reported [ , , ]. These microμraphs are in accord with the XRD spectrums which exhibit λully amor‐ phous λeatures unannealed Figure and annealed at °C ribbons Figure . “s can be seen obviously in Figure c–e, with increasinμ annealinμ temperature , , °C , the microstructure oλ Cu Zr Ni “l ribbon alloys chanμes and transλorms into irreμularly shaped λeatures which is a characteristic oλ crystalline structures. These crystalline structures belonμ to “lCu Zr, Cu Zr , Zr Cu or Cu phases obtained by XRD patterns Figure

Figure . Typical SEM imaμes λrom the cross section oλ the melt‐spun ribbon oλ Cu Zr Ni “l alloy prepared at a wheel speed oλ  ms- . a “s‐quenched and annealed at the temperatures, b °C, c °C, d °C, and e °C.

The compositional homoμeneity oλ the amorphous Cu Zr Ni “l ribbons was by measured EDX in order to conλirm initially intended composition values. The EDX analysis illustrates mean values oλ element concentrations oλ Cu Zr Ni “l alloy produced at a wheel speed oλ

103

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Metallic Glasses: Formation and Properties

 ms- in Figure . “s can be seen obviously λrom the EDX results, the peaks in the spectrum belonμ to Cu, Zr, Ni and “l elements. “s shown in Figure , the averaμe chemical composition oλ the ribbon alloy is in μood aμreement with the chemical composition values oλ Cu Zr Ni “l alloy.

Figure . EDX analysis result oλ the melt‐spun ribbon oλ Cu Zr Ni “l alloy produced at a wheel speed oλ quenched.

 ms- as‐

In order to determine the inλluence oλ annealinμ on the microhardness oλ the Cu Zr Ni “l ribbon alloys which are as‐quenched and annealed at diλλerent temperatures such as , , , , and °C, Vickers HV measurements were analysed. The λollowinμ Eq. was used λor these measurements [ ] HV =

2 P sin(q / 2) 1.8544( P ) = d2 d2

where P is the indentation λorce, d is the mean diaμonal lenμth, and . is the μeometrical λactor λor the diamond pyramid. Figure shows the variation oλ microhardness values with increasinμ annealinμ temperature λor Cu Zr Ni “l alloy prepared at wheel speed oλ  ms- . “s shown in Figure , the hardness values decrease with increasinμ annealinμ temperature. In previous works, this decline oλ the hardness values with the annealinμ temperature is μenerally reported λor Cu‐based amorphous alloys [ , – ]. The microhardness oλ as‐ quenched ribbon was calculated  HV, while it was determined –  HV λor annealed ribbons in the ranμe oλ – °C Figure . “t the temperature ranμe oλ – °C, the

Structure and Mechanical Behaviour of Cu‐Zr‐Ni‐Al Amorphous Alloys Produced by Rapid Solidification http://dx.doi.org/10.5772/63513

microhardness values oλ the Cu Zr Ni “l alloy were not chanμed distinctly and it was determined as approximately  HV. Thus, it can easily be concluded that the hiμhest microhardness value  HV oλ the Cu Zr Ni “l alloy was measured λor as‐quenched ribbon alloy.

Figure . The chanμe in Vickers microhardness values λor Cu Zr Ni “l alloy prepared by the wheel speed oλ ms- with annealinμ temperatures.



. Conclusions .

The metallic μlass Cu Zr Ni “l alloys were successλully produced by rapid solidiλication technique at wheel speeds oλ and  ms- .

.

DSC traces oλ the Cu Zr Ni “l alloys showed similar distinct μlass transition, Tμ which are around – °C. The ribbon alloys exhibited also wide super‐cooled liquid reμions, ΔTx which are – °C.

.

The activation enerμies oλ Eμ, Ex and Ep λor Cu Zr Ni “l alloy prepared at wheel speed oλ  ms- were determined . ± , . ± and . ±  kJ/mol, respectively.

.

The intermetallic “lCu Zr, Cu Zr , Zr Cu and Cu phases in the microstructure oλ Cu Zr Ni “l alloy were observed aλter annealinμ temperature oλ °C.

105

106

Metallic Glasses: Formation and Properties

.

The compositional homoμeneity oλ Cu Zr Ni “l as‐quenched ribbons was most correct‐ ly conλirmed by EDX.

.

The microhardness value oλ Cu Zr Ni “l alloy was calculated approximately  HV λor unannealed ribbons. However, it decreased with increasinμ annealinμ temperatures and was measured about  HV aλter annealinμ temperature oλ °C.

Acknowledgements We would like to thank Kahramanmaras Sutcu Imam University λor λinancial support oλ the research proμramme Project No / ‐ D . One oλ the authors C. Kursun would like to thank Council oλ Hiμher Education Y5K λor μraduate research support.

Author details Celal Kursun* and Musa Goμebakan *“ddress all correspondence to [email protected] Department oλ Physics, Faculty oλ “rt and Sciences, Kahramanmaras Sutcu Imam University, Kahramanmaras, Turkey

References [ ] “. ”outahara, H. Lassria, E.K. Hlilb, D. Fruchart, Critical behavior and its correlation with maμnetocaloric eλλect in amorphous Fe -xVx” Si x =  , and . alloys. Journal oλ Maμnetism and Maμnetic Materials. – . [ ] K. Tomolya, D. Janovszky, “. Sycheva, M. Sveda, T. Ferenczi, “. Roósz, Peculiarities oλ ball‐millinμ induced crystalline–amorphous transλormation in Cu–Zr–“l–Ni–Ti alloys. Intermetallics. – . [ ] Y. Yinμ‐Jun, K. Fu‐Wei, X. Da‐Wei, S. Jian‐Fei, S. Qinμ‐Ke, S. Jun, Formation and mechanical properties oλ bulk Cu–Ti–Zr–Ni metallic μlasses with hiμh μlass λorminμ ability. Transactions oλ Nonλerrous Metals Society oλ China. – . [ ] C.K. Kim, H.S. Lee, S.Y. Shin, J.C. Lee, D.H. Kim, S. Lee, Microstructure and mechanical properties oλ Cu‐based bulk amorphous alloy billets λabricated by spark plasma sinterinμ. Materials Science and Enμineerinμ “. – . [ ] “. Inoue, K. Hashimoto, “morphous and nano‐crystalline materials, preparation, properties and application. ”erlin Sprinμer, .

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[ ] M. Goμebakan, C. Kursun, J. Eckert, Formation oλ new Cu‐based nanocrystalline powders by mechanical alloyinμ technique. Powder Technoloμy. – . [ ] C. Kursun, M. Goμebakan, Y. Gencer, Microstructural characterization oλ rapidly solidiλied Cu Zr Ni Ti amorphous alloy. Journal oλ “lloys and Compounds. S –S . [ ] “.M. Shams El Din, R.“. Mohammed, Contribution to the problem oλ vapour‐side corrosion oλ copper–nickel tubes in MSF distillers. Desalination. – . [ ] M. Weiλanμ, L. Shulinμ, W. Jinμtanμ, Isothermal crystallization kinetics oλ an amor‐ phous CuTi alloy. Chinese Journal oλ Metal Science and Technoloμy. – . [

] M. Goμebakan, C. Kursun, J. Eckert, Formation oλ new Cu‐based nanocrystalline powders by mechanical alloyinμ technique. Powder Technoloμy. – .

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] J. Eckert, J. Das, K.”. Kim, F. ”aier, M.”. Tanμ, W.T. Wanμ, Z.F. Zhanμ, Hiμh strenμth ductile Cu‐based metallic μlass. Intermetallics. – .

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] S. Venkataraman, W. Loser, J. Eckert, T. Gemminμ, C. Mickel, P. Schulbert‐”ischoλλ, N. Wanderka, L. Schultz, D.J. Sordelet, Nanocrystal development in Cu Ti Zr Ni Si metallic μlass powders. Journal oλ “lloys and Compounds. – .

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] C.J. Hu, P.Y. Lee, Formation oλ Cu–Zr–Ni amorphous powders with siμniλicant supercooled liquid reμion by mechanical alloyinμ technique. Materials Chemistry and Physics. – .

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] E.S. Park, H.K. Lim, W.T. Kim, et al. The eλλect oλ Sn on the μlass‐λorminμ ability oλ Cu– Ti–Zr–Ni–Si metallic μlass alloys. Journal oλ Non‐Crystalline Solids. – .

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] J. Wu, Y. Pan, J. Huanμ, J. Pia, Non‐isothermal crystallization kinetics and μlass‐λorminμ ability oλ Cu–Zr–Ti–In bulk metallic μlasses. Thermochimica “cta. – .

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] L. Liu, Z.F. Wu, J. Zhanμ, Crystallization kinetics oλ Zr Cu “l Ni bulk amorphous alloy. Journal oλ “lloys and Compounds. – .

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] H.E. Kissinμer, Reaction kinetics in diλλerential thermal analysis. “nalytical Chemistry. .

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] Y.J. Yanμ, D.W. Xinμ, J. Shen, J.F. Sun, S.D. Wei, H.J. He, D.G. McCartney, Crystalliza‐ tion kinetics oλ a bulk amorphous Cu–Ti–Zr–Ni alloy investiμated by diλλerential scanninμ calorimetry. Journal oλ “lloys and Compounds. – .

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] L. Zhanμ, Z. HuaChen, Q. Zhenμ, D. Chen, Isochronal and isothermal phase transλor‐ mation oλ Cu Zr “μ “l bulk metallic μlass. Physica ”. – .

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] D. Janovszky, “. Sycheva, K. Tomolya, J. Geiμer, J. Solyom, “. Roosz, Solidiλication processes in Cu–Zr–“μ amorphisable alloy system. Journal oλ “lloys and Compounds. – .

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] F.“. Javid, N. Mattern, S. Pauly, J. Eckert, Eλλect oλ cobalt on phase λormation, micro‐ structure, and mechanical properties oλ Cu -xCoxZr x =  , , , at. pct alloys. Metallurμical and Materials Transactions “ Physical Metallurμy and Materials Science. – .

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Chapter 6

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites Devinder Singh, R.K. Mandal, R.S. Tiwari and O.N. Srivastava Additional information is available at the end of the chapter http://dx.doi.org/10.5772/64221

Abstract In the present chapter, results oλ our recent investiμations on the role oλ μallium Ga on the aluminum “l site in Zr . “l . -xGaxCu Ni metallic μlass MG composition have been discussed. The material tailorinμ and coolinμ rate eλλects on the mechanical behavior oλ Zr-based metallic μlasses and their nanocomposites have been studied. The substitution oλ Ga on the “l site in Zr–“l–Cu–Ni alloy aλλects the nucleation and μrowth characteristics oλ quasicrystals QCs and consequently chanμes the morpholoμy oλ nanoquasicrystals. The Zr . “l . -xGaxCu Ni system displayed metallic μlass λorma‐ tion in the ranμe oλ x = – . . In this process, we have come out with a new μlass composition Zr-Ga-Cu-Ni with μlass transition temperature Tμ K. The eλλect oλ coolinμ rate on the μlass λorminμ ability GF“ and mechanical properties λor this new metallic μlass composition has been discussed and compared with some other Zrbased metallic μlasses. The various indentation parameters such as microhardness, yield strenμth, strain hardeninμ constant, nature oλ shear band λormation, and so on λor the alloys have been analyzed. The study is λocused on investiμations oλ these materials to understand the structure microstructure property correlations. Keywords: metallic μlasses, quasicrystal, composites, mechanical properties, coolinμ rate

. Introduction Metallic materials are traditionally considered as crystalline in nature, possessinμ translation‐ al as well as orientation symmetry, i.e., their constituent atoms are arranμed in a reμular and periodic manner in three dimensions. However, a revolution in the concept oλ metals was

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brouμht, when metallic μlasses MGs and quasicrystals QCs have been discovered. MGs are amorphous in nature possessinμ short-ranμe orderinμ while QCs possess aperiodic lonμranμe order associated with crystalloμraphically λorbidden rotational symmetries. ”oth quasicrystal-λorminμ alloys and MGs μivinμ rise to nanoquasicrystalline phase on annealinμ have attracted attention owinμ to their promise to qualiλy λor many potential applications. . . Metallic glasses Ever since the λormation oλ λirst metallic μlass in the “u–Si system by rapid solidiλication, numerous investiμations have been carried out over the past years due to their attractive properties and technoloμical potential. In initial period oλ metallic μlass study, hiμh coolinμ rates oλ the order oλ to K/s were the usual requirement λor the λormation oλ μlassy phase. However, in the recent years, a new class oλ metallic μlass known as bulk metallic μlass ”MG has been synthesized usinμ very slow coolinμ rates. These newly developed ”MGs have μenerated immense research activity driven by both a λundamental interest in the structure and properties oλ disordered materials and their unique promise λor structural and λunctional applications. MGs have very hiμh-yield strenμth and very hiμh elastic limit compared to crystalline steel and Ti alloys Figure a . They have very hiμh λracture strenμth coupled with – % oλ elastic strain. Conventional aluminum, titanium alloys and steels can sustain – % oλ elastic strain. The μlasses have tensile yield strenμth σ∼ . GPa , i.e., a hiμh strenμth-toweiμht ratio, makinμ them a possible replacement λor “l, but with a much μreater resistance to permanent, plastic deλormation i.e., λracture touμhness . “ larμe domain oλ hiμh λracture strenμth and elastic strain can be achieved by nanocrystallization into the amorphous matrix. Figure b represents the hiμhest strenμth, speciλic strenμth and speciλic Younμ’s modulus oλ any bulk amorphous or crystalline metal.

Figure . a “morphous metallic alloys combine hiμher strenμth than crystalline metal alloys with the elasticity oλ pol‐ ymers [ ]. b Schematic representation oλ room temperature yield metals, composites, and polymers or λlexural strenμth ceramics as a λunction oλ modulus. Note the increased strenμth oλ amorphous metals over conventional crys‐ talline metals [ ].

MGs possess a number oλ very attractive properties, and in many cases, these properties are enhanced by suitable heat treatment. The ability to store a hiμh amount oλ elastic enerμy has made this material to use as a potential sprinμ material. This has led to its λirst and most visible use in the heads oλ μolλ clubs. The addition oλ ceramic second-phase particles into the material

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

improves its ductility. This composite can be used in aircraλt λrames and automobiles as armor penetrator material and medical implants. Due to larμe super-cooled liquid reμions, the workability oλ these materials is very hiμh. This property has been applied in λriction weldinμ oλ Pd-based bulk MGs [ ]. The hiμh strenμth, hardness, λracture touμhness, and λatiμue strenμth oλ MGs make them ideal λor the use as optical, die, tool, and cuttinμ materials [ , ]. “monμ the larμe number oλ multicomponent μlassy alloy systems, Zr-based MGs have outstandinμ μlass λorminμ ability GF“ . The exceptionally hiμh-yield strenμth, close to the theoretical limit, hiμh hardness, and elastic modulus oλ these MGs oλλer them potential λor structural applications. However, plastic deλormation at room temperature occurs in a hiμhly localized manner by the λormation oλ shear bands. In these MGs, the deλinitive correlations between mechanical behavior and atomic structures have not been clearly understood. . . Quasicrystals “nother important class oλ material is QCs. The breakthrouμh experiments by Shechtman et al. on rapidly solidiλied “l- % Mn alloys have created a new concept oλ nonperiodic atomic arranμements with only orientational order, which exhibit sharp diλλraction peaks with λiveλold symmetry [ ]. This new λorm oλ ordered structures havinμ orientational order and lackinμ strict translational periodicity was desiμnated as quasicrystal by Levine and Steinhardt [ ]. It may be noted that in contrast to both crystal and QCs, amorphous solids possess neither orientational nor translational order. Most λamiliar quasicrystalline systems are “l-, Ti-, and Mμ-based binary and ternary alloys, thouμh there have been a λew reports in other systems such as Cd-Mμ-Yb, “μ-In Cd , “l-Zn-Ce, and Cu-Ga-Mμ-Sc, etc. The discovery oλ the quasi‐ crystalline phases has also μenerated a μreat deal oλ interest in reμard to complex crystalline structures known as approximant phases, which have remarkable similarities with their parent quasicrystalline structures. These oλten coexist with QCs and have similar chemical composi‐ tions and similar electron diλλraction patterns Figure . Quasicrystal approximants have similar local atomic structures to QCs [ – ]. ”ecause oλ these structural similarities, the search λor other possible phases as well as intensive investiμations oλ their phase transλormation has

Figure . Selected area diλλraction patterns oλ a icosahedral quasicrystal showinμ λive-λold symmetry in “l-Mn and b pseudodecaμonal quasicrystal approximant in “l-Co-Ni alloy [ , ].

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become quite pertinent in connection with the determination oλ the phase stability oλ quasi‐ crystalline system. The successλul applications oλ QCs are very limited. QCs are corrosion resistant and have low coeλλicients oλ λriction, and thus, they can be used as a surλace coatinμ λor λryinμ pans. They can also be used in wear resistant coatinμs. “l-based quasicrystalline alloys, e.μ., “l-Mn-Ce containinμ nanoicosahedral particles may be used in surμical blades. Ti- and Zr-based QCs could be incorporated into hydroμen storaμe materials. . . Nanocomposites Quasicrystal λorminμ alloys and MGs promise to qualiλy λor many potential applications. However, bulk QCs are mostly brittle and this problem can be surmounted by producinμ μlassnanocrystal nc /nanoquasicrystal nqc composites Zr . “l . Cu Ni throuμh controlled crystallization oλ MGs. Quasicrystal evolution λrom metallic μlass systems may provide a way to produce nanostructured quasicrystalline alloys with attractive mechanical properties. The advantaμe oλ λormation oλ quasicrystalline phase throuμh devitriλication oλ MGs is also due to the λact that the microstructure can be precisely controlled. The control oλ microstructure is very important as the optimum property desiμn is related to the microstructure. It has been pointed out that Ti- and Zr-rich alloys have siμniλicantly hiμher hardness in the nanoquasi‐ crystalline state and VHN, respectively compared to the amorphous state in meltspun condition. The hardness values oλ Ti- and Zr-rich alloys increase λurther by nanoquasicrystallization oλ the amorphous phase to and VHN, respectively. Misra et al. [ ] have studied the plastic deλormation in nanostructured bulk μlass composites durinμ nanoindentation. The structural chanμes are accompanied by decrease in speciλic volume, bulk modulus and Poisson’s ratio. Small speciλic chanμes upon primary devitriλication suμμest a close relationship between the μlassy structure and the icosahedral structure.

. Effect of material tailoring on the mechanical properties Elemental substitution is widely used to λind new MGs and QCs with improved properties. In this section, the role oλ Ga on the “l site in Zr . “l . -xGaxCu Ni metallic μlass composition has been discussed. The alloy desiμn principle adopted in arrivinμ at Ga-substituted μlass compositions pertains to retaininμ the valence electron ratio e/a constant. In this respect, Ga substitution on the “l site seems to be ideal. The substitution oλ Ga in Zr . “l . -xGaxCu Ni alloys results in a chanμe λrom a two-step crystallization x = to a sinμle-step one x = . [ , ]. For x = , we have the well investiμated Zr-based alloy and λor x = . , we have come out with a new composition oλ μlass with Tμ = K [ – ]. The eλλect oλ the said material tailorinμ on the mechanical properties oλ these alloys has also been studied. The recent emphasis on nanostructured materials and synthesis oλ MGs has added a new dimension to the study oλ their indentation behavior. Indentation studies oλλer opportunities to investiμate the λunda‐ mental nature oλ deλormation in μlasses and their composites λrom a relatively small volume

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

oλ material. The indentation size eλλect ISE and shear band λormation under compression are able to throw liμht on the mechanical behavior oλ materials. . . Microstructural and structural features The Zr . “l . -xGaxCu Ni x = , . , and . MGs with a thickness oλ ∼ – μm and lenμths oλ ∼ – m have been synthesized usinμ melt spinninμ technique [ , ]. Figure a shows the macroscopic imaμe oλ the melt-spun ribbons synthesized at m/s. Figure b and the inset therein show the transmission electron microscopy TEM imaμe and correspondinμ selected area diλλraction pattern S“DP displayinμ diλλuse halos λor Zr . “l . -xGaxCu Ni x = . alloy. We note that the TEM briμht-λield microμraph displayed no discernible contrast. This clearly indicates the λormation oλ μlassy phase in the system and similar λeatures were observed λor x = and . . The μlass-nc/nqc composites are produced aλter controlled crystal‐ lization oλ melt-spun ribbons correspondinμ to compositions x = , . , and . [ – ]. The TEM microμraph oλ these composites is shown in Figure . Inset in them demonstrates the presence oλ crystalline/quasicrystalline particles embedded in the μlassy matrix. The compo‐ sition oλ the alloys with x > . consists oλ icosahedral and Zr Cu phases embedded in the μlassy matrix. The λiner μrains oλ both these phases have been observed in the Ga-bearinμ μlass composition x = . [ ]. The μrain reλinement oλ quasicrystals with respect to Ga substitution may be understood by recallinμ expression [ , ] oλ the steady-state nucleation rate Is and is reproduced below Is =Aexp[(-16 P s3 )/(3KT(DG v ) 2 )] where “’ is known as dynamical preλactor and is a λunction oλ the atomic mobility at the nucleiliquid/μlass interλace. ΔGv = drivinμ λree enerμy per unit volume λor the phase transλormation. “ccordinμ to classical theory oλ nucleation, the nucleation barrier is controlled by interλacial λree enerμy σ between the nuclei and the liquid/μlass [ ]. The decrease in σ between the icosahedral quasicrystalline nuclei and the liquid with increasinμ Ga content leads to the

Figure . a Optical imaμe showinμ the λormation oλ lonμ melt-spun ribbons synthesized at m/s. b TEM imaμe and the correspondinμ diλλraction pattern oλ as-synthesized Zr . Ga . Cu Ni alloy. Reprinted with kind permission λrom reλerences [ , ], Copyriμht and , Elsevier.

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increase in the nucleation rate. The reported interλacial enerμy per unit area oλ Ga ∼ . J/m is less than that oλ “l ∼ . J/m [ – ]. Thus, it can be said that the substitution oλ Ga reduces the interλacial enerμy between quasicrystal and remaininμ amorphous phase, thereby increas‐ inμ the nucleation rate oλ the crystalline/quasicrystalline phases. The λormation oλ icosahedral phase has been observed λor all the annealed μlasses. Thus, the icosahedral order presents predominantly in the supercooled liquid λor all the samples x = – . .

Figure . TEM microstructures and the correspondinμ diλλraction patterns oλ Zr . “l . -xGaxCu Ni alloy with x = a , x = . b , and x = . c λormed aλter heat treatment. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

The composition oλ the alloys in at.% based on electron probe microanalysis EPM“ has been λound to be Zr . “l . Cu . Ni . λor x = , Zr . “l . Ga . Cu . Ni . λor x = . , and Zr . Ga . Cu . Ni . λor x = . . The presence oλ oxyμen within the detectable limit oλ EPM“ was not λound. . . Mechanical properties In this section, we present the results oλ micro-/nanoindentation behavior oλ the three μlassy compositions and their respective composites. Figure depicts the imaμes oλ microindent λor the as-synthesized and annealed ribbons oλ x = . . It has been observed that a number oλ shear

Figure . SEM microμraphs λor the as-synthesized a and annealed ribbons b oλ x = . displayinμ shear bands. Re‐ printed with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

bands decrease μradually in the annealed condition as compared to that oλ μlassy state. “t this staμe, the indentation behavior is μoverned by the volume λraction oλ nanocrystalline/ nanoquasicrystalline phases. Thus, the μlass-nc/nqc composite possesses diλλerent indentation characteristics as is seen in the case oλ x= . that Zr Cu intermetallic phase exists alonμ with the nqc phase, which inhibits the propaμation oλ shear bands in the μlassy matrix durinμ indentation. We note λrom Table that annealed ribbons λor x = . display the hiμhest microhardness value ∼ GPa . The microhardness values λor x = . at μ load are ∼ . and ∼ . GPa λor as-synthesized and annealed samples, respectively. These are quite close to that oλ Zr-“l-Ni-Cu-“μ and Zr-“l-Ni-Cu-Nb MGs and their nqc composites [ ]. x

As-synthesized ribbons 

at.%  Microhardness GPa at ± .  

g load

Annealed ribbons 

Nanohardness GPa at ± .  

μN

Reduced

Microhardness

Nanohardness

Modulus

GPa at

g load GPa at

GPa at

± .  

± .  

μN ± .   .

.

.

.

.

.

.

GPa at μN ± .  

.

.

. .

Reduced

μN Modulus

. .

.

Table . Mechanical properties oλ as-synthesized and annealed ribbons oλ Zr . “l . -xGaxCu Ni alloys. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier .

x = , . , and .

Figure . “FM observation oλ the nanoindentation imprints λrom the as-synthesized a and annealed ribbon b oλ x = . at μN. The inset oλ b showinμ the tip imaμe oλ nanoindenter [ ].

To study the nature oλ indentation at submicroscopic scale as well as the plastic deλormation oλ composites containinμ nc/nqc phases, we now present the results oλ nanoindentation. The indentation impressions oλ ”erkovich indenter at μN λor x = . are shown in Fig‐ ures a and b . These are to compare the indentation impressions oλ the μlassy phase with that oλ their nanocomposite. The inset in Figure b shows the tip imaμe oλ the nanoindenter. We note clearly the presence oλ λine μrains in Figure b that are absent in Figure a . The size oλ the indentation impression chanμes with partial crystallization and no crackinμ occurred. The heiμht contrast around the indents is due to pileup. The contrast oλ the pileup

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in μlassy sample is much more prominent as compared to that oλ annealed ribbons. The λormation oλ pileups around the indents must be related to shear bandinμ operations and are extensively observed in amorphous alloys [ – ].

Figure . Plot oλ the indentation λorce P versus indenter displacement h obtained λrom nanoindentation tests λor the as-synthesized a and b and annealed c and d ribbons oλ x = and . respectively [ ].

Figures a and b depict the load P versus depth h behavior oλ melt-spun ribbons, whereas Figures c and d display the P versus h characteristics oλ their respective composites. The values oλ nanohardness and reduced modulus λor x = , . , and . are μiven in Table . The values oλ nanohardness λor the μlassy alloys x = – . are in the ranμe oλ ∼ – GPa. Reduced modulus is sensitive to compositions oλ the alloy as well as atomic arranμements. The reduced modulii λor the μlassy alloys x = – . lies in the ranμe oλ ∼ – GPa. These values oλ nanohardness and reduced modulus are comparable to those oλ the Zr-based alloys [ – ]. It can be seen λrom Table that composites have hiμher micro- /nanohardness values than those oλ μlassy alloys. Ramamurty et al. [ ] reported that the presence oλ nanocrystalline particles in the μlassy matrix siμniλicantly improves the stiλλness and strenμth values. The nanohardness value mean contact pressure is hiμher than that oλ Vickers hardness value [ ]. The diλλerence in the values is primarily due to two reasons i indentation size eλλect and ii actual and projected area oλ contacts, respectively, λor nano- and microhardness measure‐ ments.

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

We have observed pop-ins durinμ loadinμ marked by the arrow in Figures a and b λor the melt-spun ribbons. These pop-ins indicate displacement bursts that siμniλy the λormation oλ shear bands [ ]. The pop-ins are prominent in amorphous alloys while the presence oλ these is either less prominent or even completely suppressed in case oλ annealed alloys. Our observation is in aμreement with the results reported by others [ , , ]. The pop-ins are absent in composites and this may be attributed to the presence oλ nc/nqc μrains in the μlassy matrix. The nature oλ deλormation can be inλluenced by structural λeatures while the chemistry aλλects such behavior in a quantitative way. This is the reason why we observe similar kind oλ P versus h curves correspondinμ to x = and . aλter crystallization oλ the μlasses. The observation oλ pop-ins durinμ loadinμ cannot be attributed to the process oλ nanocrystalliza‐ tion as noted in reλerence [ ]. The break in the P versus h curves λor such a case should appear while unloadinμ. To settle this issue experimentally, the TEM investiμations oλ the indented portion oλ these samples have been done. The briμht-λield imaμes oλ the thinned specimens λor the μlassy alloys with x = and . are shown in Figures a and b . The preparation oλ these samples was done by maskinμ the indented side and thinninμ was done λrom the opposite side. The briμht-λield imaμes resemble analoμous to those shown in Figure b and do not show any reμions oλ residual contrast. The curvilinear line marked by arrows has μiven evidence oλ layer wise displacement separated by boundary. These lines must be related to the shear bandinμ operations, and thus, the possibility oλ nanocrystallization has been ruled out.

Figure . ”riμht-λield TEM imaμes oλ the indented portion oλ as-synthesized thinned specimens λor x = printed with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

and . . Re‐

The chanμe in the mechanical behavior oλ MGs and their composites can be understood on the basis oλ λree volume model. In the present case, the variation oλ λree volume with Ga substi‐ tution may increase the hardness oλ the MGs. The atomic radius oλ Ga . nm is inter‐ mediate between the atomic radius oλ “l . nm and Ni . nm and the atomic radius oλ Zr . nm and Cu . nm . Thus, the substitution oλ Ga may increase the packinμ density oλ the alloy and this would lead to the decrease in the λree volume [ ]. The hiμh resistance to plastic deλormation under applied stress may be attributed to a low λree volume [ ]. The increase in the hardness oλ MGs with alloyinμ addition has been reported recently [ ]. The shear transλormation zones STZs are the primary carriers oλ plasticity in amorphous

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materials [ , ]. The λormation oλ STZs depends upon the availability oλ λree volume. The precipitation oλ nc/nqc phases in the case oλ composites decreases the λree volume and this causes densiλication oλ the metallic μlass [ ]. This results in an increased resistance to plastic deλormation and thereλore enhancement oλ hardness oλ the metallic μlass upon structural relaxation and nanocrystallization. This observation is consistent with the results reported earlier [ , ]. In the case oλ μlass-nc/nqc composites, the hardness increases with increase in Ga addition. This may be due to the μrain reλinement oλ nanocrystals/nanoquasicrystals that produces many interλaces leadinμ to the strenμtheninμ phenomenon.

. Effect of cooling rate on the mechanical properties The absence oλ μrain boundaries and dislocations in MGs contributes to its exceptional properties [ – ]. MGs lack lonμ-ranμe order and thus, they can be considered as solids with λrozen-liquid structures composed oλ tiμhtly bonded atomic clusters and λree volume zones [ – ]. The λrozen-in excess volume is oλten interpreted as an increase oλ λree volume content in the MGs [ , ]. The λunctional and mechanical properties oλ a metallic μlass are determined by its internal atomic conλiμuration [ , ]. The diλλerent variables such as the coolinμ rate and composition aλλect the structure oλ MGs [ , , ]. “monμ these, the critical coolinμ rate is a very important λactor that plays a crucial role in determininμ the atomic structure and hence deλormation behavior oλ MGs [ – ]. The limited macroscopic plastic strain beλore λracture oλ MGs constrains their applications [ ]. Plastic deλormation oλ MGs is localized within relatively thin reμions called shear bands, resultinμ in a very low macroscopic plastic λlow limit [ , , ]. Recent investiμations show that the plastic strain oλ some monolithic MGs can be improved by enhancinμ the homoμeneity in microstructure throuμh the hiμh coolinμ rate [ ]. The hiμh coolinμ rate may result in the conλiμurationally looser atomic packinμ and thus more λree volume zones, which thereλore contribute to larμer plasticity. Chen et al. [ ] suμμested that the plasticity λor MGs can be tailored by applyinμ diλλerent coolinμ rates durinμ solidiλication. Jianμ et al. [ ] λound that the Cu-based bulk metallic μlass ”MG is havinμ hiμher hardness as compared to its ribbon counterpart oλ the same composition but synthesized at a much hiμher coolinμ rate. It has been observed that decreasinμ the coolinμ rate oλ μlass λorminμ promoted the λormation oλ denser atomic conλiμuration in the resultant alloy [ ]. The study oλ coolinμ rate eλλect on the nanomechanical response λor a Ti-based ”MG reveals that the hardness increases while the plastic deλormation μradually decreases λrom the edμe to the center oλ the sample [ ]. Recently, Huanμ et al. [ ] reported the eλλect oλ coolinμ rate on the local atomic orderinμ and the wear behavior oλ Zr-Cu-“l-“μ ”MG. These results indicate that the coolinμ rate used durinμ μlass λormation is a processinμ parameter that may be tuned to chanμe the mechanical properties oλ MGs. The eλλect oλ coolinμ rate on the mechanical behavior oλ Zr . Ga . Cu Ni metallic μlass has been studied usinμ microindentation technique. The ribbons oλ alloy have been synthesized at three coolinμ rates, correspondinμ to wheel speeds oλ , , and m/s. The diλλerent properties such as μlass λorminμ indicators, structural relaxation heat, microhardness, yield strenμth, strain-hardeninμ constant, material constant related to the resistance oλ the metal to

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

penetration and pileup parameter pertaininμ to nature oλ shear band are expected to throw liμht on the internal structure oλ μlass. They are compared and discussed with respect to the rate oλ coolinμ. This study provides some insiμhts to understand the correlation between the coolinμ rate and the mechanical behavior oλ Zr-Ga-Cu-Ni metallic μlass. . . Microstructural and structural features Figure shows the X-ray diλλraction XRD patterns oλ as-synthesized Zr . Ga . Cu Ni meltspun ribbons synthesized at diλλerent wheel speeds. It has been observed that all the patterns oλ the alloys consist oλ only broad diλλraction maxima at the position θ ≈ ° without a detectable sharp ”raμμ peak. This shows λormation oλ a μlassy phase. The XRD pattern oλ the ribbon synthesized at m/s revealinμ the presence oλ a μlassy phase exhibits μreater peak broadeninμ and lower XRD intensity as compared to the ribbons synthesized at m/s. These eλλects are more pronounced by λurther increasinμ the wheel speed to m/s. The λull width at halλ maximum FWHM was λound to be . °, . °, and . ° λor the ribbons synthesized at , , and m/s, respectively. These results indicate that the ribbon synthesized at m/s has hiμher deμree oλ short-ranμe orderinμ. The λormation oλ a μlassy phase in these samples

Figure . XRD patterns oλ as-synthesized ribbons oλ Zr . Ga . Cu Ni metallic μlass at wheel speed oλ m/s. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

,

, and

Figure . TEM microμraphs with inset showinμ the selected area electron diλλraction patterns oλ as-synthesized Zr . Ga . Cu Ni alloys synthesized at wheel speeds a , b , and c m/s. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

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was λurther investiμated by TEM. Figure and the insets therein show the TEM microμraphs and the correspondinμ selected area electron diλλraction S“ED patterns λor Zr . Ga . Cu Ni melt-spun alloys synthesized at , , and m/s, respectively. We note that all TEM microμraphs depict no discernible contrast and the correspondinμ S“ED patterns displayinμ diλλuse halos. This conλirms to the XRD results presented above.

Figure . DSC curves oλ the as-synthesized Zr . Ga . Cu Ni metallic μlass samples synthesized under diλλerent cool‐ inμ conditions. Inset hiμhliμhts the enlarμed section oλ the DSC curves below the μlass transition temperature. Re‐ printed with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

Figure shows diλλerential scanninμ calorimeter DSC scans taken at a heatinμ rate oλ K/ min λor the μlassy alloys prepared under diλλerent coolinμ conditions. “ll DSC curves exhibited one clear endothermic heat event, characteristic oλ the μlass transition to a supercooled liquid state, λollowed by only one sinμle exothermic peak between about and K. It can be seen that the DSC curves oλ all the samples are very similar with μlass transition temperature Tμ in the ranμe oλ – K and onset crystallization temperature Tx in the ranμe oλ – K. This indicates the comparable thermal stability ΔTx amonμ the samples. Table summarizes the thermal stability data λor all the investiμated samples. The crystallization enthalpies can be obtained by inteμratinμ the area covered by the crystallization peak in the DSC curve and are λound to be . , . , and . J/μ λor the ribbons synthesized at , , and m/s, respectively. The crystallization enthalpy decreases with the decreasinμ coolinμ rate and thus conλirminμ that the sample synthesized at m/s contains a larμer deμree oλ short-ranμe orderinμ or medium-ranμe orderinμ. The deμree oλ orderinμ can also be estimated by evalu‐ atinμ the crystallization λraction Vλ oλ the sample λrom the DSC curves and is μiven by [ ] Vf =

DH max - DH DH max

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

where ΔHmax is the total enthalpy chanμe when the λully amorphous alloy transλorms into a completely crystallized one and ΔH is the crystallization enthalpy oλ the examininμ sample. Considerinμ the sample synthesized at m/s to be λully amorphous with zero crystallization λraction, the crystallization λraction λor the samples synthesized at and m/s was λound to be . % and . %, respectively, suμμestinμ a neμliμible crystalline content in the samples. The small increase in the crystallization λraction results λrom the enhanced short-ranμe orderinμ. Cooling rate m/s  

Tg K  

Tx K  

Tp K  

ΔTx K  

Tμ μlass transition temperature Tx onset crystallization temperature ΔTx supercooled liquid reμion Tp exothermic peak. Table . Thermal analysis oλ the melt-spun Zr . Ga . Cu Ni ribbons synthesized at diλλerent coolinμ rate. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

“s evident λrom Figure and Table , there is a sliμht increase in the value oλ Tμ, Tx, and ΔTx with the increasinμ coolinμ rate. However, careλul analysis reveals some diλλerences around the μlass transition reμions, as more clearly seen in the inset oλ Figure . The inset oλ Figure is a local maμniλied reμion oλ the DSC curves below Tμ, illustratinμ the heat release events due to structural relaxation. The structural relaxation enthalpy associated with the exothermic peak below Tμ can be calculated by inteμratinμ the heat λlow near the μlass transition ranμe the area between the dotted lines and the curves shown in the inset oλ Figure . The relaxation enthalpy has been λound to be . , . , and . J/μ λor the μlassy ribbons synthesized at , , and m/s, respectively, as provided in Table . “ hiμher coolinμ rate has resulted in a larμer relaxation enthalpy. Slipenyuk et al. [ ] have shown that the exothermic heat release is directly related to the structural relaxation, i.e., the chanμe oλ λree volume in metallic μlasses, and can be calculated by (DH ) fv = b D vf where β is a constant, ΔH fv is the chanμe in enthalpy due to per unit λree volume, and Δvf is the chanμe oλ λree volume per atomic volume. Thus, the chanμe in λree volume oλ the μlassy ribbons synthesized at diλλerent wheel speeds may be obtained by Eq. . “ssuminμ that the λree volume per atomic volume λor the ribbon synthesized at m/s to be ρo, the substitution oλ values oλ ΔH cλ. Table and the inset in Figure into Eq. , the values oλ λree volume per atomic volume λor the ribbons synthesized at and m/s were λound to be . ρo and . ρo, respectively. Such a computation suμμests that the μlassy ribbons synthesized at lower wheel speed have less λree volume per atomic volume than the ribbons synthesized at hiμher wheel speed. The ribbons synthesized at m/s are, thereλore, believed to possess the hiμhest

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λree volume. “ccordinμ to Turnbull and Cohen [ ], the amount oλ λree volume in a metallic μlass is determined by the coolinμ rate. “ slower coolinμ rate μives suλλicient time to the atoms to attain their local ordered equilibrium positions, thereby, a more ordered atomic structure λorms durinμ coolinμ λrom the melt and thus, the obtained μlassy sample has a smaller amount oλ λree volume [ ]. The amount oλ λree volume in a metallic μlass corresponds to the atomic packinμ density and one oλ the important parameters exerts a stronμ inλluence on the me‐ chanical properties oλ the metallic μlass. Cooling rate m/s  

Hardness

n 

Log K σ GPa

VHN GPa ± .

load ± .

g

α = A/As Crystallization



Structural

enthalpy ΔH J/g   relaxation



enthalpy J/g  

.

.

  .

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

Table . Summary oλ hardness VHN , Meyer’s exponent n , material constant K , yield strenμth σ , pileup parameter α , crystallization enthalpy ΔH and structural relaxation heat oλ the Zr-based metallic μlass samples with diλλerent coolinμ rate. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

. . Mechanical properties In this section, we present the results oλ coolinμ rate eλλect on the mechanical behavior oλ Zr . Ga . Cu Ni MGs synthesized at diλλerent wheel speeds. The microhardness measure‐ ments were carried out by Vickers microhardness tester. The mean hardness reported here is

Figure . Indentation imprints at diλλerent loads λor the as-synthesized ribbons oλ Zr . Ga . Cu Ni metallic μlass at wheel speed oλ a , b , and c m/s showinμ the λormation oλ shear bands around the indents. Four indentation impressions λrom various reμions oλ the sample are superimposed. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

the averaμe oλ at least λive points on each sample. Figure shows the representative optical microμraphs oλ indents at diλλerent loads in Zr . Ga . Cu Ni MGs synthesized at , , and m/s, respectively. These microμraphs reveal that the indents are crack λree up to the load oλ μ λor all the samples. The wavy patterns around the indent reveal the μeneration and λormation oλ shear bands marked by arrows in Figure . It can be clearly seen that the number oλ visible shear bands λor the ribbons synthesized at m/s is hiμher than those observed λor the ribbons synthesized at and m/s. To λurther conλirm this, we have calculated the pileup parameter. The characteristic spiral pattern around indents constitutes pileup and is related to the λormation oλ shear bands. Pileups at which shear band reaches the surλace are extensively observed around indents in amorphous alloys [ , ]. The pileup parameter can be calculated by employinμ the λollowinμ relationship [ ]

Figure . Schematic representation oλ various quantities used λor calculation oλ pileup parameter [ ] “s and “ are the area oλ impression beλore and aλter pileup, whereas hs and h are the depth oλ impression beλore and aλter pileup.

a = A/A s where “s = fhs with f as a constant and is equal to . λor Vickers pyramid indenter. The values oλ “s and “ are the area oλ impression beλore and aλter pileup, whereas hs and h are the depth oλ impression beλore and aλter pileup. Various quantities utilized λor the determination oλ are displayed throuμh a schematic diaμram in Figure . The pileup area calculation was done by μraphical methods. The values oλ α λor the ribbons synthesized under diλλerent coolinμ conditions are reported in Table . The pileup parameter has been λound to be maximum λor the ribbons synthesized at m/s. In contrast to this, the ribbons synthesized at and m/s are havinμ relatively lower value oλ α indicates that λew shear bands are μenerated durinμ indentation. “monμ the three types oλ specimens, the ribbons synthesized at m/s contain the larμe λree volume and have the hiμhest shear band density. The hiμh λree volume content not only λavors the nucleation oλ shear bands but also helps to enhance the atomic mobility

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that can alleviate the stress concentration and thereλore prevent the metallic μlass λrom crackinμ. Thus, it can be inλerred that λor the μlassy alloy oλ λixed composition μives rise to the larμer number oλ shear bands with hiμher coolinμ rate. Furthermore, hiμher λree volume at enhanced coolinμ rates not only λacilitates permanent λlow oλ materials under compressive stresses but also contributes to enhancement in the λracture strenμth. The hardness H was calculated in GPa units by employinμ the λollowinμ relationship [ H = 1.854 ´ 9.8 ´

]

P d2

where P is the load μ and d is the diaμonal lenμth in μm. Figure a shows hardness VHN versus load μ characteristic curves λor the ribbons synthesized at , , and m/s, respec‐ tively. It can be seen λrom the load-dependent hardness curves Figure a that the hardness decreases with increase in the load due to indentation size eλλect ISE [ , ]. Table compares the values oλ microhardness and other indentation parameters oλ the ribbons synthesized at diλλerent coolinμ rates. The hardness values oλ the ribbons synthesized at , , and m/s at μ load are ∼ . , ∼ . , and ∼ . GPa, respectively, clearly suμμestinμ that λaster the coolinμ rate durinμ solidiλication, the lower the microhardness. The μlass λorminμ ability parameters and hardness values oλ Zr . Ga . Cu Ni alloy synthesized at m/s have been compared with some other known Zr-based alloys cλ. Table . Compared with the other typical MGs, Zr . Ga . Cu Ni alloy has hiμher hardness but lower Tμ and ΔTx values. The hardness λor Zr . Ga . Cu Ni alloy is . GPa that is hiμher in comparison to the majority oλ the metallic μlass alloy systems listed in Table .

Figure . a Variation oλ hardness VHN with respect to load μ λor the as-synthesized ribbons oλ Zr . Ga . Cu Ni metallic μlass synthesized at diλλerent coolinμ rates. b Loμ P versus Loμ d plots λor the as-synthesized ribbons oλ Zr . Ga . Cu Ni metallic μlass at wheel speed oλ , , and m/s. Reprinted with kind permission λrom reλerence [ ], Copyriμht , Elsevier.

Mechanical Behavior of Zr-Based Metallic Glasses and Their Nanocomposites http://dx.doi.org/10.5772/64221

Alloys Zr

Hardness GPa

Reference Raμhavan etal. [

.

Liu et al. [

Zr Cu

.

Janμ et al. [

.

Sinμh et al. [

Zr Zr

.

.

Ni ”e

ΔTx K

.

.

Cu

Tx K

Zr Pd Cu Ni “l

.

Ti

Tg K .

“l . Ni Si

“l . Cu Ni

]

.

Sun et al. [

]

Sun et al. [

]

Zr Cu

.

Sun et al. [

]

Zr Cu “l Ni

.

Jana et al. [

]

Zr

.

.

“μ . “l

“l . Cu Ni

.

Sinμh et al. [

Ni

.

Janμ et al. [

Ga . Cu Ni

.

Sinμh et al. [

Zr “l . Cu Zr

.

.

“l .

]

.

.

Ni

Ti

.

Table . Comparison oλ Tμ, Tx, ΔTx and hardness values oλ Zr with some other Zr-based metallic μlasses.

.

]

]

Zr Ti Ni Cu “l

.

Cu

.

] ] ]

Ga . Cu Ni melt-spun alloy synthesized at

m/s

The load independent hardness values permits us to compute the . % oλλset yield strenμth σ by usinμ the λollowinμ relationship [ ] σ 0 =  (VHN/3)0.1n -2 where n = Meyer’s exponent. This is determined by the slope loμ P in Kμ versus loμ d in mm curve. The intercept oλ this curve K is the material constant related to the resistance oλ the metal to penetration. Figure b shows loμ P versus loμ d curves λor the ribbons synthesized at diλλerent wheel speeds. The values oλ n and K are reported in Table . There is no siμniλicant variation observed in the values oλ n and K λor the samples. The values oλ exponent are less than as observed λor intermetallics [ ]. The yield strenμth lies in the ranμe oλ ∼ . – . GPa λor the samples and is λound to be maximum λor the ribbon synthesized at m/s. In the present case, both hardness and yield strenμth increase with decrease in the coolinμ rate and this may be attributed to the variation oλ λree volume with the coolinμ rate. “s the coolinμ rate decreases, the λree volume decreases and thus causes densiλication oλ the metallic μlass that results in an increased resistance to plastic deλormation and thereλore enhancement oλ hardness oλ the metallic μlass upon structural relaxation [ – ].

. Conclusion In this chapter, the recent proμress in the development oλ metallic μlasses, quasicrystals and their nanocomposites are discussed. The Zr . “l . -xGaxCu Ni system displayed metallic μlass λormation in the ranμe oλ x = – . . In this process, we have come out with a new composition oλ μlass without “l correspondinμ to x = . . The nanohardness and reduced

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Metallic Glasses: Formation and Properties

elastic modulus values oλ the metallic μlasses have been compared to their nanocomposites. The indentation characteristics oλ the μlass composition with x = . have shown siμniλicant improvement in reμard to hardness and elastic modulus. ”ased on transmission electron microscopic studies oλ the indented μlassy specimen, the possibility oλ deλormation-induced nanocrystallization has been ruled out. In addition to, the coolinμ rate eλλect on the μlass λorminμ ability, crystallization and mechanical behavior oλ Zr . Ga . Cu Ni metallic μlass composition is presented. “ slower coolinμ rate leads to hiμher deμree oλ structural relaxation, less λree volume content and thereλore better short-ranμe orderinμ. Such relatively ordered atomic conλiμuration and less λree volume content result in a hiμher hardness and yield strenμth λor the samples synthesized at slower coolinμ rate than those synthesized at λaster coolinμ rate. The ribbons synthesized at λaster coolinμ rate contain the larμe λree volume and have the hiμhest shear band density. The μlass λorminμ ability parameters and hardness values oλ Zr . Ga . Cu Ni alloy have shown siμniλicant improvement in comparison to some other known Zr-based alloys.

Acknowledgements The authors are thankλul to Dr. M.“. Shaz and Dr. T.P. Yadav λor many stimulatinμ discussions. One oλ the authors Devinder Sinμh μrateλully acknowledμes the λinancial support by Department oλ Science and Technoloμy DST , New Delhi, India in the λorm oλ INSPIRE Faculty “ward [IF“ -PH- ]. Sections and oλ this chapter are reproduced with permission λrom Reλerences [ ] and [ ] © , Elsevier .

Author details Devinder Sinμh *, R.K. Mandal , R.S. Tiwari and O.N. Srivastava *“ddress all correspondence to [email protected] Department oλ Physics, Panjab University, Chandiμarh, India Department oλ Metallurμical Enμineerinμ, Indian Institute oλ Technoloμy ”anaras Hindu University , Varanasi, India Department oλ Physics, Nano-Science Unit, ”anaras Hindu University, Varanasi, India

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Chapter 7

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER Vladimir S. Voitsenya, Alexandra F. Bardamid, Martin Balden, Flaviu Gostin, Sergey V. Khovrich, Vladimir G. Konovalov, Konstantin V. Kovtun, Petro M. Lytvyn, Sergey V. Ketov, Dmitri V. Luzguine-Luzgin, Sergei I. Solodovchenko, Anatoly N. Shapoval, Anatoly F. Shtan’, Vladislav N. Bondarenko, Ivan V. Ruzhkov, Ol’ga O. Skoryk and Andrei A. Vasil’ev Additional information is available at the end of the chapter http://dx.doi.org/10.5772/63885

Abstract This chapter reviews main results obtained on mirror-like samples made oλ several μrades oλ bulk metallic μlasses ”MG . Experiments were carried out under simulated conditions typical λor the operation oλ plasma λacinμ in-vessel mirrors oλ optical plasma diaμnostics in λusion reactor ITER. ”ombardment with D and T atoms radiated λrom burninμ plasma was predicted to be the main reason λor the deμradation oλ optical properties oλ such mirrors. Thereλore, to simulate the behavior oλ mirrors in ITER, mirror-like samples were subjected to bombardment by ions oλ deuterium plasma with λixed or wide enerμy distribution. The eλλects oλ ion bombardment on optical proper‐ ties, development oλ rouμhness, uptake oλ deuterium, appearance oλ blisters, and maniλestation oλ some chemical processes are presented and discussed. Keywords: amorphous mirrors, sputterinμ eλλects, deuterium uptake, chemical proc‐ esses, blister-like λeatures

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Metallic Glasses: Formation and Properties

. Introduction In the experimental λusion reactor ITER, many diλλerent methods oλ plasma diaμnostics are envisaμed to be used [ ]. “ biμ portion oλ methods are intended λor optical measurements, and these have to be based on reλlective optics, because oλ the hiμh level oλ deeply penetratinμ radiations, μammas, and neutrons, which may deμenerate the reλractive optics components. The mirrors λacinμ the burninμ plasma λirst mirrors, FM will be additionally subjected to λluxes oλ charμe-exchanμe atoms CX“, mainly D and T atoms with a wide enerμy distribution, up to several hundred eV [ ]. To overcome neμative eλλects oλ CX“ sputterinμ on mirror charac‐ teristics due development oλ surλace rouμhness , it was decided to λabricate λirst mirrors λor ITER λrom sinμle crystal SC metal, and molybdenum is the λirst candidate as a FM material. “t present, it is not known which eλλect the simultaneous irradiation oλ the FM surλace with neutrons and CX“ will have. However, there is a probability that a sinμle crystal will lose its ideal SC structure what would result in μradual development oλ rouμhness under CX“ bombardment and deμradation oλ optical properties. “ real alternative to SC mirrors can be mirrors λabricated λrom amorphous metal alloys bulk metallic μlasses, ”MGs . They do not have any arranμed structure larμer than a λew nanometer and thereλore may be more resistive under irradiation with neutrons in comparison with crystallized metals. Recent results on simulatinμ the neutron irradiation eλλects by exposinμ ”MG samples with MeV Ni+ ions did not lead to biμ deμradation oλ hardness and Younμ's modulus in the dose ranμe oλ . – dpa [ , ]. It is important to note that the structure oλ samples has continued to be amorphous, without indication oλ appearance oλ crystallization. “dditionally, due to the lack oλ crystallized structure, under lonμ-term sputterinμ, a polished ”MG mirror has to resemble a liquid under evaporation its surλace has to be smooth reμardless oλ sputterinμ time. Such assumption was mentioned in [ ] and has λound support later, aλter appearance oλ technoloμy to produce ”MG casts with size ≥ mm suλλicient λor the λabrica‐ tion oλ mirror samples to provide correspondinμ experiments. Zr-based ”MGs reveal relatively hiμh crystallization temperature compared with other ”MGs. From a practical standpoint, the μlass λorminμ ability oλ Zr-based ”MGs is very μood enablinμ the manuλacturinμ oλ λully μlassy components with thickness values in excess oλ mm. This chapter is a short review oλ main results obtained with ”MG mirror samples in experi‐ ments that partly simulate the conditions λor FM operation in ITER, that is, lonμ-term sput‐ terinμ by ions oλ deuterium plasma in some cases by ions oλ arμon plasma with enerμy λrom eV up to eV. Mirror samples were λabricated λrom λive ”MG μrades Table . The proμram oλ experiments with each ”MG μrade was not identical, and thus, new inλormation was obtained, and some new properties oλ ”MGs were λound λor both amorphous and crystallized ”MG specimens. This chapter presents results on the λollowinμ eλλects oλ lonμterm ion sputterinμ on optical properties oλ mirror-like ”MG samples, eλλects oλ deuterium adsorption, the role oλ chemical processes on ”MG surλace when the deuterium plasma is contaminated with oxyμen, and observation oλ blister-like λeatures due to deuterium exposure.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

In Section , the specimens and the experimental details reμardinμ the plasma exposures and characterization are described. Section contains the main results on chanμes due to plasma exposure, in particular deuterium adsorption, reλlectance, and erosion rate and the diλλerences due to the state, amorphous or crystallized by annealinμ. In Section , some concludinμ remarks are provided. “ppendix “ contains details on ”MG sample preparation in NSC KIPT Kharkov, Ukraine , and in “ppendix ”, the results oλ processinμ oλ an imaμe oλ laser beam aλter reλlection λrom amorphous and crystallized mirrors are shown.

. Experimental and precharacterization . . Descriptions of specimens “ list oλ all μrades oλ samples with their composition, shape, and size is presented in Table . Grade #

Composition

Size, mm

Vit -LM

Zr

. Ti

. Cu

. Ni

. Cu

. Ni

”e

Cu . Ni

”e

Vit -NSC

Zr

. Ti

Vit

Zr

.

Zr

Cu

“l “μ

Zr

Cu

. “l

Vit

Ti .

Ni

”e

.

Ø

×

.

Ø

×

.

Ø × Ø

. Nb



×

×

Table . List oλ μrades with composition μiven in at.%.

Figure . Results oλ X-ray diλλraction measurements obtained with the halλ oλ the amorphous sample oλ μrade # er curve and the halλ annealed at °K durinμ hour.

low‐

Three ”MG samples oλ grade # produced by Liquidmetal Co US“ and three billets oλ grade # have the same nominal composition. The billets oλ μrade # were casted as discs with diameter ~ and ~ mm in thickness the details oλ their λabrication in NSC KIPT are described in “ppendix “ . They were cut into

137

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Metallic Glasses: Formation and Properties

approximately equal halves λinal size Ø × mm one halλ oλ every billet was leλt amorphous and the second halλ was annealed hour at °K to have a λine-crystalline material. Thus, λor this μrade, there was a possibility to compare the behavior oλ mirrors λrom the identical material but with diλλerent structures, amorphous and λine-crystalline. The X-ray data on the structure oλ such a pair are shown in Figure . The position oλ the peaks in a diλλractoμram indicated the existence oλ the λollowinμ nano‐ crystals Zr Ni, Ti Ni, Zr Cu, and ZrxCuy with x and y exceedinμ two. ”y measurinμ the halλwidth oλ the peaks, the size oλ crystallites was rouμhly estimated to nm usinμ the λollowinμ λormula [ ] D=

0.9 l b cosq

where λ is the X-ray wavelenμth, θ is the ”raμμ anμle, and β is the λull width at halλ maximum oλ the diλλraction peak. “λter crystallization oλ three halves, all three pairs oλ μrade # were polished simultaneously. The mirror specimens oλ grade # more than pieces were cut λrom a -mm-diameter rod to discs oλ mm thickness. ”ecause oλ the small surλace area, they were not used λor some experiments, namely λor measurinμ the absorptivity oλ deuterium. The cast oλ grade # had a complicated shape thereλore, the mirror samples λive pieces had diλλerent diameter λrom Ø to Ø mm with thickness oλ mm each. The cast oλ grade # was oλ rectanμular shape aλter cuttinμ it into two samples the size oλ mirror specimens became × × mm. XRD X-ray diλλraction analysis conλirmed the μlassy state oλ all samples diλλraction patterns not shown here . “ll prepared billets were polished to a hiμh optical quality. . . Pretreatment and initial reflectance Prior to exposure experiments, all mirror samples were initially cleaned λor ≥ min with lowenerμy deuterium plasma ions, Ei ~ eV or sometimes with low-enerμy “r plasma ions, to remove the orμanic λilm appeared due to rinsinμ oλ samples in acetone and alcohol aλter the polishinμ. The reλlectance oλ mirror samples aλter this exposure was taken as their initial spectral reλlectance. The comparison oλ initial reλlectance, R λ , λor one ”MG mirror specimens oλ each μrade is presented in Figure toμether with W and Mo reλlectivity data λrom [ ]. The measurements were done in the wavelenμth ranμe – nm at normal incidence oλ the liμht by the use oλ a two-channel method described in [ ] with a homemade attachment to a standard mono‐ chromator.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

Figure . Initial reλlectance R λ measured just aλter cleaninμ by low-enerμy ions oλ deuterium or arμon plasma oλ samples oλ the λive μrades toμether with the data λor W and Mo λrom literature [ ].

The R λ values oλ some ”MG mirror samples are close to R λ values oλ tunμsten λor the wavelenμths ranμe ~ – nm and is approachinμ to the reλlectance oλ Mo mirror at > nm. Crystallization does not lead to any siμniλicant chanμe oλ initial optical properties oλ samples in the wavelenμth interval oλ measurements – nm , as can be concluded λrom comparinμ the data λor μrades # and # -C crystallized . . . Heterogeneities observed in the body of BMG samples “λter the last polishinμ procedure, some local peculiarities were discovered on smooth sample surλaces oλ the μrade # . Such peculiarities were not observed on the other μrades. Their level was a little below the main surλace, appear usually as a μroup, are rouμhly oλ oval shape, distributed over the main surλace oλ mirror sample more or less uniλormly, and are observable in both, optical microscope and scanninμ electron microscope SEM . The total relative area oλ these inhomoμeneities was estimated to be at the level oλ ~ %, so they did not have any essential inλluence on measurement oλ reλlectance, that is, the perλormance oλ the mirror. The λact oλ their elevation below the matrix supposed that the composition oλ these λeatures diλλers λrom that oλ the main volume oλ material and that their hardness aμain mechanical treatment is inλerior to the hardness oλ the bulk. Figure shows two SEM imaμes oλ the same reμion obtained by diλλerent detectors oλ SEM, detectinμ a mainly secondary electrons and b backscattered electrons. In the λirst case, the contrast oλ the imaμe is mainly determined by the surλace topoμraphy, but in the second reμime, the contrast occurs due to diλλerences in the atomic number oλ the surλace elements. In this second case, the domains with predominance oλ liμht elements appear darker than the surroundinμ thereλore, we may state that the recessed domains are enriched by the liμhter component oλ the material the averaμe atomic number here is below that oλ the matrix .

139

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Metallic Glasses: Formation and Properties

Figure . SEM imaμes λrom secondary a and backscattered electrons b oλ inhomoμeneities on the sample surλace oλ μrade # [ ].

Under ion bombardment, as a result oλ the sputterinμ process, initially lower parts oλ the surλace descended deeper below the main level and turned into shallow holes with a rather nonplanar bottom shown in Figure . The depth oλ holes increased with increase in ion λluence, which indicates i a lower resistance oλ these inclusions to ion sputterinμ and ii their volumetric character.

Figure . Photoμraph oλ interλerometer microscope aλter the layer oλ ~ μm was sputtered with deuterium plasma ions [ ]. The interλerence λrinμes indicate a depth oλ ~ nm λort the central λeature.

The diλλerence in composition oλ the material inside the inhomoμeneities and the surroundinμ matrix was conλirmed by the λollowinμ two methods an electron microprobe analysis Table and scanninμ electron microscope SEM with enerμy-dispersive X-ray spectroscopy EDX , Figure . The data oλ microprobe analysis demonstrate that the material in the holes is depleted in zirconium Zr and titanium Ti , and enriched with the other elements such as copper and nickel. From this result, it is not surprisinμ that the rate oλ sputterinμ oλ the inhomoμeneities

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

is hiμher than oλ the matrix. Takinμ into account, the sputterinμ yields, Y, oλ components oλ the mirror material [ ] one can λind λor deuterium ions at enerμy eV, Y Zr = . × − at/ ion, is about one order oλ maμnitude lower than Y ”e , Y Cu , Y Ni , and two times lower than Y Ti . “t lower ion enerμy, the diλλerence in sputterinμ yields is much μreater. The microprobe data show that aλter lonμ-term sputterinμ the composition oλ Zr in holes continues to be below that in the matrix. Element Nominal composition at.%

Matrix composition at.%

Zr

.

%

.

Ti

.

. %

.

%

.

%

Cu

.

%

.

. %

.

. %

.

. %

.

%

Ni ”e

% .

%

Inhomogeneity composition at.% .

%

. %

Table . Elemental composition oλ the mirror material λrom μrade # nominal second column and measured by means oλ microprobe method in the matrix and in the inhomoμeneity without takinμ into account ”e . Ratio to the Zr content is μiven in brackets [ ].

Figure . a SEM imaμe λrom backscattered electrons and b, c, e, λ the correspondinμ elemental maps obtained by EDX oλ the polished mirror surλace oλ a sample oλ μrade # . d Heiμht plot obtained by conλocal laser scanninμ micro‐ scopy.

Comparinμ the EDX maps λor zirconium, nickel, copper, and titanium Figure , one could see that λor inhomoμeneities oλ the hole type, the intensities oλ ZrL and TiK lines are in anticorrelation with intensities oλ CuL and NiL lines what is in aμreement with the results oλ

141

142

Metallic Glasses: Formation and Properties

microprobe analysis. The data obtained by conλocal laser scanninμ microscopy, Figure d, indicate that inhomoμeneity is below the matrix level. “s Figure shows, the alloys μrades # and # also contain heteroμeneities, however oλ another kind. The majority oλ the heteroμeneities have a dendritic appearance. On one occasion, a rather λull structure oλ such an inhomoμeneity was revealed aλter bombardment oλ one oλ μrade # samples with arμon ions, Figure a. Its shape is quite similar to the shape oλ crystals described by the authors oλ this alloy [ ] see also reλerences ibidem . Important, the surλace between these λeatures is continuinμ to be smooth.

Figure . SEM imaμes oλ the mirror surλace a sample oλ μrade # aλter sputterinμ with arμon ions oλ the layer oλ ~ μm, b sample oλ μrade # aλter sputterinμ with deuterium plasma ions oλ the layer oλ ~ μm.

No heteroμeneities oλ this kind were λound in the other two alloys μrade # and # . . . Plasma exposure and methods of surface analysis To simulate the impact oλ charμe-exchanμe atoms CX“ λlux on ”MG mirror samples in ITER, ions oλ deuterium or arμon plasma were used. The detail description oλ the experimental stand DSM- used λor perλorminμ the ion bombardment can be λound in [ , ]. The ”MG mirror samples were exposed to electron cyclotron resonance plasma ECR, λrequency oλ μenerator . GHz produced in a double-mirror maμnetic conλiμuration. Mirror specimens were λixed on a water-cooled holder just outside oλ maμnetic mirror. Durinμ the exposure, the tempera‐ ture did not exceed °C. The electron density oλ plasma was ~ m− and electron tempera‐ ture ~ eV. “ λixed neμative voltaμe in the ranμe – V was supplied to the mirror holder λor the acceleration oλ ions to the mirror surλace. In some cases, when usinμ deuterium plasma, the ion λlux was enerμy distributed between eV and eV by the combination oλ λixed neμative potential and a time-varyinμ λrequency Hz halλ-wave positive potential λor ion acceleration [ , ]. The latter was done to be closer to real enerμy distribution oλ CX“ λlux, calculated λor ITER [ ]. The mean ion current density to the sample was oλ the order oλ m“/ cm , that is, ~ ions/m .

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

In the low-enerμy ECR deuterium discharμe, there are usually not only monoatomic ions D+ but also molecular ions D + and D + [ ]. In proximity to the mirror surλace, such polyatomic ions λall apart into atoms and monoatomic ions in such a way that their enerμy, acquired durinμ passinμ the acceleratinμ voltaμe, is divided equally between λraμments. Thus, the surλace oλ mirror is bombarded by particles oλ equal mass but diλλerent enerμies E, E/ , and E/ , where E is the ion enerμy due to application oλ an acceleratinμ voltaμe. With this peculiarity oλ plasma composition and sinusoidal time variation oλ acceleratinμ voltaμe, the character oλ calculated enerμy distribution oλ projectiles bombardinμ the mirror surλace is in a qualitative accordance with the CX“ enerμy distributions measured at tokamaks [ , ] and calculated λor ITER conditions [ ]. “ll samples were exposed in several steps in DSM- to nonmass-separated plasma ions in order to study the evolution oλ investiμated properties with λluence. “λter each exposure step, the mass chanμe and the reλlectance at normal incidence oλ liμht ranμe – nm were measured. The state oλ sample surλace was analyzed by various microscopes in addition to mentioned SEM with EDX optical microscope, interλerometer microscope, conλocal laser scanninμ microscopy CLSM , atomic λorce microscopy “FM λor some ”MG samples, the state oλ surλace was studied by secondary ion mass spectrometry SIMS , electron microprobe method, and laser ablation method.

. Properties after plasma exposure . . Absorption of deuterium . . . “morphous specimens It was λound that mirror samples oλ all μrades absorb larμe amounts oλ deuterium aλter the mirror specimens were bombarded with ions oλ the deuterium plasma. ”ecause oλ deuterium absorption, a weiμht μain was observed λor all tested ”MG μrades, even iλ sometime clearly sputter erosion takes place. This is expected due to the hiμh λraction oλ hydride λorminμ elements Zr, Ti, ”e . In Table , the weiμht μain aλter each plasma exposure step λor a sample oλ μrade # is shown. For low enouμh enerμies, that is, up to eV, one may suppose that ion sputter erosion is neμliμible. Thereλore, the ratio oλ retained deuterium to the impactinμ ions can be calculated up to eV, and is shown in Figure . ”ecause the ECR discharμe does not only produce D+, but also D + and D + ions, the total λlux oλ deuterium projectiles has to be somewhat larμer than the ion λlux measured, and, correspondinμly, the portion oλ retained deuterium to impactinμ D atoms has to be noticeably lower, assumed within a λactor oλ ~ . The set oλ data λor deuterium trappinμ by diλλerent samples is shown in Figure . The ion enerμy was λixed at eV λor the samples oλ μrades # and # , and at eV λor both sides oλ sample oλ μrade # , while it was increased λrom to eV λor the sample oλ μrade # . Obviously, λor λixed ion enerμy the deuterium uptake increases linearly with ion λluence.

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Metallic Glasses: Formation and Properties

No of exposure

Accelerating voltage, Ion fluence,

Current density, Weight change, Absorption, D/ion, %

step

V

mA/cm

ions/m

μg/cm



.

.

+



.

.

+

.



.

.



.

.

+

.



.

.

+

.

.



.

.

+

.

.



.

.

+

.

.

.

. ? .

Table . Results oλ sequential exposures oλ one sample oλ μrade # in deuterium plasma.

Figure . Dependence on ion enerμy oλ the portion oλ deuterium in comparison to deuterium ion λluence retained in the amorphous mirror sample oλ μrade # [ ].

Figure . Dependence oλ deuterium retention λor ”MG samples on ion λluence λound by measurinμ weiμht μain. The details are in the text. Inserted is the sketch oλ the sample oλ μrade # aλter λinishinμ its exposure in deuterium plasma on both sides. The chipped side was exposed secondly trianμles, # - to the λluence . × ions/m .

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

“ very low level oλ absorbed deuterium λor the sample oλ μrade # at the lowest λluences and eV is most likely explained by insuλλicient cleaninμ oλ this particular sample. “ clear linear dependence takes place λrom eV up to eV λluence ranμe λrom . – × ions/ m , then sputterinμ starts to play more and more important role with increasinμ ion enerμy at ion acceleratinμ voltaμe oλ V, sputterinμ and absorption are about equal, but with λurther rise oλ ion enerμy the weiμht started to decrease showinμ that sputterinμ prevails over deuterium absorption. “n abrupt transition λrom linear dependence to saturation μives chance to estimate the thickness oλ sputtered layer, as it is described below Section . . “λter last exposure in deuterium plasma ion λluence~ . × ions/m the # μrade sample became curved with radius oλ curvature ~ cm in such a way that the side exposed to D plasma ions became convex. The experiment with exposinμ λront side oλ sample oλ μrade # data # - was also stopped at ion λluence . × ions/m aλter its bendinμ was discovered, aμain with exposed side to be convex curvature radius ~ cm . This λact indicates appearance oλ tension due to an increase in the speciλic volume oλ the nearsurλace layer exposed with deuterium plasma ions. The λact oλ bendinμ means that deuterium does not penetrate throuμh the whole thickness oλ samples, and only near-surλace layer, oλ thickness less than the total thickness oλ sample, is increasinμ in volume.

Figure . Straiμhteninμ oλ the sample oλ μrade # by exposinμ to the backside # - in Figure

.

It was decided to continue the experiments by exposinμ the back side oλ the # sample in similar conditions data # - in Figure . Durinμ sequential exposures, the sample started to straiμht‐ en μradually Figure and became plane at ion λluence ~ . × ions/m , that is, aλter approximately the same λluence to the λront side caused bendinμ oλ the sample. Three λurther exposures resulted in beμinninμ to bend the sample in opposite direction radius oλ curvature ~ cm and its partial destruction. Schematically, the shape oλ the sample aλter λull cessation oλ experiment is shown as insert in Figure . The thickness oλ the chipped oλλ part is ~ . mm. “ssuminμ the λormation oλ zirconium hydride in δ-phase ZrD~ . by all retained deuterium, the thickness oλ hydride layer would be . mm. This is excellent aμreement with the observed thickness oλ the chipped λraμment.

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Metallic Glasses: Formation and Properties

In Table , the eλλiciency oλ deuterium absorption is shown as the ratio oλ retained D atoms to the whole λluence oλ ions, λound as the product oλ measured ion saturation current and the total exposure time, because the real proportion between one-, two-, and tree-atomic ions is not known see comments in Section . . It λollows λrom Table that the eλλiciency oλ uptake depends stronμly on ion enerμy and composition, but even with identical composition diλλerence is siμniλicant λor μrades # and # . Samples

#

#

#

#

Ion enerμy, eV Ion λluence,

– ions/m− Fi

.

Mass μain, mμ/cm ND absorbed,

.

. −

D atoms/m

.

Eλλiciency oλ absorption ND/Fi

.

Portion oλ Zr in composition

.

.

.

.

.

.

.

.

.

Table . Comparison oλ deuterium uptake by diλλerent ”MG samples.

Comparinμ data oλ the table, it looks like there is a tendency λor eλλiciency oλ uptake to increase with increasinμ the portion oλ zirconium. “ laser ablation technique was applied as an attempt to see the depth distribution oλ trapped deuterium. The diameter oλ laser spot was ~ . μm, and the step in depth λor every laser shot was ~ . μm. “s shown in Figure , deuterium is conλidently reμistered even aλter the laser crater depth reached ~ μm, However, in this method, the eλλect oλ side walls oλ the laser crater cannot be λully excluded, and thereλore, any quantitative conclusion on the real depth distribution oλ implanted deuterium can only be done with a deλinite precaution.

Figure . The amplitude oλ D+ peak in ablated material as λunction oλ laser crater depth λor a specimen oλ μrade # exposed to deuterium plasma ions eV/ion up to a λluence oλ . × ions/m .

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

Samples that were not exposed to deuterium plasma did not reveal a deuterium peak in their mass spectra, as can be seen in Figure . The hydroμen peak was reμistered in every spectrum however, it is well known that such peak can be an artiλact appearinμ due to not perλect vacuum conditions in the mass spectrometer chamber.

Figure . Mass spectra oλ the liμhter masses oλ ablated material λor samples oλ μrade # a not exposed in plasma and b aλter exposure to deuterium plasma up to ion λluence . × ions/m with diλλerent ion enerμy. Peaks oλ ions H+, D + , ”e +, C +, and ”e+ can be clearly distinμuished [ ].

For another sample oλ μrade # exposed to a total ion λluence oλ ~ × m− , a second laser ablation test was carried out days aλter the exposure. Durinμ that time, the sample was stored at ambient atmospheric conditions. This sample showed a very diλλerent depth distribution, namely deuterium was not reμistered durinμ λirst λive shots, and then D+ peak appeared in the sixth shot, reached maximal amplitude in the th shot and μradually decayed to zero in the th laser shot. However, the weiμht oλ the sample did not chanμe, and thus, it may be assumed that the retained deuterium was redistributed inside the sample or that some oλ the retained deuterium was released, but the associated weiμht loss was balanced out by oxidation. . .2. Crystallized specimens Two samples oλ μrade # , one amorphous ”MG and one crystallized ”MG-C , were exposed in similar conditions λor a lonμ term to low-enerμy eV ions oλ deuterium plasma with a

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Metallic Glasses: Formation and Properties

current density oλ . m“/cm . Other important details oλ the experiment are presented in Table . Durinμ the λourth exposure total ion λluence . × ions/m , the ”MG-C was λraμmented, as shown in Figure a. On the contrary, its amorphous counterpart maintained its inteμrity and continued to μain weiμht durinμ the λollowinμ three exposure steps total ion λluence . × ions/m . The reλlectance did not appear to be aλλected noticeably not shown . No of exposure

Ion fluence,

ions/m

Weight gain for BMG, μg

Weight gain for BMG-C, μg

. . . .

Destroyed

. . . Table . The history oλ weiμht chanμe durinμ exposure oλ samples oλ μrade # to ions oλ deuterium plasma oλ ion, ion current density . m“/cm .

eV/

Figure . Photoμraphs oλ a the ”MG-C crystallized sample oλ μrade # aλter the λourth exposure to eV ions oλ deuterium plasma and b its amorphous counterpart ”MG aλter the seventh exposure to eV ions oλ deuterium plasma [ ].

Not catastrophic λor the sample, but detrimental λor the optical characteristics was the modiλication oλ the surλace oλ another crystallized mirror samples oλ μrade # , exposed to deuterium plasma ions in diλλerent reμimes, Table . “λter the last exposure on the surλace oλ this sample, there appeared deλects in the λorm oλ chips and cracks oλ diλλerent sizes and siμniλicant weiμht loss was measured due to the loss oλ some λraμments inside the vacuum chamber. The SEM imaμes oλ chip surλace at two maμniλications, presented in Figure , exhibit the characteristic λor a brittle rupture. The amorphous counterpart was unaltered.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

Such diλλerence in behavior oλ ”MG and ”MG-C samples is in a μood qualitative aμreement with the results published by Suh and “soka Kumar on amorphous and crystallized samples oλ similar composition that were subjected to cathodic charμinμ [ ]. “λter a critical charμinμ time, both amorphous and crystalline samples disinteμrated. However, the maximum hydroμen content beλore total disinteμration was up to times μreater λor the amorphous phase compared to the crystalline counterpart. No of exposure Working gas Ion energy Current density, mA/cm

Ion fluence,

“r

eV

.

.

“r

eV

.

.

“r

eV

.

D

eV

.

.

D

eV

.

.

D

eV

.

.

D

eV

.

.

D

eV

.

.

D

eV

.

ions/m

Mass loss, μg

.

.

Table . The history oλ exposure oλ one oλ the ”MG-C crystallized samples oλ μrade # to ions oλ arμon and deuterium plasma.

Figure

. SEM imaμes oλ the crystallized sample oλ μrade # aλter the last exposure no

shown in Table [

].

. . Sputtering rate The adsorption oλ deuterium makes diλλiculties when tryinμ to obtain data on sputterinμ yield by measurinμ the weiμht loss. Thereλore, a stainless steel diaphraμm diameter mm which hid the rest part oλ ”MG sample surλace oλ μrade # diameter mm was used λor determi‐

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Metallic Glasses: Formation and Properties

nation oλ the rate oλ sputterinμ by deuterium plasma ions. The depth oλ the erosion hole measured λor ion enerμies , , , and eV are presented in Table . The corre‐ spondinμ values oλ sputterinμ yield were estimated with takinμ into account the depth oλ the hole, ion current density, and exposure time. The results are presented in Figure . No. of experiment Ion energy, eV Depth of erosion, nm

Microhardness, kg/mm ±

Mass change, μg k

 n

-

.

  .

.

  .

±

Not sputter eroded surλace Table . Results oλ exposure throuμh a diaphraμm acceleratinμ voltaμe.

±

+

.

  .

±

+

.

  .

.

  .

± mm in diameter oλ the sample oλ μrade # λor λour ion

Figure . Sputterinμ yield λound λrom the depths oλ holes appeared on the surλace oλ sample oλ μrade # exposed to D+ plasma ions throuμh an mm diaphraμm.

In addition to the erosion depth, the microhardness and the optical constants oλ surλace inside each exposure spot were determined. It is seen λrom the Table that D+ ion bombardment modiλies the hardness oλ the near surλace layer. The limited inλormation on this subject, like the amount oλ trapped deuterium, its depth distribution, etc., does not allow to make a deλinite conclusion on the reason oλ this phenomenon, λor example, on its link with volumetric density oλ trapped deuterium, as was λound by the authors oλ [ ]. “λter λinishinμ the exposure to ions oλ enerμies indicated in the Table , optical indices n and k were measured with ellipsometry. Table shows that both optical indices do not depend stronμly on the chosen D+ ion enerμies and λluences however, they diλλer a little λrom the indices measured λor the initial surλace, that is, not eroded by ions oλ deuterium plasma due to protection by diaphraμm.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

”ecause oλ the small size Ø mm in diameter oλ μrade # samples the deuterium adsorption was not measured quantitatively λor these samples too low mass μain . “iminμ to obtain data on sputterinμ, one sample oλ μrade # was exposed to ions throuμh a Ni μauze with wire diameter oλ μm and mesh window width oλ × μm. The data obtained are shown in Figures and . Figure a shows a typical photoμraph made with an interλerence micro‐ scope oλ the sample oλ μrade # aλter it was bombarded with deuterium plasma ions Ei= eV, total ion λluence ~ × ions/m throuμh the μauze, and Figure b shows the result oλ processinμ oλ the interλerence picture alonμ several windows between wires oλ the μauze. “s one can see, the sputter depth is about μm. This depth corresponds to the sputterinμ yield λor not mass separated deuterium ion λlux Y≈ . atom/ion, which is about a λactor two oλ the value λound λor μrade # Figure .

Figure . Interλerence λrinμes on the surλace oλ the sample oλ μrade # aλter lonμ-term bombardment throuμh the mesh a , and the structure oλ the relieλ λound by processinμ oλ the interλerence λrinμes b [ ].

The optical and “FM data presented in Figure and Figure demonstrate similar depth oλ erosion in diλλerent windows. “lso, they do clearly indicate that sputterinμ oλ the wire itselλ occurs because the edμes oλ each wire are bombarded by ions at small anμle to the surλace, λor which the sputterinμ yield exceed considerable the yield at normal incidence [ ].

Figure . “FM data λor the same sample oλ μrade # exposed to plasma throuμh Ni μauze as in Figure a -D picture near crossinμ oλ Ni wires, b D picture oλ same data and heiμht distributions alonμ the lines shown in picture.

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Metallic Glasses: Formation and Properties

Important to note that saturation seen λor the μrade # sample in Figure is not because cessation oλ D absorption but due to λaster mass loss caused by ion sputterinμ in comparison with the mass μain due to deuterium implantation. Takinμ this into account and comparinμ the linear λit oλ deuterium absorption with the saturation level oλ mass rise measured, one can λind λor this sample that thickness oλ sputtered layer aλter last exposure shown in Figure is ~ μm, mainly due to sputterinμ by ions in the ranμe – eV. With the use oλ “r plasma ions, there was not any problem to measure the weiμht loss even aλter short exposure. The results oλ these measurements λor amorphous and crystallized specimens oλ μrade # , dependinμ on the ion enerμy, are presented in Figure in comparison with data solid line λor the bombardment oλ zirconium with “r+ ions λrom [ ]. “s seen, there is no measurable dependence in the sputter rate λor samples in diλλerent state. ”esides, there is a trend λor lower sputter erosion rate oλ both samples in comparison to zirconium. This is in qualitative aμreement with data λor stainless steel in [ ], sputterinμ yield oλ which was less than sputterinμ yields oλ each component Fe, Ni, Cr .

Figure . Sputterinμ yield dependinμ on the “r+ ion enerμy λor amorphous ”MG and crystalline ”MG-C mirror samples oλ μrade # . For comparison sputterinμ yield λor pure Zr by “r λrom Yamamura and Tawara [ ] are shown. “morphous sample ”MG- and crystallized sample ”MG-C- were made λrom same inμot.

When perλorminμ this experiment, not only mass loss but also the reλlectance at normal incidence was measured aλter each sputterinμ procedure. Thus, the behavior oλ reλlectance under lonμ-term sputterinμ λor samples ”MG and ”MG-C made λrom the same inμot oλ μrade # but with diλλerent structure was determined see next section . . . Modification of optical properties of amorphous mirrors . . . Impact of deuterium plasma ions When exposinμ ”MG specimens oλ μrades # and # in deuterium plasma, it was λound that even rather short-time bombardment with ions oλ keV enerμy ranμe, when sputter erosion

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

could be neμlected, leads to a measurable decrease oλ reλlectance in the whole spectral interval oλ measurement, both λor amorphous ”MG and λor crystallized ”MG-C mirror samples. The decrement oλ reλlectance chanμe depends on the wavelenμth the shorter the wavelenμth the deeper the drop. The reλlectance may be restored by lonμ-time exposure to low enerμy ~ eV ions oλ deuterium plasma the restoration is λull or partial, dependinμ on the exposure time to low-enerμy ions. This means that the reλlectance decrease due to keV ion enerμy impact is not connected with modiλication oλ the surλace microrelieλ but with some chemical processes on the surλace, like it was established earlier λor ”e and “l mirror samples [ , ]. We shall discuss this λact below in a special section. “s an example, Figure demonstrates the results obtained on one oλ ”MG-C samples oλ μrade # . Initially, it was sputtered with ions oλ “r plasma, what resulted in the development oλ some surλace rouμhness and correspondinμ reλlectance decrease solid circles ⇒ squares . Then, the sample was bombarded with . keV ions oλ deuterium plasma, what caused λurther reλlec‐ tance drop rhombuses however, not due to increase oλ the surλace rouμhness, because this drop was λully restored by subsequent much lonμer exposure to low-enerμy eV ions oλ the same deuterium plasma open circles . In contrast, the drop oλ reλlectance caused by “r ion sputterinμ could not be restored in similar way, that is, by exposinμ sample to low enerμy D+ ions, as it occurred due to development oλ surλace rouμhness.

Figure . Spectral dependence oλ reλlectance λor ”MG-C mirror sample oλ μrade # aλter short exposure ion λluence oλ the order ions/m to ions oλ deuterium plasma with enerμy keV ♦ and λollowinμ much lonμer exposure to eV deuterium plasma ions ο .

The correlation oλ drop and restoration oλ reλlectance λor ”e-containinμ ”MG samples oλ μrade # is in qualitative aμreement with the results oλ experiments when specimen oλ μrade # was exposed by turns to hiμh . keV and low eV enerμy ions oλ deuterium plasma [ ].

153

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Metallic Glasses: Formation and Properties

For ”MG specimens oλ μrade # , qualitatively similar eλλect was observed however, detail investiμation was not provided because oλ small size diameter mm . . .2. Impact of argon plasma ions To avoid the eλλects connected with chemistry, the consequence oλ lonμ-term erosion was studied by the use oλ “r plasma ions. In comparison with deuterium, the use oλ arμon as a workinμ μas is characterized by the λollowinμ three principle moments arμon is inert under most conditions and does not λorm conλirmed stable compounds at room temperature arμon uptake and the creation oλ a contaminatinμ layer on the metal surλace may be neμlected and the time oλ experiments can be much shortened because oλ the sputterinμ rate oλ ”MGs with “r+ ions is siμniλicantly hiμher as compared with D+ ions. The evolution oλ the microrelieλ and the reλlectance oλ ”MG mirror sample oλ μrade # was determined up to erosion depth oλ . μm by sequential exposures to “r+ ions oλ diλλerent enerμy . – . keV . Figure shows the reλlectance at three wavelenμths versus the thickness oλ eroded layer. The data demonstrate that the amorphous mirror maintains the reλlectance at about initial value even aλter stronμ sputter erosion. The constancy oλ reλlectance means that the surλace oλ amorphous mirror does not chanμe durinμ all sputterinμ procedures. This λact conλirms the supposition made by Voitsenya et al. [ ].

Figure . Dependence oλ reλlectance oλ one oλ ”MG sample mirrors oλ μrade # on thickness oλ layer eroded due to maniλold exposures to ions oλ “r plasma with diλλerent enerμy [ ].

“t the same time, the surλace oλ ”MG-C sample μrade # made λrom same inμot but crystal‐ lized, became quite rouμh aλter sputterinμ to much lower erosion depth. In Figure , the comparison oλ SEM imaμes oλ both samples are shown at the same maμniλication. One can observe that the relieλ appeared on the surλace oλ the crystallized sample oλ μrade # has approximately the lateral size oλ ~ μm, while the amorphous sample oλ μrade # stays very smooth.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

Figure . SEM photos oλ a ”MG and b ”MG-C specimens oλ μrade # by “r plasma ions to the depth . μm and . μm, respectively [ ].

made λrom the same inμot aλter sputterinμ

We have to emphasize a very siμniλicant diλλerence in behavior oλ optical characteristics oλ mirror samples oλ μrade # with amorphous ”MG and crystalline ”MG-C structures. The development oλ rouμhness on the surλace oλ ”MG-C Figure resulted in deμradation oλ its ability to transmit an imaμe. Figure shows the imaμes oλ He-Ne laser beam spot aλter reλlection a λrom etalon mirror “l λilm on quartz, that is, ideal mirror b λrom amorphous mirror oλ μrade # aλter lonμ-term bombarded with “r+ ions . μm sputtered and c λrom crystallized ”MG-C sample oλ same μrade aλter “r bombardment . μm sputtered . The reλlection oλ the crystallized ”MG-C is stronμly distorted, while the imaμe oλ the eroded ”MG has about the same quality as that λrom the ideal mirror . The results oλ comparative analysis obtained when processinμ these imaμes are presented in “ppendix ”.

Figure . Imaμe oλ the Ne–He laser beam spot aλter the beam was reλlected λrom a ideal mirror , b λrom sample oλ μrade # aλter sputterinμ . μm, and c λrom crystallized sample oλ μrade # aλter sputterinμ . μm [ ].

. . Role of chemical processes The qualitative similarity between behavior oλ reλlectance oλ ”e-containinμ ”MG mirrors and ”e mirrors exposed in similar manner suμμests that in all cases similar processes are realized on the surλace λor mirror aλter impact oλ deuterium plasma ions with hiμh ~ . keV and low ~ eV enerμies. From the study oλ ”ardamid et al [ ], the reason oλ this transλormation oλ ”eO into ”e OD and μradual rise oλ hydroxide λilm thickness aλter the sample is subjected to

155

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Metallic Glasses: Formation and Properties

keV-enerμy ions oλ the plasma that contains also some amount oλ oxyμen. “n inverse process, namely restoration oλ initial state with thin ”eO coatinμ, on the ”e surλace occurs when ion enerμy is low. To check this explanation, six ”MG samples oλ μrade # were sputter eroded by ions oλ “r plasma to the depth oλ ~ μm λor takinμ oλλ the consequences oλ previous work with them. Then, λour oλ them were bombarded by D plasma ions with enerμy oλ . keV. “λter that two oλ them were additionally lonμ-term exposed to eV ions oλ D plasma. SIMS analysis with Cs+ ions as projectiles was provided in two diλλerent points oλ the surλace oλ the three pairs oλ samples. In connection with μood equivalence oλ λour data sets λor each procedure exposure to “r ions, plus exposure to keV D ions, plus exposure to eV D ions , only results obtained λrom one point oλ one sample are shown in Figure λor simplicity.

Figure . SIMS data λor release oλ oxide ions λor all components composed oλ ”MG specimens oλ μrade # aλter a sputterinμ with “r ions, b sputterinμ by “r ions plus bombardment with . keV ions oλ deuterium plasma, c same as b plus lonμ-term exposure to eV ions oλ deuterium plasma d comparison oλ the data λor ”eO ions only [ ].

Evidently, the bombardment with keV-ranμe D plasma ions leads to increasinμ the thickness oλ uppermost oxidized layer, but the λollowinμ exposure to low-enerμy ions oλ similar D plasma results in thinninμ this layer and restoration oλ the initial depth distributions oλ all oxides. These SIMS measurements were carried out with specimens havinμ only . at.% oλ beryllium. “t the same time, the release oλ ”eO+ ions siμniλicantly exceeds the release oλ all other oxides

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

toμether. However, because the relative sensitivities oλ the various oxides in the SIMS meas‐ urements are unknown, no conclusion about depth distribution oλ diλλerent oxides in the nearsurλace layer can be obtained. It is worth to add that SIMS results are in qualitative aμreement with data oλ XPS X-ray photoelectron spectroscopy measurements on specimens with a similar composition [ ]. The authors oλ that paper have λound that the uppermost oxide λilm is mainly composed oλ ”eO in spite oλ a quite complicated chemical composition oλ the material. Similar chemical processes are probably responsible λor diλλerent ratios oλ ”e+ and ”e + peaks shown in Figure when comparinμ the output at laser ablation oλ diλλerent ions λrom ”MG samples oλ μrade # exposed and not exposed to ions oλ deuterium plasma. It is worth to note that qualitatively similar eλλects on the reλlectance drop λor keV ion enerμy ranμe and restoration λor low-enerμy ions were observed also λor ”MG samples oλ “lcontaininμ μrades # and # , althouμh not so stronμ. . . Blisters “λter the determination oλ sputterinμ rate by D plasma ions Section . , many small surλace λeatures were λound in that part oλ one ”MG mirror specimens mm in diameter oλ μrade # which was exposed to ions with enerμy eV diameter oλ diaphraμm mm, ion λluence . × ions/m . These blister-like λeatures covered approximately % oλ the irradiated surλace and μenerally had a round shape Figure a. Their size distribution diameter was obtained by measurinμ the dimensions oλ such subjects observed in the view λield oλ the microscope, ~ . mm . The mean diameter was ~ . μm with sizes ranμinμ λrom ~ up to μm one subject . The size distribution oλ these subjects is shown in Figure b. In the λollowinμ sections, we use the word blisters λor these subjects, in spite they do not look like a classical blisters, described by ”ehrisch [ ].

Figure . a SEM photo oλ the ”MG mirror surλace oλ μrade # λrom secondary electrons with characteristic surλace λeatures [ ]. The rouμh parts seen a little below the center are the structure deλects described in the part . see Fig‐ ures and b size distribution oλ blister.

157

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Metallic Glasses: Formation and Properties

Figure shows atomic λorce microscope “FM imaμes oλ the two diλλerent types oλ blisters and Figure – the correspondinμ cross-section proλiles. The heiμht oλ the λlat-top blister Figure a is ~ . μm and it is surrounded by an annular μroove oλ approximately similar depth. The blister oλ similar size, but with a black central area Figure can be seen in Figures b and b . This time, there is an irreμular-shaped dome on a λlat lid base the dome is seen as black color in Figure a.

Figure

. The “FM imaμes oλ two diλλerent kinds oλ blisters a with λlat lid and b with a dome-like lid [

Figure

. Cross-sectional proλiles oλ the blisters shown in Figure

[

].

].

“ small part oλ the blisters is not round, and some oλ them have a dark part localized closer to their edμe. Elemental maps obtained by EDX suμμest that the elemental composition oλ the dome diλλers λrom that oλ the main ”MG matrix. The elemental composition oλ blister surλaces was obtained by two methods microprobe analysis and EDX. On blisters with a uniλorm contrast and smooth lid, there was not λound any noticeable deviation oλ composition in comparison with the matrix. “t the same time those blister lids without contrast uniλormity, seen with ”SE, showed presence oλ impurities. “s an example, in Figure , EDX-maps are shown λor one oλ such blisters on the sample oλ μrade

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

# . It is seen that in darkest part oλ the surλace ”SE there are elements that are not a part oλ ”MG composition λor example, carbon, oxyμen, and iron.

Figure

. EDX-maps λor a part oλ # sample surλace with a larμe blister ~Ø

μm .

Figure . Imaμe oλ a larμe blister with indication oλ points, P to P , where measurements oλ composition were done a SE and b compo ”SE .

In Figure , SE and ”SE imaμes oλ a blister on another sample oλ μrade # are shown the blister lid is not smooth partly. Numbers – indicate the points where a microprobe analysis was done note that ”e, C, and O cannot be reμistered by used setup . The results oλ measure‐ ments in those points are presented in the Table . “μain, a small amount oλ impurity at the lid, this time, chromium, was reμistered. ”ased on results oλ EDX data λor other blisters, one may assume that in the dark parts oλ the lid there are C and O also. Several peculiarities oλ the blisters can be noted λrom presented data i the blister location on the surλace does not depend on the presence oλ deλects in the μrade # ”MG described in the Section . this is clearly demonstrated by Figure a ii it looks like there is a probability that larμe blisters are localized in those parts where there are some concentration oλ contami‐ nants, like iron, chromium, which can appear to be random iii the parts oλ blister lids that look darker in ”SE imaμes are contaminated with carbon and oxyμen iv around practically every blister, there is evidence oλ stress-induced ductile deλormation leadinμ to a depression v the inspection oλ several hundred λeatures did not result in the observation oλ any cavity such as miμht be λound λollowinμ the burstinμ oλ blisters.

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Metallic Glasses: Formation and Properties

Element

Passport data,

Matrix P , at.%

at%

Smooth part of

Coarse part of

the blister

the blister

surface P ,

surface P ,

at.%

at.%

Zr

.

.

.

.

Ti

.

.

.

.

Cu

.

.

.

.

Ni

.

.

.

.



.

.

Fe

.

Cr

.

Table . Results oλ microprobe analyses in points indicated in Figure

.

Figure . a Location oλ pricks by a nano-pin yellow crosses at the surλace oλ a larμe dome-like blister, seen in Fig‐ ure a. b The load and unloadinμ curves oλ indenter measured at locations indicated in a .

To clear up the reason oλ diλλerence between both kinds oλ blisters, the nanoindentation procedures [ ] were provided in matrix, in plane, and dome-like lids. The location oλ measurements is shown in Figures a and a. The results are presented in Figures b and b. It is seen that the dark part oλ blister lid is soλt and ductile. Three zones can be distinμuished on the load curve soλt near-surλace layer with ~ nm in thickness I , transition layer II , and more solid material comparable with hardness oλ matrix at the depth startinμ λrom nm III . These measurements demonstrate that the structure oλ the uppermost layer is very diλλerent λor this pair oλ blisters the dome-like one has a very soλt uppermost layer up to about nm the hardness oλ the material here became to be comparable with that measured λor the planerooλ λeature only at the depth ≥ nm. “t the same time, λor plane rooλ λeature, the hardness is much hiμher, composinμ ~ % oλ the matrix hardness.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

”listers were previously observed on metallic μlasses oλ several kinds and on their crystal analoμues aλter bombardment with H+ and He+ ions oλ enerμies up to keV [ – ]. “s was λound, on amorphous materials blisters appear at ion λluences approximately three times μreater than on the same materials in crystalline state. The blisters had dome-shaped lids with diameters in the ranμe . – . μm mean value . μm λor H+ ions [ ]. Such dimensions correlate well with the implantation ranμe oλ the H+ ions, thus suμμestinμ the λormation oλ classical blisters [ ] due to accumulation oλ hydroμen in the near-surλace reμion. In conclusion, we would aμain state that diλλerence in size, alonμ with the atypical shape like λlat lids, annular μrooves, the lack oλ broken lids, etc. , indicate that the observed λeatures are not blisters in common understandinμ oλ this word [ ].

Figure . a SEM photo λrom backscattered electrons oλ the part oλ the surλace shown in Figure a. scale is μm . Locations oλ pricks by a nano-pin at the surλace oλ a plain lid blister and nearby matrix are shown by lower arrow and upper arrows, correspondinμly. b The load and unloadinμ curves oλ indenter measured at locations indicated in a .

. Summary ”MG mirror samples oλ diλλerent elemental composition were investiμated in experiments simulatinμ the behavior oλ λirst mirrors under impact oλ charμe exchanμe atoms in the λusion reactor ITER. “s projectiles the ions oλ deuterium and arμon plasmas were used. ”ehavior oλ optical properties and surλace topoμraphy, uptake oλ deuterium, eλλects oλ sputterinμ on optical reλlectance in the visible spectrum, and chemical processes in a near-surλace layer were studied. It was λound that some oλ the studied properties have weak dependence on material compo‐ sition whereas the others show noticeable diλλerences. The λollowinμ statements can be made .

Initial reλlectance oλ ”MG mirror samples under study is close to reλlectance oλ W mirror λor wavelenμths exceedinμ nm. The hiμhest reλlectance was measured λor samples oλ μrade # independently on the state amorphous or crystallized.

.

In the body oλ three kinds oλ samples, structural inhomoμeneities were described, which are connected with deviation oλ local composition λrom the composition oλ matrix, as was studied in detail λor ”MG samples oλ μrade # . The composition oλ inhomoμeneities, that

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Metallic Glasses: Formation and Properties

occupied small part oλ the sample volume, was depleted with zirconium in comparison with amount oλ Zr on averaμe. The ion-sputterinμ rate oλ material in these parts was hiμher than that oλ the main matrix, which leads to shallow depressions appearinμ on the sample surλace with a rather planar bottom oλ – μm in size and the depth depended on sputterinμ time. On the samples oλ μrades # and # inhomoμeneities oλ relieλ became clearly visible aλter erodinμ a layer exceedinμ μm, either with ions oλ arμon μrade # or deuterium plasma μrade # . .

Under lonμ-term ion bombardment thickness oλ sputter eroded layer was . μm in the case oλ # μrade no noticeable chanμe oλ surλace topoμraphy and reλlectance λor matrix oλ all ”MG mirror samples was observed, excludinμ μrade # , where inhomoμeneities occupied a rather biμ portion oλ the surλace. In contrast, siμniλicant deμradation oλ the mirror surλace quality was λound λor the crystallized ”MG sample oλ μrade # sputtered to much less depth than its amorphous counterpart in similar conditions.

.

Under bombardment by deuterium plasma ions some amount oλ deuterium is absorbed, with the tendency to increase oλ absorbed portion when increasinμ the portion oλ zirco‐ nium. The hiμhest value oλ deuterium was absorbed by ”MG sample oλ μrade # . × D atoms/m . The absorbed deuterium, most likely, does not penetrate throuμh the whole thickness oλ the sample but is accumulated in the layer with a thickness oλ a λew tenths oλ millimeter note, thickness oλ samples is – mm . “t ion λluence exceedinμ × ions/m samples become bent in such a way that the λront side exposed to ions oλ deuterium plasma became convex with radius oλ curvature ~ cm λor μrade # and ~ cm λor μrade # . Followinμ exposures in similar conditions oλ the back side oλ μrade # resulted in μradual straiμhteninμ oλ the sample, and even bendinμ it in opposite direction with λurther increase in ion λluence.

.

For samples oλ μrade # in amorphous state, the deuterium ion λluence × D ions/m much exceeded the value λor their crystallized counterparts when they started to the destroy ~ × D ions/m .

.

“λter implantation oλ deuterium up to . × D atoms/m λor μrade # sample , there was not observed any indication on appearance oλ crystallized phase in amorphous samples.

.

There was not λound any measurable eλλect oλ deuterium implantation on optical prop‐ erties oλ mirror sample, either in amorphous or crystallized states.

.

On a small part oλ the surλace oλ ”MG mirror sample oλ μrade # , exposed to ions oλ deuterium plasma accelerated to the enerμy eV, the blister-like λeatures appeared with the shape diλλerent λrom what can be λound in literature.

.

When ”MG mirrors with beryllium or aluminum in their composition are exposed to deuterium plasma, chemical processes on the surλace play a deλinite role in behavior oλ their optical properties.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

In conclusion, we noted that because oλ the hiμh absorptivity oλ hydroμen isotopes, existinμ Zr-based ”MG cannot be a prospect material λor in-vessel mirrors in ITER. However, in view oλ rather rapid proμress in the development oλ amorphous materials, there is a hope that ”MG materials with a low absorptive capacity λor hydroμen isotopes will be desiμned in the λuture, what will permit their use λor λabrication oλ in-vessel mirrors operatinμ in the conditions with hiμh λluxes oλ charμe exchanμe atoms, neutrons, and μamma radiation.

Acknowledgements The authors would like to express their thanks to “.I. ”elyaeva and O.“. Galuza, with whom some oλ presented results were obtained. Appendix A: Fabrication of amorphous alloys containing beryllium The main peculiarity oλ bulk amorphous metal alloys is that they have a multicomponent mixture as the rule, – components are used λor their preparation. “monμ the most perspec‐ tive is Zr-based bulk amorphous metal alloy that demonstrate a hiμh mechanical strenμth and λracture touμhness with μood corrosive resistance. For the system Zr-Ti-Cu-Ni-”e, the maximal critical thickness known λrom literature amounts to mm at the critical coolinμ rate – K/s. This opens the way λor preparation oλ mirror samples with a standard size diameter mm and thickness – mm λor experimental modelinμ the behavior oλ mirrors in ITER. For smeltinμ oλ inμots oλ bulk metallic μlasses in Kharkov μrade # , Table , an electric arcsmeltinμ device with the nonconsumable electrode was used. The meltinμ operation was realized in water-cooled copper mold, inside its cavity oλ reλerence shape and dimensions. The alloy has the λollowinμ composition Zr Ti Cu , Ni ”e . .. The purity oλ all metal components was not less than . weiμht percent. “s the startinμ materials the iodide Zr and Ti, cathodic Ni, electrical copper, and distilled ”e were used. “ll components were mixed in the necessary proportion and placed in a mold cavity oλ the water-cooled copper mold oλ the meltinμ λacility. “λter that the meltinμ λacility was pumped λor deμassinμ the λurnace charμe and λilled with pure arμon to the pressure that ensures a stable arc burninμ. Then, arc was initiated and the metal μot λused. The λorm and the size oλ inμots are speciλied by the mold cavity shape and amount oλ the λurnace charμe. For obtaininμ a homoμeneous composition oλ an amorphous material it was molten many times with openinμ the meltinμ λacility and with λlip over oλ the inμot λor °. The initial sizes oλ inμots diameter – mm and thickness – mm were deλined by the requirements to have the standard size oλ mirror samples mm in diameter . The billets λor λabrication oλ mirror samples and specimens λor providinμ diλλerent investiμa‐ tions oλ the properties oλ the amorphous material were prepared by means oλ electric discharμe sawinμ. Figure A shows the photos oλ one oλ inμots λrom which the upper and lower caps were cut oλλ leλt , beλore it was cut λurther into two halves riμht which served as billets λor λabrication oλ mirror samples. One halλ oλ every inμot was annealed at °K durinμ one hour.

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Metallic Glasses: Formation and Properties

The crystallization was checked by usinμ the X-ray diλλractometer DRON- are shown in Figure .

, and the results

Figure A . The photos oλ the inμot with cut oλλ upper and lower caps beλore leλt and aλter riμht cuttinμ into two halves which served as billets λor the λabrication oλ mirror samples.

Appendix B: Processing of images of laser spots reflected from test mirrors To study the structure oλ a laser spot imaμe aλter laser beam reλlection λrom a μiven surλace, the samplinμ oλ briμhtness oλ pixels ” x λor every imaμe in Figure was carried out alonμ its diameter λollowinμ a horizontal line or lines inclined at °, °, and ° to this line. The spot is pixels in diameter that corresponds to mm. We removed a samplinμ component correspondinμ to a radial decrease in the spot briμhtness "a bell shape" usinμ a Fourier transλorm. The sequential combination oλ samplinμs extends the samplinμ path and improves the statistics.

Figure B . The contour diaμrams oλ the distributions ΔN/ ΔΛ⋅Δ” . The desiμnations correspond to the laser spot im‐ aμes a–c oλ Figure .

“ comparison oλ the samplinμs "by eye" shows that ”a x is characterized by very short wavelenμths and very low variations oλ the briμhtness. These parameters are hiμher λor ”b x and much hiμher λor ”c x by the reason oλ increasinμ the surλace rouμhness. The procedure described in [ ] is intended to compute the one-dimensional distribution λunction ΔN Λ /ΔΛ or, in other words, the number oλ waves N in a spectral interval ΔΛ where the wavelenμth Λ is alonμ the imaμe. We extended this λunction to two-dimensional one ΔN Λ,” / ΔΛ⋅Δ” . The λunctions a–c are marked with constant-level lines in the contour diaμram oλ Figure B . The upper limits oλ the wavelenμth axis Λ are the same λor all the spots, and those oλ the briμhtness axis ” are diλλerent.

On the Prospects of Using Metallic Glasses for In-vessel Mirrors for Plasma Diagnostics in ITER http://dx.doi.org/10.5772/63885

We observed one μroup with the very short wavelenμths and the very low oscillations oλ the briμhtness λor the laser spot a the ideal mirror . The spot b ”MG mirror has three μroups oλ the short and medium wavelenμths, and the medium variations oλ the briμhtness. “nd the spot c ”MG-C mirror has λive μroups oλ the short, medium, and lonμ wavelenμths, and the medium and hiμh variations oλ the briμhtness. The results oλ an analysis are shown in Table B . Bav, a. u. a

.

Bs, a. u. .

Bmax, a. u.

Sm, pixels

Λmax, pixels

.

b c Table B . The parameters oλ distributions ”av the averaμe briμhtness, ”s hiμhest ”, Sm the averaμe wavelenμth, Λmax the lonμest wavelenμth.

the standard deviation oλ ”, ”max the

The obtained distributions are in μood aμreement with the photos a–c oλ the laser spots. “ll the parameters in Table B are μrowinμ λrom a to c . Especially, λive μroups related to the spot c conλirm the existence oλ very lonμ irreμularities and, eventually, indicate a surλace λracture.

Author details Vladimir S. Voitsenya *, “lexandra F. ”ardamid , Martin ”alden , Flaviu Gostin , Serμey V. Khovrich , Vladimir G. Konovalov , Konstantin V. Kovtun , Petro M. Lytvyn , Serμey V. Ketov , Dmitri V. Luzμuine-Luzμin , Serμei I. Solodovchenko , “natoly N. Shapoval , “natoly F. Shtan’ , Vladislav N. ”ondarenko , Ivan V. Ruzhkov , Ol’μa O. Skoryk and “ndrei “. Vasil’ev *“ddress all correspondence to [email protected] National Science Center Kharkov Institute oλ Physics and Technoloμy , Kharkov, Ukraine Taras Shevchenko National University oλ Kyiv, Kyiv, Ukraine Max-Planck-Institut λ(r Plasmaphysik, Garchinμ, Germany Leibniz-Institute λor Solid State and Materials Research IFW Dresden, Dresden, Germany Institute oλ Semiconductor Physics oλ N“SU, Kyiv, Ukraine WPI “dvanced Institute λor Materials Research, Tohoku University, Sendai, Japan

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Metallic Glasses: Formation and Properties

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