Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology New Series / Editor in Chief: W. Martienssen
Group IV: Physical Chemistry Volume 11
Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data critically evaluated by MSIT® Subvolume E Refractory Metal Systems Part 1 Selected Systems from Al-B-C to B-Hf-W Editors G. Effenberg and S. Ilyenko
Authors Materials Science and International Team, MSIT®
ISSN
1615-2018 (Physical Chemistry)
ISBN
978-3-540-88052-3 Springer Berlin Heidelberg New York
Library of Congress Cataloging in Publication Data Zahlenwerte und Funktionen aus Naturwissenschaften und Technik, Neue Serie Editor in Chief: W. Martienssen Vol. IV/11E1: Editors: G. Effenberg, S. Ilyenko At head of title: Landolt-Börnstein. Added t.p.: Numerical data and functional relationships in science and technology. Tables chiefly in English. Intended to supersede the Physikalisch-chemische Tabellen by H. Landolt and R. Börnstein of which the 6th ed. began publication in 1950 under title: Zahlenwerte und Funktionen aus Physik, Chemie, Astronomie, Geophysik und Technik. Vols. published after v. 1 of group I have imprint: Berlin, New York, Springer-Verlag Includes bibliographies. 1. Physics--Tables. 2. Chemistry--Tables. 3. Engineering--Tables. I. Börnstein, R. (Richard), 1852-1913. II. Landolt, H. (Hans), 1831-1910. III. Physikalisch-chemische Tabellen. IV. Title: Numerical data and functional relationships in science and technology. QC61.23 502'.12 62-53136 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in other ways, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable for prosecution act under German Copyright Law. Springer is a part of Springer Science+Business Media springeronline.com © Springer-Verlag Berlin Heidelberg 2009 Printed in Germany The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Product Liability: The data and other information in this handbook have been carefully extracted and evaluated by experts from the original literature. Furthermore, they have been checked for correctness by authors and the editorial staff before printing. Nevertheless, the publisher can give no guarantee for the correctness of the data and information provided. In any individual case of application, the respective user must check the correctness by consulting other relevant sources of information. Cover layout: Erich Kirchner, Heidelberg Typesetting: Materials Science International Services GmbH, Stuttgart Printing and Binding: AZ Druck, Kempten/Allgäu
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Editors: Associate Editor:
Günter Effenberg Svitlana Ilyenko Oleksandr Dovbenko
MSI, Materials Science International Services GmbH Postfach 800749, D-70507, Stuttgart, Germany http://www.matport.com
Authors: Materials Science International Team, MSIT® The present series of books results from collaborative evaluation programs performed by MSI and authored by MSIT®. In this program data and knowledge are contributed by many individuals and accumulated over almost twenty years, now. The content of this volume is a subset of the ongoing MSIT® Evaluation Programs. Authors of this volume are: Zoya Alekseeva, Moscow, Russia
Jozefien De Keyzer, Heverlee, Belgium
Christian Bätzner, Stuttgart, Germany
Natalia Kol’chugina, Moscow, Russia
Natalia Bochvar, Moscow, Russia
Kostyantyn Korniyenko, Kyiv, Ukraine
Anatoliy Bondar, Kyiv, Ukraine
Artem Kozlov, Clausthal-Zellerfeld, Germany
Jörg Beuers, Hanau, Germany Gabriele Cacciamani, Genova, Italy Andriy Grytsiv, Vienna, Austria Lesley Cornish, Randburg, South Africa
Vasyl Kublii, Kyiv, Ukraine Viktor Kuznetsov, Moscow, Russia Nathalie Lebrun, Lille, France Hans Leo Lukas, Stuttgart, Germany Annelies Malfliet, Heverlee, Belgium
Damian M. Cupid, Freiberg, Germany Dmytro Pavlyuchkov, Jülich, Germany Tatiana Dobatkina, Moscow, Russia Oleksandr Dovbenko, Stuttgart, Germany Guenter Effenberg, Stuttgart, Germany Ol’ga Fabrichnaya, Freiberg, Germany Sergio Gama, Campinas, Brazil
Pierre Perrot, Lille, France Qingsheng Ran, Stuttgart, Germany Peter Rogl, Vienna, Austria Lazar Rokhlin, Moscow, Russia Eberhard E. Schmid, Frankfurt, Germany
Joachim Gröbner, Clausthal-Zellerfeld, Germany An Serbruyns, Heverlee, Belgium Mireille Harmelin, Paris, France
Elena Semenova, Kyiv, Ukraine
Frederick H. Hayes, Manchester, UK
Jean-Claude Tedenac, Montpellier, France
Michael Hoch, Cincinnati, USA
Vasyl Tomashik, Kyiv, Ukraine
Svitlana Ilyenko, Stuttgart, Germany
Mikhail Turchanin, Kramatorsk, Ukraine
Volodymyr Ivanchenko, Kyiv, Ukraine
Tamara Velikanova, Kyiv, Ukraine
Institutions The content of this volume is produced by MSI, Materials Science International Services GmbH and the international team of materials scientists, MSIT®. Contributions to this volume have been made from the following institutions:
The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia Degussa AG, Hanau, Germany Donbass State Mechanical Engineering Academy, Kramatorsk, Ukraine Forschungszentrum Jülich, Institut für Festkörperforschung (IFF), Institut Mikrostrukturforschung, Jülich, Germany I.M. Frantsevich Institute for Problems of Materials Science, National Academy of Sciences, Kyiv, Ukraine Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine Katholieke Universiteit Leuven, Department Metaalkunde en Toegepaste Materiaalkunde, Heverlee, Belgium G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine Materials Science International Services GmbH, Stuttgart, Germany Max-Planck-Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany Moscow State University, Department of General Chemistry, Moscow, Russia
School of Chemical and Metallurgical Engineering, The University of the Witwatersrand, DST/NRF Centre of Excellence for Strong Material, South Afrika Technische Universität Bergakademie Freiberg, Institut für Werkstoffwissenschaft, Freiberg, Germany Technische Universität Clausthal, Metallurgisches Zentrum, Clausthal-Zellerfeld, Germany Universidade Estadual de Campinas, Instituto de Fisica “Gleb Wataghin”, DFESCM, Campinas, Brazil Universita di Genova, Dipartimento di Chimica, Genova, Italy Universität Wien, Institut für Physikalische Chemie, Wien, Austria Universite de Lille I, Laboratoire de Métallurgie Physique, Villeneuve d’ASCQ, France Universite de Montpellier II, Laboratorie de Physico-chimie de la Materiere Montpellier, France University of Cincinnati, Department of Materials Science and Engineering, Cincinnati, USA
Preface The sub-series Ternary Alloy Systems of the Landolt-Börnstein New Series provides reliable and comprehensive descriptions of the materials constitution, based on critical intellectual evaluations of all data available at the time and it critically weights the different findings, also with respect to their compatibility with today’s edge binary phase diagrams. Selected are ternary systems of importance to alloy development and systems which gained in the recent years otherwise scientific interest. In one ternary materials system, however, one may find alloys for various applications, depending on the chosen composition. Reliable phase diagrams provide scientists and engineers with basic information of eminent importance for fundamental research and for the development and optimization of materials. So collections of such diagrams are extremely useful, if the data on which they are based have been subjected to critical evaluation, like in these volumes. Critical evaluation means: there where contradictory information is published data and conclusions are being analyzed, broken down to the firm facts and re-interpreted in the light of all present knowledge. Depending on the information available this can be a very difficult task to achieve. Critical evaluations establish descriptions of reliably known phase configurations and related data. The evaluations are performed by MSIT®, Materials Science International Team, a group of scientists working together since 1984. Within this team skilled expertise is available for a broad range of methods, materials and applications. This joint competence is employed in the critical evaluation of the often conflicting literature data. Particularly helpful in this are targeted thermodynamic and atomistic calculations for individual equilibria, driving forces or complete phase diagram sections. Conclusions on phase equilibria may be drawn from direct observations e.g. by microscope, from monitoring caloric or thermal effects or measuring properties such as electric resistivity, electro-magnetic or mechanical properties. Other examples of useful methods in materials chemistry are massspectrometry, thermo-gravimetry, measurement of electro-motive forces, X-ray and microprobe analyses. In each published case the applicability of the chosen method has to be validated, the way of actually performing the experiment or computer modeling has to be validated as well and the interpretation of the results with regard to the material’s chemistry has to be verified. Therefore insight in materials constitution and phase reactions is gained from many distinctly different types of experiments, calculation and observations. Intellectual evaluations which interpret all data simultaneously reveal the chemistry of the materials system best. An additional degree of complexity is introduced by the material itself, as the state of the material under test depends heavily on its history, in particular on the way of homogenization, thermal and mechanical treatments. All this is taken into account in an MSIT® expert evaluation. To include binary data in the ternary evaluation is mandatory. Each of the three-dimensional ternary phase diagrams has edge binary systems as boundary planes; their data have to match the ternary data smoothly. At the same time each of the edge binary systems A-B is a boundary plane for many other ternary A-B-X systems. Therefore combining systematically binary and ternary evaluations increases confidence and reliability in both ternary and binary phase diagrams. This has started systematically for the first time here, by the MSIT® Evaluation Programs applied to the Landolt-Börnstein New Series. The degree of success, however, depends on both the nature of materials and scientists! The multitude of correlated or inter-dependant data requires special care. Within MSIT® an evaluation routine has been established that proceeds knowledge driven and applies both, human based expertise and electronically formatted data and software tools. MSIT® internal discussions take place in almost all evaluation works and on many different specific questions the competence of a team is added to the work of individual authors. In some cases the authors of earlier published work contributed to the knowledge
base by making their original data records available for re-interpretation. All evaluation reports published here have undergone a thorough review process in which the reviewers had access to all the original data. In publishing we have adopted a standard format that presents the reader with the data for each ternary system in a concise and consistent manner, as applied in the “MSIT® Workplace Phase Diagrams Online”. The standard format and special features of the Landolt-Börnstein compendium are explained in the Introduction to the volume. In spite of the skill and labor that have been put into this volume, it will not be faultless. All criticisms and suggestions that can help us to improve our work are very welcome. Please contact us via
[email protected]. We hope that this volume will prove to be as useful for the materials scientist and engineer as the other volumes of Landolt-Börnstein New Series and the previous works of MSIT® have been. We hope that the Landolt Börnstein Sub-series, Ternary Alloy Systems will be well received by our colleagues in research and industry. On behalf of the participating authors we want to thank all those who contributed their comments and insight during the evaluation process. In particular we thank the reviewers - Pierre Perrot, Tamara Velikanova, Hans Leo Lukas, Marina Bulanova, Mikhail Turchanin, Nataliya Bochvar, Olga Fabrichnaya and Viktor Kuznetsov. We all gratefully acknowledge the dedicated scientific desk editing by Oleksandra Berezhnytska, Mariya Saltykova and Oleksandr Rogovtsov.
Günter Effenberg, Svitlana Ilyenko and Oleksandr Dovbenko
Stuttgart, June 2008
Contents IV/11E1 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume E Refractory Metal Systems Part 1 Selected Systems from Al-B-C to B-Hf-W Introduction Data Covered..........................................................................................................................................XI General ...................................................................................................................................................XI Structure of a System Report .................................................................................................................XI Introduction....................................................................................................................................XI Binary Systems ..............................................................................................................................XI Solid Phases ................................................................................................................................. XII Quasibinary Systems...................................................................................................................XIII Invariant Equilibria .....................................................................................................................XIII Liquidus, Solidus, Solvus Surfaces ............................................................................................XIII Isothermal Sections.....................................................................................................................XIII Temperature – Composition Sections ........................................................................................XIII Thermodynamics.........................................................................................................................XIII Notes on Materials Properties and Applications........................................................................XIII Miscellaneous .............................................................................................................................XIII References ...................................................................................................................................XVI General References ........................................................................................................................... XVII
Ternary Systems Al – B – C (Aluminium – Boron – Carbon) ...........................................................................................1 Al – B – Mo (Aluminium – Boron – Molybdenum) ............................................................................24 Al – B – Si (Aluminium – Boron – Silicon) .........................................................................................33 Al – C – Ti (Aluminium – Carbon – Titanium)....................................................................................41 Al – Cr – Ti (Aluminium – Chromium – Titanium).............................................................................72 Al – Fe – Nb (Aluminium – Iron – Niobium).....................................................................................109 Al – Mo – Ni (Aluminium – Molybdenum – Nickel) ........................................................................123 Al – Mo – U (Aluminium – Molybdenum – Uranium) ......................................................................144 Al – Nb – Ni (Aluminium – Niobium – Nickel).................................................................................164 Al – Nb – Si (Aluminium – Niobium – Silicon).................................................................................193 Al – Ni – V (Aluminium – Nickel – Vanadium) ................................................................................209 Al – O – Zr (Aluminium – Oxygen – Zirconium) ..............................................................................225 Al – Ta – Ti (Aluminium – Tantalum – Titanium).............................................................................242 B – C – Cr (Boron – Carbon – Chromium) ........................................................................................261 B – C – Hf (Boron – Carbon – Hafnium) ...........................................................................................282 B – C – Mo (Boron – Carbon – Molybdenum)...................................................................................306
B – C – N (Boron – Carbon – Nitrogen).............................................................................................323 B – C – Nb (Boron – Carbon – Niobium)...........................................................................................347 B – C – Si (Boron – Carbon – Silicon) ...............................................................................................367 B – C – Ta (Boron – Carbon – Tantalum) ..........................................................................................395 B – C – V (Boron – Carbon – Vanadium) ..........................................................................................408 B – C – W (Boron – Carbon – Tungsten) ...........................................................................................427 B – C – Zr (Boron – Carbon – Zirconium) .........................................................................................450 B – Cr – Mn (Boron – Chromium – Manganese) ...............................................................................476 B – Cr – Ti (Boron – Chromium – Titanium).....................................................................................485 B – Cr – Zr (Boron – Chromium – Zirconium) ..................................................................................494 B – Fe – Mo (Boron – Iron – Molybdenum) ......................................................................................500 B – Hf – Mo (Boron – Hafnium – Molybdenum) ..............................................................................514 B – Hf – W (Boron – Hafnium – Tungsten) .......................................................................................523
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Introduction Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data
Data Covered The series focuses on light metal ternary systems and includes phase equilibria of importance for alloy development, processing or application, reporting on selected ternary systems of importance to industrial light alloy development and systems which gained otherwise scientific interest in the recent years.
General The series provides consistent phase diagram descriptions for individual ternary systems. The representation of the equilibria of ternary systems as a function of temperature results in spacial diagrams whose sections and projections are generally published in the literature. Phase equilibria are described in terms of liquidus, solidus and solvus projections, isothermal and pseudobinary sections; data on invariant equilibria are generally given in the form of tables. The world literature is thoroughly and systematically searched back to the year 1900. Then, the published data are critically evaluated by experts in materials science and reviewed. Conflicting information is commented upon and errors and inconsistencies removed wherever possible. It considers those, and only those data, which are firmly established, comments on questionable findings and justifies re-interpretations made by the authors of the evaluation reports. In general, the approach used to discuss the phase relationships is to consider changes in state and phase reactions which occur with decreasing temperature. This has influenced the terminology employed and is reflected in the tables and the reaction schemes presented. The system reports present concise descriptions and hence do not repeat in the text facts which can clearly be read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient. However, a detailed bibliography of all cited references is given to enable original sources of information to be studied if required.
Structure of a System Report The constitutional description of an alloy system consists of text and a table/diagram section which are separated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carry the essential constitutional information and are commented on in the text if necessary. Where published data allow, the following sections are provided in each report: Landolt‐Bo¨rnstein New Series IV/11E1
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. Fig. 1 Structure of a system report
Introduction The opening text reviews briefly the status of knowledge published on the system and outlines the experimental methods that have been applied. Furthermore, attention may be drawn to questions which are still open or to cases where conclusions from the evaluation work modified the published phase diagram.
Binary Systems Where binary systems are accepted from standard compilations reference is made to these compilations. In other cases the accepted binary phase diagrams are reproduced for the convenience of the reader. The selection of the binary systems used as a basis for the evaluation of the ternary system was at the discretion of the assessor. DOI: 10.1007/978-3-540-88053-0_1 ß Springer 2009
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Solid Phases The tabular listing of solid phases incorporates knowledge of the phases which is necessary or helpful for understanding the text and diagrams. Throughout a system report a unique phase name and abbreviation is allocated to each phase. Phases with the same formulae but different space lattices (e.g. allotropic transformation) are distinguished by:
– small letters (h), high temperature modification (h2 > h1) (r), room temperature modification (1), low temperature modification (l1 > l2) – Greek letters, e.g., ε, ε´ – Roman numerals, e.g., (I) and (II) for different pressure modifications. In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by horizontal lines.
Quasibinary Systems Quasibinary (pseudobinary) sections describe equilibria and can be read in the same way as binary diagrams. The notation used in quasibinary systems is the same as that of vertical sections, which are reported under “Temperature – Composition Sections”.
Invariant Equilibria The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, are described by a constitutional “Reaction Scheme” (Fig. 2). The sequential numbering of invariant equilibria increases with decreasing temperature, one numbering for all binaries together and one for the ternary system. Equilibria notations are used to indicate the reactions by which phases will be
– decomposed (e- and E-type reactions) – formed (p- and P-type reactions) – transformed (U-type reactions) ¨ bergangsreaktion) is used in order to reserve the For transition reactions the letter U (U letter T to denote temperature. The letters d and D indicate degenerate equilibria which do not allow a distinction according to the above classes.
Liquidus, Solidus, Solvus Surfaces The phase equilibria are commonly shown in triangular coordinates which allow a reading of the concentration of the constituents in at.%. In some cases mass% scaling is used for better data readability (see Figs. 3 and 4). Landolt‐Bo¨rnstein New Series IV/11E1
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. Fig. 2 Typical reaction scheme
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In the polythermal projection of the liquidus surface, monovariant liquidus grooves separate phase regions of primary crystallization and, where available, isothermal lines contour the liquidus surface (see Fig. 3).
Isothermal Sections Phase equilibria at constant temperatures are plotted in the form of isothermal sections (see Fig. 4).
Temperature – Composition Sections Non-quasibinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phase fields where generally the tie lines are not in the same plane as the section. The notation employed for the latter (see Fig. 5) is the same as that used for binary and pseudobinary phase diagrams.
. Fig. 3 Hypothetical liquidus surface showing notation employed
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. Fig. 4 Hypothetical isothermal section showing notation employed
Thermodynamics Experimental ternary data are reported in some system reports and reference to thermodynamic modelling is made.
Notes on Materials Properties and Applications Noteworthy physical and chemical materials properties and application areas are briefly reported if they were given in the original constitutional and phase diagram literature.
Miscellaneous In this section noteworthy features are reported which are not described in preceding paragraphs. These include graphical data not covered by the general report format, such as lattice spacing – composition data, p-T-x diagrams, etc. DOI: 10.1007/978-3-540-88053-0_1 ß Springer 2009
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. Fig. 5 Hypothetical vertical section showing notation employed
References The publications which form the bases of the assessments are listed in the following manner: [1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead in Liquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51–56 (1974) (Experimental, Thermodyn., 16) This paper, for example, whose title is given in English, is actually written in Japanese. It was published in 1974 on pages 51- 56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and Metallurgical Institute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16 cross-references. Additional conventions used in citing are: # to indicate the source of accepted phase diagrams * to indicate key papers that significantly contributed to the understanding of the system. Standard reference works given in the list “General References” are cited using their abbreviations and are not included in the reference list of each individual system.
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General References [C.A.] [Curr.Cont.] [E] [G] [H] [L-B]
[Mas] [Mas2] [P] [S] [V-C] [V-C2]
Chemical Abstracts - pathways to published research in the world’s journal and patent literature http://www.cas.org/ Current Contents - bibliographic multidisciplinary current awareness Web resource - http://www.isinet. com/products/cap/ccc/ Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York (1965) Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York (1958) Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P., Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971); Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, Key Elements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, Subvol. a: Ac-Au … Au-Zr (1991); Springer-Verlag, Berlin. Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York (1969) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Index of Alloy Systems Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data
Index of Ternary Refractory Metal Alloy Systems Al-B-C to B-Hf-W
Al-B-C (Aluminium - Boron - Carbon) Al-B-Mo (Aluminium - Boron - Molybdenum) Al-B-Si (Aluminium - Boron - Silicon) Al-C-Ti (Aluminium - Carbon - Titanium) Al-Cr-Ti (Aluminium - Chromium - Titanium) Al-Fe-Nb (Aluminium - Iron - Niobium) Al-Mo-Ni (Aluminium - Molybdenum - Nickel) Al-Mo-U (Aluminium - Molybdenum - Uranium) Al-Nb-Ni (Aluminium - Niobium - Nickel) Al-Nb-Si (Aluminium - Niobium - Silicon) Al-Ni-V (Aluminium - Nickel - Vanadium) Al-O-Zr (Aluminium - Oxygen - Zirconium) Al-Ta-Ti (Aluminium - Tantalum - Titanium) B-C-Cr (Boron - Carbon - Chromium) B-C-Hf (Boron - Carbon - Hafnium) B-C-Mo (Boron - Carbon - Molybdenum) B-C-N (Boron - Carbon - Nitrogen) B-C-Nb (Boron - Carbon - Niobium) B-C-Si (Boron - Carbon - Silicon) B-C-Ta (Boron - Carbon - Tantalum) B-C-V (Boron - Carbon - Vanadium) B-C-W (Boron - Carbon - Tungsten) B-C-Zr (Boron - Carbon - Zirconium) B-Cr-Mn (Boron - Chromium - Manganese) B-Cr-Ti (Boron - Chromium - Titanium) B-Cr-Zr (Boron - Chromium - Zirconium) B-Fe-Mo (Boron - Iron - Molybdenum) B-Hf-Mo (Boron - Hafnium - Molybdenum) B-Hf-W (Boron - Hafnium - Tungsten)
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Aluminium – Boron – Carbon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Andriy Grytsiv, Peter Rogl
Introduction Due to the presence of carbon contaminant in aluminium borides the data on the constitution of the Al-B-C ternary system and those of the binary Al-B system have to be reviewed carefully, which was done by [1994Dus] to strictly differentiate between true aluminium borides and aluminium boron carbides. High hardness in combination with high neutron absorption, high wear resistance and impact resistance have triggered an early interest in high-strength and low-weight Al-B4C composite materials or cermets, either in bulk form with a metal binder or by reinforcing an Al-base matrix with boron carbide particles or with boron carbide-coated fibres. Understanding the phase equilibria proved of major importance in the processing of Al-B4C composites particular in finding processing criteria at temperatures high enough to promote wetting and low enough to control reactions and design microstructures. Despite much effort was spent on the synthesis and crystallographic characterization of the various ternary aluminum boron carbide compounds [1964Mat, 1965Eco, 1965Mat, 1966Gie, 1966Lip, 1969Per, 1969Wil, 1970Nei, 1977Mat, 1987Sar, 1980Ino, 1990Oka, 1992Via, 1994Kud, 1995Osc, 1996Hil1, 1996Hil2, 1997Mey], information on the equilibrium phase relations in the Al-B-C ternary system is scarce [1993Bau, 1997Via, 1998Rog]. These informations comprise early calculations of the phase equilibria disregarding the boron-rich compounds or rather assuming the phases, AlB40C4 and Al2.1B51C8, to be part of the solid solution range of “B4C” [1982Doe, 1993Kau]. Some confusion in the early experimental work on aluminum borides arose from the fact that due to contamination either from high carbon level boron starting material or from the use of graphite crucibles and substrates aluminum boron carbides were produced rather than binary aluminum borides. This is particularly true for “AlB10” [1963Wil] - shown to be “AlB24C4” or more precisely Al2.1B51C8 [1964Mat, 1967Wil, 1969Wil, 1969Per, 1990Oka] - and βAlB12 [1960Koh], later shown to be Al3B48C2 [1965Mat]. According to structural and DTA investigations [1996Hil1], Al3B48C2 exists in a tetragonal high temperature modification, which on cooling below 650˚C transforms into a bodycentered orthorhombic low temperature phase with a unique structure type. The mixture of two orthorhombic phases with coherent boundary and commensurable lattice parameters (modifications A and B), as claimed by [1965Mat, 1986Pes], thus simply explains by multiple twinning on cooling [1996Hil1]. An experimental study of the isothermal section at 1400˚C by [1993Bau] confirmed the existence of four ternary compounds Al2.1B51C8 [1964Mat, 1967Wil, 1969Wil, 1969Per], AlB40C4 [1966Gie, 1966Lip, 1970Nei], Al3B48C2 [1996Hil1] and Al3BC3 [1996Hil2]. The latter compound was first mentioned as “Al4B1–3C4” [1964Mat] and later labelled as “Al8B4C7” [1980Ino] from a cursory investigation of its crystal symmetry with X-ray single crystal Landolt‐Bo¨rnstein New Series IV/11E1
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photographs, although no details of the crystal structure were derived. The relation to a wurtzite structure was discussed [1995Osc]. A structure determination is due to [1996Hil2]. A fifth compound, Al3BC [1992Via, 1993Gon, 1997Mey, 2002Zhe], (earlier “Al4BC” [1987Sar, 1989Hal, 1990Pyz]) was reported to exist below 1000˚C [1992Via], however, was shown in the isothermal section at 1000˚C [1997Via]. From a detailed analysis (XPD, LOM, SEM, EPMA) of the Al-rich corner [1997Via] on about 30 specimens prepared from cold pressed and sintered powder compacts in the temperature region from 627 to 1000˚C, an isothermal section at 1000˚C and a tentative liquidus projection was derived assisted by a series of isothermal diffusion experiments by heating together in an alumina boat an Al-B rod and an Al-C rod. Phase equilibria at 900˚C in the Al-C rich part of the ternary Al-B-C system were established [2002Zhe] from XPD of about 45 ternary and binary alloys. Equilibrium conditions were not reached for boron-rich samples. An attempt to obtain equilibrated samples from mixtures B4C+AlB2, B4C+Al and B4C+Al4C3 were also unsuccessful. Al3BC and Al4C3 phases form very easily and are observed in all samples even after short time sintering in contrast to Al3BC3, which forms very slowly at 900˚C. On the other hand Al3BC3 was always observed in arc melted samples containing 40–60 at.% Al and 10–30 at.% B. Experimental techniques for preparation concerned (a) melting of B4C in excess of Al for the synthesis of AlB40C4 (at 1550˚C, [1970Nei]), (b) melting of boron with excess of Al in a graphite crucible for synthesis of Al3B48C2 (at 1400˚C, [1964Mat]) (c) vapor deposition at 1400 to 1600˚C for single crystals of Al3B48C2 [1967Bli] (d) hot pressing of B4C+Al in graphite dies for synthesis of Al2.1B51C8 (at 1800˚C [1966Gie, 1966Lip]), (e) infiltration of B4C by liquid Al at 1100˚C and anneal at 1000˚C to obtain Al3BC [1987Sar], (f) reaction sintering of Al+B+C powder compacts on alumina boats in sealed silica capsules 627 to 1000˚C for the synthesis of Al3BC and phase relations at 1000˚C [1992Via, 1997Via] or at 1400˚C for 10 h for the production of the single crystals of Al3B48C2 [1994Kud] (g) melting of an Al8BC mixture in alumina under argon for 160 h at 850˚C and subsequent cooling at 150K/h to RT for production of black-bluish single crystals of Al3BC [1997Mey] (h) melting of an Al40B2C3 mixture in alumina under argon at 1500˚C and subsequent cooling at 10K/h to 600˚C for production of single crystals of Al3BC3 in the form of yellow, transparent platelets [1996Hil2] and (i) Al-flux solvent method for a general production of single crystals (see i.e. [1986Kis, 1990Oka, 1996Hil1]). Samples used for the isothermal section at 1400˚C were prepared from cold compacted powder mixtures of AlB2, B4C, B and/or C, which were reaction-sintered under Ar in closed Knudsen-type graphite reactors at 1600˚C for 1h prior to 48 h heat treatment at 1400˚C [1993Bau]. Phase relations at 900˚C were studied [2002Zhe] on elemental powder compacts sintered in alumina crucibles (binary Al-B alloys) or in closed graphite crucibles (ternary alloys). The specimens were sealed in evacuated quartz ampoules and were slowly heated for 10˚/h to 720˚C (slightly above the melting point of aluminium) and kept at this temperature for 48h. After temperature was increased to 900˚C at a rate of 20˚/h, the tablets were sintered at this temperature for 1 week. Repeated repowderisation (under protective cyclohexane) and sintering at 900˚C were necessary to reach equilibrium conditions. Several studies dealt with the kinetics of wetting of B4C surfaces by liquid aluminum; detailed discussions can be found in the articles by [1979Pan] and [1989Hal]. Hot-pressing of B4.3C+Al powders at 1820˚C, 45 MPa under Ar (5 to 20 mass% Al) revealed the formation of the ternary B4C- related Al-boron carbides (solution of Al in B4C, and τ2) although the products were all thought to belong to the B4C-based solid solution [2000Liu]. With increasing Al-content (>5 mass% Al) the Al3BC3 phase evolved [2000Liu]. DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009
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Thermodynamic calculations of the Al-B-C system have been attempted by [1982Doe, 1993Wen, 1993Kau], however, are not fully consistent with experimental observations. Reviews on the constitution and on the crystal structures of the Al-B-C system have been presented by [1977Mat, 1990Luk, 1998Rog].
Binary Systems The binary systems B-C and Al-C are consistent with the critical assessments of [1998Kas, 1996Kas] and [2003Per], respectively. In spite of numerous data available from literature on the constitution of the Al-B phase diagram, contradictory results exist for the formation of aluminium diboride (Table 1). Figure 1a shows the various versions for the Al-rich part of the Al-B phase diagram. It should be noted, that recent experiments [2002Zhe] confirmed the formation of AlB2 at 900˚C, in contrast to data of [1997Via] suggesting peritectic formation at 892 ± 5˚C. In the present assessment we accept the temperature of 956 ± 5˚C for the invariant reaction L+AlB12ÐAlB2 as determined by [2000Hal]. The adopted Al-B phase diagram (Figs. 1b, 1c) is based on the assessment of [1994Dus]. The composition of the peritectic liquid at 0.55 at.% B has been confirmed by a recent thermodynamic assessment of [2001Fje]. AlB2 is still taken as a stoichiometric compound in spite of the suggestions of [1964Mat, 1999Bur, 2002Bur] for Al-deficiency in terms of Al0.9B2. Although the assessment of [1994Dus] concluded a peritectic formation of AlB12, L+(B)Ð AlB12 at 2050˚C, the thermodynamic calculation of [1993Wen] is based on congruently melting AlB12 (TM = 2150˚C).
Solid Phases The crystallographic information on all the binary and ternary phases pertinent to the Al-B-C system is listed in Table 2. Some controversy exists in the crystallographic characterization of the modifications reported for Al3B48C2. A single crystal study [1995Hil, 1996Hil1, 2000Mey] on an untwinned specimen revealed a tetragonal high temperature form (closely related to the structure of I-tetragonal boron), which on cooling undergoes a symmetry reduction resulting in microscopically twinned products that hitherto were indexed on the basis of two orthorhombic modifications, labeled A and B by [1965Mat]. The transformation was earlier proposed to be at ca. 850˚C [1960Koh, 1965Mat], whereas new results from DTA recorded 650˚C [1996Hil1]. The transition seems to be rather fast, as the low temperature modification is present in samples furnace-cooled from 1400˚C to room temperature [1993Bau]. A second point of controversy concerns the phases AlB40C4 and Al2.1B51C8 for which detailed crystallographic descriptions are available, however, AlB40C4 actually being isotypic with binary B4C, hitherto is not thoroughly established as a ternary phase independent from binary B4C. As the two structurally closely related phases AlB40C4 and Al2.1B51C8 generally are found together, a high and low temperature transition between them may be inferred [1993Bau]. Without further details the maximum solid solubility of Al in boroncarbide (“B13C2”, at 1950˚C) was reported to be 1 mass% Al (equivalent to 2 at.% Al in B4C) [1978Ekb]. Experiments to establish a possible homogeneous range for Al3BC3 (earlier: “Al8B4C7” [1980Ino], or “Al4B1–3C4” [1964Mat]), were carried out at 1830˚C by [1980Ino] resulting in a Landolt‐Bo¨rnstein New Series IV/11E1
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rather stoichiometric composition without variation of the lattice parameters. These findings were confirmed by [1993Bau, 1996Hil2]. Details of the crystal structure with linear C-B-C chains are given by [1996Hil2]. Lattice parameters of Al3BC3 were measured at room temperature up to 7.5 GPa using a multi-anvil synchrotron system with B4C anvils; for a high temperature pressure experiment the sample was placed in a graphite ampoule [2000Sol]. Al3BC3 is free of structural transitions up to 1523˚C within the pressure range 2.5 to 5.3 GPa [2000Sol]. A further ternary compound τ5 was observed after infiltration by liquid Al at 1170˚C with post heat treatment for 100 h at 800 to 1000˚C [1987Sar]. The hexagonal lattice was established by TEM; the approximate composition “Al4BC” resulted from EELS-data [1987Sar]. This phase has been also confirmed by [1989Hal, 1990Pyz]. From a detailed investigation of this Al-rich boroncarbide by X-ray powder diffraction, LOM and EPMA, [1992Via] suggested a formula of Al3BC rather than “Al4BC” and attributed a hexagonal cell; additional weak lines in the X-ray intensity pattern of Al3BC prompted a larger unit cell pffiffiffi a = a0/ 3 [1993Gon]. Although the authors of [1997Mey] recognized the larger cell, the crystal structure of Al3BC was solved for the high symmetry subcell from single crystals isolated from a sample directly reacted from the elements - however, from EPMA a composition Al2.5BC was derived (see also Table 2). Al6B-octahedra and trigonal Al5C-bipyramids are the characteristic structural elements [1997Mey]. The various data on the compositional ranges of the τ4 and τ5 phases are summarized in Fig. 2. Half filled circles correspond to the accepted stoichiometries Al3BC and Al3BC3. From the significant change of the unit cell volume of Al4C3 comparing binary and ternary alloys, a solubility of boron is suggested [1996Bid, 2002Zhe, 2000Mey]. Solubility of boron in Al4C3 was established to be 3.4 at.% at 900˚C [2002Zhe] and an interesting behavior of lattice parameters was observed. In spite of the increase of the “a” parameter and of the cell volume with boron content, the “c” parameter decreases. That may be explained by a preferential distribution of boron and carbon atoms among different crystallographic sites. A significant solubility of boron in Al4C3 was also reported by [2000Mey] to be about 9.3 at.%, however, no details on the relevant temperature were given. Furthermore these authors claim for Al4C3–xBx lattice parameters increasing with boron content. Lattice parameters of Al4C3 for samples located in three-phase regions (Al)+Al4C3+Al3BC, Al4C3+Al3BC3+Al3BC and Al4C3+Al3BC3+(C) are very close, assuming that these three-phase regions meet at the Al4C3 phase at a maximal boron solubility of Al4(C0.92B0.08)3. Insignificant solubility of carbon in AlB2 is reported by [2002Zhe] comparing lattice parameters in ternary and binary samples; AlB2 with 0.5 at.% C, heat treated at 900˚C, already contains the Al3BC phase.
Isothermal Sections Phase equilibria for the 1400˚C isothermal section are summarized in Fig. 3, revealing the existence of four ternary compounds τ1 to τ4. A small field of liquid phase exists at 1400˚C which is in equilibrium with Al3B48C2, Al2.1B51C8 and with Al3BC3 [1993Bau]. Boron-poor equilibria agree with an earlier work by [1980Ino] who reported on the two-phase equilibria Al4C3+Al3BC3 (Al7B4C8), Al3BC3 (Al7B4C8)+B4C and Al3BC3 (Al7B4C8)+C. In Fig. 3 twophase equilibria are shown to exist between the binary solid solution “B4C” and AlB40C and Al2.1B51C8. At 1400˚C all ternary compounds seem to exist at their stoichiometric compositions [1993Bau], whilst [1965Eco] claimed a homogeneity range for τ1 at 1800˚C from DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009
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“AlB48C8” to Al3B48C8. Binary AlB12 was never seen in combination with Al2.1B51C8 nor with AlB40C4 [1993Bau]. The isothermal section at 1000˚C, Fig. 4, was constructed on the basis of data from [1997Via]. Due to low interaction kinetics in the boron- and carbon-rich part of the system at 1000˚C, equilibria in this portions of the diagram are preliminary. Moreover, ternary compounds τ1 and τ2 were not included in the 1000˚C section by [1997Via], τ4 was listed as “Al8B4C7”, and no solubility of boron in Al4C3 was considered. For consistency with the present knowledge on the Al-B-C system, the ternary compounds τ1 and τ2 were introduced in Fig. 4 and the composition of τ4 was changed to Al3BC3. The solubility of boron in Al4C3 at 1000˚C was estimated to be about 4 at.%, extrapolating from data of [2002Zhe] at 900˚C. Figure 5 represents the isothermal section at 900˚C [2002Zhe] confirming the equilibrium AlB2+Al3BC, whereas [1997Via] claimed this equilibrium to be only stable below 868 ± 4˚C. Similar to 1000˚C the equilibria at 900˚C involving τ1 and τ2 are not well established due to low reactivity of the components.
Invariant Equilibria, Liquidus Surface A tentative liquidus surface for the aluminum rich portion of the diagram (Fig. 6) was proposed by [1997Via], presenting equilibria involving the τ5 phase. The invariant equilibrium U5 (L+Al3B48C2ÐAlB2+Al3BC) was reported at 868 ± 4˚C by [1997Via], but this temperature can not be accepted in respect to the observed isothermal equilibrium AlB2+Al3BC at 900˚C [2002Zhe] suggesting such transformation above 900˚C. Comparison of the reaction scheme and the isothermal section at 1000˚C (Fig. 4) with the isothermal section at 1400˚C (Fig. 3) suggests a rather complicate picture of the phase transformations in this regions mainly due to decomposition of τ5. Based on an earlier thermodynamic calculation by [1982Doe], a reaction scheme was derived [1990Luk], which gives a tentative information of the solidification behavior in the AlB-C ternary. The temperatures of the invariant equilibria were estimated and the ternary compounds τ1 to τ3 were assumed to be part of the solid solution arising from binary B4C; τ5 was not included. A more recent thermodynamic modelling of the Al-B-C phase diagram by [1993Wen] as part of the multi-component Al-B-C-N-Si-Ti system treated the ternary compounds τ1, τ2, τ3 as independent phases, however, the peritectoid formation of τ4 (Al3BC3) is in strict contradiction to the experimentally confirmed two-phase equilibrium τ4+τ5 (Al3BC3+Al3BC) [1997Via, 2002Zhe] as well as to the observed existence of τ4+τ5 in as cast alloys [2002Zhe], thereby strongly indicating direct formation of Al3BC3 from the liquid. A closed ternary miscibility gap in the Al-rich liquid is suggested from thermodynamic calculations by [1993Kau], however, hitherto without experimental confirmation [2002Zhe]. Figure 7 presents a reaction scheme for the major parts of the Al-B-C phase diagram. The reaction scheme is essentially based (i) on the tentative liquidus projection for the Al-rich part as suggested by [1997Via], (ii) on the experiments of [2002Zhe] concerning the solidification of the phases τ4, τ5 and (iii) on the thermodynamic calculation of [1993Wen] for the B-rich part, however, accepting peritectic formation of AlB12.
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Thermodynamics Enthalpies of formation and heat capacity measurements from a Calvet type automatic microcalorimeter in the temperature range 310–1200 K were reported by [1987Kis] and are listed as follows: Al3 B48 C2 : H 0 ðTÞ H 0 ð298Þ ¼ 0:7945 103 T 2 þ 0:5182T 225:1374 ðin J g1 Þ and Cp ðTÞ ¼ 0:1589 102 T þ 0:5182 ðJ g1 K1 Þ Al2:1 B51 C8 : H 0 ðTÞ H 0 ð298Þ ¼ 0:7226 103 T 2 þ 0:5411T 225:5702 1
ðin J g
:
for AlB24 C4 Þ and
Cp ðTÞ ¼ 0:1589 102 T þ 0:5182 ðJ g1 K1 for AlB24 C4 Þ Thermodynamic calculations of the Al-B-C system due to [1982Doe, 1993Wen, 1993Kau], however, are not fully consistent with experimental observations. For detailed discussion, see section Invariant Equilibria.
Notes on Materials Properties and Applications Mechanical properties of Al-B4C cermets and boron/carbon fiber-aluminium composites have been investigated by various groups [1972Bak, 1973Her, 1975Mun, 1984Via, 1985Che, 1985Hal, 1985Kov, 1985Pyz, 1985Sar, 1986Che, 1986Dub,1990Ram, 1996Pyz, 2002Ars]; the effect of reaction on the tensile behavior of infiltrated composites was reported by [2002Kou2] and size dependent strengthening in particle reinforced Al by [2002Kou1]; reaction products were studied by [2001Lee]. An increase of surface hardness of about 25 to 40 % can be achieved by impulse laser radiation on B4C/Al cermets [1988Kov]. Wetting of B4C by Al has been studied by many research teams with rather contradicting results, until the temperature and time dependent occurrence of chemical reactions/compound formation was analyzed in detail (for discussion see i.e. [1979Kis, 1979Pan, 1989Hal, 2000Kha]). The kinetics of wetting by liquid aluminium of flat, sintered boron carbide specimens with residual porosity less than 3 % were investigated by [1979Pan]. The speed of spreading of liquid aluminium at 1100˚ to 1200˚C was measured to be 0.1–0.8 mm·s–1, in accordance with r2 = f(t), where r equals the radius of the contact circle. The angle of contact was first 92˚, however, in 3 to 5 min decreased to 28˚. The slow spreading was determined by the formation of new aluminum boron carbide phases in the contact zone with a microhardness of ca. 13 GPa. The driving force Δσ = σ (cos Θ0-cosΘ) (σ = surface tension of the melt, Θ0 = contact angle of the melt, Θ = contact angle at time (t)) decreased sharply becoming zero in 4 to 5 min [1979Pan]. The contact angle of molten Al on B4C as a function of processing time for various isotherms at 5·10–3 to 10–4 Pa was also given by [1989Hal] based on sessile drops cooled to room temperature. Mechanical properties, electrical and thermal conductivity as well as their temperature dependencies were reported on the Knoop and Vickers microhardness for Al-flux grown (temperature region 1750 to 800˚C) ‘‘amber’’ single crystals Al3B48C2 and for ‘‘black’’ crystals (αAlB12, γAlB12 and AlB2.1B51C8) [1986Kis]. These studies were also performed on hot pressed specimens of various compositions x(AlB12)+(1–x)B4C and Al3B48C3 in the temperature range 24 to 827˚C [1991Kha1, 1991Kha3]. For Al3BC3 (“Al8B4C7”), Al3B48C3, Al2.1B51C8 [1991Kha2] DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009
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also examined these properties as a function of porosity and quantity of Fe-impurity. These data are summarized in Table 3 including information on flux-grown crystals Al3B48C2 and Al2.1B51C8 [1990Oka]. Both types of crystals were said to be p-type semiconductors [1986Kis]. In a ring test the strength of a powder compact of B13C2 +1 mass% Al, sintered at 1950˚C, was found to be 0.50(7) GNm–2 [1978Ekb]. [1991Kha2] reported on the kinetics of thermal densification of hot pressed powders of B4C, AlB12, Al3B48C2 and Al3BC3. Kinetics of dissolution in HCl, HNO3 and HCl-HNO3 was studied by [1998Kha] as well as the resistance of Al-boron carbides to alkali and hydrogen peroxide. [1989Hal] studied the densification kinetics of Al+B4C cermets in the range from 800 to 1400˚C in pressureless sintering as well as after applying hot isostatic pressure. The kinetic of metal depletion in post heated dense cermets B4C/Al at temperatures between 600˚C and 1000˚C was investigated by [1990Pyz]. Chemical stability against various boiling acids, oxidation resistance, IR and EPR spectra of Al-borides and Al-boron carbides (Al3B48C2, Al2.1B51C8) was studied by [1991Pri]. The spectra were taken at 77 K and 300K and for different crystal orientations relative to the magnetic field. Absorption edge and IR-active phonons in Al3B48C2 were reported by [1987Hau, 2000Wer] and IR spectra of boron carbide containing up to 1.5 at.% Al were determined between of 8 to 500 mm–1 wave numbers and for temperatures between 70 to 450 K [1997Sch]. These data seem to suggest the incorporation of Al-atoms into binary boroncarbide in form of pairs substituting the B-B-C or C-B-C chains [1997Sch]. Characteristic IR absorption bands for finely dispersed powders of Al-borides and Al-boron carbides were listed by [1998Kha]. The Seebeck-coefficients were reported to linearly increase from 260 μVK–1 for binary “B4C” to 450 μVK–1 for 1.4 at.% Al dissolved, revealing p type behavior [1997Sch]. Seebeckcoefficients, thermal and electric conductivities were further reported by [2000Liu] for B4.3Cbased samples containing 0.5, 10, 15, 20 mass% Al, highlighting the Z-value at RT of 1.04·10–6K–1 for the 5 mass% Al sample. IR and Raman spectroscopy on Al3BC3 (at RT) confirm the linear (CBC)5– unit as an isoelectronic CO2-analogon [1996Hil2, 2000Mey]. On heating in air, Al3BC3 (earlier reported as Al8B4C7), Al3B48C3 and Al2.1B51C8 show low oxidation at 500˚C (increase of mass 4 mgh–1); intensive oxidation, with a mass increase of 40 mgh–1) starts at 1280˚C for Al2.1B51C8 and at 1370˚C for Al3B48C2 [1991Pri, 1989Kha, 1991Kha4]. Oxidation in air of single crystals Al2.1B51C8 and Al3B48C2 started at about 760˚C and 710˚C, respectively [1990Oka]. The reaction products were 9Al2O3·2B2O3 for Al2.1B51C8 crystals and B2O3 for AlB40C4 specimens [1994Kud]. Whereas Al3BC3 was said to be unstable in acids [1991Kha4], more detailed experiments [1996Hil2] proved stability at room temperature against bases and dilute acids, except for HNO3 and HF. Al3BC3 was furthermore said to be stable in air up to 600˚C [1991Kha4,1996Hil2]. Al3BC is quickly attacked by dilute HCl [1997Mey]. Thermophysical properties of sintered bodies of Al3BC3 have been derived by [2000Wan]. These are linear thermal expansion in the range of 25 to 1200˚C, specific heat and thermal diffusivity via laser flash technique, Youngs modulus of 136.6 GPa, Vickers hardness of 12.1 GPa at a load of 196 N and thermogravimetric recording of growth of an oxidized layer on heating in air up to 1500˚C. Fitting a Birch-Murnaghan equation of state to the pressure dependency of the lattice parameters of Al3BC3 up to 7.5 GPa, the isothermal bulk modulus was B0 = 153 ± 6 GPa (dB0/ dp = 19 ± 4) [2000Sol]. Despite high bulk modulus the Vickers hardness of single crystals is as low as 20.7 GPa at a load of 25g and 18.2 GPa at a load of 50g [2000Sol]. Landolt‐Bo¨rnstein New Series IV/11E1
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Al3BC was successfully prepared by self propagating high temperature synthesis induced by mechanical activation of Al-B-C powder mixtures in air; mixtures low in boron (AlB0.1C) resulted in Al3BC3 under violent emission of heat [1999Tsu]. In contrast to that [2000Sav] was unable to prepare ternary aluminoborocarbides from mechanochemical synthesis. Elastic bulk and shear moduli for Al3BC (earlier reported as Al4BC) were measured by [1995Pyz] and estimated by [1999Tor].
Miscellaneous A series of patents covers the techniques to produce dense B4C/Al cermets by infiltration of the metal matrix into the porous ceramic body without wetting reactions [1976Lan, 1986Hal, 1987Pyz, 1990Pyz, 1991Pyz, 1995Pyz, 1996Pyz, 1997Du , 2000Pyz, 2001Lee]; subsequent heat treatment results in materials with designed chemistry and microstructures, flexure strength, hardness and fracture toughness. Fine microstructures were obtained via ultrarapid microwave heating [1995Rug]. B4C/Al cermets have been considered as an improved structural neutron absorber [1977Ros, 1978Boi, 1978Sur, 1986Ros, 1987Lev, 1992Bei] and for applications as friction materials for automotive brake applications [1999Cha]. Oxidation protective B4C-coatings on C-fibers in Al-matrix were reported by [1996RMi] and [1996Vin] produced C-fibres-Al composites by a squeeze casting technique. Explosive consolidation to produce Al/ B4C composites was studied by [1995Bon, 1997Yue]. Shock recovery experiments were performed on a 65 vol% B4C-Al cermet as a function of shock pressure [1989Blu].
. Table 1 Literature Data on Experimental Temperatures of Invariant Equilibrium L+AlB12ÐAlB2 T [˚C]
Technique
Heating Rate
References
Stability observation
-
1000 - 1500
[1936Hof]
Synthesis observation
-
980
[1967Ser]
DTA
4˚C/min
920
[1967Ato]
Stability observation
-
1350 - 1500
[1972Sir]
DTA
5˚C/min
1030 ± 5
[1993Ips]
Synthesis observation
-
892 ± 5
[1997Via]
DSC and Stability observation
0˚C/min*
956 ± 5
[2000Hal]
DSC
10˚C/min
914 ± 55
[2001Fje]
* DSC measurements were performed with heating rate of 5, 15 and 40˚/min., and extrapolated to 0˚C/min.
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. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
[Mas2 ]
(βB) < 2092
hR333 R3m βB
a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1096.5 c = 2386.8 a = 1097.4 c = 2387.7
[Mas2 , 1993Wer]
a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78
[Mas2 ]
(C) < 3827 (B.P.)
B4C < 2450
hP4 P63/mmc C-graphite
hR45 R3m B13C2
at 1.1 at.% C [1993Wer] linear da/dx, dc/dx at AlB31 [V-C2 ] from sample Al4B95C1, quenched from 1400˚C, contains Al3B48C2 and αAlB12 [1993Bau]
[1967Low] at 2.35 at.% Bmax (2350˚C) linear da/dx, dc/dx, [1967Low]
a = 565.1 c = 1219.6 a = 560.7 c = 1209.5 a = 560.3 c = 1209.8
from sample containing τ2, τ4, quenched from 1400˚C [1993Bau]
9 to 20 at.% C [1990Ase]
B25C
tP52 P42m B25C
a = 872.2 c = 508.0
[V-C2 ] also B51C1, B49C3; all metastable?
Al2B3 ≤ 525
hR* Al2B3 (?)
a = 1840 c = 896
at 60 at.% B [1992Var] metastable?
AlB2 ≤ 956±5
hP3 P6/mmm AlB2
a = 300.6 b = 325.2 a = 300.67 ± 0.01 b = 325.36 ± 0.02 a = 300.63 ± 0.01 b = 325.46 ± 0.01 a = 300.43 ± 0.03 b = 325.19 ± 0.06
[1994Dus], temperature from [2000Hal] [2002Zhe]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] αAlB12 ≤ 2050
γAlB12
Pearson Symbol/ Space Group/ Prototype tP216 P41212 αAlB12
oP384 P212121 γAlB12
Lattice Parameters [pm] a = 1015.8 c = 1427.0 a = 1018 c = 1434.3 a = 1016.3 c = 1425.6 a = 1015.5 c = 1426.0 a = 1014.93 ± 0.07 c = 1425.0 ± 0.5 a = 1014.4 b = 1657.3 c = 1751.0
hR21 R 3m Al4C3
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[1994Dus] ρexp. = 2.65 Mgm–3 [1991Pri] from sample Al2B92C2, quenched from 1400˚C, contains Al3B48C2 [1993Bau] from sample Al4B95C1, quenched from 1400˚C, contains Al3B48C2 and AlB31 [1993Bau] [2002Zhe] [1983Hig, 1994Dus, 2000Hig] metastable phase or ternary product stabilized by small amounts of impurity metals present in Alflux grown material ρexp. = 2.56 Mgm–3 [1991Pri]
a = 1019.5 b = 1666 c = 1769 Al4C3 < 2156
Comments/References
a = 333.8 c = 2511.7 a = 334.21 ± 0.01 c = 2503.2 ± 0.5 a = 335.78 ± 0.02 c = 2499.6 ± 0.5
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[2003Per, V-C2 ] [2002Zhe] [2002Zhe] in 57.1Al-4.3B-38.6C Al4(C0.9B0.1)3, in equilibrium with τ5 at 900˚C
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. Table 2 (continued)
Phase/ Temperature Range [˚C] * τ1, Al2.1B51C8 (eventually low temperature phase of τ2)
* τ2, AlB40C4 (eventually high temperature phase of τ1)
Pearson Symbol/ Space Group/ Prototype oC88 Cmcm Al2B51C8
hR45 R3m B4C-deriv.
Lattice Parameters [pm]
earlier labeled “AlB10” [1967Wil] or AlB24C4 [1964Mat, 1969Wil, 1970Wil] [1969Per] ρexp. = 2.54 Mgm–3
a = 569.0 b = 888.1 c = 910.0 a = 568.7 b = 887.7 c = 909.8 a = 569.0 b = 888.1 c = 910.0 a = 569.3 b = 884.7 c = 909.3 a = 567.6 b = 891.4 c = 909.5 a = 569.2 b = 889.2 c = 911.2
from sample containing τ2 and τ3, quenched from 1400˚C [1993Bau] from sample containing τ4, quenched from 1400˚C [1993Bau] from sample Al4B92C4 quenched from 1400˚C, contains Al3B48C2 (tetragonal), Al3B48C2 (A) and AlB40C4 [1993Bau] [1991Pri]
[1990Oka] ρexp. = 2.54 Mgm–3 single crystals from Al-flux
a = 564.2 c = 1236.7 a = 565.37 c = 1231.4 a = 564.8 c = 1239.9 a = 565.6 c = 1238.9
[1970Nei] ρexp. = 2.52 Mgm–3 from sample containing τ1, τ3, quenched from 1400˚C [1993Bau] from sample containing τ4 and B4C, quenched from 1400˚C [1993Bau] from sample Al4B92C4 quenched from 1400˚C, contains also Al3B48C2 (tetrag.), Al3B48C2 and Al2.1B51C8 [1993Bau] [1966Gie] for composition “Al2B48C8”
a = 563 c = 1129 a = 565 c = 1239
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Comments/References
[1966Lip] for composition “Al4B48C8”
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Al–B–C
. Table 2 (continued)
Phase/ Temperature Range [˚C] * τ3, Al3B48C2 (r) < 650
Pearson Symbol/ Space Group/ Prototype oI212 Imma Al3B48C2
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Lattice Parameters [pm] a0 = 1240.7 b0 = 1262.3 c0 = 1014.4 a = 1234 b = 1263 c = 508 a = 1232.5 b = 1261.4 c = 1016.2 a = 1233.72 b = 1262.41 c = 1016.06 a = 1232.5 b = 1264.7 c = 1016.2 a = 1230.2 b = 1262.1 c = 1016.1 a = 1229.1 b = 1262.2 c = 1015.88 a = 1233.62 b = 1262.40 c = 1015.94 a = 1239.0 ± 0.3 b = 1263.7 ± 0.3 c = 1013.6 ± 0.4 a = 1237.7 b = 1262.7 c = 507.9 to a = 1236.3 b = 1261.6 c = 510.2 a = 616.6 b = 1263.5 c = 1065.6 a = 618.1 b = 1262.2 c = 1016.1 a = 617 b = 1263 c = 1016
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Comments/References [1996Hil1], only one low temperature modification! [1965Mat], two modifications, microscopically twinned; modification A, c = c0/2 from a sample Al6B92C2 cooled from 1400˚C contains “AlB12” [1993Bau] from sample Al4B95C1 cooled from 1400˚C contains also “AlB12”, AlB31 [1993Bau] [1991Pri]
[1994Kud]
from sample Al4B92C4 cooled from 1400˚C, [1993Bau] contains Al2.1B51C8, AlB40C4 and tetragonal Al3B48C2 [1993Bau] from sample Al4B95C1, cooled from 1400˚C, see above. [2000Wer]
[1990Oka] single crystals from Al-flux modification A ; c = c0/2
[1990Oka] single crystals from Al-flux ρexp. = 2.59(2) Mgm–3 modification B, a = a0/2
[1965Mat] modification B a = a0/2
Landolt‐Bo¨rnstein New Series IV/11E1
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm] a = 616.4 b = 1262.1 c = 1016.4
* τ3, Al3B48C2 (h) > 650
* τ4, Al3BC3 < 1835
*τ5, Al3BC < 1100
[1991Pri] a = a0/2
tP52 a = 885 P42/nnm c = 508 B25C -deriv. a = 882 c = 509 a = 881.9 c = 508.25 hP42 P3c1a) Mg3BN3
hP20 P3c1 (P63/mmc for subcell) Al3BC
Comments/References
[1996Hil1] high temperature modification [1965Mat] from sample Al4B92C4 cooled from 1400˚C, contains also Al2.1B51C8, AlB40C4 and orthorhombic Al3B48C2 [1993Bau]
a = 589.97 c = 1589.0 a = 590.6 c = 1590.1 a = 590.7 c = 1591.3 a = 590.5 c = 1590.5 a = 340.1 ± 0.3 c = 1584 ± 0.2 a = 590.22 ± 0.3 c = 1589.4 ± 0.1 a = 605.0 c = 1154.0 a = 603.45
[1996Hil2] ρ = 2.66 Mgm–3 temperature from [1980Ino] [1980Ino], labelled as Al8B4C7 from sample containing τ1, quenched from 1400˚C [1993Bau] from sample containing τ2 and B4C, quenched from 1400˚C [1993Bau] pffiffiffi [2000Sol], subcell with a = a0/ 3 pressure dependence of the lattice parameters is given up to 7.5GPa [2002Zhe] [1993Gon, 1997Via]
c = 1152.02 a = 6041.9 ± 0.2 c = 1154.0 ± 0.3 a = 349.1 c = 1154.1 a = 352.0 c = 582.0
[1997Mey] from single crystals, “Al2.5BC” from EPMA [2002Zhe] pffiffiffi [1992Via], subcell with a = a0/ 3 [1987Sar], earlier “Al4BC” subcell with a = a0/ p ffiffiffi 3, c = c0/2
pffiffiffi P63/mmc for subcell with a = a0/ 3, c = c0
a)
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(101)
(001)
γAlB12
(100)
(111)
(100)
(111)
αAlB12
Al3BC3
Al2.1B51C8
Al3B48C2
29.6(1.0) 25.8(7) 22.8(8) 21.6(1.1)
19.6(5)
22.6 26 6 24.2(7) 25.0–26.9
37.6(2.0) 31.7(8) 23.7(6)
27.1(5)
26.5(5) 23.1
(0.5N, 293K) (1N, 293K) (2N, 293K) (4.9N, 293K)
(2N, 293K)
(2N, 293K) (5N, 293K) (5N, 1200K) (2N, 293K) (1N, 293K)
(0.5N, 293K) (1N, 293K) (4.9N, 293K)
(2N, 293K)
(2N 293K) (5N, 293K)
34.4(2.7) 31.0(1.5) 27.3(1.2) 23.8(9)
20.7 18.2
25.7 to 30.5
(0.5N, 293K) (1N, 293K) (2N, 293K) (4.9N, 293K)
(0.25N, 293K) (0.50N, 293K)
(1N, 293K)
33.6(1.6) (2N, 293K)
1.8(2)
1.5(3)
2.7(2)
5.3 4(1)
Fracture Toughness K1c [MPa·m1/2]
0.1
1 0.6 - 1.2
0.18 - 0.36
3.85 · 105 0.22
5.92 ·102
2.02 · 105 0.08 - 0.18
10–3 - 1
104-106 2.6·10310–6 2.6·10–6
ρ293K [Ωm]
Activation Energy ΔE [eV] 100K to 400K R=R0exp(-ΔE/2kT)
38.7 (310K) 60 (600K)
19.6 (310K)
Thermal Conductivity [Wm–1K–1]
[1986Dub] [1986Dub] [1991Pri] [1986Dub]
[1991Pri]
[2000Sol]
[1986Kis] [1986Kis] [1986Kis] [1991Pri] [1990Oka]
[1990Oka, 1994Kud] [1986Dub] [1986Dub] [1986Dub]
[1986Kis] [1986Kis] [1986Kis] [1991Pri]
References
3
Microhardness [GPa] at Various Loads and Temperatures Crystal Face Vickers Compound of Indent Knoop
. Table 3 Microhardness, Fracture Toughness, Electrical Conductivity, Activation Energies for Electrical Conductivity and Thermal Conductivity for Various Aluminium Borides and Aluminium Boron Carbides
14 Al–B–C
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. Fig. 1a Al-B-C. Various versions of the Al-rich part of the Al-B diagram
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Al–B–C
. Fig. 1b Al-B-C. Accepted Al-B phase diagram
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. Fig. 1c Al-B-C. Accepted Al-B phase diagram, enlarged Al-rich region
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Al–B–C
. Fig. 2 Al-B-C. Superposition of literature data on the homogeneity regions of τ4 and τ5 phases. Halffilled circles correspond to the accepted compositions
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. Fig. 3 Al-B-C. Isothermal section at 1400˚C; the position of Al3BC is indicated by a full circle
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Al–B–C
. Fig. 4 Al-B-C. Isothermal section at 1000˚C
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. Fig. 5 Al-B-C. Isothermal section at 900˚C
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Al–B–C
. Fig. 6 Al-B-C. Tentative liquidus surface projection
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. Fig. 7 Al-B-C. Reaction scheme
Al–B–C
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Al–B–C
References [1936Hof] [1960Koh] [1963Wil] [1964Mat] [1965Eco] [1965Mat] [1966Gie] [1966Lip] [1967Ato] [1967Bli] [1967Low] [1967Ser] [1967Wil] [1969Per] [1969Wil] [1970Nei] [1970Wil] [1972Bak] [1972Sir] [1973Her]
[1975Mun]
[1976Lan]
[1976Mon] [1977Mat] [1977Ros]
Hofmann, W., Jaeniche, W., “Contribution to the Knowledge of the Aluminium-Boron System” (in German), Z. Metallkd., 1, 1–5 (1936) (Phase Diagram, Crys. Structure, 13) Kohn, J.A., Eckart, D.W., “Aluminium Boride, AlB12”, Anal. Chem., 32, 296–298 (1960) (Crys. Structure, Experimental, 6) Will, G., “On the Crystal Structure of AlB10”, J. Am. Chem. Soc., 85, 2335–2336 (1963) (Crys. Structure, Experimental; 7) Matkovich, V.I., Economy, J., Giese Jr. R.F., “Presence of Carbon in Aluminium Borides”, J. Am. Chem. Soc., 86, 2337–2340 (1964) (Crys. Structure, Experimental, 14) Economy, J., Matkovich, V.I., Giese, Jr.R.F., “Crystal Chemistry of α-Boron Derivatives”, Z. Kristallogr., 122, 248–258 (1965) (Review, Crys. Structure, 26) Matkovich, V.I., Giese, Jr.R.F., Economy, J., “Phases and Twinning in C2Al3B48”, Z. Kristallogr., 122, 108–155 (1965) (Crys. Structure, Experimental, 7) Giese, Jr.R.F., Economy, J., Matkovich, V.I., “Topotactic Transition in C4AlB24”, Acta Crystallogr., 20, 697–698 (1966) (Crys. Structure, Experimental, 7) Lipp, A., Ro¨der, M., “On an Aluminium Bearing Boron Carbide” (in German), Z. Anorg. All. Chem., 343, 1–5 (1966) (Crys. Structure, Experimental,13) Atoda, T., Higashi, I., Kobayashi, M., “Process of Formation and Decomposition of Aluminium Borides”, Sci. Papers Inst. Phys. Chem. Res., 61, 92–99 (1967) (Phase Diagram, Crys. Structure, 8) Bliznakov, G., Peshev P., Niemyski, T., “On the Preparation of Crystalline Aluminium Borides by a Vapour Deposition Process”, J. Less-Common Met., 12, 405–410 (1967) (Experimental, 14) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Serebryanskii, V.T., Epel’baum, V.Z., Zhdanov, G.S., “Equilibrium Diagram of the Aluminium - Boron System”, Russ. J. Inorg. Chem., 12(9), 1311–1316 (1967) (Phase Diagram, 33) Will, G., “Crystal Structure Analysis of AlB10 by the Convolution Molecule Method”, Acta Crystallogr., 23, 1071–1079 (1967) (Crys. Structure, 11) Perrotta, A.J., Townes, W.D., Potenza, J.A., “Crystal Structure of C8Al2.1B51”, Acta Crystallogr., 25B, 1223–1229 (1969) (Crys. Structure, Experimental, 11) Will, G., “The Crystal Structure of C4AlB24”, Acta Crystallogr., 25B, 1219–1222 (1969) (Crys. Structure, Experimental, 11) Neidhard, H., Mattes, R., Becher, H.J., “On the Preparation and Structure of an Aluminium Bearing Boron Carbide”, Acta Crystallogr., 26B, 315–317 (1970) (Crys. Structure, Experimental, 11) Will, G., “On the Existence of AlB10: a Critical Review of the Crystal Structures of AlB10 and C4AlB24”; Electrochem. Technol., 3(1-2), 119–126 (1970) (Crys. Structure, Experimental, 11) Baker, A.A., Braddick, D.M. Jackson, P.W., “Fatigue of Boron-Aluminium and Carbon-Aluminium Fibre Composites”, J. Mater. Sci., 7, 747–62 (1972) (Mechan. Prop., Experimental, 18) Sirtl, E., Woerner, L.M., “Preparation and Properties of Aluminium Diboride Single Crystals”, J. Cryst. Growth, 16, 215–218 (1972) (Crys. Structure, Phase Diagram, 15) Herring, H.W., Lytton, J.L., Steele, J.H., “Experimental Observations of Tensile Fracture in Unidirectional Boron Filament Reinforced Aluminium Sheet”, Metall. Trans. A, 4(3), 807–817 (1973) (Experimental, Mechan. Prop., 9) Munir, Z.A., Veerkamp, G.R., “Investigation of the Parameters Influencing the Microstructure of HotPressed Boron Carbide”, California Univ., Davis (USA). Dept. of Engineering,. 95 pp. (1975) (Mechan. Prop., Crys. Structure, 32) Lange, R.G., Munir, Z.A., “Sintering Kinetics of Pure and Doped Boron Carbide. Final Technical Report”, California Univ., Davis (USA). Dept. of Mechanical Engineering, 35 pp. (1976) (Experimental, 0) Mondolfo, L.F., “Aluminium - Boron System” in “Aluminium Alloys: Structure and Properties”, Butterworths, London, pp. 228–230 (1976) (Review, Phase Diagram, 29) Matkovich, V.I., Economy, J., “Structural Determinants in Higher Borides” in “Boron and Refractory Borides”, Matkovich, V.I. (Ed.), Springer Verlag, Berlin, 78–95 (1977) (Crys. Structure, Review, 36) Roszler, J.J., “Production of Neutron Shielding Material. Patent; B4C+Al in Al Boxes”, US Patent Document 4,027,377/A/, (1977)
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Al–B–C [1978Boi]
[1978Ekb] [1978Sur] [1979Kis]
[1979Pan]
[1980Ino] [1982Doe]
[1983Hig] [1984Sig] [1984Via] [1985Che]
[1985Hal]
[1985Kov]
[1985Pyz]
[1985Sar]
[1986Che] [1986Dub] [1986Hal] [1986Kis] [1986Pes]
[1986Ros] [1987Hau]
3
Boiko, Yu.V., Gol’tsev, V.P., Gorobtsov, V.G., Kavkhuta, G.A., Strelkov, G.I., Khrenov, O.V., Yuzhanin, M.I., “Development and Investigation of Properties of Disperse Boron-Containing Materials for Control Rods of a Nuclear Reactor” (in Russian), Vest. Akad. Navuk BSSR, Ser. Fiz.-Energ. Navuk, 3, 5–8 (1978) (Mechan. Prop., Experimental) Ekbom, L.B., “Effect of Increased Boron Content on the Sintering Behavior and Mechanical Properties of Boron Carbide”, Keram. Z., 183–189 (1978) (Experimental, Mechan. Prop., 6) Suri, A.K., Gupta, C.K., “Studies on the Fabrication of Aluminium Bonded Boron Carbide Rings”, J. Nucl. Mater., 74(2), 297–302 (1978) (Experimental, 4) Kislyi, P.S., Kozina, G.K., Bodnaruk, N.I., “Wetting and Impregnation of Boron Carbide with Copper, Aluminum, and Their Alloys” (in Russian), Adgez. Rasplav. Pajka Mater., 4, 54–57 (1979) (Experimental) Panasyuk, A.D., Oreshkin, V.D., Maslennikova, V.R., “Study of the Kinetics of the Reactions of Boron Carbide with Liquid Aluminium, Silicon, Nickel and Iron”, Sov. Powder Metall. Met. Ceram., 199(7), 487–490 (1979), translated from Poroshk. Metall., 199(7), 79–83 (1979) (Experimental, 9) Inoue, Z., Tanaka, H., Inomata, Y., “Synthesis and X-Ray Crystallography of Aluminium Boron Carbide”, J. Mater. Sci., 15, 3036–3040 (1980) (Crys. Structure, Experimental, 7) Do¨rner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, Univ. Stuttgart (1982) (Experimental, Thermodyn., 126) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Crys. Structure, 17) Sigworth, G.K., “The Grain Refining of Aluminium and Phase Relationships in the Al-Ti-B System”, Mater. Trans. 15A, 277–282 (1984) (Experimental, Phase Diagram, Thermodyn. Calculation, 28) Viala, J. C., Bouix, J., “Elaboration of Aluminum-Matrix Composite Materials Reinforced with Inorganic Fibers”, Mater. Chem. Phys., 11(2), 101–123 (1984) (Mechan. Prop., Experimental, 41) Chernyshova, T.A., Tsirlin, A.M., Gevlich, S.O., Rebrov, A.V., Obolenskii, A.V., “Effect of Surface Condition on the Strength of Coated Boron Fibers”, Sov. Powder Metall. Met. Ceram., 24(3), 210–213 (1985), translated from Poroshk. Metall., 24(3), 39–43 (1985) (Mechan. Prop., Experimental, 9) Halverson, D.C., Pyzik, A.J., I.A. Aksay, I.A., “Processing and Microstructural Characterization of B4CA1 Cermets”, “Composites and Advanced Ceramic Materials”, Anon. Proc. 9th Annu. Conf., American Ceramic Society, Inc., Columbus, OH, 736–744 (1985) (Mechan. Prop., Experimental, 14) Koval’chenko, M.S., Laptev, A.V., Zhidkov, A.B., “Annealing Effect on Structure and Properties of Hot Pressed Cermets Based on Boron Carbide” (in Russian), Poroshk. Metall., 24(9), 51–54 (1985) (Mechan. Prop., Experimental, 6) Pyzik, A. J., Aksay, I. A., “Processing, Microstructure, and Mechanical Properties of Boron CarbideAluminum Alloys Composites”, Anon. Abst. 38th Annu. Pacific Coast Regional Meeting American Ceramic Society, American Ceramic Society, Columbus, OH, (1985) (Mechan. Prop., Experimental, 0) Sarikaya, M., Pyzik, A.J., Ilsay, I.A., Snowden, W. E., “Effect of Secondary Phases on the Properties of B4C-A1 Composites.”, Anon. Abst. of the 38th Annu. Pacific Coast Regional Meeting American Ceramic Society, American Ceramic Society, Columbus, OH, (1985) (Mechan., Prop., Experimental, 0) Chernyshova, T.A., Rebrov, A.V., “Interaction Kinetics of Boron Carbide and Silicon Carbide with Liquid Aluminium”, J. Less-Common Met., 117, 203–207 (1986) (Kinetics, Experimental, 4) Dub, S.N., Prikhna, T.A., Il’nitskaya, O.N., “Mechanical Properties of the Al-B-C Compounds Crystals” (in Russian), Sverkhtverd. Mater., 6, 12–18 (1986) (Mechan. Prop., Experimental, 22) Halverson, D.C., Pyzik, A.J., Aksay, I.A., “Boron-Carbide-Aluminum and Boron-Carbide-Reactive Metal Cermets”, US patent document 4,605,440/A/, (1986) Kisly, P.S., Prikhna T.A., Golubyak, L.S., “Properties of High-Temperature Solution Grown Aluminium Borides”, J. Less-Common Met., 117, 349–353 (1986) (Experimental, 10) Peshev, P., Gyurov, G., Khristov, M., Gurin, V.N., Korsukova, M. M., Solomkin, F.Yu., Sidorin K.K., “Preparation and some Properties of Aluminium Carboboride Single Crystals”, J. Less-Common Met., 117, 341–348 (1986) (Crys. Structure, Mechan. Prop., Optical Prop., Experimental, 16) Roszler, J.J., “Process for the Manufacture of a Material Shielding Against Neutrons” (in German), DE Patent Document 2643444/C2/, (1986) Haupt, H., Werheit, H., Siejak, V., Gurin, V.N., Korsukova, M.M., “Absorption Edge and IR-active Phonons of Al3B48C2, “Boron, Borides and Related Compounds”, Proc. 9th Int. Sympos., Werheit, H. (Ed.), Univ. Duisburg, Germany, 387–389 (1987) (Experimental, 2)
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3 [1987Kis]
[1987Lev]
[1987Pyz] [1987Sar]
[1988Kov]
[1989Blu]
[1989Hal] [1989Kha]
[1990Ase]
[1990Luk]
[1990Oka]
[1990Pyz]
[1990Ram] [1991Kha1]
[1991Kha2]
[1991Kha3]
[1991Kha4]
[1991Pri]
Al–B–C Kisly, P.S., Prikhna, T.A., Gontar A.N., Podarevskaya, O.V., “Structure and Properties of Monocrystals of the Al-B-C System Compounds” in “Boron, Borides and Related Compounds”, Proceedings 9th Int. Sympos., Werheit, H. (Ed), Univ. Duisburg, Germany, 273–274 (1987) (Thermodyn., Crys. Structure, Phys. Prop., Experimental, 1) Levinskas, D., “Evaluation of Boron Carbide Coatings”, Western Region American Nuclear Society Student Conference: Nuclear Technology for the Year 2000, American Nuclear Society, La Grange Park, IL., NM(USA), 68–71 (1987) (Experimental, 0) Pyzik, A.J., Aksay, I.A., “Multipurpose Boron Carbide-Aluminum Composite and its Manufacture via the Control of the Microstructure”, US patent document 4,702,7707 A/, 27, (1987) Sarikaya, M., Laoui, T., Milius D.L., Aksay, I.A., “Identification of a New Phase in the Al-B-C Ternary by High-Resolution Transmission Electron Microscopy”. Proc. 45th Ann. Meeting of the Electron Microscopy Society of America, Bailey, G.N., (Ed.), San Franc. Press, USA, 168–169 (1987) (Crys. Structure, Experimental, 4) Koval’chenko, M.S., Paustovskij, A.V., Bolejko, B.M. Zhidkov, A.V., “Laser Surface Hardening of Cermets on the Base of Boron Carbide”(in Russian), Poroshk. Metall., 5, 77–80 (1988) (Mechan. Prop., Experimental, 6) Blumenthal, W.R., Gray, G.T., “Structure-Property Characterization of Shock-Loaded B4C-Al”, Inst. Phys. Conf. Ser. No 102: Session 7, Paper Presented at Int. Conf. Mech. Prop. Materials at High Rates of Strain, Oxford, 363–370 (1989) (Experimental, 8) Halverson, D.C., Pyzik, A.J., Aksay, I.A., Snowden, W.E., “Processing of Boron Carbide-Aluminium Composites”, J. Am. Ceram. Soc., 72(5), 775–80 (1989) (Experimental, 33) Kharlamov, A.I., Duda, T.I., Lojchenko, S.V., Fomenko, V.V., “Preparation and Properties of Aluminium Boridocarbide Powder of Al8B4C7 Composition”, 12thUkrainian Republic Conference on Inorganic Chemistry, Vol. 1, Simferopol’, Ukr. SSR, 44pp. (1989) (Mechan. Prop., Experimental, 0) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. NATO Advanced Research Workshop, Manchester, U.K., 1989, published as ASI-Series, Ser. E: Appl. Sci., Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental,14) Lukas, H.L., “Aluminium-Boron-Carbon” in “Ternary Alloys. A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams”, Petzow, G., Effenberg, G., (Eds.), Vol. 3, VCH, Weinheim, 140–146 (1990) (Review, Phase Diagram, 14) Okada, S., Kudou, K., Hiyoshi, H., Higashi, I., Hamano, K., Lundstro¨m, T., “Preparation of AlC4B24 and Al3C2B48 Crystals”, J. Int. Ceram. Soc. Jpn., 98, 1342–1347 (1991), translated from Nippon Seramikkusu Kyokai Gakujutsu Ronbunshi, 98(12), 1330–1336 (1990) (Experimental, Crys. Structure, 24) Pyzik, A.J., Williams P.D., McCombs, A., “New Low Temperature Processing for Boron Carbide/ Aluminium Based Composite Armor”, Final Report, US-Army Research Office, DAAL 0388 C0030, 1990 (Experimental, 14) Ramesh, K. T., Ravichandran, G., “Dynamic Behavior of a Boron Carbide-Aluminum Cermet: Experiments and Observations”, Mech. Mater., 10(1-2) 19–29 (1990) (Experimental, 22) Kharlamov, A.I., Loichenko, S.V., “Electronic Transport Properties of Hot-pressed Boron-rich Compounds of the Al-B-C System” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 94–97 (1991) (Experimental, 5) Kharlamov, A.I., Loichenko, S.V., “Investigation: The Process of Densification of Boron-Rich Compounds of the Al-B-C System” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 473–481 (1991) (Experimental, 2) Kharlamov, A.I., Murzin, L.M., Loichenko, S.V., Duda, T.I., “Electrical Conductivity and Seebeck Coefficient of Hot-Pressed Specimens of Aluminium Borides and Carboborides”, Sov. Powder Metall. Met. Ceram., 9(345), 770–773 (1991), translated from Poroshk. Metall., 9(345), 62–65 (1991) (Experimental, Electr. Prop., 7) Kharlamov, A.I., Duda, T.I., Fomenko, V.V., “Preparation and Properties of High-Dispersive Powders of Aluminium Dodecaboride and Carboborides” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. (Eds.), Albuquerque, USA, 1990, AIP, New York, 512–515 (1991) (Experimental, 0) Prikhina, T.A., Kisly, P.S., “Aluminium Borides and Carboborides” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 590–593 (1991) (Experimental, 11)
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Al–B–C [1991Pyz] [1992Bei] [1992Var] [1992Via] [1993Bau]
[1993Gon]
[1993Ips] [1993Kau] [1993Wen] [1993Wer]
[1994Dus] [1994Kud]
[1995Bon] [1995Hil] [1995Osc] [1995Pyz]
[1995Rug] [1996Bid] [1996Hil1]
[1996Hil2]
[1996Kas] [1996Pyz] [1996RMi]
[1996Vin]
3
Pyzik, A.J., Nilson, R.T., “B4C/A1 Cermets and Method for Making Same”, US Patent Document 5,039,633, (1991) Beidler, C.J., Hauth III, W.E., Goel, A., “Development of a B4C/A1 Cermet for Use as an Improved Structural Neutron Absorber”, J. Testing and Evaluation, 20(1), 67–70 (1992) (Experimental, 6) Vardiman, R.G., “Microstructures in Aluminium, Ion Implanted with Boron and Heat Treated”, Acta Metall. Mater., 40, 1029–35 (1992) (Crys. Structure, Eperimental, 7) Viala, J.C., Gonzales, G., Bouix, J., “Composition and Lattice Parameters of a New Aluminium-Rich Boron Carbide”, J. Mater. Sci. Lett., 11, 711–714 (1992) (Crys. Structure, Experimental, 9) Bauer, J., Bittermann, H., Rogl, P., “Phase Relations and Structural Chemistry in the Ternary System Aluminium - Boron - Carbon”, COST-507, Annual Report, (1993) (Crys. Structure, Phase Diagram, Experimental, 12) Gonzalez, G., Esnouf, C., Viala, J.C., “Structural Study of a New Aluminium Rich Borocarbide Formed by Reaction at the B4C/Al Interface”, Mater. Sci. Forum, 126-128, 125–128 (1993) (Crys. Structure, Experimental, 4) Ipser, H., privat communication (1993) (Experimental) Kaufmann, L., private communication (1993) (Thermodyn.) Wen, H., “Thermodynamic Calculations and Constitution of the Al-B-C-N-Si-Ti System” (in German), Thesis, Univ. Stuttgart, 1–183 (1993) (Calculation, Phase Diagram, Thermodyn., 223) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489–511 (1993) (Crys. Structure, Experimental, 51) Duschanek, H., Rogl, P., “The System Al-B”, J. Phase Equilib., 15(5), 543–52 (1994) (Crys. Structure, Phase Diagram, Experimental, #, 78) see also ibid, 16(1), 6 (1995) Kudou, K., Okada, S., Hikichi, H., Lundstro¨m, T., “Preparation and Properties of Si-doped Al3C2B48Type Crystals” (in Japanese), J. Soc. Mater. Sci., Jpn., 43(485), 223–228 (1994) (Experimental, Crys. Structure, Phys. Prop., 20) Bond, G.M., Inal, O.T., “Shock-Compacted Aluminium/Boron Carbide Composites”, Compos. Eng. 5(1), 9–16 (1995) (Experimental, 18) Hillebrecht, H., Meyer, F., “B48A13C2 - a Filled Variant of Tetragonal Boron”, Z. Kristallogr., Suppl. 10, 101 (1995) (Crys. Structure, Experimental, 2) Oscroft, R.J., Roebuck P.H.A., Thompson, D.P., “Characterisation and Range of Composition for Al8B4C7”, Br. Ceram. Trans., 94(1), 25–26 (1995) (Experimental, 11) Pyzik, A.J., Beaman, D.R., “Al-B-C Phase Development and Effects on Mechanical Properties of B4C/ Al-Derived Composites”, J. Am. Ceram. Soc., 78(2), 305–312 (1995) (Crys. Structure, Mechan. Prop., Experimental, 25) Ruginets, R., Fischer, R. “Microwave Sintering of Boron Carbide Composites”, Am. Ceram. Soc. Bull., 74(1), 56–58 (1995) (Experimental) Bidaud, E., research at Univ. Wien, unpublished (1996) Hillebrecht, H., Meyer, F.D., “The Structure of B48Al3C2 - A Filled and Distorted Variant of Tetragonal Boron (I)” in “Boron, Borides and Related Compounds”, Proc. 12th Int. Symp., Baden/Wien, paper PA.4, 59 (1996) (Crys. Structure, Experimental, 6) Hillebrecht, H., Meyer, FD., “Synthesis, Crystal Structure, and Vibrational Spectra of Al3BC3, a Carbidecarboborate of Aluminium with Linear (C=B=C)5– Anions”, Angew. Chem., 35(21), 2499–2500 (1996), translated from Angew. Chemie, 108(21), 2655–2657 (1996) (Crys. Structure, Experimental, 17) Kasper, B., “Phase Equilibria in the B-C-N-Si System”, Thesis, Max Plank Institute-PML, Stuttgart, (1996) (Phase Diagram, Thermodyn.) Pyzik, A.J., Deshmukh, U.V., Dunmead, S.D., Ott, J.J., Allen, T.L., Rossow, H.E., “Light Weight Boron Carbide/Aluminium Cerments”, United States Patent: 5,521,016, (1996) R’Mili, M., Massardier, V., Merle, P., Vincent, H., Vincent, C., “Mechanical Properties of T300/A1 Composites. Embrittlement Effects due to a B4C Coating”, J. Mater. Sci., 31, 4533–4539 (1996) (Mechan. Prop., Experimental, 12) Vincent, H., Vincent, C., Berthet, M. P., Bouix, J., Gonzalez, G., “Boron Carbide Formation from BCl3CH4-H2 Mixtures on Carbon Substrates and in a Carbon-Fibre Reinforced Al Composite”, Carbon, 34 (9), 1041–1055 (1996) (Crys. Structure, Mechan. Prop., Experimental, 25)
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3 [1997Du] [1997Mey] [1997Sch]
[1997Via] [1997Yue] [1998Kas]
[1998Kha]
[1998Rog]
[1999Bur] [1999Cha]
[1999Tor]
[1999Tsu] [2000Hal]
[2000Hig] [2000Kha]
[2000Liu] [2000Mey]
[2000Pyz] [2000Sav] [2000Sol] [2000Wan]
Al–B–C Du, W.F., Watanabe, T., “High-Toughness B4C-AlB12 Composites Prepared by Al Infiltration”, J. Eur. Ceram. Soc., 17, 879–884 (1997) (Mechan. Prop., Experimental, 15) Meyer, F.D. Hillebrecht, H., “Synthesis and Crystal Structure of Al3BC, the First Boridecarbide of Aluminium”, J. Alloy. Compd., 252, 98–102 (1997) (Crys. Structure, Experimental, 30) Schmechel, R., Werheit, H., Robberding, K., Lundstro¨m, T., Bolmgren, H., “IR-active Phonon Spectra of B-C-Al Compounds with Boron Carbide Structure”, J. Solid State Chem., 133, 254–259 (1997) (Experimental, 11) Viala, J.C., Bouix, J., Gonzalez, G., Esnouf, C. “Chemical Reactivity of Aluminium with Boron Carbide”, J. Mater. Sci, 32, 4559–4573 (1997) (Phase Diagram, Experimental, 39) Yu¨cel, O., Tekin, A., “The Fabrication of Boron-Carbide-Aluminium Composites by Explosive Consolidation”, Ceram. Int., 23, 149–152 (1997) (Experimental, Mechan. Prop., 3) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Kharlamov, A.I., Kirillova, N.V., Loichenko, S.V., Fomenko, V.V., “Properties of Aluminium Borides and Borocarbides”, Russ. J. Appl. Chem., 71(5), 743–749 (1998), translated from Zh. Prikl. Khim, 71(5), 717–724 (1998) (in Russian), (Crys. Structure, Kinetics, Mechan. Prop., Experimental, 13) Rogl, P., “Al-B-C (Aluminium-Boron-Carbon)”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12170.2.20, (1998) aslo published in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G., (Ed.), ASM-Intl, MSI, 3–15 (1998) (Assessment, Crys. Structure, Experimental, Phase Diagram, 50) Burkhardt, U., Grin, Y., “Refinement of the Aluminium Diboride Crystal Structure” in “Borides and Related Compounds”, Abst. 13th Int. Symp. on Boron, Dinar (France), 13pp., (1999) (Crys. Structure, 3) Chapman, T.R., Niesz, D.E., Fox, R.T., Fawcett, T., “Wear-resistant Aluminum - Boron - Carbide Cermets for Automotive Brake Applications”, Wear, 236, 81–87 (1999) (Mechan. Prop., Experimental, 9) Torquato, S., Yeong, C.L.Y., Rintoul, M.D., Milius, D.L., Aksay, I.A., “Elastic Properties and Structure of Interpenetrating Boron Carbide/Aluminum Multiphase Composites”, J. Am. Ceram. Soc., 82(5), 1263–1268 (1999) (Mechan. Prop., 32) Tsuchida, T., Kan, T., “Synthesis of Al3BC in Air from Mechanically Activated Al/B/C Powder Mixtures”, J. Eur. Ceram. Soc., 19, 1795–1799 (1999) (Crys. Structure, Experimental, 12) Hall, A., Economy, J., “The Al(L)+AlB12ÐAlB2 Peritectic Transformation and its Role in the Formation of High Aspect Ratio AlB2 Flakes”, J. Phase Equilib., 21(1), 63–69 (2000) (Phase Diagram, Experimental, 21) Higashi, I., “Crystal Chemistry of α-AlB12 and γ-AlB12”, J. Solid State Chem., 154, 168–176 (2000) (Crys. Structure, Experimental, 18) Kharlamov, A. I., Nizhenko, V.I., Kirillova, N.V., Floka, L.I., “Wettability of Hot-Pressed Samples of Boron-Containing Aluminium Compounds by Liquid Metals and Alloys” (in Russian), Zh. Prikl. Khim., 73(6), 884–888 (2000) (Experimental, 14) Liu, C.H., “Structure and Properties of Boron Carbide with Aluminum Incorporation”, Mater. Sci. Eng. B, B72, 23–26 (2000) (Phys. Prop., Crys. Structure, Experimental, 10) Meyer, F.D., Hillebrecht, H., “Ternary Phases in the System Al/B/C” in “High Temperature Materials Chemistry”, Vol. 15, Part 1, K. Hilpert et al. (Eds.), Proc. 10th Intl. IUPAC Conf., Forschungszentrum Ju¨lich, Germany, Published by Schriften des Forschungszentrums Juelich, 161–164 (2000) (Crys. Structure, 5) Pyzik, A.J., Deshmukh, U.V., Krystosek, R. D., “Aluminum-Boron-Carbon Abrasive Article and Method to Form Said Article”, US Patent: 6,042,627, (2000). Savyak, M., Uvarova, I., Timofeeva, I., Isayeva L., Kirilenko, S., “Mechanochemical Synthesis in Ti-C, Ti-B, B-C, B-C-A1 Systems”, Mater. Sci. Forum, 343-346, 411–416 (2000) (Experimental, 4) Solozhenko, V.L., Meyer, F.D., Hillebrecht, H., “300-K Equation of State and High-Pressure Phase Stability of Al3BC3”, J. Solid State Chem., 154, 254–256 (2000) (Crys. Structure, Experimental, 11) Wang, T., Yamaguchi, A., “Some Properties of Sintered Al8B4C7”, J. Mater. Sci. Letter., 19, 1045–1046 (2000) (Calculation, Crys. Structure, 6)
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Al–B–C [2000Wer] [2001Fje]
[2001Lee]
[2002Ars] [2002Bur]
[2002Kou1] [2002Kou2]
[2002Zhe] [2003Per]
[Mas2] [V-C2]
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Werheit, H., Schmechel, R., Meyer, F.D., Hillebrecht, H., “Interband Transitions and Optical Phonons of B48Al3C2”, J. Solid State Chem., 154, 75–78 (2000) (Optical Prop., Experimental, 10) Fjellstedt, J., Jarfors, A.E.W., El-Benawy, T., “Experimental Investigation and Thermodynamic Assessment of the Al-rich Side of the Al-B System”, Mater. Des., 22(6), 443–449 (2001) (Thermodyn, Phase Diagram, Experimental, 14) Lee, K.B., Sim, H.S., Cho, S.Y., Kwon, H., “Reaction Products of Al-Mg/B4C Composite Fabricated by Pressureless Infiltration Technique”, Mater. Sci. Eng. A, 302, 227–234 (2001) (Crys. Structure, Phase Diagram, Experimental, 17) Arslan, G., Kara, F., Turan, S., “Mechanical Properties of Melt Infiltrated Boron Carbide-Aluminium Composites”, Key Eng. Mater., 206-213(2), 1157–1160 (2002) (Experimental, Mechan. Prop., 5) Burkhardt, U., Gurin, V., Borrmann, H., Schnelle, W., Grin, Y., “On the Electronic and Structural Properties of Aluminium Diboride Al0.9B2” in “Boron, Borides and Related Compound”, Abst. 14th Int. Symp., (ISBB’02), Saint Petersburg, O4, (2002) (Crys. Structure, 3) Kouzeli, M., Mortensen, A., “Size Dependent Strengthening in Particle Reinforced Aluminium”, Acta Mater., 50, 39–51 (2002) (Mechan. Prop., Experimental, 59) Kouzeli, M., Marchi, C. S., Mortensen, A., “Effect of Reaction on the Tensile Behavior of Infiltrated Boron Carbide-Aluminum Composites”, Mater. Sci. Eng. A, A337, 264–273 (2002) (Experimental, Mechan. Prop., 51) Zheltov, P., Grytsiv, A., Rogl, P., Velikanova, T.Ya., Research at Univ. Wien (unpublished) (2002) (Phase Diagram, Crys. Structure) Perrot, P., “Aluminium-Carbon”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published, (2003) (Phase Diagram, Crys. Structure, Assessment, 19) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Boron – Molybdenum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Qingsheng Ran, updated by Lazar Rokhlin, Tatiana Dobatkina, Elena Semenova, Natalia Kol’chugina
Introduction The Al-B-Mo system is of interest mainly because of the application of low-alloyed Mo compositions in manufacturing parts for operation under conditions of high-temperature gases and erosion wear [1988Buk, 1994Buk, 1995Buk, 1997Buk, 1999Buk, 2003Buk1, 2003Buk2]. The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Ran], which is updated by the present work. An isothermal section at 1000˚C was presented by [1965Rie] obviously from the determination of phase regions using X-ray diffraction. Metals were cold pressed; samples were prepared by reacting at 900˚C and annealing at 1200˚C for 15 h. While cooling, equilibrium was said to be frozen in at 1000˚C. [1942Hal] prepared a ternary compound “Mo7Al6B7” by thermal reaction of MoO3, B2O3, Al and S and attempted a first indexation of its powder pattern. The existence of the phase was later confirmed by [1965Rie, 1966Jei], which assumed that its formula is MoAlB. Moreover, an Al-stabilized molybdenum boride with the CrB type structure was found. To investigate the structure of this phase more accurately, [1966Jei] heated a sample of the composition Mo:Al:B = 1:6:1 to about 1800˚C and cooled it for 40 minutes to 1000˚C. The charge was subsequently treated in hot NaOH solution and single crystals were obtained from the residue. In performing X-ray diffraction, the Debye-Scherrer, Weissenberg, and rotating crystal methods were used. In [1977Gur], crystals of solid solutions of boron (to 1 mass%) in MoAl4(5) were separated from Al-B-Mo alloys (the alloys were prepared preliminarily by melting followed by spontaneous solidification) by dissolution of Al matrix in HCl. The MoAlB (orthorhombic) compound and another compound of the Al-B-Mo system with a rhombohedral crystal structure were reported by [1976Hig]. Reviews [1972Cha, 1976Hig] included the aforementioned data on the phases existing in the ternary Al-B-Mo system. Investigations into the Al-B-Mo phase relations, structure identifications are given in Table 1.
Binary Systems The binary systems Al-B [2004Mir] (Fig. 1), B-Mo [1992Rog] (Fig. 2), and Al-Mo [2005Sch] are accepted in the present assessment. Landolt‐Bo¨rnstein New Series IV/11E1
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Solid Phases The solid phases of the Al-B-Mo system are described in Table 2 that is presented mainly with allowance for the binary phase diagrams accepted in this assessment. Two ternary phases were found in the system; these are τ1 (MoAlB) and τ2 (Mo9AlB10). The homogeneity ranges of the phases are taken from the isothermal section at 1000˚C given by [1965Rie]. The aluminium solubility in B-Mo phases and boron solubility in Al-Mo phases are taken also from this section. Virtually the same boron solubility in MoAl4 was confirmed by [1977Gur].
Isothermal Sections The only isothermal section at 1000˚C (Fig. 3) is constructed based on data by [1965Rie] corrected with allowance for the binary phase diagrams accepted in the present evaluation. In accordance with the binary B-Mo system, the existing MoB4 compound is added in Fig. 3; two three-phase regions containing the MoB4 phase in the section are given tentatively. The MoAl4 and Mo4Al17 compounds are added to the section after the accepted binary Al-Mo system. They are located between the Al rich liquid and Mo3Al8 that are in equilibrium with the τ1 phase. Therefore, it was reasonable to accept the existence of the MoAl4 and Mo4Al17 compounds in equilibrium with the τ1 phase. This is indicated in the section in Fig. 3.
Notes on Materials Properties and Applications Owing to the importance of molybdenum-based alloys containing small amounts of Al and B for the manufacturing of parts for operation under conditions of high-temperature gases and erosion wear, various mechanical properties of them have been investigated at room and high temperatures [1988Buk, 1994Buk, 1995Buk, 1997Buk, 1999Buk, 2000Buk, 2003Buk1, 2003Buk2]. Table 3 shows the mechanical properties and their correlation with the structural state.
Miscellaneous The possibility of forming composites with Al matrix strengthened by molybdenum borides is considered in [1996Bie].
. Table 1 Investigations of the Al-B-Mo Phase Relations, Structures and Thermodynamics Reference [1942Hal]
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
Interaction of molybdenum boride with 74 mass% Mo, 18 mass% Al, 8.5 aluminium followed by dissolution of Al mass% B / near Mo7Al6B7 matrix / X-ray diffraction, chemical analysis, density measurements
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied 1000˚C / 0–80 at.% Al, 0–100 at.% Mo, 0–100 at.% B /
[1965Rie]
Cold pressing of powders followed by heating, X-ray diffraction
[1966Jei]
Heating of components to 1800˚C followed 1000˚C / MoAlB by slow cooling, chemical extraction / X-ray diffraction
[1977Gur] Melting followed by chemical extraction after solidification / X-ray diffraction, chemical analysis
1400–1000˚ / 55 mass% Al, 43 mass% Mo, 0.7 mass% B / MoAl4-based solid solution of B
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βAl)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25˚C, 20.5 GPa [Mas2]
(αAl) < 660.452
cF4 Fm 3m Cu
a = 404.96
at 25˚C [Mas2]
(βB) < 2092
hR333 R3m βB
a = 1093.30 c = 2382.52
[1993Wer]
a = 1093.02 c = 2381.66
pure B [1976Lun]
a = 1006.1 c = 1421.0
presumably metastable phase, possibly stabilized by impurities, preparation at 1100 to 1400˚C [1971Amb]
(B) (tetr.)
tP192 derived from αAlB12
(αB)
hR36 R 3m αB
(Mo) < 2623
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cI2 Im 3m W
a = 490.8 c = 1256.7 a = 314.70
MSIT1
presumably metastable phase, preparation below 1000˚C [1971Amb] pure B, single crystal [1994Cha] at 25˚C [Mas2] dissolves up to 19.5 at.% Al and up to 1 at.% B [2005Sch, Mas2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
αAlB12 < 2094
oP384 P212121 αAlB12
a = 1662.3 b = 1754.0 c = 1018.0
[2004Mir] [V-C2]
γAlB12 ≤ 1550
oP396 P212121 γAlB12
a = 1662.3 b = 1754.0 c = 1018.0
[1983Hig] Metastable or metal impurity stabilized [1994Dus]
AlB2 < 972 - 213
hP3 P6/mmm AlB2
a = 300.5 c = 325.7
[2004Mir] [V-C2]
MoAl12 < 712
cI26 Im 3 WAl12
a = 757.3
[2005Sch] 92.3 at.% Al
MoAl5(h2) 846 - (750 - 800)
hP12 P6322 WAl5
a = 491.2
[2005Sch] 83.3 at.% Al
a = 493.3 c = 4398
[2005Sch] 83.3 at.% Al
MoAl5(h1) hP60 (750–800) - 648 P321 MoAl5(h1) MoAl5(r) ≲ 648
hP36 R 3c MoAl5(r)
a = 495.1 c = 2623
[2005Sch] 83.3 at.% Al
Mo5Al22(h) 964 - 831
oF216 Fdd2 Mo5Al22(h)
a = 7382 ± 3
[2005Sch] 81.5 at.% Al
Mo4Al17 < 1034
mC84 C2 Mo4Al17
a = 915.8 ± 0.1 b = 493.23 ± 0.08 c = 2893.5 ± 0.5 β = 96.71 ± 0.01˚
[2005Sch] 81 at.% Al
MoAl4(h) 1177 - 942
mC30 Cm WAl4
a = 525.5 ± 0.5 [2005Sch] b = 1776.8 ± 0.5 79 at.% Al c = 522.5 ± 0.5 β = 100.88 ± 0.06˚
Mo1–xAl3+x 1260 - 1154
cP8 Pm 3n Cr3Si
a = 494.5 ± 0.1
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[2005Sch] 76.5 at.% Al
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
MoAl3(h) 1222 - 818
mC32 C2/m MoAl3
a = 1639 ± 1.0 [2005Sch] b = 359.4 ± 0.1 75 at.% Al c = 838.6 ± 0.4 β = 101.88 ± 0.07˚
Mo3Al8 < 1555 ± 10
mC22 Cm Mo3Al8
a = 920.8 ± 0.3 [2005Sch] b = 363.78 ± 0.03 72.7 at.% Al c = 1006.5 ± 0.3 β = 100.78 ± 0.05˚
ζ1(h), Mo2Al3(h) 1570 - 1490
Unknown
-
[2005Sch] 63 at.% Al
ζ2(h), MoAl(h) 1750 - 1535
cP2 Pm 3m CsCl
a = 309.8
[2005Sch] 50–52 at.% Al
Mo3Al ≲ 2150
cP8 Pm 3n Cr3Si
a = 495.0
[2005Sch] 23–29 at.% Al
Mo2B < 2280
tI12 I4/mcm Al2Cu
a = 554.7 c = 473.9
[Mas2] [V-C2] 33 at.% B
αMoB < 2180
tI16 I41/amd MoB
a = 310.3 c = 1695
[Mas2] [V-C2] 48–50 at.% B
βMoB 2600 - 2180
oC8 Cmcm CrB
MoB2 2375 - 1517
hP3 P6/mmm AlB2
a = 303.9 c = 305.5
[Mas2] [V-C2] 62–66 at.% B
Mo2B5 < 2140
hR21 R3m Mo2B5
a = 300.7 c = 2091.0
[Mas2] [V-C2] 67–69 at.% B
MoB4 < 1807
hP20 P63/mmc WB4
a = 520.3 c = 634.5
[Mas2] [V-C2] 79 at.% B
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[Mas2] [V-C2] 48–51 at.% B
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
* τ1, MoAlB ≲ 1000˚C
oC12 Cmcm BCU
a = 321.2 b = 1398.5 c = 310.2
[1966Jei, V-C2] 31–34.5 at.% B and 33–35.5 at.% Al
* τ2, Mo9AlB10
oC8 Cmcm CrB
a = 316.3 b = 847.5 c = 308.2
[V-C2] 48–50 at.% B and 5–6.5 at.% Al
. Table 3 Investigations of the Al-B-Mo Materials Properties Reference
Method / Experimental Technique
Type of Property
[1988Buk]
Tensile creep tests at 1500–2000˚C. Creep rate and stress-rupture strength for Mobased compositions low-alloyed with Al and B
[1994Buk]
Tensile creep tests 1500–2000˚C.
Correlation between short-term, long-term, and low-cycle strength characteristics of Mo-based compositions low-alloyed with Al and B were established.
[1995Buk]
Tensile tests 20–2000˚C.
Ultimate strength and relative elongation of Mo-based compositions low-alloyed with Al and B are determined.
[1997Buk]
Fatigue tests under conditions of rigid alternating bending at 950–1550˚C.
Effect of heat treatment and welding on fatigue resistance of Mo-based compositions lowalloyed with Al and B is studied.
[1999Buk]
Tensile tests at (0.5–0.8) Tm.
Correlation between short-term, long-term static, creep resistance, and few-cycle strength characteristics of Mo-based compositions lowalloyed with Al and B were established.
[2000Buk]
Tensile and fatigue tests under conditions of rigid alternating bending.
Empirical correlations between fatigue strength, conventional yield strength, and grain size were established.
[2003Buk1] Tensile tests at 20–2000˚C and fatigue tests.
Correlations between the structural state and short-term and low- and high-cycle fatigue were established.
[2003Buk2] Tensile tests at 20–2000˚C.
Effect of heat treatment on ultimate strength, yield strength, relative elongation, and relative uniform deformation is considered.
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. Fig. 1 Al-B-Mo. The Al-B binary system
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. Fig. 2 Al-B-Mo. The B-Mo binary system
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. Fig. 3 Al-B-Mo. Isothermal section at 1000˚C
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References [1942Hal] [1965Rie]
[1966Jei] [1971Amb] [1972Cha]
[1976Hig] [1976Lun] [1977Gur]
[1983Hig] [1988Buk]
[1990Ran]
[1992Rog]
[1993Wer]
[1994Buk]
[1994Cha]
[1994Dus] [1995Buk]
[1996Bie]
Halla, F., Thury, W., “On Borides of Molybdenum and Tungsten” (in German), Z. Anorg. Allg. Chem., 249, 229–237 (1942) (Crystal Structure, Experimental, 7) Rieger, W., Nowotny, H., Benesovksy, F., “Complex Borides of the Transition Metals (Mo, W, Fe, Co, Ni)” (in German), Monatsh. Chem., 96(3), 844–851 (1965) (Crys. Structure, Experimental, Phase Diagram, #, 10) Jeitschko, W., “The Crystal Structure of MoAlB” (in German), Monatsh. Chem., 97, 1472–1476 (1966) (Crys. Structure, Experimental, 10) Amberger, E., Ploog, G., “Formation of Lattices of Pure Boron” (in German), J. Less-Common Met., 23, 21–31, (1971) (Crys. Structure, Experimental, 17) Chaban, N.F., Kuz’ma, Yu.B., “Metallides in the Period IV Transition Metal-Aluminum (Gallium) Boron Ternary System” (in Russian), Stroenie, Svoistva i Primenenie Metall., 102–107 (1972) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, 18) Higashi, I., Takahashi, Y., Atoda, T., “Crystal Growth of Borides and Carbides of Transition Metals from Al Solutions”, J. Cryst. Growth, 33, 207–211 (1976) (Experimental, 16) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Gurin, V.N., Korsukova, M.M., Popov, V.I., Elizarova, O.V., Belozov, N.N., Kuz’ma, Yu.B., “Solid Solutions of B, C, Si in Aliminides of Transient Metals” (in Russian), Akad. Nauk SSSR, Nauka, Moscow, 39–42 (1977) (Crys. Structure, Experimental, Phase Diagram, 6) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Experimental, Crys. Structure, 21) Bukhanovskii, V.V., Kharchenko, V.K., Polishchuk, E.P., Kravchenko, V.S., Galinzovskaya, T.D., Onoprienko, A.A., “Influence of Production Factors on the High-Temperature Strength Characteristics of Molybdenum Alloys”, Strength Met., 20(6), 818–824 (1988), translated from Probl. Prochn., (6), 102–108 (1988) (Experimental, Kinetics, Morphology, 18) Ran, Q., “Al-B-Mo (Aluminium-Boron-Molybdenum)” in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.19496.1.20, (1990), also published in “Ternary Alloys”, Petzow, G., Effenberg, G. (Eds.), Vol. 3, VCH Verlagsgesellschaft, Weinheim, Germany, 189–191 (1990) (Phase Diagram, Phase Relations, Review, 5) Rogl, P., “The system B-N-Mo” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), Ohio, USA: ASM, Materials Park, 64–67 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Experimental, Review, *, 10) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Correlations Between Short-Term, Long-Term Static, and Low-Cycle Strength Characteristics of Low-Alloy Molybdenum Alloys at High Temperatures”, Strength Met., 28(12), 895–902 (1994), translated from Probl. Prochn., (12), 43–51 (1994) (Experimental, Mechan. Prop., 21) Chakrabarti, D.J., Laughlin, D.E., “B-Cu (Boron-Copper)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 74–78 (1994) (Review, Phase Diagram, Crys. Structure, Thermodyn., 24) Duschanek, H., Rogl, P., “The Al-B (Aluminum-Boron) System”, J. Phase Equilib., 15, 543–552 (1994) (Assessment, Crys. Structure, Phase Relations, Phase Diagram, Review, Thermodyn., 78) Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Mechanical Characteristics of Molybdenum Alloys of the Systems the Mo-Al-B and Mo-Zr-B in the Temperature Range 290–2270 K”, Strength Met., 27(11-12), 688–695 (1995), translated from Probl. Prochn., 27(11-12), 70–80 (1995) (Experimental, Mechan. Prop., 19) de Bie, J.E., Froyen, L., Lust, P., Delaey, L., “Solidification of In-Situ Aluminium Composites”, Mater. Sci. Forum, 215-216, 435–442 (1996) (Morphology, Phase Relations, Theory, 10)
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Al–B–Mo [1997Buk]
[1999Buk] [2000Buk]
[2003Buk1]
[2003Buk2]
[2004Mir]
[2005Sch]
[Mas2] [V-C2]
4
Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Effect of Heat Treatment and Welding on the Fatigue Resistances of Molybdenum Alloys of the Mo-Zr-B and Mo-Al-B Systems”, Met. Sci. Heat Treat., 39(5-6), 247–250 (1997), translated from Metalloved. Term.. Obrab. Met., (6), 16–19 (1997) (Experimental, Mechan. Prop., 8) Bukhanovskii, V.V., “Correlation Between the Strength and Creep of Molybdenum and Tungsten Alloys”, Russ. Metall. (Engl. Transl.), (5), 91–98 (1999) (Experimental, Mechan. Prop., 22) Bukhanovskii, V.V., “Correlations between the Characteristics of Fatigue Resistance, Short-Term Strength, and Structure of Low-Molybdenum Alloys”, Strength Met., 32(4), 361–367 (2000), translated from Probl. Prochn., (4), 75–85 (2000) (Experimental, Mechan. Prop., 14) Bukhanovsky, V., Mamuzic, I., Borisenko, V., “Interrelation between the Structural State of Material and Mechanical Properties of Low-Alloyed Molybdenum Alloys”, Metallurgia, 42(3), 159–166 (2003) (Experimental, Kinetics, Mechan. Prop., 25) Bukhanovsky, V.V., Mamuzic, I., Borisenko, V.A., “The Effect of Thermal Treatment on the Mechanical Properties of Low-Alloyed Molybdenum Alloys Over the Wide Range of Temperatures”, Metallurgia, 42(1), 9–14 (2003) (Experimental, Morphology, Mechan. Prop., 20) Mirkpvic, D., Gro¨bner, J., Schmid-Fetzer, R., Fabrichnaya, O., Lukas, H.L., “Experimental Study and Thermodynamic Re-Assessment of the Al-B System,” J. Alloys Compd., 384, 168–174 (2004) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calculation, #, 26) Schuster, J.C., “Al-Mo (Aluminum-Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 30.12123.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 61) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Boron – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Hans Leo Lukas
Introduction The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Luk], which is updated by the present work. The only experimental ternary information on this system is the existence of a ternary phase Al3SiB48 [1969Lam, 1972Lam] (Table 1). There is evidence that this is a high temperature phase, which is easily quenchable to room temperature by rapid or moderate cooling. The previously reported βAlB12 phase is assumed to be identical with this phase, the structure of which is closely related to that of pure B. A phase with the same structure also appears as Al3B48C2 [1964Mat]. [1981Por] estimated interaction parameters of B and Si in the (Al) solid solution. [1982Doe] calculated the whole ternary phase diagram based on assessments of the binary subsystems. Since that time newer assessments for all three binary systems are available. Especially for the Al-B system good calorimetrically measured enthalpy data are available for the intermediate phases [2001Mes] increasing the quality of the new assessment of [2004Mir]. The diagrams given here are from calculations using the dataset of [2005Gro], which is composed of the binary assessments of Al-B [2004Mir], Al-Si [1996Gro] and B-Si [1998Fri]. They differ markedly from those of [1982Doe] although the reason is only a slightly more negative Gibbs energy of mixtures of the same overall compositions of AlB12 + Si compared to those of liquid-Al + B-Si phases, whereas in the dataset of [1982Doe] the contrary was assumed. The tie-lines between AlB12 and Si-Al melt, as calculated by [2005Gro], agree with experimental findings of [2005Yos1] [1999Wan, 2000Wan, 2003Nog, 2005Hui] investigated the influence of B-additions in the magnitude of 100 to 1000 mass-ppm to the morphology of the eutectic microstructure of hypoeutectic Al-Si alloys. The two newest ones did not find significant effects of B.
Binary Systems For the thermodynamic calculations the following assessments were merged by [2005Gro] to a ternary dataset: Al-B [2004Mir], Al-Si [1996Gro] and B-Si [1998Fri]. The phase diagrams calculated from these binary datasets agree well with those compiled in [Mas2], except Al-B, where [Mas2] assumes an additional phase AlB10, which was proved to be stable only as carbon containing ternary phase [1964Mat, 1994Dus].
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Solid Phases One ternary phase, Al3SiB48, was reported by [1969Lam]. The authors got either this phase or αAlB12 on heating samples between 1400 and 1600˚C, indicating the upper temperature limit to be in this range. This phase was formerly described as binary phase βAlB12, however, [1964Mat] proved, that it is stable only with ternary additions, and gave a formula Al3B48C2. [1969Lam] showed, that carbon can be replaced in this phase by silicon, but not by germanium. The only stable binary Al-B phases are AlB2 and αAlB12 [1964Mat, 1994Dus]. Three more stable binary intermediate phases exist in the B-Si system, see Table 2. The crystal structures of the unary phases of pure boron are differently described in literature [V-C2]. The most likely interpretation is the existence of a phase hR36 below ca. 1100˚C and a high temperature phase B described by the Pearson symbols hR105, hR111 or hR141 regarding the number of atomic positions per rhombohedral unit cell. The hexagonal lattice parameters agree fairly well, a = 1092–1096, b = 2381–2389 pm. Regarding the partial occupations of the different structure descriptions, the number of atoms per rhombohedral unit cell is nearly equal for the different descriptions, 105–107. All well characterized solid phases are summarized in Table 2. Some of the crystal structures contain sites occupied only partially, there two Pearson symbols are given, one counting the sites, the other one counting the atoms per unit cell.
Invariant Equilibria Figure 1 shows the reaction scheme derived from calculations using the dataset of [2005Gro]. The calculated temperatures and phase compositions are given in Table 3. The scheme and especially the temperatures have to be taken as tentative. The temperature limits of Al3SiB48 are not known, there is only evidence that it forms between 1400 and 1600˚C and that it transforms at lower temperatures to another phase for which no details are known. Phase Al3SiB48 (τ) was added to the dataset of [2005Gro] as stoichiometric phase with linearly temperature dependent Gibbs energy forming peritectoidally at 1500˚C and neglecting the possible transformation. Below 1500˚C this phase appears everywhere in equilibrium with the three phases AlB12, SiB36 and SiB6.
Liquidus, Solidus and Solvus Surfaces The liquidus surface resulting from the calculation is shown in Fig. 2. It is remarkably different from that calculated from the dataset of [1982Doe] due to a newer assessment of the Al-B system [2004Mir] based on precise calorimetric measurements of the enthalpies of the AlB intermediate phases [2001Mes]. The isotherms are less dependent on the thermodynamic quantities than the curves of double saturation and therefore they should be taken as reasonably well defined.
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Isothermal Sections The calculated isothermal sections at 1300 and 500˚C are shown in Figs. 3 and 4 respectively.
Temperature – Composition Sections Figure 5 shows a partial calculated Temperature - Composition section at 0.1 at.% B. Note that here the field L+AlB2+AlB12 is so narrow, that it is degenerated to a single line.
Thermodynamics There are no experimental data on ternary thermodynamic quantities. Some binary data, however, are very important for ternary calculations: the calorimetrically measured enthalpies of formation of AlB2 and AlB12 [2001Mes] as well as the enthalpy change during the invariant reaction L + AlB12 Ð AlB2 [2004Mir] A thermodynamic dataset was assessed by [2005Gro] by merging the binary datasets of [2004Mir] (Al-B), [1996Gro] (Al-Si) and [1998Fri] (B-Si) without using ternary terms.
Notes on Materials Properties and Applications Wang et al. [1999Wan, 2000Wan] reported a good refining effect of B on hypoeutectic Al-Si alloys, but [2003Nog] and [2005Hui] could not find a significant effect of B on the microstructure of the eutectic in Al-Si alloys. [2005Yos2] demonstrated, that “temperature gradient zone melting” (TGZM) with a thin Al-Si-melt may help to remove B-impurities from semiconductor grade Si material.
. Table 1 Investigations of the Al-B-Si Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1969Lam], Single-crystal X-ray diffraction [1972Lam]
Identification of the ternary phase τ, Al3SiB48
[2005Yos1] Distribution coefficient of B between Al-Si melt and solid Si
1100–1300˚C
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. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
pure Al at 25˚C [V-C2]
(βB)(h) < 2092
hR315(a,b)
a = 1096 c = 2389
hR333(a), hR315(b)
a = 1093 c = 2381 a = 1093 c = 2382
three different complete structure descriptions, cited including atomic positions by [V-C2]
hR423(a), hR321(b) R3m βB
Lattice Parameters [pm]
Comments/References
(αB)(r)
hR36(c) R3m αB
a = 491.1 c = 1257.3
[V-C2] presumably metastable phase
(Si) < 1414
cF8 Fd3m C diamond
a = 543.06
at 25˚C, [V-C2]
γAlB12
oP396(a), oP377(b) P212121 γAlB12
a = 1662.3 b = 1754.0 c = 1018.0
[1983Hig] Metastable or metal impurity stabilized [1994Dus]
αAlB12 ≲ 2094
tP216(a), tP189(b) P41212 or P43212 αAlB12
a = 1016.1 c = 1428.3
[1977Hig]
AlB2 972 - ~213
hP3(a), hP2.9(b) P6/mmm AlB2
a = 300.9 c = 326.2
[1964Mat]
SiB36 < 2137
hR339(a), hR314(b)
a = 1101 [V-C2] c = 2390 a = 1099 to 1113 93.3 to 97 at.% B [1984Ole] c = 2383 to 2400
SiB6 < 1850
oP340(a), oP314.5(b) a = 1439.7 b = 1831.8 Pnnm c = 991.1 B6Si
R3m B36Si
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
B3Si < 1270
hR45 R3m B4C
* τ, Al3SiB48
tP68(a), tP52(b) P42/nnm B25C ?
Lattice Parameters [pm]
Comments/References
a = 632 to 635 [1984Ole] c = 1269 to 1275
(c)
a = 891 c = 505
[1969Lam] metastable at room temperature? Pearson symbol (a) and prototype [V-C2]
Pearson symbol showing number of sites per unit cell (for hR hexagonal unit cells) Pearson symbol showing number of atoms per unit cell (for hR hexagonal unit cells) (c) Number of atoms in Pearson symbol is valid for hexagonal unit cell. (a)
(b)
. Table 3 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
L + (βB) Ð αAlB12 + SiB36
2013
U1
L
L + SiB36 Ð αAlB12 + SiB6
αAlB12 + SiB36 + SiB6 Ð τ
L + SiB6 Ð αAlB12 + (Si)
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≈1500
1373
U2
P1
U3
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Al
B
Si
2.6
91.2
6.2
(B)
0.0
97.9
2.1
αAlB12
7.7
92.3
0.0
SiB36
0.0
96.7
3.3
L
1.9
63.0
35.1
SiB36
0.0
94.1
5.9
αAlB12
7.7
92.3
0.0
SiB6
0.0
86.2
13.8
αAlB12
7.7
92.3
0.0
SiB36
0.0
94.1
5.9
SiB6
0.0
86.2
13.8
τ
5.8
92.3
1.9
L
2.7
8.2
89.1
SiB6
0.0
85.5
14.5
αAlB12
7.7
92.3
0.0
(Si)
0.0
1.0
99.0
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. Table 3 (continued) Composition (at.%) Reaction
T [˚C]
Type
Phase
SiB6 + (Si) Ð SiB3, αAlB12
1270
D1
SiB6
L + αAlB12 Ð AlB2 + (Si)
677
U4
B
Si
0.0
85.5
14.5
(Si)
0.0
0.7
99.3
SiB3
0.0
73.8
26.2
αAlB12
7.7
92.3
0.0
81.4
0.1
18.5
L αAlB12
L Ð (Al) + (Si), AlB2
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7.7
92.3
0.0
AlB2
33.3
66.7
0.0
(Si)
0.0
0.01
99.99
L
87.92
0.03
12.05
(Al)
98.5
0.0
1.5
(Si)
0.0
0.0
100.0
AlB2
33.3
66.7
0.0
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. Fig. 1 Al-B-Si. Reaction scheme
Al–B–Si
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. Fig. 2 Al-B-Si. Liquidus surface projection
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. Fig. 3 Al-B-Si. Isothermal section at 1300˚C
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. Fig. 4 Al-B-Si. Isothermal section at 500˚C
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. Fig. 5 Al-B-Si. Partial temperature-composition section at 0.2 at.% B
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References [1964Mat] [1969Lam]
[1972Lam]
[1977Hig] [1981Por]
[1982Doe]
[1983Hig] [1984Ole] [1990Luk]
[1994Dus] [1996Gro] [1998Fri]
[1999Wan] [2001Mes]
[2000Wan] [2003Nog] [2004Mir]
[2005Gro]
[2005Hui] [2005Yos1]
[2005Yos2]
Matkovich, V.I., Economy, J., Giese, R.F., “Presence of C in Al Borides”, J. Am. Ceram. Soc., 86, 2337–2340 (1964) (Crys. Structure, Experimental, 14) Lamikhov, L.K., Neronov, V.A., Rechkin, V.N., Samsonova, T.I., “On the β-AlB12 Phase in the Si-Al-B System” in org. Mater. (Engl. Trans.), 5, 1034–1036 (1969), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 5(7), 1214–1217 (1969) (Experimental, Morphology, 12) Lamikhov, L.K., Neronov, V.A., Rechkin, V.N., Samsonova, T.I., “β-AlB12 Phase in the Si-Al-B System” (in Russian), in “Metalloterm. Metody Poluch. Soedin. Splavov”, Kornilov, A.A. (Ed.), Nauka, Sib. Otdel. Novosibirsk, 52–57 (1972) (Review, Crys. Structure, 12) Higashi, I., Sakurai, T., Atoda, T., “Crystal Structure of α-AlB12”, J. Solid state Chem., 20, 67–77 (1977), (Experimental, Crys. Structure, 21) Portnoy, K.I., Bogdanov, V.I., Mikhailov, A.V., Fuks. D.L., “Interaction Parameters in Interstitial Solid Solutions Based on Aluminium”, Russ. J. Phys. Chem. (Engl. Transl.), 55(4), 583–584 (1981), translated from Zh. Fiz. Khim., (55), 1041–1043 (1981) (Experimental, 10) Doerner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, Uni. Stuttgart, Inst. Metallkunde, (1982) (Calculation, Thermodyn., 126) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Experimental, Crys. Structure, 21) Olesinski, R.W., Abbaschian, G.I., “The B-Si System”, Bull. Alloy Phase Diagrams, 5, 478–484 (1984) (Review, Phase Diagram, 51) Lukas, H.L., “Aluminium - Boron - Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.20852.1.20, (1990) (Crys. Structure, Phase Diagram, Assessment, 6) Duschanek, H., Rogl, P., “The Al-B (Aluminum-Boron) System”, J. Phase Equilib., 15, 543–552 (1994) (Assessment, Crys. Structure, Phase Relations, Review Thermodyn., 78) Gro¨bner, J., Lukas, H.L., Aldinger, F., “Thermodynamic Calculation of the Al-Si-C System”, Calphad, 20, 247–254 (1996) (Calculation, Phase Diagram, Thermodyn., 37) Fries, S.G., Lukas, H.L., “B-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Vol. 2”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Luxembourg, 126–128 (1998) (Calculation, Phase Diagram, Thermodyn., 0) Wang, Li., Bian, Xiufang, Sun, Yimin, “Refining Effect of Boron on Hypoeutectic Al-Si Alloys”, Chin. J. Nonfer. Met., 9(4), 714–718 (1999) (Experimental, Morphology, 13) Meschel, S.V., Kleppa, O.J., “Thermochemistry of Alloys of Transition Metals and Lanthanide Metals with some IIIB and IVB Elements in the Periodic Table”, J. Alloys Compd., 321, 183–200 (2001) (Experimental, Thermodyn., 82) Wang, Li, Bian, Xiufang, “Refining Effect of Boron on Hypoeutectic Al-Si Alloys”, J. Mater. Sci. Technol., 16(5), 517–520 (2000) (Experimental, Morphology, 7) Nogita, K., Dahle, A.K., “Effects of Boron on Eutectic Modification of Hypoeutectic Al-Si Alloys”, Scr. Mater., 48(3), 307–313 (2003) (Kinetics, Morphology, 16) Mirkovic, D., Gro¨bner, J., Schmid-Fetzer, R., Fabrichnaya, O., Lukas, H.L., “Experimental Study and Thermodynamic Re-assessment of the Al-B System”, J. Alloys Comp., 384, 168–174, (2004) (Experimental, Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 26) Groebner, J., Mirkovic, D., Schmid-Fetzer, R., “Thermodynamic Aspects of Grain Refinement of Al-Si Alloys Using Ti and B”, Mater. Sci. Eng. A, 395(1-2), 10–21 (2005) (Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 56) Huiyuan, G., Yanxiang, L., Xiang, C., Xue, W., “Effects of Boron on Eutectic Solidification in Hypoeutectic Al-Si Alloys”, Scr. Mater., 53(1), 69–73 (2005) (Experimental, Morphology, 8) Yoshikawa, T., Morita, K., “Thermodynamic Properties of B in Molten Si and Phase Relations in the Si-Al-B System”, Metal. Trans., 46(6), 1335–1340 (2005) (Experimental, Phase Relations, Thermodyn., 16) Yoshikawa, T., Morita, K., “Removal of B from Si by Solidification Refining with Si-Al Melts”, Metal. Trans. B, 36(6), 731–736 (2005) (Experimental, Phase Relations, 18)
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Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Carbon – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Lesley Cornish, Gabriele Cacciamani, Damian M. Cupid, Jozefien De Keyzer
Introduction Al-C-Ti alloys are of interest for many different reasons, mainly as alloys in themselves, dispersion-hardened alloys with TiC, and also as grain refiners for commercial Al-based alloys, giving improved properties. Alloys are based on Ti-6Al/TiC [2002Zha2], which has high strength, low density and good elastic modulus. Ti3Al and γαffl (TiAl) phases are being developed for lightweight structural applications, especially at elevated temperatures. The system is also important for composite manufacture; using TiC in (Al), and also using the ternary carbides. Ti2AlC is recognized as a “machinable ceramic”, its layered structure produces a combination of metallic and ceramic properties. Additionally, there is interest for coatings: Ti-Al-C PVD coatings could compete with Ti-Al-N, and in a different application, TiC has potential for a barrier diffusion layer for silicon semiconductor devices. Phase boundary information on this system comes from the work of [1954Thy, 1980Sch, 1987Zak], [1989Cam1, 1989Cam2, 1991Cam1, 1991Cam2, 1992Pie, 1994Pie] mostly using optical and electron microscopy together with and X-ray diffraction. Other publications [1963Jei, 1964Jei, 1964Now, 1976Ivc1, 1976Ivc2, 1980Pea] covered X-ray diffraction studies, crystal structures, lattice parameters and physical properties of the two ternary carbides Ti2AlC and Ti3AlC, with general agreement. One of the first reviews was by [1965Mol], followed by [1983Kub, 1983Sri]. The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Hay], which is updated by the present work. Recent reviews by [1994Pie, 2000Ban, 2006Rag] (where [1994Pie] also undertook experimental work) mainly include the experimental information obtained after the original MSIT review of [1990Hay]: [1990Via, 1991Cam1, 1992Pie, 1994Pie, 1994Zha]. Information of some other recent articles that were not considered there ([1990Hay]) are included in this review [1991Cam2, 1995Via, 2000Ria, 2000Tze, 2000Van, 2003Ge1, 2005Zho2]. An overview of the investigations considering phase equilibria, solid phases and thermodynamics of the system is given in Table 1.
Binary Systems For the Al-Ti system, the MSIT evaluation from [2004Sch] is available. More recently, a new review of [2006Sch] and a thermodynamic assessment of [2007Wit] appeared which show differences in two main regions. [2006Sch] and [2007Wit] consider the equilibria between (αTi), (βTi) and Ti3Al. Starting from the same experimental literature information, [2004Sch] concluded that Ti3Al decomposes congruently and there is no invariant equilibrium involving Landolt‐Bo¨rnstein New Series IV/11E1
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these three phases, while [2006Sch] and [2007Wit] concluded that there are two peritectoid reactions between these three phases and a small two-phase field between (βTi) and Ti3Al in the 1150˚C-1200˚C temperature range. According to [2004Sch], the different phase equilibria may be related to very small differences in Gibbs energy, and consequently in driving forces, to give the possible phase transitions in that range. Thus, it is very difficult to decide which interpretation of the experimental data is actually correct. The second difference considers the equilibria between γ(TiAl), TiAl2 and one-dimensional antiphase domain structures or “long period structures (LP)” stable at high temperature in the range between 65 and 75 at.% Al. In this case, the assessments by [2006Sch] and [2007Wit] are based on more recent experimental data and give a more complete discussion of the stable equilibria in this region. Thus the phase diagram calculated by [2007Wit] is accepted in the present evaluation since it is based on experimental data available, the majority of which were assessed in [2006Sch], being complemented by own key experiments. The adopted Al-Ti phase diagram is shown in Fig. 1 from [2007Wit]. The Al-C phase diagram is accepted from the review of [2004Per]. The C-Ti phase diagram (Fig. 2) is taken from the recent thermodynamic calculation of [2003Fri]. It leads to a better fit of high temperature heat capacity measurements than all previous studies. The phase diagram of [2003Fri] is reproduced with the same accuracy as given by [1999Dum] and gives a better description of the experimental data than previous assessments. However, it should be noticed that the liquidus lines are not supported by any experimental data.
Solid Phases The solid phases are given in Table 2. The different morphologies of TiAl3 are deemed to be related to the different morphologies observed at different cooling rates and Ti contents [1982Bla, 1993Sve]. Similarly, TiAl2 has been reported to have metastable structures in ascast alloys [2001Bra]. The Ti2Al5 phase [2004Sch] is interpreted here as a one dimensional antiphase domain structure or “long period structure (LP)” stable at high temperature in the range between 65 and 75 at.% Al. The Ti5Al11 phase has been included in Table 2, although it is not accepted here as a separate phase, and is presumed to be related to ζ, Ti2Al5, with which it shares similar lattice parameters. According to [1989Cam1, 1989Cam2], carbon solubilities in the binary aluminides are small (~1 at.% C) and the absence of fine scale carbide precipitates after ageing the alloys with 1 at.% C for 15 days at 750˚C and 450˚C respectively shows the absence of a significant decrease of C solubility in (αTi) or Ti3Al [1989Cam1, 1989Cam2]. Ternary carbides are stated to form above 1250˚C by peritectic reactions from TiC1–x and liquid. All workers have reported minimal solubility of aluminium in TiC1–x [1984Ker, 1980Sch, 1989Cam1, 1989Cam2, 1990Cam, 1991Cam1, 1991Cam2, 1994Pie, 2000Ria]. [1980Sch] confirmed the existence of the hexagonal ternary H-phase Ti2AlC and of the cubic Perovskite-(P-) phase Ti3AlC reported by earlier workers [1963Jei, 1964Jei, 1964Now, 1976Ivc1, 1976Ivc2, 1980Pea]. A cubic phase that was earlier noted in a Ti-10Al-1C (mass%) alloy [1954Thy] probably was the cubic Perovskite-phase Ti3AlC; this was confirmed by [2000Ria], although [1994Pie] reported not being able to find a match. [1994Pie] found a third ternary phase, N, Ti3AlC2. It is presumed that the reason for this phase not being observed earlier was that longer sintering times and higher pressures [2006Li] were needed. DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
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The ternary carbides have less C than their stoichiometric compositions [1980Sch, 1992Pie, 1994Pie, 2000Tze] and the range for N is at least Ti3Al1.1C1.8 [2000Tze] to Ti3Al1.1C2.3 [2005Kho]. The relationship of TiC1–x and Ti2AlC1–x to Al4C3 was discussed by [1982Now]. The unusual diffusion path reported by [1995Via] between (Al) and TiC could possibly have been due to the then unrealized N phase.
Quasibinary Systems The section between γ (TiAl) and TiC1–x by [1987Zak] is not a true quasibinary since both peritectic reactions have a (non-linear) three-phase field.
Invariant Equilibria The invariant reactions are not fully derived yet, and more work needs to be done. The agreed invariant equilibria of [1991Cam1, 1994Pie] are given in Table 3. Ti2AlC melts incongruently at 1625±10˚C; Ti3AlC also melts incongruently at 1580±10˚C, whereas Ti3AlC2 decomposes in the solid state [1994Pie]. [2004Hwa] observed that ternary carbides formed from peritectic reaction with TiC1–x ; this confirmed the reaction: L + TiC1–x Ð H of [1987Zak], which [1994Pie] reported having a maximum at 1625±10˚C. [1992Che] deduced that cubic Ti3AlC is a metastable transition phase which is ultimately replaced by the more stable H (Ti2AlC) phase, which agrees with the reactions of [1994Pie]. According to [2005Zho2], the lower stability limit of Ti3AlC2 is ~1300˚C (decomposing to give TiC1–x and Al vapor) and [2000Tze] reports the upper stability limit as at least 1400˚C, if not 1450˚C. The most Al rich equilibrium in Table 3 is from [1990Via, 1993Sve], where the latter deduced that the formation kinetics of Al4C3 were slow, and the reaction agreed with [1998Fra]. [1990Via] gave alternative Al rich reactions: L + Al4C3 + TiAl3 Ð (Al) or L + TiAl3 Ð Al4C3+ (Al), both being very near to the melting point of aluminium.
Liquidus, Solidus and Solvus Surfaces [1991Cam1] derived a “semi-schematic” liquidus surface. The volume fraction of H was less than the projection suggested; this was thought to be due Al substituting for C. A schematic liquidus surface and partial reaction scheme was also derived by [1994Pie]. They are not presented here, since they contain extra Al-Ti phases which are not accepted in this review. However, the liquidus surface is dominated by TiC1–x , then graphite. The reactions forming the other phases all lie very close to the Al-Ti binary. [1990Via] gave an experimental liquidus projection in the Al rich corner, whereas [1998Fra] produced one from thermodynamic calculations. The major difference was that [1990Via] gave the lowest temperature reaction, which was dependent on cooling conditions. [1993Sve] also attempted to rationalize the different morphologies of TiAl3 observed with composition.
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Isothermal Sections [2003Ge1] drew an isothermal section at 1300˚C showing the third ternary phase, N (Ti3AlC2), which was involved in the three-phase triangles: (N + H + TiC1–x), (N + TiC1–x + TiAl3) and (N + H + Ti5Al11), but the last phase (as well as its Al-Ti phase diagram) is not accepted here. Additionally, the phase compositions did not agree with the boundaries due to Al loss. Similarly, the isothermal section at 1300˚C of [1992Pie, 1994Pie] is not presented. The 1100˚C isothermal section of [1994Zha] is not presented because it contains Ti2Al5 which is not accepted here, although it mainly agrees with the other sections presented. [2000Ban] gave isothermal sections at 1300, 1100 and 1000˚C, but these are not reproduced because they are incompatible with the Al-Ti binary accepted here, and show a radical change in the positions of the P and H ternary phases over the temperature ranges (especially 1100˚C). [1993Tia] established a metastable diagram between 900 and 1050˚C. The isothermal section at 1250˚C was drawn by [1991Cam1], and is shown in Fig. 3 after modifications to be consistent with the accepted Al-Ti binary. [2000Ria] modified the 1050˚C isothermal section of [1990Cam], with additional results that retained the same phase fields, but changed the shape of some of them, notably ((αTi) + P + TiC1–x). This is redrawn in Fig. 4 to be consistent with the Al-Ti binary system accepted here, the positions of the P and H, and keeping their alloys in the respective phase fields. The redrawn P and H phases in Figs. 4 and 1 are also consistent with those of [1990Via], but the isothermal sections of [1990Via] are not given here because negligible solubility is given for all binary phases. The results of [1954Thy] and [1989Cam1, 1989Cam2, 1991Cam1] indicate that both (αTi) and (βTi) solid solutions are in equilibrium with TiC, thus generating the tie-triangles ((αTi) + (βTi) + TiC) and ((αTi) + TiC + P) (Fig. 5) at 1000˚C. This is taken rather than the Ti rich corner of [1980Sch], which showed that at 1000˚C, (βTi) solid solution is in equilibrium with the Perovskite carbide phase Ti3AlC (P) generating the two tie-triangles ((βTi) + TiC + P) and ((αTi) + (βTi) + P). This decision was made because [1980Sch] showed those relationships as dotted, and did not give the sample compositions. Their drawn phase relationships also contradict their statement that (βTi) was not retained in the ternary. Additionally, seven experimental alloys of [1994Zha] agree with the findings and interpretation of all previous workers, except for [1980Sch]. There is agreement that γαffl (TiAl) is in equilibrium with Ti2AlC at 1000˚C [1980Sch, 1987Zak, 1989Cam1, 1991Cam1]. The isothermal section at 750˚C (Fig. 6) is taken from [2000Ria], but redrawn to be consistent with the Al-Ti binary accepted here (Fig. 1), the positions of the P and H, as well as the low carbon solubilities reported in the Al-Ti phases, except for (αTi) [1954Thy, 1980Sch, 1987Zak, 1989Cam1, 1989Cam2, 1991Cam1, 1991Cam2, 1992Pie, 1994Zha]. This is consistent with the isothermal section at 1000˚C in terms of reduced ternary element solubilities in the phases at lower temperatures. The section of [2000Ria] is preferred to that of [1991Cam1] because it was derived using alloys further into the ternary system (up to 15 at.% C) [2000Ria], instead of only up to 3 at.% C [1991Cam1]. The redrawing of the diagram brings the only contentious alloy of [1991Cam1] to a position nearly on the boundary between the ((αTi) + TiAl3 + P) and ((αTi) + P + TiC1–x), which could explain why P was not found. Although it appears that TiC1–x [1991Cam1] was re-precipitated within the grains, it would have been the primary phase [1994Pie] and [1995Via] has shown that it takes a long time to disappear. In the isothermal sections presented, the compositions of the P and H phases were modified to agree with those of the 1000˚C isothermal section since this was the widest ranges, and DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
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otherwise the other sections had compositions which fell outside this range, especially those of [2000Ria] which used higher Ti content alloys. Experimental work on liquid (Al) by [1993Sve] showed that there were three-phase regions: ((Al) + TiAl3 + TiC1–x) and ((Al) + Al4C3 + TiC1–x). Several calculations have been undertaken on the Al rich corner: [1998Fra] showing the equilibrium between (Al) with TiC1–x and TiAl3 and [2000Van] which confirmed [1998Fra], while studying grain refiners.
Temperature – Composition Sections Isopleths for the concentration range 0–1 mass% C at 2, 4, 6, 8 and 10 mass% Al were determined by [1954Thy] and boundaries between the following phase fields were located: (αTi), (βTi), ((αTi) + (βTi)), ((βTi) + TiC1–x), ((αTi) + (βTi) + TiC1–x) and a field containing more than one phase one of which is TiC1–x. The solubility of C in (αTi) was found to increase with increasing Al contents rising to over 1 mass% C at 1150˚C and 10 mass% Al. These are reproduced in Figs. 7–11) from [1990Hay] who redrew them slightly (but without compromising the original experimental results) to be consistent with isothermal sections, and they are redrawn slightly to be consistent with the adopted Al-Ti binary. A section at 0.2 at.% C [1997Li, 1999Li] agrees with the accepted binary. [1987Zak] derived a γ(TiAl)TiC1–x section (Fig. 12).
Thermodynamics [1986Ban] determined the Gibbs free energy of formation of the compounds Al4C3 and TiC as a function of temperature in the range 300–1800 K. The same was calculated by [1991Rap] in order to demonstrate that the formation of Al4C3 is metastable with respect to TiC and it is due to kinetic reasons. Thermodynamics and kinetics of formation of TiC by a reaction between Al-Ti melts and a carbonaceous gas have been studied by [1991Sah]. More recently [2006Ton] used thermodynamic analysis to show that it is feasible to produce TiC-Al (Ti) nanocomposite powders using Al and Ti as starting materials and CH4 as reacting gas in a thermal plasma environment. Based on thermodynamic results previously available in literature, [1991Yok] presented a three-dimensional potential phase diagram for the Al-C-Ti system at 700˚C. [1993Jar] investigated thermodynamics and kinetics of the reactions taking place during infiltration of graphite fibres by molten Al-Ti alloys. A temperature range for safe processing of Al-TiC composites has been suggested by [1993Mit] based on thermodynamic considerations. [1993Sve] determined by a theoretical thermodynamic analysis the C and Ti concentration in Al-rich liquid in equilibrium with solid phases at temperatures between 700 and 1100˚C. A thermodynamic discussion of the solid/liquid phase equilibria in the Al rich corner has been reported also by [1998Fra]. Limiting values of thermochemical data of formation of the Al-C-Ti ternary compounds are reported in [1994Pie], based on binary and ternary Al-C-Ti equilibria and thermochemical data for various reactions previously studied in literature. The authors concluded that the dataset so obtained was sufficient to calculate tie lines in agreement with known phase equilibria but a full thermodynamic assessment of the system is needed. Landolt‐Bo¨rnstein New Series IV/11E1
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Thermodynamics of formation of TiC in DC-casting and melt spun Al-C-Ti alloys has been discussed by [1994Ton]. Thermodynamics and kinetics of formation of TiC in Al-C-Ti have also been investigated by [2000Zha2]. Thermodynamics and kinetics of precipitation of TiC in Al-C-Ti melts at 1100˚C have been investigated by [1999Jar]. Low temperature heat capacity of Ti3Al1.1C1.8 has been measured by [1999Ho]. Heat capacity, thermal expansion coefficients and thermal conductivity of the Ti2AlC phase have been measured by [2002Bar] at room to high temperature: these are reported in Table 4. Low temperature heat capacity of the same phase has been measured by [2006Dru].
Notes on Materials Properties and Applications The work on materials properties is summarized in Table 5. The system is of interest because of the effect of different additions to the grain refining of many commercial Al-based alloys [1949Cib], with the benefits of improved mechanical properties, reduced ingot cracking, improved feeding and increased casting speeds, improved homogeneity and reduced porosity and better deformation behavior. Grain refiners based on Al-C-Ti can be added in master alloy form, and are an alternative to conventional Al-B-Ti grain refiners [1994May] and are less susceptible to agglomeration. [1976Jon] calculated solubility limits of liquid aluminium in different systems, including Al-C-Ti. Epitaxy was found to be important; thus TiC was seen as being more effective than TiAl3, having more epitaxial faces [1972Cis]. However, TiC is surrounded by Al4C3, then Ti3AlC, which reduces the grain refinement effectiveness, and so heating liberates the beneficial TiC [1986Ban]. Several studies were made on the formation of the grain refining precipitates [1977Ivc, 1989Cam1, 1991Fro, 1991Sah, 2000Bri1, 2000Zha1, 2002Zha1, 2003Zha2], to improve the methods [2000Zha1, 2005Zha, 2005Kum] and to assess the efficiency [2000Bri2, 2000Rom, 2000Sch, 2000Van, 2002Gaz, 2002Tro, 2002Xia, 2002Liu, 2003Zha2]. Aluminium alloys grain refinement has also been achieved by synthesizing Al-Ti-C launders and moulds in an electromagnetic field [2005Xia]. There are also several alloys based on Ti-6Al/TiC1–x [2002Zha2] which have high strength, low density and good elastic modulus. Aluminium dispersion-hardened alloys are also important and different alloys in different conditions have been studied: up to 50 mass% TiC1–x [1986Bat]. The Ti3Al and γ (TiAl) phases are being developed for lightweight structural applications, especially at elevated temperatures. γ (TiAl) is disadvantaged by low temperature ductility and toughness, and these are being improved by precipitation [1991Mab]. Composites are also being developed by precipitating Ti3AlC in HIPped γαffl (TiAl) alloys [1992Che, 1994Whi], although the precipitate long-term stability was questioned [1992Che]. Optimization of the lamellar microstructural has been attempted [2001Qin]. Improvements have also been made by indirect-extruding TiAl-Mn-Mo-C alloys to obtain a fine microstructure, and rheocasting TiAl alloys and composites [1997Ich]. The system is also important for composite manufacture; composites have been made by infiltrating graphite fibres with molten Al-Ti, using the phase diagram [1993Jar]. In-situ internal carburization has also been carried out to fabricate the composites [1991Jar, 1999Bir, 2001Zha]. [1995Mor] made in-situ TiC1–x particulate reinforced aluminum composites by reacting graphite particles in a liquid Al-Ti alloy and showed that the Young’s modulus and tensile strength increased proportionally to the TiC1–x volume fraction. Ductile Al/TiC DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
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MMCs (metal matrix composites) have also been made [1993Mit, 2003Arp] and the effects of the TiC1–x dendrite spacing were studied [2002Ria]. Bulk [2004Hwa, 2004Wan3, 2005Jia, 2005Li, 2006Xia] and porous [1992Vol] Al-C-Ti cermets have also been manufactured using the SHS (self-propagating high-temperature synthesis) sintering technique [1992Vol, 1997Lee, 1997Ye, 2002Mei, 2006Xia] and phase evolution studied [2005Rog]. Variations of thermal explosion/quick pressure (TE/QP) methods have also been used [2001Yan, 2005Ma], and “hot shock” [1998Tan]. Combustion synthesis, also known as self-propagating high-temperature synthesis (SHS), offers advantages of low processing cost, energy-efficiency, and high production rate, and has been used to manufacture TiC1–x [1993Cho, 2006Xia]. Addition of synthesized master alloys to molten aluminium has also been achieved [2006Sel], and a dipping exothermic reaction process (DERP) has also been used [2003Son, 2004Son2]. [1996Tom] produced ternary composites by combustion synthesis with titanium, aluminium, and graphite powders. At low C-contents, binary intermetallics were formed, whereas Ti2AlC and TiC1–x were formed at higher C-contents. Other elements have also been added to the materials [2003Yan]. A useful review of these and other composites is [2000Tjo]. Some of the composites are being developed for cutting tools [1967Hod, 1976Ivc1], using the intermetallic phases, and TiC1–x shows the most promise. An extension of this are complex carbides (or nitrides) M2AlC (or N), where M = Cr, Ti, V, Nb or Ta [1996Cha]. Also, the system has been considered for coatings on Ti-based metal matrix composites containing alumina fibres [1994Din]. Studies have been made of solid state reactions at graphite/TiAl or Ti3Al interfaces, showing that whilst graphite moulds are safe for these alloys, carbon fibres would have to be coated to prevent Ti diffusion [1995Via]. Ti2AlC is recognized as a “machinable ceramic” (layered structure) with good mechanical properties and low thermal expansion, thus giving potential high temperature applications (furnace apparatus). It has been manufactured by self-propagating high-temperature synthesis (SHS) [2001Lop], combustion synthesis [2004Kex], “thermal explosion” (a form of SHS) [2002Kho1, 2002Kho2], hot pressing [2004Hon, 2005Pen], spark plasma sintering [2003Zho, 2004Mei, 2005Zho1], two-stage pressureless sintering followed by hot pressing [2005Wan1]. Similarly, interest has also been shown in Ti3AlC [2001Lop] and Ti3AlC2 [2000Tze, 2001Lop, 2003Ge1, 2003Ge2, 2004Son1, 2005Kho, 2005Zho2, 2006Li], with similar manufacturing processes. Ti3AlC pins twins in γ (TiAl) [1997Chr]. Ti-Al-C PVD coatings have been manufactured as possible competitors to Al-N-Ti coatings [1991Kno] and for coatings for steels [2005Wal]. Another coating application is as a diffusion barrier for materials in contact structures for silicon semiconductor devices. Aluminium can be used (rather than the more effective and expensive PtSi), but needs a barrier diffusion layer, and this can be TiC1–x [1982Wit].
Miscellaneous The formation mechanism for TiC1–x from the ball-milled elemental components was studied by [2003Wan2]. Various reaction mechanisms for combustion reactions to yield TiC1–x [1993Cho] (with results of the first TiC formation supported by [1995Via]), Ti3AlC2 [2002Ge], and TiC / γ(TiAl) composites [2002Hwa] have been proposed. [1993Cho] studied the effect of Al addition (0-40 mass%) on the combustion reaction between Ti and C to form TiC1–x using combustion wave velocity data and differential thermal analysis. Their results Landolt‐Bo¨rnstein New Series IV/11E1
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indicate the reaction between titanium and aluminium initiates the reaction between titanium and carbon through a series of reaction steps. [2002Hwa] also used combustion synthesis to create TiC1–x/Al composites, but reported the presence of ternary carbides in the final microstructure. In-situ processing of aluminium matrix metal matrix composites based on TiC has been investigated [1995Mor, 1996Tom, 1993Ber]. [1996Mat] performed electronic structure calculations for C-containing γ (TiAl) and showed that C substitutes Al sites. [1996Ily] undertook electronic structure calculations on Ti1–xAlxC (x = 0.17), and showed that this compound has a bandgap of 2.18 eV. Ab-initio electronic structure calculations were performed on C-containing γ (TiAl) [1996Mat] and on theoretical Ti1–xAlxC compounds (NaCl lattice where Al and Ti occupy the same lattice sites) by [1996Ily, 1997Ily1, 1997Ily2, 1998Mat]. The wetting angle of molten Al on TiC1–x was measured using the sessile drop method by [1966Yas] at 117˚. [1996Sob] used the same method to investigate the wettability of Al-rich Al-Ti alloys (Al-0.5Ti to Al-10Ti (mass%)) and showed that Ti additions lead to better wettability of graphite because of reactions between titanium and carbon. Gibbs free energy calculations [1994Ton, 1996Tom, 1995Mor] have been used to predict/ determine reactions and phase stabilities for aluminium-based metal matrix composites with Ti and C. [1997Wu] ball milled titanium, aluminium, and graphite powders in the atomic ratio Al-0.23Ti-0.23C and produced an amorphous phase that further transformed to an fcc metastable phase. Annealing at 600˚C for 45 minutes yielded the equilibrium phases (Al), Al3Ti and TiC1–x(?).
. Table 1 Investigations of the Al-C-Ti Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/ Composition/Phase Range Studied
[1954Thy]
Optical metallography and XRD
0-1 mass% C at 2, 4, 6, 8 and 10 mass% Al
[1963Jei]
Powder metallurgy, XRD
Ti2AlC
[1964Jei]
Powder metallurgy, XRD
Ti3AlC
[1964Now]
Powder metallurgy, XRD
Ti2AlC
[1972Now]
Powder metallurgy, XRD
Ti2AlC
[1976Ivc2]
Powder metallurgy, XRD
Ti2AlC and Ti3AlC
[1977Gur]
Metallography, XRD, magnetic properties TiAl2.97C0.03 measurements
[1980Pea]
Dimensional analysis
Ti2AlC
[1980Sch]
XRD
1000˚C; 30 alloys of nonspecified composition
[1984Ker]
Metallography, SEM and EDS.
Solubility of Al in TiC1–x at 1600˚C
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. Table 1 (continued)
Reference
Temperature/ Composition/Phase Range Studied
Method/Experimental Technique
[1987Zak]
Metallography and X-ray diffraction
600˚C - 1400˚C, TiAl-TiC, up to 25 mass% TiC0.98
[1993Tia]
XRD, TEM
(Ti0.50Al0.50)99.5C0.5
[1989Cam1, 1989Cam2, 1990Cam, 1991Cam1, 1991Cam2]
Light microscopy, SEM, TEM, X-ray diffraction, X-ray ED spectroscopy, EDX and hardness measurements
750˚C, 1000˚C, 1250˚C, 15 < at.% Al < 55, 0.5 < at.% C <3
[1999Far]
TEM, HRTEM
Ti3AlC2
[1992Aye]
SEM, WDS and EDS on HIPped ingots
Ti2AlC1–x
[1992Pie, 1994Pie]
X-ray diffraction
1100˚C, 900˚C and 700˚C
[1993Sve]
Light optical microscopy, SEM, EDS
1100˚C, 900˚C and 700˚C / Low C and up to 2.22 at.% Ti / Liquid (Al), Al4C3 and TiC1–x
[1995Via]
Diffusion couples, X-ray diffraction
1000˚C
[1997Li, 1999Li]
EPMA
Ti46Al-0.2C
[2000Ria]
SEM, TEM, X-ray diffraction
750˚C, 1000˚C
[2005Wan3]
XRD
Ti3AlC2, Ti2AlC and TiC1–x
[2006Lin]
XRD, TEM with EDS, HRTEM
Ti2AlC
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βAl)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25˚C, 20.5 GPa [Mas2]
(αAl) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25˚C [Mas2]
(C)d
cF8 Fd3m C (diamond)
a = 356.69
at 25˚C, 60 GPa [Mas2]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(C)gr < 3827
hP4 P63/mmc C (graphite)
a = 246.12 c = 670.90
at 25˚C [Mas2] sublimation point
(ωTi)
hP3 P6/mmm ωTi
a = 462.5 c = 281.3
at 25˚C, HP at 1 atm [Mas2]
(βTi) 1670 - 882
cI2 Im3m W
a = 330.65
pure Ti [Mas2]
-
Possible ordering from A2 (βTi) to B2 (β0) [2000Ohn]
a = 295.06 c = 468.35
at 25˚C [Mas2]
cP2 β0, (βTi,Al) ~1427 - ~1104 Pm3m CsCl (αTi) < 882
hP2 P63/mmc Mg
α2, Ti3Al < 1193 (up to 10 GPa at RT)
hP8 P63/mmc Ni3Sn
γ, TiAl < 1456
a = 578.83 c = 463.7 a = 580.6 c = 465.5 a = 574.6 c = 462.4
tP4 P4/mmm AuCu
a = 400.0 c = 407.5 a = 398.4 c = 406.0
D019 ordered phase [1997Sah] [1980Sch]
at 22 at.% Al [L-B] at 38 at.% Al [L-B] L10 ordered phase, has c/a > 1 50 at.% Al at 1000˚C [2001Bra] 62 at.% Al at 1100˚C [2001Bra]
Ti1–xAl1+x > 1170
tP4 P4/mmm AuCu
a = 403.0 c = 395.5
Ti3Al5 < 809
tP32 P4/mbm Ti3Ga5
a = 1129.3 Ti3Al5, stable below 810˚C [2001Bra] (a = 399.3 substructure) c = 403.8
ζ, Ti2Al5 1432 - 976
tP28 P4/mmm Ti2Al5
a = 393.6 c = 413.54
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Disordered Ti1–xAl1+x has c/a < 1 [2001Bra], and was reported as a separate high temperature phase
[1986Mii] Presumed ordering effect
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm] a = 398.0 c = 2436.0
Comments/References
TiAl2 < 1224
tI24 I41/amd HfGa2
Ti5Al11 ≥ 995
CuAu a = 395.3 superstructure c = 410.4
[2001Bra] 1300˚C, but not recognized here as a separate phase
ε(h), TiAl3 (h) 1396 - 734
tI8 I4/mmm TiAl3(h)
[V-C2]
a = 384.6 c = 859.4
[1980Mii, 1986Mii] by TEM
a = 384.9 1000˚C [2001Bra] c = 860.9 (c = 430.5 substructure)
ε(l), TiAl3 (l) < 932
tI32 I4/mmm TiAl3(l)
a = 387.7 76 at.% Al at 640˚C [2001Bra] c = 3382.82 (c = 422.9 substructure)
TiC1–x
cF8 Fm3m NaCl
a = 432.8
[V-C2]
Al4C3
hR21 R3m Al4C3
a = 333.8 c = 2511.7
[V-C2]
* H, Ti2AlC1–x
hP8 P63/mmc Cr2AlC
a = 305.6 c = 1362.3
[V-C2]
* P, Ti3AlC1–x
cP5 Pm3m CaTiO3
a = 415.6
[V-C2]
* N, Ti3AlC2
hP12 P63/mmc CMo
a = 307.53 c = 1857.8
[1994Pie], isotypic with Ti3SiC2
Metastable / high pressure phases TiAl3 (m)
cP4 Pm3m AuCu3
Ti3Al (I) hP16 15 to > 41 GPa P63/mmc TiNi3
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a = 397.2
Splat cooled; 85 at.% Al [1994Bra]
a = 531.2 c = 960.4
[1997Sah] at 16 GPa; Not found at 0–35 GPa, 25-2250˚C [2004Sch]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
TiAl2 (m)
oC12 Cmmm ZrGa2
a = 394.21 c = 401.61
~66 to 67 at.% Al - metastable in as-cast alloys [2001Bra]
Ti52Al48 (m)
cP20 P4132 βMn
a = 690
Precipitated in amorphous thin film after 1 h at 525˚C [1999Abe]
. Table 3 Invariant Equilibria Composition (at.%) T [˚C]
Reaction
Type
Phase
Al
C
Ti
L + TiC1–x + H Ð P
1580±10
P
-
-
-
-
L + TiC1–x Ð (βTi) + P
1430 < T < 1580
U
-
-
-
-
L + P Ð (βTi) + H
Lower than the above
U
-
-
-
-
L + (βTi) Ð (αTi) + H
<1460
U
-
-
-
-
L + TiC1–x Ð TiAl3 + Al4C3
812
U
L
99.6281
0.0019
0.0019
. Table 4 Thermodynamic Properties of Single Phases
Phase
Temperature Range [˚C]
Property, per mole of atoms [J, mol, K]
Comments
Ti2AlC 30-1300
CP = 58.1+0.1·T–7·10–5·T2+ 1.8·10–8·T 3
[2002Bar]
Ti2AlC 30-1000
ktot = 49–0.01·T
thermal conductivity [2002Bar]
Ti2AlC 30-1000
TEC = 8.7·10
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–1
(K )
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. Table 5 Investigations of the Al-C-Ti Materials Properties Reference
Method / Experimental Technique
Type of Property
[1967Hod]
Machining tests; microhardness measurements
Cutting tool ranking; microhardness
[1976Ivc1]
Electrical and thermal; microhardness measurements
Coefficients of thermal expansion; emf; electrical and thermal conductivities; microhardness of constituent phases
[1976Jon]
Optical microscopy, SEM and calculation
Grain refining capability
[1977Ivc]
Optical microscopy, SEM and DTA
Precipitate characterization and formation
[1982Wit]
Sputter coater, Rutherford backscattering and XRD
Coating characterization
[1986Ban, 2000Rom]
Powder metallurgy, optical microscopy, SEM and TEM
Precipitate characterization and formation for grain refining
[1986Bat]
Mechanical properties
Yield strength, elongation and hardness
[1989Cam1, 1989Cam2]
Optical microscopy, SEM and TEM with EDX, hardness measurements
Precipitate characterization and formation; hardness
[1991Fro]
Mechanical alloying, DSC, DTA, optical microscopy, and SEM with EDX and WDS
Precipitate characterization and formation
[1991Jar, 1992Jar, 1999Bir]
Crucible furnace, optical microscopy and XRD
Manufacture of TiC-(Al) composite; precipitate formation and characterization
[1991Kno]
PVD-coatings, SEM with EDX
Phase proportions
[1991Mab]
Reactive synthesis, arc-melting, compression testing, optical microscopy, TEM and XRD
Phase proportions; compressive strengths
[1991Sah]
Carrier gas, calorimetry and TEM Precipitate formation and characterization
[1992Che]
Powder metallurgy, reaction sintering, SEM, and TEM with EDX and EELS
Precipitate characterization and formation
[1992Cho, 1996Sun, 1999Zha, 2002Fra, 2002Kho2, 2003Ge1, 2003Ge2, 2005Ma]
SHS combustion chamber, XRD and SEM
Characterization and reaction temperatures
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. Table 5 (continued) Method / Experimental Technique
Reference
Type of Property
[1992Vol, 1997Lee, 1997Ye, Powder metallurgy, combustion 2002Hwa] reaction chamber (reaction sintering), optical microscopy, XRD, WDS and SEM
Formation and characterization of TiC precipitates
[1993Den]
Atom probe
Partitioning between γTiAl and Ti3Al3
[1993Ber, 1993Cho, 1993Jar, 1995Mor, 1995Via 1996Tom, 2002Ge, 2002Hwa]
Reaction between graphite and liquid Al-Ti alloy
Aluminum matrix metal matrix composites (In-situ)
[1993Cho, 2006Xia]
Combustion chamber, XRD and SEM with EDS
Manufacture and characterization of TiC-(Al) cermet
[1993Jar]
Infiltration, SEM with EDS
Precipitate characterization and formation
[1993Mit]
Casting; extrusion; microhardness tests; tensile tests; SEM, TEM and XRD
Formation and characterization of (Al)/TiC; microhardnesstemperature relationships; Young’s modulus, yield stress and elongation
[1993Tia]
TEM, mechanical tests
P/γTiAl orientation relationships; 0.2% proof stress
[1994Din]
HIP, XRD
Phase distribution
[1994May]
Optical microscopy, SEM, Precipitate characterization, (extraction replica) TEM with EDX formation and distribution and XRD
[1994Ton]
Rapid quenching
Precipitate formation, distribution and characterization
[1994Whi]
HIP, optical microscopy, SEM, TEM and XRD, compression, creep and tensile tests
Volume fraction; yield strength, UTS, elongation and ductility (RA); compressive strain, compressive creep curves
[1995Nuk]
Infiltration of preform, DTA and XRD
Formation reactions of TiC
[1996Mat, 1996Ily, 1997Ily1, Ab initio and structure 1997Ily2, 1998Mat, calculations 2001Zho1, 2001Zho2, 2003Wan1]
Electronic structure calculations
[1996Sob]
Sessile drop
Wettability of substrate with molten Al-Ti alloys
[1997Chr, 2005App]
TEM
Characterization of mechanically tested samples
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. Table 5 (continued) Reference
Method / Experimental Technique
Type of Property
[1997Wu]
Mechanical alloying of titanium, aluminium, and carbon powder
Nanocrystalline phase formation
[1998Tan, 2000Bri1]
Reactive synthesis, DSC, optical microscopy, SEM and XRD
Formation, characterization and distribution of precipitates
[1999Far]
HRTEM
Characterization of Ti2AlC and Ti3AlC2
[2000Bar1]
HIP, XRD, mechanical tests
Characterization; coefficients of expansion
[2000Bri2, 2000Van]
Optical microscopy, XRD and SEM
Assessment of grain refining efficiency
[2000Fin]
Mechanical properties
Elastic properties of Ti3AlC2
[2000Ken]
Powder metallurgy, XRD, SEM
Formation and characterization of precipitates
[2000Zha1]
XRD, chemical analysis, SEM with Precipitate characterization, EDS formation and distribution
[2000Tze]
HIP, XRD, SEM with EDX and hardness measurements
Formation, characterization and hardness
[2001Lop]
Mill, combustion reaction and XRD
Characterization and phase proportions
[2001Qin]
Arc melting, directional casting, optical microscopy, XRD, SEM, TEM and thermodynamic calculations
Characterization and microstructure stability
[2001Yan]
Mill, press, XRD, optical microscopy SEM and TEM
Formation, characterization and alloy compositions to avoid deleterious TiAl3
[2001Zha]
Arc melting, XRD and SEM with EDX
Formation and characterization
[2002Bar]
Electrical and mechanical properties
Electrical and thermal characterization of Ti2AlC
[2002Gaz]
Metallography, LiMCA and PoDFA
Particle kinetics characteristics
[2002Ge]
Combustion, DSC
Formation kinetics of Ti3AlC2
[2002Kho1]
High energy attrition mill, XRD, SEM with EDS
Formation and characterization
[2002Mei]
SPS, XRD and EPMA
Formation and characterization
[2002Ria]
Arc melting, optical microscopy, Characterization and dendrite SEM, image analysis and XRD arm spacing of TiC
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. Table 5 (continued) Method / Experimental Technique
Reference
Type of Property
[2002Tro]
Thermoelectric (TEP) measurements, metallography and SEM with EDX
Formation, characterization and effect on grain refining
[2002Wan1, 2002Wan2, 2004Wan1, 2004Wan2, 2006Lee]
Oxidation studies
Synthesis and oxidation of Ti3AlC2
[2002Xia, 2002Liu]
Metallography, SEM
Characterization and distribution of particles; effectiveness of grain refining
[2002Zha1]
Melt spinning, XRD and SEM
Formation and characterization; effect on grain refining
[2002Zha2]
Arc melting, XRD, SEM with EDS, Formation, characterization; nanoindenter microhardness and modulus
[2003Arp]
Pressure infiltration, DTA, XRD and SEM
(Al)/TiC composite manufacture and characterization
[2003Guo1, 2003Guo2, 2004Guo]
Combustion chamber studies
Formation of TiC1–x, Ti3AlC2
[2003Hua]
Reactions of activated Al-Ti-C powder under laser
Formation and characterization
[2003Guo1, 2003Guo3]
Combustion chamber, XRD, SEM Formation and characterization of Ti2AlC1–x
[2003Son, 2004Son2]
Mill, “DERP” thermal analysis furnace, optical microscopy, XRD, SEM with EDS and TEM
Formation and characterization
[2003Sun]
Bulk modulus studies
Ti2AlC bulk modulus
[2003Wan2]
Ball milling, DSC and XRD
Formation reactions for TiC for grain refiners
[2003Zha1]
Melt spinning, XRD, SEM and TEM
Formation and characterization
[2003Zha2]
DTA, XRD, SEM
Formation and characterization
[2003Zho]
SPS, HIP, XRD and SEM
Formation and characterization of Ti3AlC2
[2004Hon, 2005Hon]
Press, XRD, SEM, density and hardness measurements
Characterization, density and hardness
[2004Hwa, 2004Wan3, 2005Wan4]
SHS combustion chamber, XRD and SEM with WLS and EDX
Precipitate characterization, formation and distribution
[2004Kex, 2005Rog]
Ball mill, combustion synthesis reactor, XRD and SEM with EDS
Formation and characterization
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. Table 5 (continued) Reference
Method / Experimental Technique
Type of Property
[2004Mei, 2005Zho1, 2005Zho2]
Spark plasma sintering (SPS), XRD and SEM with EDS.
Manufacture, characterization and density of Ti2AlC
[2004Son1]
Mill, hydrostatic press, SEM and TEM with EDS
Formation and characterization of Ti3AlC2
[2001Zho1]
XRD, electronic measurements
Structure and electrical properties of Ti3AlC2
[2005Bao]
Electrical and mechanical properties
Thermal shock of Ti3AlC2
[2005Jia, 2006Xia]
SHS combustion chamber, XRD, Precipitate characterization, DTA and SEM with WLS and EDX formation and distribution
[2005Kho]
SHS combustion chamber, press, Characterization, formation and XRD, SEM with EDX, EPMA and distribution image analysis
[2005Kum]
XRD, SEM with EDX
Characterization of master alloy, and grain-refined alloys
[2005Li]
SHS combustion chamber, XRD, XRF, SEM with EDX and thermodynamic calculations
Precipitate characterization, formation and distribution
[2005Pen]
Press, XRD and SEM
Characterization and density
[2005Wal]
Sputtering chamber, XRD and XPS
Characterization
[2005Wan1, 2005Wan2]
Sintering, pressing, mechanical Manufacture, characterization, tests and XRD, SEM with EDX and three-point bend tests, harnesses density measurements and fracture toughness
[2005Xia]
Al-C-Ti launders and moulds, electromagnetic field
Precipitate characterization and distribution
[2005Zha]
Equal-channel angular press, SEM and microhardness tester
Formation, precipitate characterization and distribution; hardness
[2006Li]
Mechanical alloying, XRD and SEM
Formation and characterization
[2006Sel]
Powder metallurgy, optical Formation, characterization, yield microscopy, SEM and mechanical stress, UTS, ductility and Young’s measurements modulus.
[2006Zho]
Uniaxial compression and indentation by hemispherical nanoindenor
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. Fig. 1 Al-C-Ti. Adopted Al-Ti phase diagram
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. Fig. 2 Al-C-Ti. Adopted C-Ti phase diagram
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. Fig. 3 Al-C-Ti. Partial isothermal section at 1250˚C
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. Fig. 4 Al-C-Ti. Partial isothermal section at 1050˚C
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. Fig. 5 Al-C-Ti. Isothermal section at 1000˚C
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. Fig. 6 Al-C-Ti. Partial isothermal section at 750˚C
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. Fig. 7 Al-C-Ti. Isopleth at 2 mass% Al for the concentration range 0–1 mass% C, plotted in at.%
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. Fig. 8 Al-C-Ti. Isopleth at 4 mass% Al for the concentration range 0–1 mass% C, plotted in at.%
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. Fig. 9 Al-C-Ti. Isopleth at 6 mass% Al for the concentration range 0–1 mass% C, plotted in at.%
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. Fig. 10 Al-C-Ti. Isopleth at 8 mass% Al for the concentration range 0–1 mass% C, plotted in at.%
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. Fig. 11 Al-C-Ti. Isopleth at 10 mass% Al for the concentration range 0–1 mass% C, plotted in at.%
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. Fig. 12 Al-C-Ti. Vertical section γ(TiAl) - TiC1–x
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Kerans, R.J., Mazdiyasni, K.S., Ruh, R., Lipsitt, H.A., “Solubility of Metals in Substoichiometric TiC(1–x)”, J. Am. Ceram. Soc., 67, 34–38 (1984) (Experimental, Mechan. Prop., Morphology, Phase Diagram, Phase Relations, 29) Banerji, A., Reif, W., “Development of Al-Ti-C Grain Refiners Containing TiC”, Metall. Trans. A, 17, 2127–2137 (1986) (Calculation, Experimental, Morphology, Phase Relations, Thermodyn., 28) Baturinskaya, N.L., Kalchuk, N.A., Sezonenko, Yu.D., Chernyi, V.G., “Cast Dispersion-Hardemed Aluminum Alloys with a High Titanium Carbide Content”, Russ. Metall., 3, 145–147 (1986) (Experimental, 3) Miida, R., Jap. J. Appl. Phys., 25, 1815–1824 (1988), as quoted in [1990Sch] Zakharov, A.M., Lashkova, L.A.K; Semeryakova, S.G., “Polythermal Section (γ)TiAl-TiC of the Ti-Al-C System”, Russ. Metall. (Engl. Transl.), (4), 197–198 (1987), translated from Izv. Akad. Nauk SSSR, Met., (4), 196–197 (1987) (Crys. Structure, Experimental, Phase Diagram, 7) Cam, G., Flower, H.M., D. R. F. West, D.R.F., “The Alloying of Titanium Aluminides with Carbon”, Mater. Res. Soc. Symp. Proc., 133, 663–668 (1989) (Experimental, Morphology, Phase Relations, 7) Cam, G., Flower, H.M., West, D.R.F., “Phase Transformations in Ti-Al-C Alloys”, Proc. Euromat ’89, Aachen, 1–6 (1989) (Experimental, Morphology, Phase Relations, 7) Cam, G., Ph.D. Thesis, Imperial College, University of London, 1990, as quoted in [2000Ria] Hayes, F. “Al-C-Ti (Aluminium-Carbon-Titanium)” in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14870.1.20, (1998), also published in “Ternary Alloys”, Petzow, G., Effenberg, G. (Eds.), Vol. 3, VCH Verlagsgesellschaft, Weinheim, Germany, 557–566 (1990) (Phase Diagram, Review, 15) Schuster, J.C., Isper, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81 (6), 389–396 (1990) (Phase Diagram, Experimental, 33) Viala, J.C., Vincent, C., Vincent, H., Bouix, J., “Thermodynamic Approach of the Chemical Interaction between Aluminium and Carbon-Titanium” (in French), Mater. Res. Bull., 25, 457–464 (1990) (Experimental, Phase Diagram, Phase Relations, 20) Cam, G., Flower, H.M., West, D.R.F., “Constitution of Ti-Al-C Alloys in the Temperature Range 1250750˚C”, Mater. Sci. Technol., 7, 505–511 (1991) (Experimental, Morphology, Phase Diagram, Phase Relations, 15) Cam, G., Flower, H.M., West, D.R.F., “Solid State Phase Transformations in Titanium-AluminumCarbon and Titanium-Aluminum-Niobium-Carbon Alloys”, Mater. Sci. Technol., 7(7), 587–591 (1991) (Crys. Structure, Experimental, Morphology, Phase Relations, 10) Froyen, L., Delaey, L., “In Situ Formation of Aluminium Composites”, COST 507 Leuven Proc., Commission of the European Communities, Part A, B1 (1991) (Experimental, 1) Jarfors, A., Fredriksson, H., Froyen, L., “On the Thermodynamics and Kinetiks of Carbides in the Aluminium-rich Corner of the Al-Ti-C Phase Diagram”, Mater. Sci. Eng. A, 135, 119–123 (1991) (Assessment, Experimental, Kinetics, Phase Diagram, Thermodyn., 5) Knotek, O., Lo¨ffler, F., Wolkers, L., “On the Phase Stability and Transition of Metastable TiAlC and TiAlN-PVD-coatings”, COST 507 Leuven Proc., Part A, D7, (1991) (Experimental, 4) Mabuchi, H., Harada, K., Tsuda, H., Nakayama, Y., “Fabrication of Ti2AlC/TiAl Composites Using Combustion Reaction Process”, ISIJ Int., 31(10), 1272–1278 (1991) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Thermodyn., 20) Rapp, R.A., Zheng, X., “Thermodynamic Consideration of Grain Refinement of Aluminium Alloys by Titanium and Carbon”, Metall. Trans. A, 22A, 3071–3075 (1991) (Calculation, Thermodyn., 7) Sahoo, P., Koczak, M.J., “Analysis of in Situ Formation of Titanium Carbide in Aluminum Alloys”, Mater. Sci. Eng. A, 144(1-2), 37–44 (1991) (Experimental, Kinetics, Thermodyn., 23) Yokokawa, H., Sakai, N., Kawada, T., Dokiya, M., “Chemical Potential Diagram of Al-Ti-C System: Al4C3 Formation on TiC Formed in Al-Ti Liquids Containing Carbon”, Metall. Trans. A, 22A, 3075–3076 (1991) (Calculation, Phase Diagram, Thermodyn., 11) Ayer, R., Ray, R., Scanlon, J.C., “Analytical Microscopy of a γ-TiAl Composite Containing a Carbide Phase”, Scr. Metall. Mater., 26(9), 1337–1342 (1992) (Crys. Structure, 19) Chen, S., Beaven, P.A., Wagner, R., “Carbide Precipitation in γ-TiAl Alloys”, Scr. Metall. Mater., 26(8), 1205–1210 (1992) (Electronic Structure, Experimental, Morphology, 10) Choi, Y., Mullins, M.E., Wijayatilleke, K., Lee, J.K., “Fabrication of Metal Matrix Composites of TiC-Al Through Self-Propagating Synthesis Reaction”, Metall. Trans. A, 23(9), 2387–2392, 1992 (Experimental, Morphology, 14)
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Al–C–Ti Jarfors, A.E.W., Fredriksson, H., “The Reaction Between Titanium Carbide and Aluminum Carbide in Liquid Aluminum; A Microgravity Study”, Eur. Space Agency, ESA SP-333, Proc. VIIIth European Symposium on Materials and Fluid Sciences in Microgravity, Vol. 20, 849–854 (1992) Pietzka, M.A.O., “Structural Chemistry, Phase Equilibria and Chemical Analysin in the Ti-Al-C and Ti-Al-N Systems” (in German), Dipl., Uni. Wien, 1–52 (1992) (Phase Diagram, 35) Vol’ne, B.M., Evstigneev, V.V., “Structure Formation in the SHS-System Titanium Aluminum Carbon”, Combustion Explosion and Shock Waves, 28(2), 173–178 (1992) (Crys. Structure, Experimental, Morphology, 8) Bergman, A., Jarfors, A., Liu, Z., Fredriksson, H., “In-situ Formation of Carbide Composites by Liquid/Solid Reactions”, Key Eng. Mater., 79-80, 213–233 (1993) (Calculation, Experimental, Morphology, Phase Diagram, Phase Relations, 9) Choi, Y., Rhee, S.W., “Effect of Aluminium Addition on the Combustion Reaction of Titanium and Carbon to Form TiC”, J. Mater. Sci., 28(24), 6669–6675 (1993) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Thermodyn., 17) Frisk, K., “A Revised Thermodynamic Description of the Ti-C System”, Calphad, 27(3), 367–373 (2003) (Phase Diagram, Phase Relations, Thermodyn., Assessment, 22) Denauin, A., Naka, S., Huguet, A., Menand, A., “Atom-Probe Investigation of the Partitioning of Interstitial Elements in Two-Phase γ + α2 Ti-Al-Based Alloys”, Scr. Metall. Mater., 28(9), 1131–1136 (1993) (Crys. Structure, Experimental, 8) Jarfors, A.E.W., Svendsen, L., Wallinder, M., Fredriksson, H., “Reactions During Infiltration of Graphite Fibers by Molten Al-Ti Alloy”, Metall. Trans. A, 24A(11), 2577 (1993) (Experimental, Mechan. Prop., Phase Diagram, Thermodyn., 17) Mitra, R., Fine, M.E., Weertman, J.R., “Chemical Reaction Strengthening of Al/TiC Metal Matrix Composites by Isothermal Heat Treatment at 913 K”, J. Mater. Res., 8(9), 2370–2379 (1993) (Crys. Structure, Experimental, Morphology, Phase Relations, Thermodyn., 25) Svendsen, L., Jarfors, A., “Al-Ti-C Phase Diagram”, Mater. Sci. Technol., 9, 948–957 (1993) (Experimental, Phase Diagram, Thermodyn., 17) Tian, W.H., Harada, K., Nakashima, R., Sano, T., Nemoto, M., “Crystal Structures and Morphologies of Carbide and Nitride Precipitates in TiAl” (in Japanese), J. Jpn. Inst. Met., 57(11), 1235–1243 (1993) (Crys. Structure, Experimental, Phase Diagram, 17) Braun, J., Ellner, M., Predel, B., Z. Metallkd., 85, 855–862 (1994) as quoted in [2001Bra] Dinsdale, A., “Summary of the Proceedings of the CALPHAD XXIII-CAMSE 94 Meeting”, Calphad, 18 (4), 337–368 (1994) (Abstract, Calculation, Phase Diagram, 52) Mayes, C.D., McCartney, D.G., Tatlock, G.J., “Observations on the Microstructure and Perfomance of an Al-Ti-C Grain-Refining Master Alloy”, Mater. Sci. Eng. A, 188, 283–290 (1994) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 29) Pietzka, M.A., Schuster, J.C., “Summary of Constitutional Data on the Aluminium-Carbon-Titanium System”, J. Phase Equilib., 15(4), 392–400 (1994) (Calculation, Experimental, Phase Diagram, Phase Relations, Thermodyn., 23) Tong, X., Shen, N., Liu, B., “Formation and Distribution of TiC Phase in Rapidly Quenched Al-3.18Ti-0.65C Alloy” (in Japanese), Acta Metall. Sin., Ser. A, 30(4), 155–159 (1994) (Experimental, Thermodyn., 2) Whittenberger, J.D., Ray, R., Farmer, S.C., “Elevated-Temperature Deformation Properties of In-Situ Carbide Particle Strengthened Ti-48Al Materials”, Intermetallics, 2(3), 167–178 (1994) (Crys. Structure, Experimental, Morphology, Thermodyn., 19) Zhang, M.X., Chang, Y.A., “Phase Diagrams of Ti-Al-C, Ti-Y-O, Nb-Y-O, and Nb-Al-O at 1100˚C”, J. Phase Equilib., 15(5), 470–472 (1994) (Experimental, Phase Diagram, Phase Relations, 12) Morimoto, H., Nomura, M., Ashida, Y., “In-Situ TiC Particulate Reinforced Aluminum Composites Fabricated by Reaction Between Graphite Particles and Liquid Al-Ti Alloys” (in Japanese), J. Jpn. Inst. Met., 4, 429–436 (1995) (Crys. Structure, Experimental, Morphology, Phys. Prop., Thermodyn., 24) Nukami, T., Flemings, M.C., Metall. Mater. Trans., 24A (1995) as quoted in [2000Tjo] Viala, J.C., Peillon, N., Clochefert, L., Bouix, J., “Diffusion Paths and Reaction Mechanisms in the High-Temperature Chemical Interaction Between Carbon and Titanium Aluminides”, Mater. Sci. Eng. A, 203, 222–237 (1995) (Experimental, Interface Phenomena, Kinetics, Phase Diagram, Phase Relations, Thermodyn., 22)
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Chang, W-S., Muddle, B.C., “Precipitation of (Ti,V)2Al(C,N,) in Multiphase Al-Ti-V Alloys”, Mater. Sci. Eng. A, 207, 64–71 (1996) (Crys. Structure, Experimental, 18) Ilyasov, V.V., Nikiforov, I.Ya., Ilyasov, Yu.V., “Electronic Energy Structure of the Ternary Carbide System Ti-Al-C”, Phys. Status Solidi B, 198, 687–693 (1996) (Calculation, Crys. Structure, Mechan. Prop., Phys. Prop., 22) Matar, S.F., Etourneau, J., “Investigation of the Electronic Structure of Carbon-Containing TiAl”, J. Alloys Compd., 233(1-2), 112–120 (1996) (Crys. Structure, Experimental, Thermodyn., 10) Sobczak, N., Gorny, Z., Ksiazek, M., Radziwill, W., Rohatgi, P., “Interaction Between Porous Graphite Substrate and Liquid or Semi-liquid Aluminum Alloys Containing Titanium”, Mater. Sci. Forum, 217-222, 153–158 (1996) (Crys. Structure, Experimental, Morphology, Thermodyn., 17) Sun, X., Mei, B., Yu, A.R., Liao, G., “Thermal Explosion Synthesis in Ti-C-Al System”, Acta Metall. Sin., 32(10), 1102–1106, 1996 (Experimental, Morphology, 3) Tomoshige, R., Matsushita, T., “Production of Titanium-Aluminum-Carbon Ternary Composites with Dispersed Fine TiC Particles by Combustion Synthesis and Their Microstructure Observations” (in Japanese), J. Ceram. Soc. Jpn., 104(2), 94–100 (1996) (Crys. Structure, Experimental, Morphology, Thermodyn., 14) Christoph, U., Appel, F., Wagner, R., “Dislocation Dynamics in Carbon-Doped Titanium Aluminide Alloys”, Mater. Sci. Eng. A, 240, 39–45 (1997) (Experimental, Morphology, Thermodyn., 15) Ichikawa, K., Kinoshita, Y., “Rheocasting Techniques Applied to Intermetallic TiAl Alloys and Composites”, Mater. Sci. Eng. A, 239-240, 493–502 (1997) (Experimental, Mechan. Prop., Phase Relations, 16) Ilyasov, V.V., Nikiforov, I.Ya., “Computer Simulation of the Electronic Structure and Chemical Bond in the Ternary System Ti1–xAlxC”, Phys. Solid State, 39(2), 185–188 (1997) (Calculation, Phys. Prop., 22) Ilyasov, V.V., Nikiforov, I.Ya., Ilyasov, Yu.V., “Ti L-Spectrum XANES and Electron Structure of the System Ti-Al-C”, J. Phys. IV, France, 7(C2), 281–282 (1997) (Experimental, 9) Li, J., Hao, S., “α (α2) and γ Phase Equilibria in Ti-Al-C and Ti-Al-B Ternary Systems”, Trans. Nonferrous Met. Soc. China, 7(2), 63–66 (1997) (Experimental, Morphology, Phase Relations, 10) Lee, W.-C., Chung, S.-L., “Ignition Phenomena and Reaction Mechanisms of the Self-Propagating High-Temperature Synthesis Reaction in the Titanium-Carbon-Aluminum System”, J. Am. Ceram. Soc., 80(1), 53–61 (1997) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Thermodyn., 20) Sahu, P.Ch., Chandra Shekar, N.V., Yousuf, M., Govinda Rajan, K., “Implications of a Pressure Induced Phase Transition in the Search for Cubic Ti3Al”, Phys. Rev. Lett., 78(6), 1054–1057 (1997) (Crys. Structure, Experimental, 20) Wu, N.Q., Wu, J.M., Li, Z.Z., Wang, G.-X., “Formation of Nanocrystalline F.C.C. Phase by Mechanically Driven Crystallization”, Mater. Trans., JIM, 38(3), 255–259 (1997) (Experimental, 16) Ye, L.L., Liu, Z.G., Li, S.D., Quan, M.X., Hu, Z.Q., “Thermochemistry of Combustion Reaction in AlTi-C System During Mechanical Alloying”, J. Mater. Res., 12(3), 616–618 (1997) (Experimental, 19) Frange, N., Frumin, N., Levin, L., Polak, M., Dariel, M.P., “High-Temperature Phase Equilibria in the Al-Rich Corner of the Al-Ti-C System”, Metall. Mater. Trans. A, 29, 1341–1345 (1998) (Calculation, Experimental, Phase Relations, Thermodyn., 27) Matar, S.F., Petitcorps, Y.L., Etourneau, J., “Ab Initio Study of the Chemical Role of Carbon within TiAl Alloy System: Application to Composite Materials”, Comput. Mater. Sci., 10, 314–318 (1998) (Abstract, Experimental, 7) Tanaka, H., Tomoshige, R., Imamura, K., Chiba, A., Kato, A., “Hot-shock Consolidation and Mechanical/thermal Properties of Ti-Al-C Composites Using Explosive Shock Energy and Combustion Synthesis” (in Japanese), J. Ceram. Soc. Jpn., 106(7), 676–681 (1998) (Experimental, Interface Phenomena, Morphology, 17) Abe, E., Ohnuma, M., Nakamura, M. “The Structure of a New ε-Phase Formed During the Early Stage of Crystallization of Ti-48 at.% Al Amorphous Film”, Acta Mater., 47(13), 3607–3616 (1999) (Crys. Structure, Experimental, 25) Birol, Y., “In Situ Processing of TiCp-Al Composites by Reacting Graphite with Al-Ti Melts”, J. Mater. Sci., 34(7), 1653–1657 (1999) (Experimental, Morphology, 19)
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[1999Far]
[1999Ho]
[1999Jar] [1999Li] [1999Zha] [2000Ban]
[2000Bar1] [2000Bri1]
[2000Bri2]
[2000Fin]
[2000Ken]
[2000Ohn]
[2000Ria]
[2000Rom] [2000Sch] [2000Tjo]
[2000Tze]
[2000Van]
[2000Zha1]
Al–C–Ti Dumitrescu, L.F.S., Hillert, M., Sundman, B., “A Reassessment of Ti-C-N based on a Critical Review of Available Assessments of Ti-N and Ti-C”, Z. Metallkd., 90(7), 534–541 (1999) (Calculation, Phase Relations, Phase Diagram, Thermodyn., 38) Farber, L., Levin, I., Barsoum, M.W., El-Raghy, T., Tzenov, T., “High-Resolution Transmission Electron Microscopy of some Tin+1Axn Compounds (n=1, 2; A=Al or Si; X=C or N)”, J. Appl. Phys., 86(5), 2540–2543 (1999) (Crys. Structure, Experimental, 23) Ho, J.C., Hamdeh, H.H., Barsoum, M.W., El-Raghy, T., “Low Temperature Heat Capacities of Ti3Al1.1C1.8, Ti4AlN3, and Ti3SiC2”, J. Appl. Phys., 86(7), 3609–3611 (1999) (Experimental, Thermodyn., 15) Jarfors, A.E.W., “Peritectic-like Precipitation of Titanium Carbide in Al-Ti-C Melts at 1373 K”, Mater. Sci. Technol., 15, 481–494 (1999) (Experimental, Phase Relations, Phys. Prop., Theory, Thermodyn., 29) Li, J., Zong, Y., Hao, Sh., “Effects of Alloy Elements (C, B, Fe, Si) on the Ti-Al Binary Phase Diagram”, J. Mater. Sci. Technol., 15(1), 58–62 (1999) (Experimental, Phase Relations, 13) Zhang, E., Zeng, S., Yang, B., Ma, M., “Formation Process of SHS TiC in Al-Ti-C System”, Intl. J. SelfPropagating High-Temperature Synthesis, 8(1), 59–68 (1999) (Experimental, Morphology, 11) Bandayopadhyay, D., Sharma, R.C., Chakraborti, N., “The Ti-Al-C System (Titanium-AluminiumCarbon)”, J. Phase Equilib., 21(2), 195–198 (2000) (Assessment, Crys. Structure, Phase Diagram, Phase Relations, 11) Barsoum, M.W., Ali, M., El-Raghy, T., “Processing and Characterization of Ti2AlC, Ti2AlN and Ti2AlC0.5N0.5”, Metall. Trans. A, 31A, 1857–1865 (2000) (Crys. Structure, Experimental, 36) Brinkman, H.J., Zupanic, F., Duszczyk, J., Katgerman, L., “Production of Al-Ti-C Grain Refiner Alloys by Reactive Synthesis of Elemental Powders: Part I. Reactive Synthesis and Characterization of Alloys”, J. Mater. Res., 15(12), 2620–2627 (2000) (Crys. Structure, Experimental, Morphology, 16) Brinkman, H.J., Zupanic, F., Duszczyk, J., Katgerman, L., “Production of Al-Ti-C Grain Refiner Alloys by Reactive Synthesis of Elemental Powders: Part II. Grain Refining Performance of Alloys and Secondary Processing”, J. Mater. Res., 15(12), 2628–2635 (2000) (Crys. Structure, Experimental, Morphology, Thermodyn., 9) Finkel, P., Barsoum, M.W., El-Rghy, T., “Low Temperature Dependencies of the Elastic Properties of Ti4AlN3, Ti3Al1.1C1.8, anf Ti3SiC2”, J. Appl. Phys., 87(4), 1701–1703 (2000) (Experimental, Mechan. Prop., 22) Kennedy, A.R., Weston, D.P., Jones, M.J., Enel, C., “Reaction in Al-Ti-C Powders and its Relation to the Formation and Stability of TiC in Al at High Temperatures”, Scr. Mater., 42(12), 1187–1192 (2000) (Experimental, Thermodyn., 12) Ohnuma, I., Fujita, Y., Mitsui, H., Ishikawa, K., Kainuma, R., Ishida, K., “Phase Equilibria in the Ti-Al Binary System”, Acta Mater., 48, 3113–3123 (2000) (Calculation, Experimental, Phase Relations, Thermodyn., 37) Riaz, S., Flower, H.M., West, D.R.F., “Phase Relationships Involving TiC and Ti3AlC (P Phase) in Ti-Al-C System”, Mater. Sci. Technol., 16, 984–992 (2000) (Crys. Structure, Experimental, Phase Relations, 16) Romankiewicz, F., “Grain Refinement of Aluminium with AlTi6C0,1” (in German), Z. Metallkd., 91 (10), 822–825 (2000) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 7) Schneider, W., “Grain Refinement of Al Wrought Alloys with Newly Developed AlTiC Master Alloy” (in German), Z. Metallkd., 91(10), 800–806 (2000) (Mechan. Prop., Phase Relations, Review, 14) Tjong, S.C., Ma, Z.Y., “Microstructural and Mechanical Characteristics of in Situ Metal Matrix Composites”, Mater. Sci. Eng., 29, 49–113 (2000) (Crys. Structure, Mechan. Prop., Morphology, Review, Thermodyn., 229) Tzenov, N.V., Barsoum, M.W., “Synthesis and Characterization of Ti3AlC2”, J. Am. Ceram. Soc., 83(4), 825–832 (2000) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., Mechan. Prop., Phys. Prop., 27) Vandyoussefi, M.V., Worth, J., Greer, A.L., “Effect of Instability of TiC Perticles on Grain Refinement of Al and Al-Mg Alloys by Addition of Al-Ti-C Inoculants”, Mater. Sci. Technol., 16, 1121–1128 (2000) (Calculation, Experimental, Phase Relations, 30) Zhang, B.-Q., Fang, H.-S., Li, J.-G., Ma, H.-T., “An Investigation on Microstructures and Refining Performances of Newly Developed Al-Ti-C Grain Refining Master Alloys”, J. Mater. Sci. Lett., 19(16), 1485–1489 (2000) (Crys. Structure, Experimental, 44)
DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
MSIT1
Landolt‐Bo¨rnstein New Series IV/11E1
Al–C–Ti [2000Zha2] [2001Bra] [2001Lop]
[2001Qin]
[2001Yan] [2001Zha]
[2001Zho1] [2001Zho2]
[2002Bar]
[2002Fra] [2002Gaz]
[2002Ge]
[2002Hwa]
[2002Kho1] [2002Kho2] [2002Liu]
[2002Mei] [2002Ria]
[2002Tro]
[2002Wan1] [2002Wan2]
6
Zhang, Z., Liu, X., Bian, X., “Thermodynamics and Kinetic of Forming TiC in Al-Ti-C System” (in Japanese), Acta Chim. Sin., 36(10), 1025–1029 (2000) (Experimental, Morphology, 12) Braun, J., Ellner, M., “Phase equilibria Investigations on the Aluminium-Rich Part of the Binary System Ti-Al”, Metall. Mater. Trans. A, 32A, 1037–1047 (2001) (Phase Diagram, Experimental, 34) Lopacinski, M., Puszynski, J., Lis, J., “Synthesis of Ternary Titanium Aluminum Carbides Using Selfpropagating High-temperature Synthesis Technique”, J. Am. Ceram. Soc., 84(12), 3051–3053 (2001) (Experimental, 11) Qin, G.W., Oikawa, K., Sun,Z.M., Sumi, S., Ikeshoji, T., Wang, J.J., Guo, S.W., Hao, S.M., “Discontinuous Coarsening of the Lamellar Structure of γ-TiAl-Based Intermetallic Alloys and its Control”, Metall. Mater. Trans. A, 32, 1927–1938 (2001) (Crys. Structure, Experimental, Kinetics, Phase Relations, Thermodyn., 46) Yang, B., Chen, G.X., Zhang, J.H., “Effect of Ti/C Additions on the Formation of Al3Ti of in Situ TiC/ Al Composites”, Mater. Design, 22(8), 645–650 (2001) (Experimental, Morphology, 22) Zhang, E., Wang, H., Zeng, S., “Microstructure Characteristics of in Situ Carbide Reinforced Titanium Aluminide (Ti3Al) Matrix Composites”, J. Mater. Sci. Lett., 20, 1733–1735 (2001) (Experimental, Interface Phenomena, Morphology, 7) Zhou, Y.C., Wang, X.H., Sun, Z.M., Chen, S.Q., “Electronic and Structural Properties of the Layered Ternary Carbide Ti3AlC2”, J. Mater. Chem., 11, 2335–2339 (2001) (Calculation, Crys. Structure, 29) Zhou, Y., Sun, Zh., Wang, X., Chen, Sh., “Ab Initio Geometry Optimization and Ground State Properties of Layered Ternary Carbides Ti3MC2 (M=Al, Si and Ge)”, J. Phys.: Condens. Matter, 13(44), 10001–10010 (2001) (Calculation, Crys. Structure, Electronic Structure, 17) Barsoum, M.W., Salama, I., El-Raghy, T., Golczewski, J., Porter, W.D., Wang, H., Seifert, H.J., Aldinger, F., “Thermal and Electrical Properties of Nb2AlC, (Ti,Nb)2AlC and Ti2AlC”, Metall. Mater. Trans. A, 33 (9), 2775–2779 (2004) (Experimental, Thermodyn., 28) Fras, E., Wierzbinski, S., Janas, A., Lopez, H.F., “SHSB Processing and Properties of Al/TiC “in-situ” Composites”, Arch. Metall., 46(4), 407–423 2001 (Experimental, Morphology, 23) Gazanion, F., Chen, X.-G., Dupuis, C., “Studies on the Sedimentation and Agglomeration of Al-Ti-B and Al-Ti-C Grain Refiners”, Mater. Sci. Forum, 396-402, 45–52 (2002) (Experimental, Thermodyn., 12) Ge, Z.B., Chen, K.X., Zhou, H.P., Guo, J.M., “Kinetic Process of Combustion Synthesis of Ternary Carbide Ti3AlC2”, Key Eng. Mater., 224-226, 539–544 (2002) (Experimental, Morphology, Thermodyn., 15) Hwang, C.-C., Chung, S.-L., “Combustion Synthesis in the Ti+C/Ti+Al System - Influence of Reactant Composition”, J. Mater. Sci. Lett., 21(6), 447–450 (2002) (Experimental, Phase Relations, Thermodyn., 15) Khoptiar, Y., Gotman, I., “Ti2AlC Ternary Carbide Synthesized by Thermal Explosion”, Mater. Lett., 57 (1), 72–76 (2002) (Crys. Structure, Experimental, Morphology, Phase Relations, 11) Khoptiar, Y., Gotman, I., Gutmanas, E.Y., “SHS of Ti2AlC and Ti3AlC2 Machinable Ceramics”, Intl. J. Self-Propagating High-Temperature Synthesis, 11(4), 335–44, 2002 (Experimental, Morphology, 16) Liu X.F., Wang, Z.Q., Zhang, Z.G., Bian X.F., “The Relationship Between Microstructures and Refining Performances of Al-Ti-C Master Alloys”, Mater. Sci. Eng. A, 332(1-2), 70–74, 2002 (Experimental, Morphology, 13) Mei, B., Miyamoto, Y., “Investigation of TiAl/Ti2AlC Composites Prepared by Spark Plasma Sintering”, Mater. Chem. Phys., 75, 291–295 (2002) (Experimental, Morphology, Phase Relations, 10) Riaz, S., Flower, H.M., West, D.R.F., “Characteristics of TiC Dendrites in as Solidified Ti-Al-C Alloys”, Mater. Sci. Technol., 18(8), 941–943 (2002) (Crys. Structure, Experimental, Morphology, Phase Relations, 7) Tronche, A., Vandyoussefi, M., Greer, A.L., “Instability of TiC Particles in Aluminium Melts Inoculated with an Al-Ti-C Grain Refiner”, Mater. Sci. Technol., 18(10), 1072–1078 (2002) (Experimental, Phase Relations, 32) Wang, X.H., Zhou, Y.C., “Synthesis and Oxidation of Bulk Ti3AlC2”, Key Eng. Mater., 224-226, 785–790 (2002) (Experimental, Interface Phenomena, 12) Wang, X.H., Zhou, Y. C., “Intermediate-Temperature Oxidation Behavior of Ti2AlC in Air”, J. Mater. Res., 17(11), 2974–2981 (2002) (Experimental, Kinetics, 24)
Landolt‐Bo¨rnstein New Series IV/11E1
MSIT1
DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
35
36
6 [2002Xia]
[2002Zha1]
[2002Zha2]
[2003Arp]
[2003Fri] [2003Ge1]
[2003Ge2]
[2003Guo1]
[2003Guo2]
[2003Guo3] [2003Hua]
[2003Son]
[2003Sun]
[2003Wan1]
[2003Wan2]
[2003Yan]
[2003Zha1] [2003Zha2]
[2003Zho]
Al–C–Ti Xiangfa, L., Zhenqing, W., Zuogui, Z., Xiufang, B., “The Relationship Between Microstructures and Refining Performances of Al-Ti-C Master Alloys”, Mater. Sci. Eng. A, 332, 70–74 (2002) (Crys. Structure, Experimental, Morphology, 16) Zhang, Z., Bian, X., Wang, Z., Liu, X., Wang, Y., “Microstructures and Grain Refinement Perfomance of Rapidly Solidified Al-Ti-C Master Alloys”, J. Alloys Compd., 339, 180–188 (2002) (Experimental, Phase Relations, 25) Zhang, E., Zeng, S., Wang, B., “Preparation and Microstructure of In Situ Particle Reinforced Titanium Matrix Alloy”, J. Mater. Proc. Tech., 125-126, 103–109 (2002) (Experimental, Mechan. Prop., Morphology, 12) Arpon, R., Narciso, J., Louis, E., Garcia-Cordovilla, C., “Interfacial Reactions in Al/TiC Particulate Composites Propduced by Pressure Infitration”, Mater. Sci. Technol., 19, 1225–1230 (2003) (Experimental, Interface Phenomena, Phase Relations, Thermodyn., 22) Frisk, K., “A Revised Thermodynamic Description of the Ti-C System”, Calphad, 27(4), 367–373 (2003) (Assessment, Phase Diagram, Phase Relations, Thermodyn., 22) Ge, Z., Chen, K., Guo, J., Zhou, H., Ferreira, J.M.F., “Combustion Sinthesis of Thernary Carbide Ti3AlC2 in Ti-Al-C System”, J. Eur. Ceram. Soc., 23(3), 567–574 (2003) (Experimental, Phase Diagram, 14) Ge, Z.B., Chen, K.X., Guo, J.M., Zhou, H.P., Ning, X.S., “Formation Mechanism of Ternary Carbide Ti3AlC2 by Combustion Synthesis”, J. Inorg. Mater., 18(2), 427–432 (2003) (Experimental, Morphology, 9) Guo, J.M., Chen K.X., Ge, Z.B., Zho, H.P., Ning, X.S., “Effects of TiC Addition on Combustion Synthesis of Ti(2)AlC Powders” (in Chinese), Acta Metall. Sin., 39(3), 315–319 (2003) (Experimental, Morphology, 18) Guo, J.M., Chen, K.X., Ge, Z.B., Liu, G.H., Zhou, H.P., Ning, X.S., “Effects of Carbon Addition on Combustion Synthesis of Ti3AlC2 Powders” (in Chinese), Acta Metall. Sin., 39(4), 409–413 (2003) (Experimental, Morphology, 17) Guo, J., Chen, K., Ge, Z., Zhou, H., Ning, X., “Combustion Synthesis Ternary Carbide Ti2AlC1–x”, Rare Metal Mater. Eng., 32(12), 1029–1032 (2003) (Experimental, Morphology, 19) Huang, Y., Ma, N., Zou, D., Liang G., Su, J., “Chemical Reactions in Activated Al-Ti-C Powder Mixture under Action of Laser Beam”, Rare Metal Mater. Eng., 32(12), 999–1002 (2003) (Experimental, Morphology, 12) Song, I.-H., Kim, D.K., Hahn, Y.-D., Kim, H.-D., “The Effect of a Dilution Agent on the Dipping Exothermic Reaction Process for Fabricating a High-Volume TiC-reinforced Aluminum Composite”, Scr. Mater., 48(4), 413–418 (2003) (Crys. Structure, Experimental, Morphology, Phys. Prop., 12) Sun, Zh., Ahuja, R., Li, S., Schneider, J.M., “Structure and Bulk Modulus of M2AlC (M = Ti, V, and Cr)”, Appl. Phys. Lett., 83(5), 899–901 (2003) (Calculation, Crys. Structure, Electronic Structure, Kinetics, Mechan. Prop., 15) Wang, J.Y., Zhou, Y.C., “First-Principles Study of Equilibrium Properties and Electronic Structure of Ti3Si0.75Al0.25C2 Solid Solution”, J. Phys.: Condens. Matter, 15(35), 5959–5968 (2003) (Calculation, Crys. Structure, Electronic Structure, Thermodyn., 21) Wang, Z., Liu, X., Zhang, J., Bian, X., “Reaction Mechanism in the Ball-Milled Al-Ti-C Powders”, J. Mater. Sci. Lett., 22(20), 1427–1429 (2003) (Crys. Structure, Experimental, Kinetics, Phase Relations, 6) Yang, B., Wang, F., Zhang, J.S., “Microstructural Characterization of in Situ TiC/Al and TiC/Al-20Si5Fe-3Cu-1Mg Composites Prepared by Spray Deposition”, Acta Mater., 51(17), 4977–4989 (2003) (Crys. Structure, Experimental, Interface Phenomena, Kinetics, Morphology, Phase Relations, Thermodyn., 65) Zhang, Z., Bian, X., Wang, Y., Liu, X., Wang, Z., “TEM Observation of a Rapidly Solidified Al-Ti-C Alloy”, J. Alloys Compd., 349, 121–128 (2003) (Crys. Structure, Experimental, 26) Zhang, B.Q., Fang, H.S., Lu, L., Lai, M.O., Ma, H.T., Li, J.G., “Synthesis Mechanism of an Al-Ti-C Grain Refiner Master Alloy Prepared by New Method”, Metall. Mater. Trans. A, 34(8), 1727–1733 (2003) (Experimental, Thermodyn., 19) Zhou, A., Wang, C.-A., Hunag, Y., “Synthesis and Mechanical Properties of Ti3AlC2 by Spark Plasma Sintering”, J. Mater. Sci., 38(14), 3111–3115 (2003) (Experimental, Mechan. Prop., 14)
DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
MSIT1
Landolt‐Bo¨rnstein New Series IV/11E1
Al–C–Ti [2004Guo]
[2004Hon]
[2004Hwa]
[2004Kex] [2004Mei]
[2004Per]
[2004Sch]
[2004Son1]
[2004Son2] [2004Wan1]
[2004Wan2]
[2004Wan3]
[2005App] [2005Bao]
[2005Hon]
[2005Jia]
[2005Kho]
[2005Kum]
[2005Li] [2005Ma]
6
Guo, J.M., Chen, K.X., Zhou, H.P., Ning, X.S., “Effects of TiAl3 Addition in Ti-Al-C System on Combustion Synthesis of Ti3AlC2 Powders” (in Chinese), Acta Metall. Sin., 40(1), 109–112 (2004) (Experimental, Morphology, 14) Hong, X., Mei, B., Zhu, J., Zhou, W., “Fabrication of Ti2AlC by Hot Pressing of Ti, TiC, Al and Active Carbon Powder Mixtures”, J. Mater. Sci., 39(5), 1589–1592 (2004) (Crys. Structure, Experimental, Kinetics, Morphology, Phase Relations, 13) Hwang, C.-C., Chung, S.-L., “A Study of Combustion Synthesis Reaction in the Ti+C/Ti+Al System”, J. Mater. Sci., 39(6), 2073–2080 (2004) (Crys. Structure, Experimental, Morphology, Phase Relations, 18) Kexin, C., Junming, G., Renli, F., Ferreira, J.M.F., “Combustion Synthesis Ternary Carbide Ti2AlC1–x Powders”, Mater. Sci. Forum, 445-456, 191–195 (2004) (Experimental, Morphology, Thermodyn., 11) Mei, B., Zhou, W., Zhu, J., Hong, X., “Synthesis of High-Purity Ti2AlC by Spark Plasma Sintering (SPS) of the Elemental Powders”, J. Mater. Sci., 39(4), 1471–1472 (2004) (Crys. Structure, Experimental, Kinetics, Morphology, 7) Perrot, P., “Al - C (Aluminium - Carbon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, Document ID: 20.15738.1.20, (2004) (Phase Relations, Phase Diagram, Assessment, 19) Schmid-Fetzer, R, “Al-Ti (Aluminum-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.15634.1.20 (2004) (Phase Diagram, Phase Relations, Review, 85) Song, I.-H., Kim, D.K., Hahn, Y.-D. Kim, H.-D., “Investigation of Ti3AlC2 in the In-Situ TiC-Al Composite Prepared by the Exothermic Reaction Process in Liquid Aluminum”, Mater. Lett., 58(5), 593–597 (2004) (Experimental, Morphology, 14) Song, I.H., Kim, D.K., Hahn, Y.D., Kim, H.D., “Synthesis of In-Situ TiC-Al Composite by Dipping Exothermic Reaction Process”, Met. Mater. Int., 10(3), 301–306 (2004) (Experimental, Morphology, 26) Wang, X.H., Zhou, Y.C., “Oxidation Behavior of TiC-Containing Ti3AlC2 Based Material at 500-900˚C in Air”, Mater. Res. Innovat., 7(6), 381–390 (2004) (Experimental, Interface Phenomena, Kinetics, Phase Relations, 25) Wang, X.H., Zhou, Y.C., “Improvement of Intermediate-Temperature Oxidation Resistance of Ti3AlC2 by Pre-Oxydation at High Temperatures”, Mater. Res. Innovat., 7(4), 205–211 (2004) (Crys. Structure, Experimental, Interface Phenomena, Kinetics, Morphology, Phase Relations, 10) Wang, H.Y., Jiang, Q.C., Li, X.L., Zhao, F., “Effect of Al Content on the Self-Propagating Hightemperature Synthesis Reaction of Al-Ti-C System in Molten Magnesium”, J. Alloys Compd., 366(1-2), L9–L12 (2004) (Crys. Structure, Experimental, Morphology, 13) Appel, F., “An Electron Microscope Study of Mechanical Twinning and Fracture in TiAl Alloys”, Philos. Mag., 85(2-3), 205–231 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, 74) Bao, Y.W., Wang, X.H., Zhang, H.B., Zhou, Y.C., “Thermal Shock Behavior of Ti3AlC2 from Between 200 and 1300˚C”, J. Eur. Ceram. Soc., 25(14), 3367–3374 (2005) (Electr. Prop., Experimental, Mechan. Prop., Morphology, Phys. Prop., Thermodyn., 22) Hongxiang, Z., Zhenying, H., Mingxing, A., Yang, Z., Zhili, Z., Shibo, L., “Tribophysical Properties of Polycrystalline Bulk Ti3AlC2”, J. Am. Ceram. Soc., 88(11), 3270–3274 (2005) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 35) Jiang, Q.C., Wang, H.Y., Zhao, Y.G., Li, X.L., “Solid-State Reaction Behavior of Al-Ti-C Powder Mixture Compacts”, Mater. Res. Bull., 40(3), 521–527 (2005) (Crys. Structure, Experimental, Morphology, Phase Relations, 13) Khoptiar, Y., Gotman, I., Gutmanas, E.Y., “Pressure-Assisted Combustion Synthesis of Dense Layered Ti3AlC2 and its Mechanical Properties”, J. Am. Ceram. Soc., 88(1), 28–33 (2005) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phase Relations, 11) Kumar, G.S.V., Murty, B.S., Chakraborty, M., “Development of Al-Ti-C Grain Refiners and Study of Their Grain Refining Efficiency on Al and Al-7Si Alloy”, J. Alloys Compd., 396(1-2), 143–150 (2005) (Crys. Structure, Experimental, Interface Phenomena, Morphology, 22) Li, P., Kandalova, E.G., Nikitin, V.I., “In Situ Synthesis of Al-TiC in Aluminum Melt”, Mater. Lett., 59(19-20), 2545–2548 (2005) (Experimental, Morphology, Phase Relations, Thermodyn., 18) Ma, M., Liu, R., Zhao, H., Yu, Y., “In-Situ TiCp/Al Composites Prepared by TE/QP Method”, J. Mater. Sci. Technol., 21(5), 652–656 (2005) (Experimental, Morphology, Phys. Prop., 13)
Landolt‐Bo¨rnstein New Series IV/11E1
MSIT1
DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009
37
38
6 [2005Pen] [2005Rog]
[2005Wal] [2005Wan1]
[2005Wan2]
[2005Wan3] [2005Wan4] [2005Xia]
[2005Zha]
[2005Zho1]
[2005Zho2]
[2006Dru]
[2006Lee]
[2006Li]
[2006Lin]
[2006Rag] [2006Sch]
[2006Sel]
[2006Ton]
[2006Xia]
Al–C–Ti Peng, C., Wang, C.-A., Huang, Y., “Synthesis of High Purity Ti3AlC2 Bulk Material by Hot-Pressing”, Key Eng. Mater., 280-283, 1369–1372 (2005) (Experimental, Morphology, 7) Rogachev, A.S., Gachon, J.-C., Grigoryan, H.E., Vrel, D., Schuster, J.C., Sachkova, N.V., “Phase Evolution in the Ti-Al-B and Ti-Al-C Systems During Combustion Synthesis: Time Resolved Study by Synchrotron Radiation Diffraction Analysis”, J. Mater. Sci., 40(9-10), 2689–2691 (2005) (Crys. Structure, Experimental, Phase Relations, 7) Walter, C., Martinez, C., El-Raghy, T., Schneider, J.M., “Towards Large Area MAX Phase Coatings on Steel”, Steel Res., 76(2-3), 225–228 (2005) (Phase Relations, Phys. Prop., Review, 14) Wang, C.-A., Zhou, A., Peng, C., Huang, Y., “Fabrication and Properties of Ti3AlC2-Based Ceramics by Two-Step Method”, Key Eng. Mater., 280-283, 1365–1368 (2005) (Experimental, Mechan. Prop., Phase Relations, 6) Wang, H.Y., Zhao, F., Jiang, Q.C., Wang, Y., Ma, B.X., “Effect of Mg Addition on the Self-Propagating High Temperature Synthesis Reaction in Al-Ti-C System”, J. Mater. Sci., 40(5), 1255–1257 (2005) (Crys. Structure, Experimental, Morphology, Phase Relations, 12) Wang, Ch.-A., Zhou, A., Qi, L., Huang, Y., “Quantitative Phase Analysis in the Ti-Al-C Ternary System by X-ray Diffraction”, Powder Diffr., 20(3), 218–223 (2005) (Experimental, 10) Wang, H., Xia, T., Zhao, W., Liu T., “Preparation of Al-Ti-C Master Alloy by SHS”, Rare Met. Mater. Eng., 34(12), 2009-12, 2005 (Experimental, Morphology, 13) Xiangfa, L., Lina, Y., Jianwen, L., Zhenqing, W., Xiangfa, B., “A New Technique to Refine Pure Aluminum by Al-Ti-C Mold”, Mater. Sci. Eng. A, 399, 267–270 (2005) (Experimental, Interface Phenomena, 13) Zhang, Z., Watanabe, Y., Kim, I., Liu, X., Bian, X., “Microstructure and Refining Performance of an Al-5Ti-0.25C Refiner before and after Equal-Channel Angular Pressing”, Metall. Mater. Trans. A., Suppl., 36A(3), 837–844 (2005) (Experimental, Morphology, 30) Zhou, W.B., Mei, B.C., Zhu, J.Q., Hong, X.L., “Rapid Synthesis of Ti2AlC by Spark Plasma Sintering Technique”, Mater. Lett., 59(1), 131–134 (2005) (Crys. Structure, Experimental, Mechan. Prop., Phys. Prop., 15) Zhou, W., Mei, B., Zhu, J., Hong, X., “Synthesis of High-Purity Ti3SiC2 and Ti3AlC2 by Spark Plasma Sintering (SPS) Technique”, J. Mater. Sci., 40(8), 2099–2100 (2005) (Crys. Structure, Morphology, Experimental, Phys. Prop., 5) Drulis, M.K., Drulis, H., Gupta, S., Barsoum, M.W., El-Raghy, T., “On the Heat Capacities of M2AlC (M=Ti, V, Cr) Ternary Carbides”, J. Appl. Phys., 99(9), 093502 (2006) (Crys. Structure, Electronic Structure, Experimental, Phys. Prop., Thermodyn., 34) Lee, D.B., Park, S.W., “High-temperature Oxidation of Ti3AlC2 Between 1173 and 1473 K in Air”, Mater. Sci. Eng. A, 434(1-2), 147–154 (2006) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Phase Relations, Thermodyn., 15) Li, S.-B., Zhai, H.-X., Bei, G.P., Zhou, Y., Zhang, Z.L., “Formation of Ti3AlC2 by Mechanically Induced Self-Propagating Reaction in Ti-Al-C System at Room Temperature”, Mater. Sci. Technol., 22(6), 667–672 (2006) (Crys. Structure, Experimental, Morphology, 27) Lin, Z.J., Zhuo, M.J., Zhou, Y.C., Li, M.S., Wang, J.Y., “Microstructural Characterization of Layered Ternary Ti2AlC”, Acta Mater., 54(4), 1009–1015 (2006) (Crys. Structure, Experimental, Morphology, 33) Raghavan, V., “Al-C-Ti (Aluminum-Carbon-Titanium)”, J. Phase Equilib. Diffus., 27(2), 148–149 (2006) (Crys. Structure, Phase Diagram, Phase Relations, Review, 14) Schuster, J.C., Palm, M., “Reassessment of the Binary Aluminum-Titanium Phase Diagram”, J. Phase Equilib. Diffus., 27(3), 255–277 (2006) (Assessment, Calculation, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 272) Seleuk, C., Kennedy, A.R., “Al-TiC Composite Made by the Addition of Master Alloys Pellets Synthesised from Reacted Elemental Powders”, Mater. Lett., 60(28), 3364–3366 (2006) (Crys. Structure, Experimental, Morphology, 15) Tong, L., Reddy, R.G., “In-Situ Synthesis of TiC-Al (Ti) Nanocomposite Powders by Thermal Plasma Technology”, Metall. Mater. Trans. B, 37B, 531–539 (2006) (Calculation, Crys. Structure, Experimental, Morphology, Thermodyn., 23) Xiao, G.Q., Fan, Q.C., Gu, M.Z., Jin, Z.H., “Microstructural Evolution During the Combustion Synthesis of TiC-Al Cermet with Larger Metallic Particles”, Mater. Sci. Eng. A, 425(1-2), 318–325 (2006) (Crys. Structure, Experimental, Morphology, 33)
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Zhou, A.G., Barsoum, M.W., Basu, S., Kalidindi, S.R., El-Raghy, T., “Incipient and Regular Kink Bands in Fully Dense and 10 vol.% Porous Ti2AlC”, Acta Mater., 54(6), 1631–1639 (2006) (Experimental, Mechan. Prop., Morphology, 23) Witusiewicz, V.T., Bondar, A.A., Hecht, U., Rex, S., Velikanova, T.Ya., “The Al-B-Nb-Ti System. III. Thermodynamic Re-Evaluation of the Constituent Binary System Al-Ti”, in press, J. Alloys Compd., doi:10.1016/j.jallcom.2007.10.061 (2007) (Phase Relations, Phase Diagram, Experimental, Thermodyn., Calculation, Review, 89) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Chromium – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Natalia Bochvar, Tatiana Dobatkina, Ol’ga Fabrichnaya, Volodymyr Ivanchenko, Damian M. Cupid
Introduction Aluminum and chromium are well known alloying elements used in titanium industry to produce Ti-based alloys [1956Zwi, 1974Zwi, 1990Lam]. Titanium aluminides based on α2 (Ti3Al), γ (TiAl) and TiAl3 are currently being considered as candidate materials for elevated temperature applications as a result of their low density, good corrosion resistance, modulus and attractive high temperature properties [1989Kim, 1991Dur, 1991Moh, 1994Gao, 1995Whi, 1996Liu, 2000Wel, 2000Yan, 2003Zha, 2004Wan, 2005App]. However, these alloys are brittle at temperatures lower than 400 K [2005App]. Alloying with Cr raises their plasticity and resistance to high temperature oxidation [1991Par]. Reviews of experimental studies in the Al-Cr-Ti system were performed by [1953Han, 1992Hay, 2005Rag]. Review of experimental studies up to 1990 is presented by [1992Hay]. Phase equilibria in the Ti rich corner were extensively investigated by [1953Tay, 1958Kor, 1958Tag, 1960Enc]. Partial isothermal sections presented by [1992Hay] in the Ti rich corner and in the range of 600–1200˚C are mainly based on the results of experimental studies of [1960Enc]. [2005Rag] reviewed more recent experimental studies on phase equilibria [2001Fuj, 2000Kai1, 1996Jew1, 1996Jew2, 1996Jew3, 1994Sok]. In the review of [2005Rag], a schematic liquidus surface based on studies of [1997Mab2] and [2001Ich], a reaction scheme and isothermal sections between 497–1200˚C in more extended composition ranges than in [1992Hay] are presented. The Calphad method was used in works [1988Gro, 1999Sha, 2001Kau] to establish a thermodynamic database and to calculate phase diagrams. However, the details of modeling were not reported in these works. The only experimental data for thermodynamic values were obtained by [1999Jac] by vapour pressure measurements. Experimental and theoretical studies of the Al-Cr-Ti system are summarized in Table 1.
Binary Systems The binary diagrams were taken from [2008Cor] for the Al-Cr, from [2007Wit] for the Al-Ti and from [2000Zhu] for the Cr-Ti. The lattice parameters for the binary Al-Ti compounds were taken from [2006Sch] and for the binary Cr-Ti compounds from [2008Iva].
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Solid Phases Three ternary phases exist in the Al-Cr-Ti system. A ternary phase with the ordered structure L12, denoted τ1, was formed by stabilisation of the binary metastable TiAl3(m) compound by Cr addition [1992Nic, 1993Nak, 1994Kla, 1997Par, 2005Rag]. The τ1 phase exhibits a homogeneity range and dissolves from 7 to 15 at.% Cr at 1200˚C, whereas the content of Ti changes from 23 to 27 at.% [1992Nic, 2001Kau]. According to experimental studies of [1997Mab1, 1997Mab2, 2001Ich, 2003Bar] and calculation of Al-Cr-Ti liquidus surface by [1999Sha] phase τ1 most probably forms by a peritectic reaction. The temperature of τ1 formation was about 1350˚C according to [2001Ich]. In the review of [2005Rag] formation of τ1 by peritectic reaction at 1370˚C was postulated. The ternary phase Ti(Cr,Al)2, (C14 ordered phase), denoted τ2, forms in the Al-Cr-Ti system in the solid state by Al substitution of Cr atoms in the structure of the hightemperature γTiCr2 phase. However, according to the available experimental data, τ2 and γTiCr2 never form a single phase (continuous solid solution). Therefore, τ2 and γTiCr2 are considered here as different phases in spite of having the same structure. The high temperature limit of formation of the τ2 phase was not determined exactly. According to [2001Fuj] and [1997Xu], the maximum temperature of stability of τ2 is between 1200 and 1050˚C. The τ2 phase is stable down to room temperature [1978Jac, 1992Nic, 1993Nak, 1994Kla, 1997Par, 1997Mab1, 1997Mab2]. The τ2 phase has a narrow homogeneity range which extends along the Al-Cr side. The Al content in this phase varies from 3 to 20 at.% at 1050˚C [1997Xu], from 5 to 42 at.% at 1000˚C [1996Jew1, 1997Jew] or from 5 to 40 at.% at 800˚C [1996Jew1, 1997Jew]. Deviation of the Ti content from stoichiometric content (33.3 at.% Ti) was insignificant: 1 at.% at 5 at.% Al and 3 at.% at 35 at.% Al at 1000˚C and 3 at.% at 40 at.% Al at 800˚C [1996Jew1, 1997Jew]. It should be mentioned that with an increase in temperature from 1000 to 1050˚C the τ2 phase shrinks due to the extension of the β phase. The ternary phase denoted τ3 based on the ordered B2 phase (β0), forms in the Al-Cr-Ti system in the solid state. The temperature of formation of the τ3 phase was not determined. One can assume that the high temperature stability limit of the τ3 phase is between 1050˚C, where the phase does not exist [1997Xu], and 1000˚C where it was indicated by [1997Jew, 2000Kai1, 2000Kai2]. The τ3 phase is stable down to room temperature [1992Zhe, 1994Zha, 1999Sha, 2001Sun]. It has a large homogeneity range at 1000˚C, which reduces as temperature decreases [1996Jew1, 1997Jew]. This phase was also indicated by [1964Ram, 1964Sch]. It should be mentioned that the τ3 phase has the same structure as the ordered B2 phase (β0) stable in the Al-Ti binary system in the temperature range of 1159–1427˚C. However, these phases are considered as different phases because they never form a continuous solid solution. According to [1997Jew], the TiAl, Ti3Al, and TiAl2 binary phases extend into the ternary system. The αTiCr2 and βTiCr2 phases also extended into the ternary system dissolving 1 and 4 at.% Al, respectively [1996Jew1, 1997Xu]. At 1000˚C, the solubility of Cr in the TiAl phase increased from about 2 at.% to about 8 at. % while the Al content varied from about 47 at.% to 58 at.%. At 800˚C, the maximum solubility of Cr in TiAl was about 4.5 at.% and remained fairly constant within the Al composition range. The solubility of Cr in the Ti3Al phase is about 2.5 at.% in the temperature range from 800 to 1000˚C. The solubility of Cr in TiAl2 phase is 3.5 at.% at 1000˚C and about 6 at.% at 800˚C [1997Jew]. The solubility of Cr in TiAl3(l) is 4 at.% at 497˚C [1994Sok]. DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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The solubility of Ti in the Cr5Al8 phase is about 10 at.% at room temperature [1992Nic, 1993Nak, 1997Mab1]. The binary CrAl7 phase dissolves up to 1.03 at.% Ti according to [1960Zol] or 2 at.% Ti according to [1994Sok]. The crystal structure data of solid phases in the Al-Cr-Ti system are presented in Table 2.
Invariant Equilibria No temperatures or compositions for invariant reactions with liquid have been reported for the Al-Cr-Ti system. The methods of alloy preparation, the experimental technique and phase range studied are presented in Table 1. [2003Bar] experimentally established two univariant eutectic equilibria. The eutectic transformations L Ð τ1 + β and L Ð τ1 + Cr5Al8 are observed in the temperature range 1285–1255˚C and 1260–1225˚C, respectively. Investigating solid state phase equilibria at temperatures of 1000 and 800˚C, [1996Jew1, 1996Jew2] assumed the existence of invariant reactions τ2 + τ1 Ð β + TiAl2, TiAl + τ1 Ð τ2 + TiAl2 and β + Cr5Al8 Ð Cr2Al + τ1 at a temperature of around 890˚C. [1997Jew] assumed one more invariant reaction β + τ3 Ð τ2 + TiAl at a temperature of 1100±10˚C. For this reaction to occur two three-phase reactions should exist. One of them, τ2 Ð τ3 + β was established by [1996Jew1] at approximately 1120˚C and an Al content of 25 at.%. Existence of the other reaction τ3 Ð β + TiAl has not been confirmed. A reaction scheme was suggested in the present evaluation based on the configuration of the liquidus surface from the thermodynamic calculations of [1999Sha], as well as taking into account the nature and positions of the various threephase equilibria presented in the isothermal sections of [1997Jew]. The reaction scheme is shown in Fig. 1. The invariant three-phase reaction (βTi) Ð αTiCr2 +Ti3Al suggested by [1992Hay] was introduced into the reaction scheme. The temperatures of the invariant reactions indicated in the reaction scheme [1992Hay] were not established experimentally, but were evaluated based on the binary phase diagrams and isothermal sections.
Liquidus, Solidus and Solvus Surfaces Liquidus isotherms in the Ti rich corner (50–100 mass% Ti) were presented by [1958Kor] based on DTA results. Liquidus and solidus isotherms for τ1 and a partial liquidus surface showing equilibria between τ1 and the surrounding phases were constructed by [2001Ich] based on experimental studies. However, the Cr2Al and τ2 phases shown on the liquidus surface of [2001Ich] are not expected to be in equilibrium with liquid [2005Rag]. The boundaries of the liquidus surface projection presented in Fig. 2 are accepted according to [1999Sha]. Solid lines show boundaries of phase fields obtained by thermodynamic calculations, while dashed lines indicate the results of experiments carried out for three alloys Ti40Al-10Cr, Ti-50Al-10Cr, Ti-52Al-20Cr (at.%) (their compositions are shown by filled circles [1999Sha]). A part of the univariant reaction L Ð τ1 + β, between two empty circles in Fig. 2 is based on experimental data of [2003Bar] being in agreement with [1999Sha]. The positions of the invariant reactions in the binary Al-Ti and Al-Cr systems were corrected to be in agreement with binary diagrams of Al-Ti [2007Wit] and Al-Cr [2008Cor] accepted in the present evaluation. Additionally, phase equilibria were added in the Al rich corner, near the Al-Cr side. Isotherms at 1350, 1300, and 1250˚C in the primary crystallisation field of the τ1 phase are presented according to experimental data of [2001Ich]. Landolt‐Bo¨rnstein New Series IV/11E1
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[2005Rag] included a primary crystallisation field of the Ti1–xAl1+x solid solution in the liquidus surface. This phase is not considered in the Al-Ti binary system accepted in the present evaluation according to [2007Wit]. Therefore, the Ti1–xAl1+x phase is not shown in Fig. 2 and there are no reactions involving Ti1–xAl1+x in the reaction scheme (Fig. 1). Also, the difference in the liquidis surface of [2005Rag] and the one presented here is that the primary crystallization field of Ti2Al5 (the same as Ti5Al11) in [2005Rag] is much wider and there is an univariant reaction between liquid, Ti5Al11 and bcc (β) phases, while in the present evaluation, such reaction does not occur. The other difference between the present work and [2005Rag] is in the formation of τ1 by a peritectic reaction: [2005Rag] postulated the formation of τ1 from L+Ti5Al11+β, while here we accept its formation from L+TiAl+Ti2Al5 based on the calculations of [1999Sha]. The liquidus presented here is in a good agreement with the liquidus isotherms for τ1 obtained by [2001Ich], while the isotherm at 1350˚C [2001Ich] is outside the primary crystallisation field of the τ1 phase of [2005Rag]. It should be stressed that experimental information about liquidus in the Al-Cr-Ti system is very limited and the liquidus surface presented by [2005Rag] is rather tentative.
Isothermal Sections Phase equilibria investigations in the Al-Cr-Ti system were performed in the following works at various temperatures: [2000Kai2] at 1300˚C, [1992Hay, 1992Nic, 1992Zhe, 1994Kla, 1998Has, 2000Kai2, 2001Fuj] at 1200˚C, [1993Nak, 1997Mab1, 1997Mab2, 1997Par] at 1150˚C, [1992Hay] at 1100˚C, [1997Xu] at 1050˚C, [1991Mab, 1992Hay, 1993Nak, 1993Hao, 1995Bra, 1995Jew, 1995Hao, 1996Jew1, 1996Jew2, 1997Jew, 2000Kai2, 2001Fuj, 2003Nis1] at 1000˚C, [1974Zwi] at 980˚C, [1988Gro, 1992Hay, 1995Bra, 1995Jew, 1996Jew1, 1996Jew2, 1997Jew] at 800˚C, [1974Zwi] at 760˚C, [1992Hay] at 600˚C and [1994Sok] at 497˚ C. Methods of melting, conditions of homogenization followed by thermal treatment are given in Table 1. [1974Zwi] constructed a phase diagram in the Ti rich region, however, the Ti3Al phase was not considered and therefore these equilibria are not very probable. The isothermal section at 497˚C constructed by [1994Sok] is not exactly correct because the binary Cr2Al11 compound shown in this section decomposes at 785˚C according to accepted binary system [2008Cor]. The samples investigated in the works of [1991Mab, 1992Nic, 1992Zhe, 1993Nak, 1994Kla, 1997Mab1, 1997Mab2, 1997Par, 2003Nis1] were not quenched in water after heat treatment, but were furnace cooled or cooled in air. Using furnace or air cooling methods, the equilibrium phases at high temperature are probably not retained at room temperature. Therefore it can be assumed that phase equilibria established in these works belong to temperatures closer to room temperature. A partial isothermal section in the middle part of the system at 1300˚C is presented in Fig. 3 according to [2000Kai2]. Tie lines show equilibria between (αTi) and β, (αTi) and TiAl, and β and TiAl phases, respectively. A partial isothermal section in the Al rich corner at 1200˚C is presented in Fig. 4 according to [2001Fuj]. [2001Fuj] confirmed the existence of the three-phase field β + TiAl + (αTi), indicated by [1992Hay, 1998Has, 2000Kai2]. An isothermal section at 1050˚C constructed by the diffusion couple method [1997Xu] is presented in Fig. 5. The boundary of the β to τ3 transition is constructed according to the DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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data of [2000Kai1]. It should be mentioned that the phase interactions at 1050˚C (Fig. 5) and 1000˚C (Fig. 6) are different. As temperature decreases from 1050˚C, the single-phase β region contracts, whereas τ2 becomes increasingly stable. At 1050˚C, both the τ1 + τ2 two-phase field and the τ3 single-phase field are absent but are present at 1000˚C. Isothermal sections at 1000 and 800˚C are presented in Figs. 6 and 7 according to [1997Jew]. [1997Jew] used results obtained by [1995Jew, 1996Jew1], which studied equilibria adjacent to the Cr-Ti side with the Al content changing between 30 and 60 at.%, and the Al rich region up to 40 at.% Ti and 80 at.% Cr investigated by [1996Jew2]. In these studies, two three-phase fields in the Al-corner were reliably established: Cr2Al + Cr5Al8 + τ1 and TiAl3 + Cr5Al8 + τ1. Other fields were shown by dashed lines and need additional investigations. The boundary between β and τ3 stability fields is shown according to [2000Kai1]. Results of [1997Jew] at 1000˚C are in general agreement with [1993Hao, 1995Hao], who studied Ti rich region of the system by the diffusion couple method. However, the authors of [1993Hao, 1995Hao] did not indicate the existence of the ternary phase τ3 having the B2 structure. Threephase equilibria τ3 + τ2 + TiAl and τ1 + τ2 + TiAl established by [1997Jew] at 1000˚C were confirmed by [2001Fuj], and the Ti3Al + TiAl + τ3 equilibrium was confirmed by [2000Kai2, 2001Fuj]. The isothermal section at 800˚C was corrected to be consistent with the accepted binary systems Cr-Ti [2000Zhu] and Al-Ti [2007Wit]. According to the accepted binary diagrams, the βTiCr2 and Ti2Al5 compounds decompose by eutectoid reactions at temperatures of 804 and 990˚C respectively and therefore should be absent at 800˚C. Results of [1997Jew] are in partial agreement with the Ti rich corner of the isothermal section at 800˚C calculated by [1988Gro], particularly in the existence of the three-phase field αTi + β + Ti3Al. However [1988Gro] did not indicate the existence of the ternary compound τ3 in contrast with [1997Jew]. Phase compositions of the Al-30at.% Ti-15 at.% Cr alloy heat treated in air at 1000˚C and 800˚C for 100 h and cooled in air for 5minutes were studied by [1995Bra]. It was concluded that the τ1 phase is stable at 1000˚C, but decomposes at 800˚C (this was not confirmed in later studies of [1995Jew, 1996Jew1, 1996Jew2, 1997Jew]). A partial isothermal section at 600˚C in the Ti rich region is presented in Fig. 8 according to [1992Hay]. [2001Kau] calculated isothermal sections at 1500, 1300, 1200 and 1000˚C. The melting temperature of the τ1 phase was also calculated at 1284˚C. The calculated isothermal sections are in general agreement with the experimental studies. The differences concern the temperature of formation of the ternary phases τ1 and τ2. According to [2003Bar] and [2003Mor], the melting temperature of τ1 should be above 1284˚C and according to [2001Fuj], the temperature of formation of τ2 should be below 1200˚C.
Temperature – Composition Sections A partial vertical section at the constant content of Cr (5 mass%) was presented by [1953Bus] based on microstructural investigations. [1958Tag] constructed partial vertical sections at Cr/ Al=1, at constant Cr content (3 at.%) and at constant Ti content (83 at.%) based on equilibrium studies of Ti rich alloys. These sections are not shown in the present evaluation because the Al-Ti binary diagram was not known in details at the time and the TiAl3 phase was not detected in [1958Tag]. Vertical sections from Al-25Ti (at.%) to Cr between 6 and 28 at.% Cr were constructed in the works of [1997Mab1, 1997Mab2, 2001Ich] based on the results of Landolt‐Bo¨rnstein New Series IV/11E1
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DTA investigations and microstuctural analysis of alloys quenched from temperatures of 1300, 1250, 1150 and 1000˚C. In the present work, the vertical section at Ti content of 25 at.% is however not presented as it is in contradiction to the liquidus surface and reaction scheme. According to the present evaluation, the τ2 phase does not appear in the liquidus, while equilibria between liquid and τ2 are shown in the vertical section at 25 at.% Ti. [2001Kau] calculated a vertical section from Al to the composition of Ti-50Cr (at.%). However [2001Kau] did not take into account the existence of the ternary phase τ3 and did not distinguish the two phases with the C14 structure: the binary phase γTiCr2 and the ternary phase τ2.
Thermodynamics Al and Ti activity of Al-Cr-Ti alloys were calculated by [1999Jac] from vapour pressure measurements in the range of 950–1400˚C in two two-phase regions (Ti3Al+TiAl and TiAl +τ2). The results are presented in Figs. 9 and 10. Calculations based on the Calphad method were published by [1988Gro], [1999Sha] and [2001Kau]. [1988Gro] presented a calculated partial isothermal section in the Ti rich composition area at 800˚C. In the calculations of [1999Sha] and [2001Kau], ordering in the τ3 phase was not taken into account and the γTiCr2 and τ2 Laves phases were not distinguished. Details of phase modelling were not reported in these papers. [1999Sha] calculated the isothermal section at 1000˚C and the liquidus surface. Kaufman [2001Kau] calculated isothermal sections in the range of 1000–1500˚C and a phase fraction diagram for the Al15Cr4Ti7 composition which is close to the composition of the τ1 phase. According to the calculations of [2001Kau], the τ1 phase melted in a narrow temperature range at around 1284˚C.
Notes on Materials Properties and Applications The addition of Cr to binary titanium aluminides, in particular TiAl and TiAl3, has proven beneficial to properties. Specifically, when Cr was substituted up to 3 at.% Al in TiAl, the yield stress and fracture strain in bending increased, with the increase attributed to solid solution strengthening [1991Hua, 1991Mas, 1993Has, 1993Kaw, 1994Li, 2001Sun]. Cr also stabilizes the cubic phase, nominally Ti25Cr8Al67 when substituted for Al in TiAl3. This cubic phase has much greater compressive and bending fracture strains than the tetragonal TiAl3 phase [1990Mab, 1990Zha, 1991Mik, 1993Nak, 1993Wri, 2001Mil, 2001Yam, 2002Bra]. This cubic phase also has improved oxidation resistance over TiAl3 at 1200˚C, with Cr promoting the formation of a pure Al2O3 scale while suppressing the formation of TiO2 [1991Par]. However, this attempt at overcoming the brittleness of TiAl3 at ambient temperature results in a relatively poor creep strength at high temperatures. Therefore, the methods of mechanical alloying and high energy ball milling were used to prepare the ternary (Al,Cr)3Ti L12intermetallic compound additionally strengthened by small volume fractions of incoherent Y2O3 dispersoids [1997Hei, 2001Sur]. Specifically, at room temperature oxide dispersion strengthening with 3 vol.% Y2O3 increases the Vickers hardness level by about 500 units compared to 200 for the single-phase Ti25Cr8Al67 compound and retained it up to 500˚C on the level of 480 HV. A successful introduction of the material into the industrial scale depends particularly on its room temperature ductility and the efficiency of the chosen production DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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process. [2000Mab] tried to raise the plasticity and yield stress of (Al,Cr)3Ti L12-intermetallic compound by alloying with V. The Vickers hardness increases slightly by the addition of vanadium, and is nearly independent on the Cr+V contents, although the higher titanium alloys have larger hardeness. The bend ductility at ambient temperature is not improved by the addition of vanadium. The best ductility is about 0.5% (plastic bend strain) for the Al-25Ti12Cr-3V alloy. [2003Bar] shows that the transition from single-phase L12 alloys to eutectic L12+β(Cr) microstructures is accompanied by enhancement of both strength and plasticity, while retaining a high elastic modulus. The primary creep strains of various TiAl alloys reported in the literature were analyzed through modeling by [2003Zha]. [2002Cal] showed, that mechanical alloying followed by spark plasma sintering can be used to produce satisfactory materials in the TiAl-Cr and Al3Ti-Cr systems. Remarkably high flow stresses were found. Such strengths have the same order of magnitude as extrapolated values on the basis of the HP slope ky for single phase materials. Deformation mechanisms based on dislocation activity are most likely dominating in the case of TiAl-X alloys. The fracture stresses for nanocrystalline TiAl3-Cr are very high and within an order of magnitude of the theoretical shear strength of Ti. Deformation of microcrystalline TiAl3-Cr produces yield stresses above 1 GPa and some ductility in compression. Intermetallic Ti-63Al-7Cr coatings can be used for protection from high temperature oxidation of Ti based as well as γ-TiAl based alloys [1997Ley]. [2003Lee] reported that Al21Ti-15Cr alloy would be the most effective among many potential L12-based alloys as a suitable coating material for TiAl alloy because the phase stability of this alloy. This result is similar to one published by [2006Fox], who reported that the best oxidation stability at 750–1100˚C is exhibited by the Ti25 Cr20Al55 alloy containing near 70–75 at.% of L12 and 25–30% Laves phases. Further improvement of this alloy’s oxidation stability results from synergistic doping by a small amount (0.1–0.2 at.%) of active metals such as Hf and Yas well as Si (above 1%). [2003Nis1, 2003Nis2] proposed three layer coatings formed by two-step Cr and Al diffusion for TiAl based alloys. [2001Zho] studied the TiAl based coatings with high Cr content. It was shown that they exhibited excellent oxidation resistance at temperatures ranging from 900 to 1000˚C. Excellent oxidation resistance can be attributed to formation of protective alumina on the coatings. Oxide formed on the surface of Ti-50Al-15Cr coating after oxidation consisted of A12O3, whereas oxide formed on the surface of Ti-50Al-10Cr coating after oxidation is composed of a mixture of a large portion of A12O3 and a small portion of TiO2 that decrease the oxidation resistance. A review of advanced coatings on high temperature applications, including TiAl based and TiAl3 based alloys has been performed by [2006Nar]. Papers concerned with investigations of material properties are listed in Table 3.
Miscellaneous Interest in the Al-Cr-Ti system was focused on Ti rich alloys. When alloyed with Al and other transition metals, Ti rich alloys can exhibit martensitic transformation from α to β when quenched from the single-phase β-region [1953Bus]. Several studies were concentrated on the nature, kinetics and mechanism of this transformation [1953Bus, 1956Boe, 1956Zwi, 1963Wei]. The results were summarized in the form of T-T-T diagrams for β to α transformation. The growth kinetics of a discontinuous coarsening of lamellar structure in Ti-44Al-0.5Cr (at.%) intermetallic compounds was studied by [1998Jun]. The growth rate is linear during Landolt‐Bo¨rnstein New Series IV/11E1
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the initial period of the reaction and reaches a maximum near the eutectoid temperature. The migrating grain boundary plane maintains no definitive geometrical angular relationship to the lamellar interface. The discontinuous coarsening occurs through the grain boundary diffusion and its activation energy is estimated to be about 230 kJ·mol–1. This kinetic analysis suggests that the driving force for grain boundary migration is the elastic strain energy rather than the chemical and interfacial free energies, despite the comparatively larger value of the last two. Microstructural instability in a fully lamellar Ti-46.5Al-4Cr (at.%) during short-term creep was studied by [2000Cha] under creep tests at 175 MPa and 800˚C. Within Cr depleted regions α2 lamellae transform to γ by different mechanisms. In contrast, Cr enriched regions exhibit a transformation from the ordered hexagonal α2 to a cubic ordered B2 phase through a metastable Cr enriched α2 phase. [2001Lee] studied the general behavior of the CCTs for the Cr-containing alloy (Ti-45Al-2Cr (at.%)) at cooling rates from about 0,5 to about 500˚C·s–1 are similar; the α → γL lamellar reaction predominates at slow cooling rates, and only the α → α2 ordering reaction occurs at fast cooling rates; no massive transformation occurs at intermediate cooling rates. The addition of Cr to the Ti-45,5Al (at.%) binary alloy results in a simple displacement of the CCT curves of the α → γL lamellar reaction towards somewhat lower temperatures, in partial agreement with the depression of a transus temperature. Peculiarities of massive transformation in TiAl-based alloys were studied by [2001Xia]. [2000Wel] used mechanical spectroscopy and creep testing to measure activation enthalpies of viscoelastic relaxation in Ti-46.4Al-4Cr (at.%) alloy. They were determined as 3.9 and 3.6 eV, respectively, which are in the range of those obtained for self-diffusion. Thus, it can be assumed that grain boundary controlled diffusional creep is a dominating process at 800˚C and low stresses. The results indicate that the mechanical properties of Ti-46.5Al-4Cr (at.%) sheet material at 700–800˚C and high stresses are controlled volume diffusion assisted climb of dislocations. [2001Ish] studied opportunity to reduce the hydrogen desorption temperature of Ti3AlHx by substituting Ti with Cr. It was shown that the hydrogen capacity was largely reduced, but the hydrogen desorption temperature was unchanged or increased. Therefore Cr is the worst substitutional element to Ti3Al. The solidification-path of various Al-Cr-Ti alloys, which crystallized (Al,Cr)3Ti as the primary phase, were analyzed by [2001Mor, 2003Mor] using the data of Al-Cr-Ti phase diagram, k-values - functions of composition, and diffusion coefficient obtained by experiment. [1999Yan, 2000Yan] reviewed recent theoretical and experimental investigations of sublattice substitution of alloying element (Cr,V, Mn, Fe, Nb and others) in TiAl and Ti3Al phases. Electronic structure calculations for TiAl3-Cr alloys were performed by [1996Liu] and for TiAl-Cr alloys by [1996Woo] and obtained results were related to mechanical properties. [2007Hes] investigated pre-deformation microstructures, including lamellar, duplex and feathery structures, and microsctructures obtaining after hot forging of Ti-47Al-2Cr alloy. TEM investigations of TiAl based alloys were performed by [1994Zhe, 1994Gao, 1997Mas, 2005App].
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Al–Cr–Ti
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. Table 1 Investigations of the Al-Cr-Ti Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied
Reference
Method / Experimental Technique
[1953Bus]
Homogenezation at 1000˚C, isothermal heat treatment, quenching, optical microscopy
Vertical section 550–1050˚C at constant Ti 5 mass% and Al in the range of 1–5 mass% Ti-3Al-5Cr (mass%), T-T-T diagram
[1953Tay]
XRD, metallography
980, 760˚C, 60–100% Ti isothermal sections
[1958Kor]
Arc-melting, heat treatment in vacuum, DTA, XRD, metallography, hardness, electrical resistivity
Liquidus isotherms 1400–1690˚C, isothermal sections at 600–1200˚C; 60–100% Ti
[1958Tag]
Phase equilibrium studies, metallography, XRD
600, 720, 850 and 1000˚C, 70–100 at.% Ti; isothermal sections and vertical section at Cr/Al=1, constant Cr 3.5 at.%, constant Ti 83 at.%
[1960Enc]
Phase equilibrium studies, metallography, XRD
600–1400˚C, 60–100 at.% Ti
[1960Zol]
X-ray analysis
Al alloy with 0.2% Ti and 2% Cr (mass%). The solubility of Ti in CrAl7
[1963Luz]
Homogenezation at 930˚C, annealing, electrical resistivity and hardness measurements
600–700 and 800˚C, solubiliy of Cr in Ti-6Al (mass%) alloy
[1964Ram]
XRD
As-cast and after annealing at 700˚C; AlCr2-AlTi2 (Ti: 7–50 at.%)
[1974Zwi]
Dilatometery
Annealing at 980˚C for 4 h and 760˚C for 10 d, Ti rich region, isothermal sections
[1978Jac]
X-ray
Ti(AlxCr1–x)2 at 0≤ x≤0.2, hydrogen sorption at 25˚C
[1984Sup]
XRD
Homogenezation at 1000˚C, heat treatment at 800˚C, TiCr2-TiAl2 solid solutions
[1988Gro]
Compacting and homogenizing in vacuum. Electron microprobe method, CALPHAD calculation
Annealing at 1100˚C for 4 h, then long period at 800˚C and quenching in water. Ti rich region, tie lines and calculated isothermal section.
Landolt‐Bo¨rnstein New Series IV/11E1
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. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method / Experimental Technique
[1990Zha]
Arc melting after cold isostatically pressed at 350 MPa powders and then hot isostatically pressed (HIPping) at 1200˚C and 175 MPa, 4 h in argon. Optical microscopy, SEM, energy dispersive spectrometry (EDS), X-ray diffraction, microhardness, DTA
[1991Mab]
XRD, TEM, SEM/EDX, DTA, microhardness Al-(20 to 80)Cr-50Ti (at.%). Isothermal section at 1000˚C
[1991Moh]
Arc-melting, homogenization at 1200˚C. 1300˚C, Ti-51Al-2.1Cr; site occupancy equilibration, quenching X-ray of Cr microanalysis, TEM
[1991Par]
Arc-melting, cyclic oxidation furnace at 1200˚C, gravimetry, XRD, SEM
1200˚C, Al67Cr8Ti25; oxidation resistance
[1991Spa]
Bulk ingots, arc melting, homogenization at 1200˚C for 16 h, X-ray diffraction, EPMA.
After homogenization. Ti27.6Cr6.8Al65.6 alloy.
[1992Dur]
Engel-Brewer theory
Al67Ti25Cr8 calculation of electron concentration
[1992Kim]
Arc-melting, annealing at 1050˚C in vacuum, forging at 1200˚C In situ XRD, optical microscopy, EPMA
1200˚C, Ti-48Al-2Cr (at.%)
[1992Mor]
HIP-treatment at 1140˚C and 100 MPa, spray deposition XRD, TEM
Al67Ti25Cr8 ordering
[1992Nic]
Bulk alloys, arc-melting from pure elements. HIPping at 1200˚C and 172 MPa, 2 h, in argon. X-ray diffraction, electron microprobe, EDS, hardness.
As-HIPed alloys. Al-Cr-50Ti (at.%). Isothermal section at 1200˚C.
[1992Win]
Bulk alloys, induction melting in argon, Ti25Cr8Al67. Crystal structure and lattice then homogenization at 1050˚C - 72 h in parameter of as-cast and homogenizated samples. vacuum. Optical microscopy, scanning electron microscopy, X-ray diffraction, EDS
[1992Zhe]
Bulk alloys, arc-melting, homogenization Annealing 1150–1200˚C for 24–168 1000˚C for 168 h in vacuum. Optical h in argon and cooling in air. microscopy, X-ray diffraction, SEM, TEM, Ti-44.3Al-3Cr (at.%). Phase composition. EDAX, microhardness.
[1993Ers]
Calculation, using the Linear Muffin Tin Orbital Method
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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As-HIPed alloys. Ti25Cr8Al67, Ti25Cr9Al66, stress-strain behavior
TiAl, Ti25Cr25Al50, Ti50Cr25Al25
Landolt‐Bo¨rnstein New Series IV/11E1
Al–Cr–Ti
7
. Table 1 (continued) Reference
Method / Experimental Technique
Temperature/Composition/Phase Range Studied
[1993Hao]
Diffusion couple technique, optical microscopy, EPMA.
[1993Nak]
1.Powder mixtures obtained at about As-homogenizated alloys. The range 260 MPa and sintered in vacuum at 1000 Al-(20 to 80)Cr-50Ti (at.%), isothermal or 1150˚C for 12 or 24 h. sections. 2.Bulk alloys, arc-melting, homogenization at 1150˚C for 2 days in vacuum. X-ray diffraction, optical microscopy, SEM, EDS.
[1994Boh]
Annealing in vacuum, XRD, SEM/EDX (backscattered)
700–1000˚C, Ti-48Al-2Cr (at.%)
[1994Kla]
Bulk alloys, arc melting from pure elements, HIPping at 1200˚C, 172 MPa, 2 h, X-ray diffraction, EPMA, backscattered electron imaging, EDS, hardness.
As-HIPed alloys. The range Al-(20 to 80) Cr-50Ti (at.%), isothermal section.
[1994Sok]
Arc melting from pure elements under Al-corner at up to 25 at.% Ti and 20 at.% Ar atmosphere. X-ray and microstruture Cr. Quenching in ice water after analyses annealing at 497˚C for 1000 h. Isothermal section.
[1994Zha]
Bulk alloys, arc-melting, homogenization As homogenizated alloy: Ti-42Al-3Cr at 1100˚C for 30 h and then 1050˚C for (at.%), phase composition 30 h. TEM, SEM, X-ray diffraction
[1994Zhe]
TEM
[1995Bra]
Bulk alloys, arc melting. X-ray diffraction, Heat treating at 1200˚C for 100 hours SEM, microhardness. in a 95%Ar-5%H2 and then furnace cooling. Heat treating in air at 800˚C for 100 hours and then air cooling to room temperature. Two alloys: Ti40Cr15Al45 and Ti30Cr15Al55. Phase composition.
[1995Hao]
Diffusion couples. EPMA, X-ray diffraction
The range Ti-(0 to 70)Cr-(0–75) Al (at.%). Isothermal section at 1000˚C.
[1995Jew]
Bulk alloys and diffusion couple technique. X-ray diffraction, optical microscopy, SEM, EDX.
Annealing at 800˚C for 750 h or 1000˚C for 650 h and then water quenching. Cr rich region: Cr-(0 to 70)Ti-(0–30)Al (at.%). Isothermal sections, phase equilibria at 800 or 1000˚C.
[1995Kog]
Arc-melting, homogenization at 1100˚C, 1100˚C; 17 alloys Al-(25–28)Ti-(9–14)Cr annealing, quenching, XRD (at.%)
Landolt‐Bo¨rnstein New Series IV/11E1
Annealing 1000˚C for 200 h, 1000˚C 100 h in vacuum. The range Ti-(20–80)Cr-50Al (at.%). Isothermal section
25˚C, Ti-45Al-3Cr (at.%)
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature/Composition/Phase Range Studied
[1995Nak]
Arc-melting, XRD, optical metallography, 1200˚C, Ti-48Al-2(Cr, Mo, Mn) (at.%) SEM, TEM, EPMA
[1996Jew1]
Thirteen bulk alloys, arc-melting, homogenization in vacuum for 2 h at 1380˚C or 1270˚C and furnace cooling. Optical microscopy, SEM, EDX, X-ray diffraction, DSC for enthalpy determination.
Annealing at 800˚C for 1000 h or 1000˚C for 500 h, then water quenching. Cr-Ti rich alloys with 0 to 30 at.% Al. Isothermal section, phase equilibria at 800 and 1000˚C, tie lines.
[1996Jew2]
Eleven bulk alloys, arc-melting, homogenization in evacuated ampules at 1140 or 1250˚C for 3 to 4 h. Diffusion couple technique. Optical microscopy, SEM, EDX, X-ray diffraction, DSC measurements.
Annealing 800˚C for 100 to 1000 or or 1000˚C for 100 to 500 h, respectively, and water quenching. Al-(0 to 80)Cr-40Ti (at.%) alloys. Isothermal section, phase equilibria at 800 and 1000˚C.
[1997Mab1], [1997Mab2]
Compact samples from powders, sintering in vacuum for 48 h at 1200˚C, then blowing-air quenching. Button ingots, arc-melting, homogenization at 1200˚C in vacuum for 2 d, then aircooling. X-ray diffraction, optical microscopy, SEM, EDX, DTA, microhardness, bend testing.
After homogenization (isothermal section) or heat treatment: 1300˚C for 2 h, 1250˚C for 4 h, 1150˚C for 48 h, 1000˚C for 144 h, then water quenching (vertical section). Al-(0 to 70)Cr-(0 to50)Ti (at.%). Isothermal section at 1150˚C and vertical section for Al-25Ti-(0–28)Cr (at.%).
[1997Fan]
Mechanical alloying for 40 h of milling. X- Heat treating at 700˚C and cooling to ray diffraction, TEM, DSC. room temperature. The Al-25Ti-8Cr (at.%) alloy.
[1997Jew]
Twelve bulk alloys, arc-melting, homogenization in vacuum at temperature below solidus for 2 to 4 h, then furnace cooling. Optical microscopy, SEM, EDX, X-ray diffraction, DSC measurements.
Annealing 800˚C for 100, 150 and 1000 h or 1000˚C for 100,160 and 500 h, then water quenching. Isothermal section at 800 and 1000˚C, temperature of invariant reaction.
[1997Par]
Six button ingots, arc melting, homogenization at 1150˚C for 48 h in vacuum, then furnace cooling. X-ray diffraction, EDX, optical microscopy, SEM, compression tests.
After homogenization. Al-(21 to 30)Ti-(6, 15, 23)Cr (at.%). Isothermal section at 1150˚C.
[1997Xu]
Diffusion couple TiAl3-Cr and diffusion triples Cr-Ti-TiAl3. Optical microscopy, SEM, EDX.
Annealing at 1050˚C for 496 or 532 h. Isothermal section at 1050˚C.
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Al–Cr–Ti
7
. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method / Experimental Technique
[1998Has]
Bulk samples, arc melting. Diffusion couple technique. Homogenization at 1050˚C (bulk) and 1100˚C (diffusion couple) for 345.6 ks in vacuum, then furnace cooling. Optical microscopy, X-ray diffraction, EPMA, DTA. Calculation with Thermocalc program.
Heat treatment at 1200˚C for 172.8 ks, then ice water quenching. Ti rich region. Calculation of isothermal section at 1200˚C and phase equilibria between αTi, β and TiAl phases at 1200˚C accoding to experimental data.
[1998Jun]
Button ingots, arc melting, homogenization at 1330˚C for 3 h and air cooling. Optical microscopy, X-ray diffraction, TEM.
Aging treatment at 650 to 1200˚C. The Al-44Al (at.%) alloy with 0.5 at.% Cr. Lattice parameters of Ti3Al and TiAl phases at various temperatures.
[1999Hao]
Arc-melting, homogenization at 900˚C, quenching, electron microscopy/EDAX
900˚C, Ti-(46–53)Al-(1–5)Cr (at.%)
[1999Jac]
Vapour-pressure measurements (Knudsen cell)
1000–1400˚C, Ti-47Al-2Cr, Ti-47Al-13Cr, Ti-51Al-12Cr
[1999Nic]
Arc-melting, HIP at 1200˚C and 172 MPa, 1200˚C, Al-25Ti-(9,12)Cr, EPMA, XRD Al-(23,27)Ti-10Cr (at.%)
[1999Sha]
Bulk ingot, arc melting. Melt-spun ribbon. EPMA, TEM, EDXA, calculation phase equilibria.
[2000Kai1]
Bulk ingot, arc melting, homogenization Annealing at 1000˚C for 168 h and ice 1200˚C for 3 h. Diffusion couple. SEM, water quenching. Ti-(15 and 18 to EDS, TEM, optical microscopy. 27)Al-(3 and 5)Cr (at.%). Phase equilibria at 1000˚C in the (βTi) region.
[2000Cha]
Powder metallurgical process, heat treatment in vacuum, creep-testing, TEM,TEM/EDX, SEM
800˚C, Ti-46.5Al-4(Cr,Nb, Ta, B) Influence of creep to microsctructure instability
[2000Kai2]
Bulk ingot, arc melting. EPMA.
Annealing at 1000˚C for 168 or 504 h, at 1200˚C for 168 h, at 1300˚C for 24 h, then ice water quenching. Ti rich region: Ti-(35 to 47)Al-(0.5 to12)Cr (at.%). Phase equilibria at 1300, 1200, 1000˚C between (αTi), β, TiAl.
[2000Sha]
Bulk ingots, arc melting, melt-spun ribbons. EPMA, TEM, EDX.
After casting and after annealing at 450˚C for 100 h. Ti-(20 to 52)Al-(7.5 to 20)Cr (at.%). Metastable β-based structures.
Landolt‐Bo¨rnstein New Series IV/11E1
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Ti-40Al-10Cr (at.%), Ti-50Al-10Cr (at.%), Ti-52Al-20Cr (at.%). Phase equilibria at 1000˚C, calculated liquidus surface projection and isothermal section at 1000˚C.
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature/Composition/Phase Range Studied
[2001Fuj]
X-ray diffraction, TEM, optical microscopy
Heat treatment at 1200˚C for 4 d, 1000˚C for 7 d, the ice water quenching. 13 alloys in Ti rich region with (38 to 60) at.% Al, (0 to 21) at.% Cr. Phase equilibria at 1200 and 1000˚C.
[2001Ich]
DTA, XRD, isothermal heat treatment
1250–1350˚C, 50–100 at.% Al (liquidus and solidus τ1 and vertical section at Ti 25 at.%)
[2001Kan]
Calculation of phase equilibrium by cluster variation method.
Calculation TiAl/Ti3Al and αTi/Ti3Al phase boundaries at 1000˚C in Al-Cr-Ti phase diagram.
[2001Kau]
CALPHAD method
Isothermal sections at 1500, 1300, 1200, 1000˚C, phase fraction diagram for Al15Cr4Ti7 and vertical section from Al to Ti-50 at.% Cr
[2001Lee]
Arc-melting, homogenezation at 1200˚C 900–1200˚C, Ti45Al-2Cr, kinetic study in vacuum, TEM
[2001Sun]
Bulk ingots, arc melting, HIPping at Annealing 900˚C for 8 h and air cooling. 1200˚C for 3 h, heat treatment at 1200˚C Ti52Al48–xCr (x = 0, 1, 2, 4, 6). Phase equilibria between Ti3Al, TiAl, τ3. for 12 h, air cooling. X-ray diffraction, XPS, TEM, SEM, EPMA, tensile test, creep test.
[2001Yam], [1999Yam]
Arc-melting, homogenezation at 1177˚C 1177˚C, in vacuum, quenching; XRD, optical Al-(9.24–12.32)Cr-(25.38–26.71)Ti (at.%) microscopy, SEM, EPMA
[2003Bar]
Bulk ingots, arc melting. Optical As-cast alloys. Al-(14 to 27) at.%Ti-(11 to microscopy, SEM, X-ray diffraction, DTA, 35) at.%Cr. Univariant equilibria with hardness, compression test, Young’s liquid phase. modulus.
[2003Lee]
Arc-melting, homogenization at 1150˚C 800˚C; Al-21Ti-23Cr, Al-21Ti-15Cr, in vacuum 10–1 Pa, XRD, TEM, SEM/EDX, Al-30Ti-15Cr optical microscopy
[2003Mor]
Progressive type of solidification equation based on experimental data on phase diagram, partitioning coefficients, diffusion coefficients
[2003Nis1], [2003Nis2]
Bulk ingot, arc melting. EPMA, SEM, X-ray Coated by Cr at 1300˚C and Al at diffraction, oxidation resistance studies 1000–1300˚C and air quenching. Melted and hardness measurements. samples were annealing at 1000˚C for 86.4 ks. Phase equilibria between τ1, τ2, β, TiAl.
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Calculation of solidification path (1250–1350˚C) of various Al-Cr-Ti alloys with Ti content of 15–29 at.% and Cr content of 6–20 at.%
Landolt‐Bo¨rnstein New Series IV/11E1
Al–Cr–Ti
7
. Table 1 (continued) Reference
Method / Experimental Technique
Temperature/Composition/Phase Range Studied
[2003Ram]
Cr and Ti magneton sputtering on Al 600, 700˚C heat treatment, target - thin films EPMA, SEM, XRD, TEM, TiAl-(0.6–4.6)Cr gravimetry
[2004Nay]
Mechanical alloying for 40 h of milling. X- After milling. The Cr8Ti27Al67 alloy. Lattice parameters. ray diffraction, TEM.
[2004Wan]
Skul melting, XRD, SEM/BSE, EPMA
Ti-46Al-8Cr-2.65(Nb,W,B)
[2007Hes]
Arc-melting, vacuum heat treatment, cooling in air, TEM, SEM and optical microscopy
1100˚C homogenization, 1400˚C heat treatment, Ti-47Al-2Cr (at.%)
. Table 2 Crystallographic Data of Solid Phase Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(Al) < 660.452
cF4 Fm3m Cu
β, (βTi,Cr) < 1863
cI2 Im3m W
Lattice Parameters [pm]
dissolves up to 0.8 at.% Ti at 665.7˚C [2008Cor] and 0.35 at.% Cr at 661.5˚C [2007Wit] at 25˚C [Mas2]
a = 404.96
0 to 100 at.% Cr at 1410 to 1359˚C [2008Iva] dissolves up to 40 at.% Al at 1000˚C [1999Sha] for Ti6.5Cr68.6Al24.9, quenched from 800˚C [1997Jew] for Ti12.9Cr58.0Al29.1, quenched from 800˚C [1997Jew] dissolves up to 46 at.% Al at 1350˚C [2008Cor] at 25˚C [Mas2] dissolves up to 49.4 at.% Al at 1491˚C [2007Wit] [V-C2]
a = 294.4 a = 294.6 (Cr) < 1863
a = 288.48
(βTi) 1670 - 882 a = 330.65 (αTi) < 882
Landolt‐Bo¨rnstein New Series IV/11E1
hP2 P63/mmc Mg
Comments/References
dissolves up to 51.4 at.% Al at 1456˚C [2007Wit] and 0.6 at.% Cr [2008Iva] a = 295.06 c = 468.35
at 25˚C [V-C2]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C] CrAl7 < 790
mC104 C2/m V7Al45
Cr2Al11 940 - 785
orthorhombic
oC584 Cmcm
Lattice Parameters [pm]
a = 2519.6 b = 757.4 c = 1094.9 β = 128.7˚C
Comments/References Cr2Al13 at room temperature 13.5 at.% Cr [2008Cor]
CrAl5 16.9 to 19.2 at.% Cr quenched from 920˚C [2008Cor]
a = 1240 b = 3406 c = 2020 a = 1252.1 b = 3470.5 c = 2022.3 a = 1260 b = 3460 c = 2000
single crystal “εCrAl4” [2008Cor]
CrAl4 < 1030
hP574 P63/mmc MnAl4
a = 1998 c = 2467 a = 2010 c = 2480
at room temperature, 20.9±0.3 at.% Cr [2008Cor] at Cr rich border at 1000˚C [2008Cor]
γCr4Al9 1170 - ≈1060
-
-
30 to 34 at.% Cr [2008Cor]
βCr4Al9 ≲ 1060
cI5 I43m Cu4Al9
a = 912.3
30 to 35 at.% Cr at Al rich border at 920˚C [2008Cor]
αCr4Al9 ≲700
hR156 R3m Cr4Al9
a = 1291 c = 1567.7
32.8 to 35 at.% Cr [2008Cor]
βCr5Al8 cI52 ≈1350 - ≈1100 I43m Cu5Zn8
35 to 42 at.% Cr, a = 910.4 to 904.7 quenched from liquid [2008Cor]
αCr5Al8 ≲ 1160
hR78 R3m Cr5Al8
a = 1281.3 c = 795.1
Cr2Al < 910
tI6 I4/mmm MoSi2
≈ 65.5 to ≈ 71.4 at.% Cr a = 300.5 to 302.8 [2008Cor] c = 864.9 to 875.5
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
35 to 42 at.% Cr [2008Cor]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C]
Lattice Parameters [pm]
Comments/References
X ≲ 400
Cr5Al3 or Cr3Al super lattice
-
≈ 75 to ≈ 80 at.% Cr possibly metastable [2008Cor]
w
-
-
In quenched alloy of 60 to 100 at.% Cr, like metastable wTi [2008Cor]
γTiCr2 1359 - 1269
hP12 P63/mmc MgZn2
a = 493.2 c = 800.50
Laves phase C14 at 25˚C for Ti1.12Cr2 [2008Iva]
βTiCr2 1271 - 804
hP24 P63/mmc MgNi2
αTiCr2 < 1223
cF24 Fd3m MgCu2
ωTi Cr2 < 450
hP3 P3m1 ωTiCr2
β0 1427–1159
cP2 Pm3m CsCl
Laves phase C36. Dissolves up to 4 at.% Al [1996Jew1] at 25˚C [2008Iva]
a = 493.2 c = 1601.0
Laves phase C15. Dissolves up to 1 at.% Al [1997Xu] at 25˚C for TiCr1.9 [2008Iva]
a = 693.20
from 3 to 9 at.% Cr, metastable phase [2008Iva] at 4.6 at.% Cr
a = 461.6 c = 282.7
high temperature ordered B2 phase [2007Wit]
-
hP8 α2, Ti3Al < 1293 P63/mmc (up to 10 GPa) Ni3Sn
from 18 at.% Al at 600˚C to 38.5 at.% Al at 1119˚C D019 ordered phase (α2-Ti3Al) [2007Wit] dissolves up to 2.5 at.% Cr at 800 to 1000˚C [1997Jew] at 25 at.% Al [2006Sch]
a = 576.5 c = 462.5
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Al–Cr–Ti
. Table 2 (continued) Phase/ Temperature Range [˚C] γ, TiAl < 1456
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
tP4 P4/mmm AuCu
from 46.7 at.% Al at 1119˚C to 62.1 at.% Al at 1432 at.% Al [2007Wit] L10 ordered phase (γ-TiAl) Dissolves up to 8 at.% Cr at 1000˚C and up to 4.5 at.% Cr at 800˚C [1997Jew] at 50.0 at.% Al [2006Sch]
a = 400.0 c = 407.5 a = 400.6 c = 406.1 a = 398.1 c = 406.6 a = 400.5 c = 406.3 a = 399.0 c = 406.1 η, TiAl2 < 1224
ζ, Ti2Al5 1432 - 976
tI24 I41/amd HfGa2
tetragonal super structure of AuCu-type
ε (h), TiAl3(h) 1396 - 734
tI8 I4/mmm TiAl3(h)
ε (l), TiAl3(l) < 932
tI32 I4/mmm TiAl3(l)
TiAl3(m)
cP4 Pm3m AuCu3
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
Comments/References
for Ti51.0Cr1.7Al47.3, quenched from 1000˚C [1997Jew] for Ti36.6Cr8.0Al55.4, quenched from 1000˚C [1997Jew] for Ti50.9Cr2.0Al47.1, quenched from 800˚C [1997Jew] for Ti48.2Cr4.2Al47.6, quenched from 800˚C [1997Jew] dissolves about 2.3 at.% Cr at 976˚C [2007Wit] for TiAl2 [2006Sch]
a = 397.0 c = 2497.0 a = 396.1 c = 2415.2 a = 395.2 c = 2411.7
for Ti34.3Cr3.5Al62.2, quenched from 1000˚C [1997Jew] for Ti34.2Cr3.1Al62.7, quenched from 800˚C [1997Jew] from 65 at.% Al at 1432˚C to 70.4 at.% Al at 1396˚C [2007Wit]
a = 392.30 c = 1653.48 a = 390.53 c = 2919.21
at 68.75 at.% Al [2006Sch] at 71.4 at.% Al [2006Sch] from 74.2 at.% Al at 1396˚C to 76.5 at.% Al at 734˚C [2007Wit] [2006Sch]
a = 384.9 c = 861
from 74.6 at.% Al at 932˚C to 75.7 at.% Al at 665.7˚C [2007Wit] [2006Sch]
a = 387.7 c = 3383.2 a = 397.2 a = 396.7
by splat cooling [2006Sch] by mechanical alloying [2006Sch]
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7
. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Ti3Al5 < 810
tP32 P4/mbm Ti3Al5
* τ1, (Ti1–x+yCrx) Al3–y ≲ 1370
cP4 Pm3m AuCu3
Lattice Parameters [pm] a = 1129.3 c = 403.8
[2006Sch]
L12 ordered phase stabilized by Cr at 7 to 13 at.% Cr at 1200˚C [1992Nic], at 0.28 ≤ x≤ 0.48 and 0≤y≤0.68 at x = 0.32, y = 0.32 for Ti25Cr8Al67, annealed at 1200˚C [1992Nic] at x = 0.32, y = 0,32 for Ti25Cr8Al67, as-cast [1990Zha] at x = 0.272, y = 0.376 for Ti27.6Cr6.8Al65.6, annealed at 1200˚C [1991Spa] at x = 0.32, y = 0.32 for Ti25Cr8Al67 annealed 800˚C [1992Win] at x = 0.32, y = 0.32 for Ti25Cr8Al67, mechanical alloying for 40 h, annealed 700˚C [1997Fan] at x = 0.32, y = 0.32 for Ti25Cr8Al67 mechanical alloying for 40 h [2004Nay] at x = 0.476, y = 0.588 for Ti27.8Cr11.9Al60.3, quenched from 1000˚C [1997Jew] at x = 0.42, y = 0.656 for Ti30.9Cr10.5Al58.6, quenched from 1000˚C [1997Jew]
a = 395.4
a = 395.4 a = 395.8
a = 396.3
a = 396.3**
a = 397.3**
a = 394.5
a = 395.1
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Comments/References
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Al–Cr–Ti
. Table 2 (continued) Phase/ Temperature Range [˚C] * τ2, Ti(Cr,Al)2 < 1200
* τ3 < 1050
Pearson Symbol/ Space Group/ Prototype hP12 P63/mmc MgZn2
cP2 Pm3m CsCl
Lattice Parameters [pm]
Comments/References C14 ordered phase stabilized by Al
a = 495** c = 805** a = 503** c = 820.7** a = 504.42 c = 824.62 a = 505.5 c = 827.8 a = 504.7 c = 825.3 a = 504.3 c = 824.3 a = 504.5 c = 824.1 a = 502.8 c = 821.3 a = 492.6 c = 803.2 a = 493.2 c = 803.6 a = 495.6 c = 809.0
for Ti33..3Cr56.7Al10 at 1000˚C [1996Jew1] for Ti33..3Cr36.7Al30 at 1000˚C [1996Jew1] for Ti31.5Cr26.7Al41.8, quenched from 1000˚C [1997Jew] for Ti33.6Cr29.3Al37.1, quenched from 1000˚C [1997Jew] for Ti35.8Cr30.2Al34.0, quenched from 1000˚C [1997Jew] for Ti31.6Cr29.3Al39.1, quenched from 800˚C [1997Jew] for Ti33.3Cr27.6Al39.0, quenched from 800˚C [1997Jew] for Ti34.2Cr34.5Al31.3, quenched from 800˚C [1997Jew] for Ti33..3Cr61.2Al5..5 as-cast [1978Jac] for Ti33..3Cr56.7Al10 as-cast [1978Jac] for Ti33..3Cr53.4Al13..3 as-cast [1978Jac] B2 ordered phase stabilized by Cr for Ti56.9Cr7.6Al35.5, quenched from 1000˚C [1997Jew] for Ti52.2Cr13.3Al34.5, quenched from 1000˚C [1997Jew] for Ti46.5Cr21.6Al31.9, quenched from 1000˚C [1997Jew] for Ti53.9Cr14.1Al53.9, quenched from 800˚C [1997Jew] for Ti46.9Cr20.8Al32.3, quenched from 800˚C [1997Jew] for Ti55Cr3Al42, annealed at 1050˚C [1994Zha] for Ti50Cr25Al25 [1964Sch]
a = 316.3 a = 315.0 a = 312.8 a = 315.4 a = 313.3 a = 319.7 a = 312
Note: ** - Lattice parameters are derived from figures
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Al–Cr–Ti
7
. Table 3 Investigations of the Al-Cr-Ti Materials Properties Reference
Method / Experimental Technique
Type of Property
[1963Luz]
Resistivity measurements, hardness
Solubility of Cr in Ti-6 mass% Al, electrical and mechanical properties
[1978Jac]
Sivert’s method
Hydrogen capacity of Ti(Cr1–xAlx) (0≤x≤0.2) at 7 MPa, 293 K and 4 MPa, 80 K
[1984Sup]
Faradey microbalance method
The temperature dependence of magnetic susceptibility of TiCr2 - TiAl2 alloys
[1990Mab]
Hot-pressed compact specimens, TEM, X- Compression tests at temperatures from ray diffraction, Vickers microhardness, 300 to 1300 K of Ti25Cr8Al67 with L12 structure compression tests at elevated temperature
[1990Zha]
Light microscopy, SEM, DTA, compressing tests, Vickers hardness
[1991Bea]
Arc-melting, HIP at 1220˚C, hot-rolling at 1300˚C, Ti48Al-2Cr 1260˚C or forging at 1025, heat treatment, optical microscopy, SEM
[1991Hua]
Light microscopy, AEM, 4-point bending Effect of alloying with Cr on the and tensile test microstructure and plasticity of Ti-(45–54) Al-(1–4)Cr (at.%) alloys prepared as meltspun ribbon, consolidated by cold compaction, HIPped and extruded
[1991Mab]
X-ray diffractometry, TEM, SEM, mechanical testing
[1991Mae]
Mechanical properties of Ti-35.5Al-1.5Cr Tensile tests at room temperature and 1073 K with a strain rate of 8.3·10–5·s–1. (at.%) at room temperature and 800˚C Creep rupture tests were carried out on the as cast materials at 800˚C in air with the applied stress between 180–250 MPa.
[1991Mas]
Temperature and strain rate Mechanical properties of γ (TiAl)-based dependencies of the tensile properties alloy cointaining bcc phase of were examined by tensile tests in Instron composition of Ti-47Al-3Cr (at.%) type machine. Compression tests for isothermally forged alloys were carried out below 1073 K at 7.5·10–4 sec–1 initial strain rate. The deformed and fractured specimens were characterized by optical microscopy, SEM, TEM and XRD (diffraction and photograph).
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Mechanical properties of as-casted and HIPped alloys of Ti25Cr8Al67 and Ti25Cr9Al66
Mechanical properties of Ti25Cr8Al67 at 400–1000˚C
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Al–Cr–Ti
. Table 3 (continued) Reference
Method / Experimental Technique
[1991Mik]
Hardness (HDP), compression and bend testing, ultrasonic vawe velocity measuremenrs, dilatometry
Mechanical properties, including elastic moduli; average coefficient of thermal expansion for temperature interval 200–1200˚C foras cast and HIPped Ti25Cr8Al67
[1991Nic]
Hardness (HDP), compression and bend testing, ultrasonic wave velocity measuremenrs
Hardness, crack loading, Young’s modulus of Ti25Cr8Al67
[1991Par]
Tests in cyclic oxidation furnace in static air at 1200˚C for 200 cycles: 1 h immersion at 1200˚C followed by 20 min cool down to rome temperature
Oxidation resistance of Ti25Cr8Al67 at 1200˚C
[1992Miz]
Tensile tests at 1200˚C at an initial strain Annealing time dependence of tensile strength and elongation in isothermally rate 5.4 · 10–4 s–1 forged Ti-46.1Al-3.1Cr (at.%) at 1200˚C. Annealing time from 0 to 170 h
[1992Mor]
Hardness, toughness testing in a range of Mechanical properties loads
[1992Nic]
Hardness (DPH)
Hardness profile and load necessary to couse cracks for 25Ti-Al-xCr 0<x<20 at.% in as cast and HIPped states
[1992Zhe]
Microhardness
Microhardness of phases in Ti-44.3Al-3Cr (at.%) alloy
[1993Has]
Tensile test, strain rate 10–4·s–1
Yield stress, fracture stress, elongation and Yong’s modulus at ambient temperature of Ti-48.7Al-1.3Cr (at.%) alloy annealed at 1350˚C for 1 h
[1993Kaw]
Four-point bend tests at room temperature at strain rate 10–3·s–1
The effect of Ti/Al ratio and Cr additions on mechanical properties in TiAl-base alloys
[1993Nak]
Vickers microhardness using a 25 g load, three-point bend tests at ambient temperature using an Instron mashine at a constant cross-head deflection rate of 0.1 mm·min–1
Mechanical properties of sintered specimens of Ti25Cr8Al67, Ti25Cr9Al66 and melted button ingots of Ti25CrxAl75–x, where x=8, 13, 14 and 16 at.% as a function of phase composition, and porosity
[1993Wri]
Hardness (HDP)
Hardness as a function of aging temperature for 1 h heat treatments of Ti27.5Cr8Al27.5 after solution treatment for 1 h at l000˚C and oil quenched
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Type of Property
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. Table 3 (continued) Reference
Method / Experimental Technique
Type of Property
[1994Kla]
Hardness (HDP)
Constructing of contour map of hardness values and contour map of cracking resistance load for Al-Ti50Al50Ti50Cr30Al20-Cr80Al20 composition region
[1994Li]
Conventional tensile tests were carried out at room temperature in air using a servohydraulic Instron model 8500 testing machine equipped with a digital control unit
Tensile properties of as HIPped Ti-44Al-2Cr alloy
[1995Bra]
Heat treating in air at 800 and 1000˚C for High temperature oxidation resistance of 100 h Ti-45Al-15 Cr and Ti-55Al-15Cr (at.%) alloys
[1995Whi]
Compression, creep test
Mechanical properties at elevated temperatures for alloys Al-25Ti-9(Cr, Fe, Mn)
[1996Jew1] Hardness (HDP)
Variation of hardness of the Ti(Cr,Al)2 phase with composition
[1997Hei]
The hardness at temperatures up to 500˚C of oxide dispersion strengthened (3 vol.% Y2O3) Ti25Cr8Al67 produced by mechanical alloying and hot compressing at 800˚C
The elevated temperature Vickers hardness using 200 N load
[1997Mab2] Microhardness, bend test
Interrelations, between starting phase composition, porosity and ductility of TiAl3 alloyed with Cr
[1997Ley]
Interrupted weight gain tests in laboratory air at 750 and 900˚C
Oxidation behavior of Ti-Al-Cr coatings deposited by magnetron sputtering of Ti-(44 -66)Al-(4 -22)Cr (at.%) alloys
[1998Jim]
Strain-rate-change tests at strain rates ranging from 10 –6 to 10–3 s–1 in a universal testing machine. The mechanical tests were performed in compression in the temperature range from 800 to 1100˚C.
Mechanical properties at elevated temperatures of Ti-47Al-3Cr (at.%) alloy
[2000Cha]
Creep-test
Mechanical properties of Ti-46.5Al-4(Cr, Nb,Ta,B)
[2000Coe]
Hardness tests were carried out with loads of 70 and 300 mN in an ultramicroindentation device with a Vickers indenter
Mecanical properties of thing films TiAl-(0.6–4.6)Cr (at.%) produced by magnetron sputtering
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Al–Cr–Ti
. Table 3 (continued) Reference
Method / Experimental Technique
Type of Property
[2000Mab]
Vickers hardness, bend tests, compressive tests
Mechanical properties of button ingots of Al-25Ti-5Cr-4V and Al-xTi-yCr-3V(x = 25, 27, 29, y = 6, 8, 10, 12, 14) alloys prepared by arc-melting
[2000Ohi]
Vickers hardness
Hardness and microstructural variations during aging of Al-Cr-Ti ternary Ll0- and Ll2-phase alloys
[2000Wel]
Subresonance torsion apparatus
Mechanical properties of Ti-46.5Al-4(Cr, Nb,Ta,B) consisting of TiAl and Ti3Al up to 1000˚C
[2001Ish]
Sorbtion and desorption of hydrogen
The effect of sustituting of Ti for Cr on the hydrogen capacity and desorption temperature of Ti75–xCrxAl25 (x = 0, 15, 25)
[2001Mil]
Young’s modulus, Vickers hot hardness, The Vickers hot hardness at load p = stress-strain curves of TiAl3 and 2.34 N, four-point bending tests at a Ti28Cr12Al61 cross-head speed of 0.2 mm· min–1, compression tests with a strain rate of 10–3·s–1, indentation method, using nine trihedral diamond indenters with tip angles of 45, 50, 55, 60, 65, 70, 75, 80 and 85˚C.
[2001Sun]
Tensile tests at room temperature and at 900˚C in air using an MTS testing machine, a nominal engineering strain rate of 2·10–3·s–1 was used for the tensile tests; creep tests were carried out at 800˚C
[2001Zho]
Isothermal oxidation tests at 900, 950 and Effect of Ti-50Al-15Cr and Ti-50Al-10Cr 1000˚C in static air. A thermal balance (at.%) coatings on the oxidation behavior was used to measure the weight change of TiAl alloys continuously. Cyclic oxidation tests using a cyclic oxidation furnace in static air at 950 and 1000˚C. Each cycle consisted of an hour immersion in static air in the furnace followed by a 20 min cool down to room temperature.
[2002Bra]
Vickers hardness was found with a Leco Effect of alloying with Cr on the M-400 hardness tester (1 kg, 20 s), three- mechanical properties of TiAl3, point-bend testing was adapted from the (Ti31Cr3Al66, Ti28Cr6Al66, Ti25Cr9Al66) ASTM standard C1161–94
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
MSIT1
Effect of alloying with Cr on the microstructure, tensile properties and creep resistance of Ti52Al48–xCr alloys (x = 0, 0.5, 1.0, 2.0, 4.0, and 6.0 at.%
Landolt‐Bo¨rnstein New Series IV/11E1
Al–Cr–Ti
7
. Table 3 (continued) Reference
Method / Experimental Technique
Type of Property
[2002Cal]
Compression tests were made as a function of temperature in a vacuum atmosphere (<10–3 Pa). Most compression tests were carried out at a constant crosshead rate of 0.002 mm·s–1 and others at 0.0033 mm·s–1.
Evaluation of the mechanical properties of Ti50Al50, Ti50Al48Cr2, Al67Ti25Cr8 at temperatures ranging from 298 to 773 K. Nanocrystalline alloys have been produced by means of mechanical milling and spark plasma sintering.
[2003Bar]
The Vickers hardness at load 98 N; Young’s modulus was determined in four point bending tests at a cross-head speed of 0.2 mm/min, compression tests with a strain rate of 10–3·s–1
The analysis of mechanical properties under transition from single-phase L12 aloys to eutectic microstructures L12+β (Cr)
[2003Lee]
Compressive tests at room temperature using an Instron testing machine (Model 4206) at a strain rate of 1·10–1·s–1.Fracture toughness was obtained through threepoint bend test using an Instron testing machine (Model 4206) at a strain rate of 1·10–4·s–1 for single edges notched bend specimens.
Examination of the relationship between mechanical properties and phase stability of Al-21Ti-23Cr, Al-21Ti-15Cr and Al-30Ti-15Cr alloys proposed as coating materials with high oxidation resistanse
[2003Nis1] [2003Nis2]
Oxidation tests were carried out in air under a thermal cycling condition between 900˚C and room temperature.
Protective properties of three layer coatings for TiAl-based alloys
[2003Ram]
Microidentation
Microhardness of heat treated TiAl-based thing films alloyed with Cr up to 3 at.%
[2004Wan]
Strain-rate testing
Mechanical properties Ti-46Al-8Cr-2.65 (Nb,W,B) at 800–1100˚C
[2006Fox]
The isothermal interrupted oxidation tests in air, over a wide range of temperatures from 750 to 1100˚C
Oxidation resistance of TiAl3-based alloys alloyed with Cr
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Al–Cr–Ti
. Fig. 1a Al-Cr-Ti. Partial reaction scheme, part 1
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. Fig. 1b Al-Cr-Ti. Partial reaction scheme, part 2
Al–Cr–Ti
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Al–Cr–Ti
. Fig. 1c Al-Cr-Ti. Partial reaction scheme, part 3
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. Fig. 1d Al-Cr-Ti. Partial reaction scheme, part 4
Al–Cr–Ti
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Al–Cr–Ti
. Fig. 2a Al-Cr-Ti. Tentative liquidus surface projection (solid lines - calculations, dashed lines experiments)
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. Fig. 2b Al-Cr-Ti. Tentative liquidus surface projection, enlarged view of the Al corner
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Al–Cr–Ti
. Fig. 3 Al-Cr-Ti. Partial isothermal section at 1300˚C
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. Fig. 4 Al-Cr-Ti. Partial isothermal section at 1200˚C
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Al–Cr–Ti
. Fig. 5 Al-Cr-Ti. Tentative isothermal section at 1050˚C
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. Fig. 6 Al-Cr-Ti. Isothermal section at 1000˚C
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Al–Cr–Ti
. Fig. 7 Al-Cr-Ti. Isothermal section at 800˚C
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. Fig. 8 Al-Cr-Ti. Partial isothermal section at 600˚C
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Al–Cr–Ti
. Fig. 9 Al-Cr-Ti. Activities of Al in Ti-47Al-2Cr, Ti-47Al-13Cr and Ti-51Al-12Cr alloys according to the experimental data of [1999Jac]
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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. Fig. 10 Al-Cr-Ti. Activities of Ti in Ti-47Al-2Cr, Ti-47Al-13Cr and Ti-51Al-12Cr alloys according to the experimental data of [1999Jac]
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Al–Cr–Ti
References [1953Bus] [1953Han]
[1953Tay] [1956Boe] [1956Zwi]
[1958Tag] [1958Kor]
[1960Enc] [1960Zol]
[1963Luz]
[1963Wei] [1964Ram] [1964Sch] [1974Zwi] [1978Jac]
[1984Sup]
[1988Gro]
[1989Kim] [1990Lam]
[1990Mab] [1990Zha]
Busch, L.S., “Transformation Characteristics of 3% Al - 5% Cr Titanium Alloy”, J. Metals, 5, 146–149 (1953) (Morphology, Experimental, Phase Diagram, Phase Relations, Mechan. Prop., 2) Hansen, M., McPherson, D.J., Rostoker, W., “Constitution of Ti Alloy Systems”, ADC Techn. Rep. 53–41, Armour Res. Found. Illinois Inst. Techn., WADC TR 53–41, 129–136 (1953) (Phase Diagram, Phase Relations, 3) Taylor, J.L., Duwez, P., “Constitution of Titanium Rich Ti-Cr-Al Alloys”, J. Metals, 5, 253–256 (1953) (Experimental, Phase Diagram, Phase Relations, 17) Boehm, H., Westphal, H., “Transformation of β into α and ω in Titanium Alloys”, Z. Metallkd., 47, 558–563 (1956) (Morphology, Experimental, Phase Diagram, Phase Relations, Mechan. Prop., 6) Zwicker, U., “The Titan-Aluminium-Chromium and Titan-Aluminium-Vanadium Systems and Technical Alloys with 5% Cr and 3% Al Also 6% Al and 4% V” (in German), Z. Metallkd., 47, 535–548 (1956) (Experimental, Mechan. Prop., Phase Diagram, 12) Tagunova, T.V., “The Ti-Cr-Al System”, Zh. Neorg. Khim., 3, 815–819 (1958) (Experimental, Phase Diagram, Phase Relations, 3) Kornilov, I.I., Mikheev, V.S., Chernova, T.S., “Investigations of Phase Diagram of the Ti-Cr-Al System” (in Russian), Zh. Neorg. Khim., 3(3), 786–796 (1958) (Morphology, Experimental, Phase Diagram, Phase Relations, 6) Ence, E., Farrar, P.A., Margdin, H., “Binary and Ternary Diagrams of the Ti-Al-Cr and Ti-Al-V Systems”, Wright Air Develop. Division, Techn. Rep., 60–316 (1960) Zoller, H., “The Influence of Zinc, Magnesium, Silicon, Copper, Manganese, Titanium on the Primary Crystallization of Al7Cr” (in German), Schweizer Archiv, 478–491 (1960) (Experimental, Phase Diagram, 33) Luzhnikov, L.P., Novikova, V.M., Mareev, A.P., “Solubility of β Stabilisers in α-Ti”, Met. Sci. Heat Treat., 5(2), 78–81 (1963) translated from Metallov. Term. Obrab. Met., (2), 13–16 (1963) (Experimental, Phase Diagram, Phase Relations, 4) Weigand, H.H., “Transformation in α / β Ti Alloys Containing Al”, Z. Metallkd., 54, 43–49 (1963) (Experimental, Phase Relations, Morphology, 16) Raman, A., Schubert, K., “The Occurence of Zr2Cu- and Cr2Al-Type Intermetallic Compounds”, Z. Metallkd., 55, 798–804 (1964) (Experimental, 23) Schubert, K., Raman, A., Rossteutscher, W., “Structural Data of some Metallic Phases (10)” (in German), Naturwissenschaften, 51, 506–507 (1964) (Crys. Structure, 0) Zwicker, U., “Titanium and Titanium Alloys” (in German), Springer, Berlin, 576–585 (1974) (Phase Diagram, Phase Relations, Review, 300) Jacob, I., Shaltiel, D., “A Note on the Influence of Al on the Hydrogen Sorption Properties of Ti(AlxB1–x)2 (B = Cr, Mn, Fe, Co)”, Mater. Res. Bull., 13, 1193–1198 (1978) (Crys. Structure, Experimental, 10) Suprunenko, P.A., Markiv, V.Ya., Tsvetkova, T.M., “Magnetic and X-Ray Diffraction Study of Laves Phases in the Ternary Systems (Ti, Zr, Hf)-Cr-Al”, Russ. Metall., (1), 207–210 (1984) translated from Izv. Akad. Nauk SSSR, Metally, (1), 207–210 (1984) (Experimental, 11) Gros, J.P., Ansara, I., Allibert, M., “Prediction of α/β Equilibria in Titanium-Based Alloys Containing Al, Mo, Zr, Cr. II”, Les Editions de Physique. Avenue du Hoggar, Zone Industrielle de Courtaboeuf, B.P. 112, F-91944 Les Ulis Cedex, France. 1988; Accession Number: 90(9):72–359; Conference: Sixth World Conference on Titanium. III, Cannes, France, 6–9 June 1988, 1559–1564 (1988) (Phase Diagram, Phase Relations, 0) Kim, Y.-W., “Intermetallic Alloys Based on γ Titanium Aluminide”, JOM, 24–30 (1989) (Crys. Structure, Mechan. Prop., Phase Diagram, Phase Relations, Review, 61) Lampman, S., “Wrought Titanium and Titanium Alloys”, Metals Handbook, Tenth Edition. Vol. 2. Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, 2, 592–633 (1990) (Experimental, Interface Phenomena, Mechan. Prop., Phase Diagram, 36) Mabuchi, N., Hirukawa, K., Tsuda, H., Nakayama, Y., “Formation of Structural L12 Compounds in the Ternary Al-Ti-Cr System”, Scr. Metall., 24(3), 505–508 (1990) (Crys. Structure, Experimental, 20) Zhang, S., Nic, J.P., Mikkola, D.E., “New Cubic Phases Formed by Alloying Al3Ti with Mn and Cr”, Scr. Metal. Mater., 24(1), 57–62 (1990) (Crys. Structure, Experimental, Mechan. Prop., 13)
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Al–Cr–Ti [1991Bea]
[1991Dur]
[1991Hua]
[1991Mab]
[1991Mae] [1991Mas]
[1991Mik] [1991Moh] [1991Nic]
[1991Par]
[1991Spa]
[1992Dur] [1992Hay] [1992Kim]
[1992Miz] [1992Mor] [1992Nic]
[1992Win] [1992Zhe]
[1993Ers]
[1993Hao]
7
Beaven, P.A., Pfullmann, Th., Rogalla, J., Wagner, R., “Rhase Transformations and Microstructural Development in TiAl-Based Alloys”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 151–156 (1991) (Experimental, Phase Relations, 9) Durlu, N., Inal, O.U., Yost, F. G., “L1(2)-Type Ternary Titanium Aluminides of the Composition Ti25X8Al67: TiAl3-Based or TiAl2-Based?”, Scr. Metall. Mater., 25(11), 2475–2479 (1991) (Crys. Structure, Review, 30) Huang, S.C., Hall, E.L., Shih, D.S., “Microstructure and Ductility of TiAl Alloys Modified by Chromium Additions”, ISIJ Intl., 31(10), 1100–1105 (1991) (Electronic Structure, Experimental, Mechan. Prop., 17) Mabuchi, H., Nakayama, Y., “Development of Al-Ti-X Ternary L12 Intermetallic Compounds” (in Japanese), Bull. Jpn. Inst. Met., 30(1), 24–30 (1991) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, 36) Maeda, T., Okada, M., Shida, Y., “Ductility and Strength in Mo Modified TiAl”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 555–560 (1991) (Experimental, Phys. Prop., 15) Masahashi, N., Mizuhara, Y., Matsuo, M., Hashimoto, K., Kimura, M., Hanamura, T., Fujii, H., “Ternary Alloying of γ Titanium-Aluminides for Hot-Workability”, Mater. Res. Soc. Symp. Proc.: HighTemp. Ordered Intermetallic Alloys IV, 213, 795–800 (1991) (Experimental, Phys. Prop., 9) Mikkola, D.E., Nic, J.P., Zhang, S., Milligan, W.W., “Alloying of Al3Ti to Form Cubic Phases”, ISIJ Intl., 31(10), 1076–1079 (1991) (Crys. Structure, Experimental, Mechan. Prop., 16) Mohandas, E., Beaven, P.A., “Site Occupation of Nb, V, Mn and Cr in γ-TiAl”, Scr. Metall. Mater., 25 (9), 2023–2027 (1991) (Crys. Structure, Experimental, 15) Nic, J.P., Zhang, S., Mikkola, D.E., “Alloying of Al3Ti with Mn and Cr to Form Cubic L12 Phases”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 697–702 (1991) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, 12) Parfitt, L.J., Smialek, J.L., Nic, J.P., Mikkola, D.E., “Oxidation Behavior of Cubic Phases Formed by Alloying Al3Ti with Cr and Mn”, Scr. Metall. Mater., 25, 727–731 (1991) (Electrochem., Experimental, Phys. Prop., 11) Sparks, C.J., Porter, W.D., Schneibel, J H., Oliver, W.C., Golec, C.G., “Formation of Cubic L12 Phases from Aluminum Titanium (Al3Ti) and Aluminum Zirconium (Al3Ar) by Transition Metal Substitutions for Aluminum”, Mater. Res. Soc. Symp. Proc., 186, 175–80 (1991) (Crys. Structure, Experimental, 15) Durlu, N., Inal, O.T., “L12-Type Ternary Titanium Aluminides as Electron Concentration Phases”, J. Mater. Sci., 27(12), 3225–3230 (1992) (Assessment, Crys. Structure, 41) Hayes, F.H., “The Al-Cr-Ti System (Aluminum-Chromium-Titanium)”, J. Phase Equilib., 13(1), 79–86 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 90) Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems at High Temperatures”, Mater. Sci. Eng. A, 152A(1–2), 54–59 (1992) (Abstract, Crys. Structure, Experimental, Phase Diagram, 12) Mizuhara, Y., Masahashi, N., “The Phase Stability of γ Titanium Aluminides with the β Phase”, Scr. Metall. Mater., 27, 1079–1084 (1992) (Experimental, Mechan. Prop., Morphology, 5) Morris, D.G., Gunter, S., “Ordering Ternary Atom Location and Ageing in Ll2 Trialuminide Alloys”, Acta Metall. Mater., 40(11), 3065–3073 (1992) (Experimental, 23) Nic, J.P., Klansky, J.L., Mikkola, D.E., “Structure/Property Observations for Al-Ti-Cr Alloys Near the Cubic (Al,Cr)3Ti Phase”, Mater. Sci. Eng. A, 152(1/2), 132–137 (1992) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, 16) Winnicka, M.B., Varin, R.A., “Microstructure and Ordering of L12 Titanium Trialuminides”, Metall. Trans. A, 23A(11), 2963–2972 (1992) (Crys.Structure, Experimental, Morphology, 24) Zheng, Y., Zhao, L., Tangri, K., “Microstructure Evolution During Heat Treatment of a ChromiumBearing Ti3Al + TiAl Alloy”, Scr. Metall. Mater., 26(2), 219–224 (1992) (Experimental, Mechan. Prop., Morphology, 19) Erschbaumer, H., Podloucky, R., Rogl, P., Temnitschka, G., Wagner, R., “Atomic Modelling of Nb, V, Cr and Mn Substitutions in γ-TiAl. I: c/a Ratio and Site Preference”, Intermetallics, 1, 99–106 (1993) (Calculation, Crys. Structure, 31) Hao, S.M., Liu, X.J., Zheng, Z.Z., Zeng, N.H., “A Study on the Isothermal Section of Phase Diagram in Ti-Al-Cr System at 1000˚C”, Proc. 7th Nat. Symp. Phase Diagrams, 14–16 (1993) (Experimental, Phase Diagram, 0)
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7 [1993Has]
[1993Kaw]
[1993Nak] [1993Wri]
[1994Boh]
[1994Gao] [1994Kla] [1994Li]
[1994Sok]
[1994Zha] [1994Zhe] [1995Bra] [1995Jew] [1995Hao]
[1995Kog] [1995Nak] [1995Whi]
[1996Jew1] [1996Jew2] [1996Jew3] [1996Liu]
[1996Woo]
Al–Cr–Ti Hashimoto, K., Masao, K., “Effects of Third Element Addition on Mechanical Properties of TiAl”, Struct. Intermet.: 1st Int. Symp. Struct. Intermetallics, Champion Pa. Sept., The Miner., Mater. and Mater. Soc., Daralia, R., Lewandowski, J.J., Liu, C.T., Martin, P.L., Mirakle, D.B., Nathal, M.V. (Eds.), 309–318 (1993) (Experimental, Mechan. Prop., Phase Diagram, 18) Kawabata, T., Tamura, T., Izumi, O., “Effect of Ti/Al Ratio and Chromium, Niobium, and Hafnium Additions on Material Factors and Mechanical Properties in TiAl”, Metall. Trans. A, 24(1), 141–150 (1993) (Crys. Structure, Experimental, Mechan. Prop., 20) Nakayama, Y., Mabuchi, H., “Formation of Ternary L12 Compounds in Al3Ti-Base Alloys”, Intermetallics, 1, 41–48 (1993) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, 40) Wright, R.N., Erickson, A.E., O’Brien, M.H., Rabin, B.H., “Characterization of Al2Ti Precipitates in Cubic Al64.5Ti27.5Cr8”, Scr. Metall. Mater., 28(10), 1293–1297 (1993) (Electronic Structure, Experimental, Mechan. Prop., 19) Bohnenkamp, U., Wang, G.-X., Jewett, T.J., Dahms, M., “A Quantitative Investigation of Phase Formation During Annealing of Cold-Extruded Elemental Powder Mixtures Ti-48 at.% Al and Ti-48 at.% Al-2 at.% Cr”, Intermetallics, 2, 275–283 (1994) (Crys. Structure, Experimental, Phase Diagram, Thermodyn., 20) Gao Y., Zhu J., Cai Q.G., “Microstructure of the Ordered β Phase in Ti-Al Alloy with Chromium and Vanadium.”, Scr. Met. Mater., 31(5), 571–575 (1994) (Experimental, Morphology, 12) Klansky, J.I., Nic, J.P., Mikkola, D.E., “Structure/Property Observations for Al-Ti-Cr Intermetallic Alloys”, J. Mater. Res., 9(2), 255–258 (1994) (Crys. Structure, Experimental, Phase Diagram, 13) Li, Y.G., Loretto, M.H., “Microstructure and Fracture Behaviuor of Ti-44Al-xM Derivatives”, Acta Metall. Mater., 42(9), 2913–2919 (1994) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 12) Sokolovskaya, E.M., Kazakova, E.F., Poddyakova, E.I., Portnoi, V.K., Temirbaeva, A.A., “Phase Composition of Alloys of the Al-Cr-Ti System at 770 K” (in Russian), Vestn. Mosk. Univ. Ser. 2: Khim., 35(5), 453–454 (1994) (Phase Diagram, Phase Relations, Experimental, 5) Zhang, B., Wang, J., Wan, X., Chen, W., “A Study on the β and ω Phases in a Ti-Al-Cr Alloy”, Scr. Metall. Mater., 30, 399–404 (1994) (Crys. Structure, Experimental, 13) Zhang, J.G., Li, Q., Liu, Z.G., Feng, D., Frommeyer, G., “Complex Stacking Fault Energy of Cr-Alloyed γ-TiAl”, Phys. Lett. A, 196, 125–127 (1994) (Crys. Structure, Experimental, 11) Brady, M.P., Smialek, J.L., Terepka, F., “Microstructure of Alumina-Forming Oxidation Resistant AlTi-Cr Alloys”, Scr. Metall. Mater., 32(10), 1659–1664 (1995) (Experimental, Phase Relations, 17) Jewett, T.J., Dahms, M., “Investigation of the γ-Ti(Cr,Al)2 Phase at 800˚C and 1000˚C”, Scr. Metall. Mater., 32(10), 1533–1539 (1995) (Experimental, Phase Relations, 18) Hao, S., Zeng, L., “Isothermal Section of Phase Diagram in Ti-Al-Cr Ternary System at 1000˚C”, Acta Metall. Sin. (China), 31(4), B152-B158 (1995) (Experimental, Morphology, Phase Diagram, Phase Relations, 5) Kogachi, M., Kameyama, A., “Composition Dependence of site Occupancies in Ternary L12 Compound Al3Ti-Cr”, Intermetallics, 3, 327–334 (1995) (Crys. Structure, Experimental, 29) Nakai, K., Ono, T., Ohtsubo, H., Ohmori, Y., “Phase Stability and Decomposition Processes in Ti-Al Based Intermetallics”, Mater. Sci. Eng. A, 192-193, 922–929 (1995) (Experimental, Phase Relations, 21) Whittenberger, J.D., Kumar, K.S., DiPietro, M.S., Brown, S.A., “Characteristics of Elevated Temperature Deformation in Several L12-Modified Al3Ti-Based Alloys”, Intermetallics, 3, 221–232 (1995) (Experimental, Mechan. Prop., 32) Jewett, T., Dahms, M., “Stability of the Ti(Cr,Al)2 Phase”, Z. Metallkd., 87(4), 254–261 (1996) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 33) Jewett, T.J., Ahrens, B., Dahms, M., “Phase Equlibria Involting the τ-L12 and TiAl2 Phases in the Ti-AlCr System”, Intermetallics, 4, 543–556 (1996) (Experimental, Phase Relations, 31) Jewett, T.J., Ahrens, B., Dahms, M., “Stability of the L12 Phase at 800˚C C in the Ti-Al-Cr-System”, Scr. Mater., 34(3), 395–399 (1996) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 13) Liu, S., Hu, R., Zhao, D., Wang, C., Luo, P., Pu, Z., “Bonding Characteristics of the Intermetallic Compound Al3Ti+Cr”, J. Mater. Sci. Technol., 12, 180–184 (1996) (Crys. Structure, Electronic Structure, Experimental, 30) Woodward, C., MacLaren, J.M., “Planar Fault Energies and Sessile Dislocation Configurations in Substitutionally Disordered Ti-Al with Nb and Cr Ternaty Additions”, Philos. Mag. A, 74(2), 337–357 (1996) (Calculation, Crys. Structure, 43)
DOI: 10.1007/978-3-540-88053-0_7 ß Springer 2009
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Al–Cr–Ti [1997Jew] [1997Fan] [1997Hei] [1997Mab1]
[1997Mab2]
[1997Mas] [1997Ley]
[1997Par]
[1997Xu] [1998Jim]
[1998Jun]
[1998Has] [1999Yan]
[1999Jac] [1999Hao] [1999Nic] [1999Sha] [1999Yam]
[2000Cha]
[2000Coe]
[2000Kai1]
7
Jewett, T.J., Ahrens, B., Dahms, M., “Stability of TiAl in the Ti-Al-Cr System”, Mater. Sci. Eng. A, 225, 29–37 (1997) (Crys. Structure, Experimental, Phase Relations, 26) Fan, G.J., Song, X.P., Quan, M.X., Hu, Z.Q., “Mechanical Alloying and Thermal Stability of Al67Ti25M8 (M = Cr,Zr,Cu)”, Mater. Sci. Eng. A, 231, 111–116 (1997) (Crys. Structure, Experimental, 22) Heilmaier, M., Saage, H., Eckert, J., “Formation of ODS L12-(Al, Cr)3Ti by Mechanical Alloying”, Mater. Sci. Eng. A, 239-240, 652–657 (1997) (Experimental, Mechan. Prop., 18) Mabuchi, H., Tsuda, H., Matsui, T., Morii, K., “185. Microstructure and Mechanical Properties of Ternary L12 Intermetallic Compound in Al-Ti-Cr System”, Metall. Abstr. Light Met. Alloys, 30, 208–210 (1996-1997) (Experimental, Phase Diagram, Phase Relations, 0) Mabuchi, H., Tsuda, H., Matsui, T., Morii, K., “Microstructure and Mechanical Properties of Ternary L12 Intermetallic Compound in Al-Ti-Cr System”, Mater. Trans., JIM, 38(6), 560–565 (1997) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 40) Masahashi, N., Mizuhara, Y., “APFIM Study of β and γ Microduplex TiAl Intermetallic Alloys”, Mater. Sci. Eng. A, 223, 29–35 (1997) (Experimental, Phys. Prop., 24) Leyens, C., Schmidt, M., Peters, M., Kaysser, W.A., “Sputtered Intermetallic Ti-Al-X Coatings: Phase Formation and Oxidation Behavior”, Mater. Sci. Eng. A, 239-240, 680–687 (1997) (Experimental, Morphology, 17) Park, J.Y., Oh, M.H., Wee, D.M., Miura, S., Mishima, Y., “L12 (Al,Cr)3Ti-Based Two-Phase Intermetallic Compounds-I. Plastic Deformation Behavior”, Scr. Mater., 36(7), 795–800 (1997) (Experimental, Phase Relations, 12) Xu, H., Jin, Z., Wang, R., “Study of the Phase Equilibria of Al-Cr-Ti System at 1050˚C”, Scr. Mater., 37(10), 1469–1473 (1997) (Experimental, Phase Relations, 25) Jimenez, J.A., Wesemann, J., Frommeyer, G., “High-Temperature Deformation Behavior of the Intermetallic Ti-47 at.% Al-3 at.% Cr Alloy”, Metall. Mater. Trans. A, 29, 1425–1430 (1998) (Electronic Structure, Experimental, Mechan. Prop., 29) Jung, J.Y., Park, J.K., “Growth Kinetics of Discontinous Coarsening of Lamellar Structure in Ti-44 at.% Al (-0.5 at.% Cr) Intermetallic Compounds”, Acta Mater., 46(12), 4123–4130 (1998) (Crys. Structure, Experimental, Kinetics, Thermodyn., 31) Hashimoto, K., Kimura, M., Mizuhara, Y., “Alloy Design of γ Titanium Aluminides Based on Phase Diagrams”, Intermetallics, 6, 667–672 (1998) (Experimental, Phase Relations, 14) Yang, R., Hao, Y.L., “Estimation of (γ + α2) Equilibrium in Two-Phase Ti-Al-X Alloys by Means of Sublattice Site Occupancies of x in TiAl and Ti3Al”, Scr. Mater., 41(3), 341–346 (1999) (Calculation, Phase Relations, 13) Jacobson, N.S., Brady, M.P., Mehrotra, G.M., “Thermodynamics of Selected Ti-Al and Ti-Al-Cr Alloys”, Oxid. Met., 52(5/6), 537–556 (1999) (Experimental, Phase Relations, Thermodyn., 48) Hao, Y.L., Xu, D.S., Cui, Y.Y., Yang, R., Li, D., “The Site Occupancies of Alloying Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47(4), 1129–1139 (1999) (Crys. Structure, Experimental, 41) Nic, J.P., Mikkola, D.E., “Site Occupancy in Ternary L12 Ordered Alloys as Determined by Diffarction: Observations on (Al,Cr)3Ti Alloys”, Intermetallics, 7, 39–47 (1999) (Crys. Structure, Experimental, 20) Shao, G., Tsakiropoulos, P., “Solidification Structures of Ti-Al-Cr Alloys”, Intermetallics, 7, 579–587 (1999) (Experimental, Phase Relations, 15) Yamamoto, Y., Hashimoto, K., Kimura, T., Nobuki, M., Kohno, N., “L12 Single Phase Region in Al-Ti Base Ternary and Quaternary Systems at 1450 K” (in Japanese), J. Jpn. Inst. Met., 63(10), 1317–1326 (1999) (Crys. Structure, Experimental, Phase Relations, 15) Chatterjee, A., Dehm, G., Scheu, C., Clemens, H., “Onset of Microstructural Instability in a Fully Lamellar Ti-46.5 at.% Al-4 at.% (Cr,Nb,Ta,B) Alloy During Short-term Creep”, Z. Metallkd., 91 (9), 755–760 (2000) (Experimental, Mechan. Prop., Phase Relations, 21) Coelho, C., Ramos, A.S., Trindade, B., Vieira, M.T., Fernandes, J.V., Vieira, M., “Characterisation of Modified Sputtered (TiAl)-Based Intermetallic Materials Doped with Silver and Chromium”, Key Eng. Mater., 188, 37–44 (2000) (Experimental, Mechan. Prop., 13) Kainuma, R., Ohnuma, I., Ishukawa, K., Ishida, K., “Stability of B2 Ordered Phase in the Ti rich Portion of Ti-Al-Cr and Ti-Al-Fe Ternary Systems”, Intermetallics, 8, 869–875 (2000) (Crys. Structure, Experimental, Phase Relations, 19)
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7 [2000Kai2]
[2000Mab]
[2000Ohi]
[2000Sha] [2000Wel] [2000Yan]
[2000Zhu] [2001Ich]
[2001Ish]
[2001Fuj]
[2001Kan]
[2001Kau] [2001Lee]
[2001Mil]
[2001Mor] [2001Sun]
[2001Sur] [2001Xia]
[2001Yam]
[2001Zho]
Al–Cr–Ti Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among α (hcp), β (bcc) and γ (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855–867 (2000) (Crys. Structure, Experimental, Phase Relations, 29) Mabuchi, H., Tsuda, H., Matsui, T., Morii, K., “Effects of Vanadium Addition on the Material Properties of L12-Type Ordered Alloys in the Al-Ti-Cr System”, J. Jpn. Inst. Met., 64(11), 1041–1047 (2000) (Crys. Structure, Experimental, Phase Relations, 39) Oh-ishi, K., Kumamoto, H., Horita, Z., Nemoto, M., “Phase Separation in TiAl(L10)-(Al,Cr)3Ti(L12) System” (in Japanese), J. Jpn. Inst. Met., 64(8), 609–615 (2000) (Experimental, Mechan. Prop., Phase Relations, 14) Shao, G., Tsakiropoulos, P., “On the w Phase Formation in Cr-Al and Ti-Al-Cr Alloys”, Acta Mater., 48, 3671–3685 (2000) (Calculation, Crys. Structure, Experimental, Phase Relations, 39) Weller, M., Chatterjee, A., Haneczok, G., Clemens, H., “Internal Friction of γ-TiAl Alloys at High Temperature”, J. Alloys Compd., 310, 134–138 (2000) (Experimental, Mechan. Prop., 15) Yang, R., Hao, Y., Song, Y., Guo, Z.X., “Site Occupancy of Alloying Additions in Titanium Aluminides and its Application to Phase Equilibrium Evaluation”, Z. Metallkd., 91(4), 296–301 (2000) (Crys. Structure, Phase Relations, Review, 38) Zhuang, W., Shen, J., Liu, Y., Shang, S., Du, Y., Schuster J.C., “Thermodynamic Optimisation of the Cr-Ti System”, Z. Metallkd., 91, 121–127 (2000) (Thermodyn., Phase Diagram, 53) Ichimaru, M., Mori, N., Miura, Y., “Al-Ti-Cr Ternary Phase Diagram and Solidification Structure of (Al,Cr)3Ti Alloys”, J. Jpn. Inst. Light Met., 51(12), 640–645 (2001) (Experimental, Morphology, Phase Diagram, Phase Relations, 9) Ishikawa, K., Hashi, K., Suzuki, K., Aoki, K., “Effect of Substitutional Elements on the Hydrogen Absorption-Desorption Properties of Ti3Al Compounds”, J. Alloys Compd., 314, 257–261 (2001) (Experimental, Phase Relations, 9) Fujita, T., Ikeda, H., Tanaka, S., Horita, Z., “Construction of Ti-Al-Cr Phase Diagram Using Quantitative X-Ray Microanalysis in Analytical Electron Microscope” (in Japanese), J. Jpn. Inst. Met., 65(5), 382–388 (2001) (Crys. Structure, Experimental, Phase Relations, 20) Kang, S.Y., Onodera, H., “Analyses of HCP/D019 and D019/L10 Phase Boundaries in Ti-Al-X (X = V, Mn, Nb, Cr, Mo, Ni and Co) Systems by the Cluster Variation Method”, J. Phase Equilib., 22, 424–430 (2001) (Calculation, Phase Relations, 15) Kaufman, L., “Calculation of Multicomponent Phase Diagrams for Niobium Alloys”, Niob. Sci. Tech., 107–120 (2002) (Calculation, Phase Diagram, 20) Lee, C.Y., Park, J.K., “Effect of Cr and Si Additions on the Continuous-Cooling-Transformation Kinetics of γ-Based Ti-45 at.% Al Alloy”, Philos. Mag. A, 81(10), 2415–2426 (2001) (Crys. Structure, Experimental, Kinetics, 14) Milman, Yu.V., Miracle, D.B., Chugunova, S.I., Voskoboinik, I.V., Korzhova, N.P., Legkaya, T.N., Podrezov, Yu.N., “Mechanical Behaviour of Al3Ti Intermetallic and L12 Phases on Its Basis”, Intermetallics, 9, 839–845 (2001) (Crys. Structure, Experimental, Mechan. Prop., 36) Mori, N., Ogi, K., “Analisis of Solidification-Path of Al-Ti-Cr Alloys by Progressive-Type Solidification Equations” (in Japanese), J. Jpn. Inst. Met., 65(9), 848–851 (2001) (Experimental, Phase Relations, 12) Sun, F.-S., Cao, C.-X., Kim, S.-E., Lee, Y.-T., Yan, M.-G., “Alloying Mechanism of β Stabilizers in a TiAl Alloy”, Metall. Mater. Trans. A, 32A, 1573–1589 (2001) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 37) Suryanarayana, C., “Mechanical Alloying and Milling”, Prog. Mater. Sci., 46(1-2), 1–184 (2001) (Crys. Structure, Experimental, Kinetics, Phase Relations, Review, Thermodyn., 932) Xia, Q., Wang, J.N., Yang, J., Wang, Y., “On the Massive Transformation in TiAl-based Alloys”, Intermetallics, 9, 361–367 (2001) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, Thermodyn., 29) Yamamoto, Y., Hashimoto, K., Kimura, T., Nakamura, M., Kohno, N., “Residual Strain in Mechanically Ground Powders of Al-Ti Base Ternary and Quaternary Compounds with Ll2 Single-Phase at 1450K”, Mater. Trans., JIM, 42(7), 1392–1399 (2001) (Crys. Structure, Experimental, Phase Relations, 18) Zhou, Ch., Yang, Y., Gong, Sh., Xu, H., “Effect of Ti-Al-Cr Coatings on the High Temperature Oxidation Behavior of TiAl Alloys”, Mater. Sci. Eng. A, 307, 182–187 (2001) (Crys. Structure, Experimental, Kinetics, 13)
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[2003Ram] [2003Zha] [2004Nay] [2004Wan] [2005App] [2005Rag] [2006Fox]
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Brandt, C., Inal, O. T., “Mechanical Properties of Cr, Mn, Fe, Co, and Ni Modified Titanium Trialuminides”, J. Mater. Sci., 37(20), 4399–4403 (2002) (Crys. Structure, Experimental, Mechan. Prop., 17) Calderon, H.A., Garibay-Febles, V., Umemoto, M., Yamaguchi, M., “Mechanical Properties of Nanocrystalline Ti-Al-X Alloys”, Mater. Sci. Eng. A, 329-331, 196–205 (2002) (Experimental, Mechan. Prop., Morphology, 26) Barabash, O.M., Milman, Yu.V., Miracle, D.V., Karpets, M.V., Korzhova, N.P., Legkaya, T.N., Mordovets, N.M., Podrezov, Yu.N., Voskoboinik, I.V., “Formation of Periodic Mecrostructures Involving the L12 Phase in Eutectic Al-Ti-Cr Alloys”, Intermetallics, 11(9), 953–962 (2003) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, 26) Lee, J.K., Oh, M.W., Oh, M.H., Wee, D.M., “Phase Stability of L12-Based Alloys in Al-Ti-Cr Systems”, Intermetallics, 11(8), 857–865 (2003) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phase Relations, 24) Mori, N., Ogi, K., “Solidification-Path and Microstructure of Al-Ti-Cr Alloy Analyzed by ProgressiveType Solidification Equation”, Mater. Trans., JIM, 44(11), 2334–2338 (2003) (Calculation, Phase Relations, 15) Nishimoto, T., Izumi, T., Hayashi, S., Narita, T., “Two-Step Cr and Al Diffusion Coating on TiAl at High Temperatures”, Intermetallics, 11, 225–235 (2003) (Experimental, Transport Phenomena, 38) Nishimoto, T., Izumi, T., Hayashi, Sh., Narita, T., “Effect of Coating Layer Structures and Surface Treatments on the Oxidation Behavior of a Ti-50 at.% Al Alloy”, Intermetallics, 11(5), 459–466 (2003) (Experimental, Interface Phenomena, Kinetics, Mechan. Prop., 46) Ramos, A.S., Vieira, M.T., “Properties of γ-TiAl-M (M=Ag, Cr) Sputtered Films”, Mater. Sci. Forum, 426-432, 1843–1848 (2003) (Electronic Structure, Experimental, Mechan. Prop., 20) Zhang, W.J., Deevi, S.C., “The Controlling Factor in Primary Creep of TiAl-Based Alloys”, Intermetallics, 11(2), 177–185 (2002) (Experimental, Mechan. Prop., 34) Nayak, S.S., Murty, B.S., “Synthesis and Stability of L12-Al3Ti by Mechanical Alloying”, Mater. Sci. Eng. A, 367(1-2), 218–224 (2004) (Crys. Structure, Experimental, Morphology, Phase Relations, 25) Wang, Y., Wang, J.N., Yang, J., “Superplastic Behavior of a High-Cr TiAl Alloy in its Cast State”, J. Alloys Compd., 364(1-2), 93–98 (2004) (Experimental, Mechan. Prop., 10) Appel, F., “An Electron Microscope Study of Mechanical Twinning and Fracture in TiAl Alloys”, Philos. Mag., 85(2-3), 205–231 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, 74) Raghavan, V., “Al-Cr-Ti (Aluminum-Cromium-Titanium)”, J. Phase Equilib. Diffus., 26, 349–356 (2005) (Review, Phase Diagram, 35) Fox-Rabinovich, G.S., Wilkinson, D.S., Veldhuis, S.C., Dosbaeva, G.K., Weatherly, G.C, “Oxidation Resistant Ti-Al-Cr Alloy for Protective Coating Applications”, Intermetallics, 14(2), 189–197 (2006) (Crys. Structure, Experimental, Interface Phenomena, Kinetics, Morphology, Phase Relations, 20) Narita, T., Izumi, T., Nishimoto, T., Shibata, Y., Thosin, K.Z., Hayashi, S., “Advanced Coatings on High Temperature Applications”, Mater. Sci. Forum, 522-523, 1–14 (2006) (Phase Diagram, Phase Relations, Review, 15) Schuster, C., Palm, M., “Reassessment of the Binary Aluminium-Titanium Phase Diagram”, J. Phase Equilib. Diffus., 27(3), 255–277 (2006) (Review, Phase Diagram, Crys. Structure, 375) Heshmanu-Manesh, S., Ahmadabadi, M.N., Ghasemiarmaki, H., Jafarian. H.R., “Effect of Initial Microstructure and Further Thermomechanical Processing on Microstructural Evolution in a Ti-47Al2Cr Alloy”, J. Alloys Compd., 436, 200–203 (2007) (Experimental, Morphology, 10) Witusiewicz, V.T., Bondar, A.A., Hecht, U., Rex, S., Velikanova, T.Ya., “The Al-B-Nb-Ti System. III. Thermodynamic Re-Evaluation of the Constituent Binary System Al-Ti”, in press, J. Alloys Compd., DOI:10.1016/j.jallcom.2007.10.061 (2007) (Phase Relations, Phase Diagram, Experimental, Thermodyn., Calculation, Review, 89) Cornish, L., Saltykov, P., Cacciamani, G., Velikanova, T., “Al-Cr (Aluminium-Cromium), MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2008) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 51) Ivanchenko, V., “Cr-Ti (Chromium-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2008) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 22)
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Al–Cr–Ti Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Iron – Niobium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Sergio Gama, updated by Annelies Malfliet
Introduction [1966Ram] first reported an isothermal section at 1000˚C, followed by [1970Bur] who provided an 800˚C isotherm. There is reasonable agreement between the two works. Interestingly, however, both investigators found a ternary phase, called μ’, although the position and solubility range of this phase is drawn differently in both sections. Some diffraction lines of this phase are equal to those of the μ phase, but the exact crystal structure is unknown and further investigation is required. No other ternary phases were identified. More recently, a new version of the isothermal section at 1000˚C according to [1993Bej2] was presented in [1999Mot]. By comparing this section with the previous section at this temperature of [1966Ram], the μ’ phase is not present, while a new ternary Nb rich phase is observed. [1987Rag] provided a critical evaluation of the Al-Fe-Nb system, presenting a 800˚C isothermal section based on both [1966Ram] and [1970Bur], but containing the metastable Nb3Fe2 phase not considered by the former investigators. In the previous MSIT evaluation of [1992Gam] the isothermal section at 1000˚C determined by [1966Ram] is redrawn to be consistent with the accepted binaries. The Al-Fe-Nb system is also part of the review by [1990Kum] who highlights the inconsistencies in different studies of ternary Al-refractory metal-X systems. Most investigations on the applications of Al-Fe-Nb alloys are restricted to two composition regions. First, compositions around Fe-25Al with small percentages of Nb have been studied to improve the mechanical properties of this intermetallic through the formation of a two-phase microstructure of Fe3Al and λ, Nb(Fe1–xAlx)2, [1991Ran, 1997Par, 1999Den, 2004Mor2, 2005Mor, 2005Mot]. Secondly, Al-rich compositions are of interest due to the possibility to prepare nanocomposites with an amorphous matrix containing metallic nanocrystals through rapid quenching followed by the proper heat treatment. Investigations regarding the phase relations, structures and thermodynamics of the Al-Fe-Nb system are summarized in Table 1.
Binary Systems The Al-Fe system is accepted from [2006MSIT] which is mainly based on [1993Kat], except for the Fe rich corner, where the data of [2001Ike] are taken. For the Al-Nb system, the liquidus has not yet been fully determined and some inconsistencies in the invariant temperatures can be found by comparing different experimental studies. The accepted phase diagram is from [Mas2], mainly based on the work of [1980Jor], except for the Al-rich side which is based on [1966Lun]. The Fe-Nb system, presented in Fig. 1, is accepted from [1993Bej1]. This diagram Landolt‐Bo¨rnstein New Series IV/11E1
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is based on earlier work of the same author [1991Bej], where Nb3Fe2 is observed as a metastable phase around 61 at.% Nb.
Solid Phases The binary phases are described in Table 2. None of the Al-Fe binary compounds have appreciable solubility of Nb. For the Fe3Al phase, a solubility of 1 and 2 at.% Nb at 1300˚C is reported by respectively [2004Mor2] and [1988Dim], 1.6 at.% at 1150˚C by [2003Ste], 1 and 0.87 at.% at 1000˚C by respectively [1988Dim] and [2003Ste] and 0.5 at.% at 700˚C by [1988Dim]. For the binary Al-Nb phases, about 10 at.% Fe substitutes either Al or Nb in the σNb2Al phase at 1000˚C [1966Ram] and 2 at.% at 800˚C [1970Bur]. The NbAl3 phase takes up only little amount of Fe at 1000˚C [1966Ram]. The solubility range of this phase drawn at 5 at. % Fe in the preliminary study of [1966Ram], leads to the suggestion of [1989Sub] that Fe substitutes for Nb in the latter phase although no clear experimental details are given to validate this. For the Fe-Nb phases, a solubility of about 20 and 50 at.% Al in respectively μ, Nb19(Fe1–xAlx)21 and λ, Nb(Fe1–xAlx)2, at 1000˚C is found by [1966Ram]. This large solubility in both phases, whereby Al mainly substitutes for the Fe atoms, is confirmed in the investigation of the ternary system by [1970Bur] at 800˚C, reported as 26 at.% Al in μ, Nb19(Fe1–xAlx)21 and 56 at.% Al in λ, Nb(Fe1–xAlx)2. However, [1986Bla] concludes based on diffraction intensities of NbxFe2Al1–x alloys, that Al replaces the Nb atoms at the 8f positions in the hexagonal λ, Nb(Fe1–xAlx)2, phase, thereby replacing up to half the Nb atoms. [1966Ram] was the first to describe a ternary phase, called μ’, with basic stoichiometry Nb2FeAl existing in alloys kept for 168 h at 1000˚C. Most diffraction lines are equal to those of the μ, Nb19(Fe1–xAlx)21 phase, but some important ones are missing and new ones are observed. Although this indicates that μ’ is structurally related to the μ phase, the exact crystal structure is undetermined. [1970Bur] claims to observe this μ’ phase also in water quenched alloys kept for 1000 h at 800˚C, albeit at a different composition range than [1966Ram]. Therefore, as [1990Kum] states, it is not clear whether the same phase is observed in both investigations. In addition, the μ’ phase is completely absent in the isothermal section at 1000˚C of [1993Bej2] and as the solubility of the μ phase is extended to 27 at.% Al, this could indicate the metastable formation of the μ’ phase as observed by [1966Ram] from the μ phase. [1993Bej2] also observed a new ternary phase containing 5–10 at.% Fe and 19–25 at.% Al, but no crystallographic data are available.
Liquidus, Solidus and Solvus Surfaces Based on the substitution of Fe by Al in both λ, Nb(Fe1–xAlx)2, and FeAl, it is suggested in [1999Mot] that the eutectic valley between these phases should also follow a line of constant Nb. For Al-Fe-Nb alloys with 22.8, 41.2 and 45.0 at.% Al, the eutectic composition has been determined to contain respectively 9.7, 9.1 and 8.0 at.% Nb. Based on these measurements, a possible liquidus in this region of the phase diagram is constructed in Fig. 2. [2005Mot] found in addition that the coupled zone below the eutectic line is asymmetrically situated towards the λ, Nb(Fe1–xAlx)2, phase as a film of FeAl surrounds the primary λ, Nb(Fe1–xAlx)2, dendrites DOI: 10.1007/978-3-540-88053-0_8 ß Springer 2009
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in a hypereutectic alloy, while no λ, Nb(Fe1–xAlx)2, film surrounds the FeAl primary phase in a hypoeutectic alloy.
Isothermal Sections Two isothermal sections of the Al-Fe-Nb system are proposed by [1966Ram] and [1970Bur] at respectively 1000 and 800˚C, mainly based on metallographic and X-ray techniques. By comparing the two isothermal sections which differ only by 200˚C, two remarks should be mentioned. First, there is a significant difference in solubility of Fe in the binary Al-Nb phases. Secondly, the position and solubility range of the ternary μ’ phase in both sections is different. In his review, [1990Kum] poses the question whether these two μ’ phases are the same or different phases and states that further investigation of this part of the ternary system is recommended. [1987Rag] proposed an isothermal section at 800˚C based on both the results of [1966Ram] and [1970Bur] with the Al-Fe, Al-Nb and Fe-Nb binaries of respectively [1982Kub], [1981Ell] and [1982Kub]. In this combined section, a mean solubility of 5 at.% Fe in σ (Nb2Al) was assumed and the position of the μ’ phase was taken from [1970Bur]. To be in agreement with the binary Fe-Nb from [1982Kub], the metastable Nb3Fe2 phase was added. In [1999Mot] a more recent version of the isothermal section at 1000˚C as determined by [1993Bej2] is presented. This section is similar to the section proposed by [1966Ram] except for the region around μ’ and the ternary τ phase. The 1000˚C isothermal section presented in Fig. 3 is mainly taken from [1999Mot] as determined by [1993Bej2]. This section is preferred with respect to [1966Ram] as the latter only aimed for a rough estimation of the phase diagram by using metallographic and XRD techniques. Although no detailed information is available about the experimental work of [1993Bej2], from other works of this author in the same period [1991Bej, 1993Bej1, 1999Mot] it is known that this author had at least the opportunity to perform electron microprobe analysis to have a more precise composition analysis. In Fig. 4, the extension of the λ, Nb (Fe1–xAlx)2, phase in the two-phase region λ+FeAl is adjusted to the experimental results from [1988Dim]. The latter homogenized a Nb2Fe73Al25 alloy for 168 h at 1000˚C and found a twophase microstructure with 1Nb-74Fe-25Al (at.%) as the composition for the matrix and 25Nb-56Fe-19Al (at.%) as the composition for the λ, Nb(Fe1–xAlx)2 particles. [2004Mor2] confirmed that at high temperature the composition of the binary λ, NbFe2 phase in equilibrium with Fe3Al is close to NbFe3 and through the steady substitution of Fe by Al in the ternary system it evolves towards NbAl3. Figure 4 shows the isothermal section at 800˚C after [1970Bur], with some additional lines in the Al corner based on the binary systems.
Thermodynamics Thermodynamic data of the Al-Fe-Nb system are reported in Table 3. [1971Dub] used isothermal calorimetry based on an aluminothermic reaction to determine the heat of solution of aluminium in Fe-Nb alloys. The heat of solution is 105 ± 13 kJ·mol–1 for 1.3–6.3 mass% Al in a Fe-38.2Nb alloy and 100 ± 13 kJ·mol–1 for 1.9–9.5 mass% Al in a Fe-67.6Nb alloy. [2001Sud] used calorimetric data to estimate the enthalpy of mixing in the ternary Al-Fe-Nb system. Their calculations show that the addition of aluminium to Fe-Nb results in a minimum of the mixing enthalpy between –25 and –20 kJ·mol–1 located at the binary Al-Fe. Landolt‐Bo¨rnstein New Series IV/11E1
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Notes on Materials Properties and Applications Hardness tests have been performed on Nb1–xFe2Alx by [1986Bla] and on Nb33.4Fe33.6Al33.0 by [1996Mac]. According to [1962Min], the addition of Nb increases the hardness, electrical resistivity and improves the high temperature mechanical properties of the Fe3Al phase.The temperature dependence of the hardness of Nb2Fe68Al30, Nb2Fe73Al25, Nb5Fe75Al20 and Nb2Fe83Al15 has been measured by [2006Mor1, 2006Mor2]. The hardening is related to the change in state of order and, in particular, to the formation of a two-phase D03 order-plusdisorder microstructure. Above the temperature for onset of precipitation, the loss of Nb in solid solution by the formation of metastable L21 needles and stable λ, Nb(Fe1–xAlx)2 blocky precipitates in these alloys initially has no major effect on hardening, but after long anneals the coarsening of these precipitates results nevertheless in softening. Stress-strain curves for Nb2Fe73Al25 have been measured at different temperatures and strain rates by [2004Mor2, 2006Mor1]. The rapid formation and coarsening of the λ, Nb(Fe1–xAlx)2 phase above 700˚C is detrimental for the high temperature strength of these alloys [2005Mor]. This limits the use of these alloys to 700–800˚C, and even lower temperatures (< 600˚C) are recommended for longterm applications. [1991Ran] used combustion synthesis to prepare two-phase microstructures consisting of Fe3Al and Nb2Fe for which sufficiently high strength can be retained up to 700˚C. Compression tests on these alloys indicate the possibility of superplasticity around 900˚C. The brittle-to-ductile transition for the two-phase alloys Nb5.1Fe62.2Al32.7 and Nb16.7Fe53.9Al29.4 has been measured by [1996Mac]. [1997Par] studied the effect of Nb addition on the precipitation strengthening and fracture mechanisms in Fe-25Al alloys. [1999Den] has performed calculations on the density of the valence electrons to describe the effect of Nb additions on the grain boundary and bulk bond strength in Fe-25Al. [1998Aud2] investigated the mechanical and corrosion behavior of partially and fully amorphous Al-rich Al-Fe-Nb alloys as an improved light material. Results indicate that fully amorphous alloys exhibit low Young’s moduli, high microhardness and good corrosion behavior. [1993Kai] studied the corrosion kinetics in H2/H2O/H2S atmosphere in the temperature range 700–980˚C. With increasing Nb addition to Fe-3Al (mass%) the corrosion rate reduces with a most pronounced reduction from 30 to 40 mass% Nb. The addition of 10 mass% Al to Nb-70Fe (mass%) leads at 700 and 800˚C to a significant decrease in corrosion rate. [1997Cho] has observed that the addition of Nb to Fe-25Al (at.%) accelerates both the granular and the intergranular corrosion. Mo¨ssbauer spectroscopy studies on NbxFeAl1–x for 0.1 ≤ x ≤ 0.3 indicate that the FeAl and λ phase (Nb(Fe1–xAlx)2) are paramagnetic down to 85 K [1993Mah]. The Debye temperature for the FeAl phase is 112 ± 5˚C. The magnetic behavior of the λ, Nb(Fe1–xAlx)2 phase in the Al-Fe-Nb system has been calculated through the LAPW method by [2002Ima]. Information on the type of property and the used technique or method in these investigations is given in Table 4.
Miscellaneous [1990Car] performed ab initio electronic band calculations to predict the effect of Fe addition on the relative stability between the tetragonal D022 and cubic L12 structure for the NbAl3 phase. Although he assumes in contradiction to [1989Sub] that Fe substitutes for Al in the DOI: 10.1007/978-3-540-88053-0_8 ß Springer 2009
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tetragonal D022 structure, their conclusions that Fe is not an efficient stabilizer of the cubic structure, are in agreement. [1999Mek] used a quasichemical method based on statistico-thermodynamical theory in combination with electronic theory in the pseudopotential approximation to determine the effect of Nb on the ordering in the FeAl phase. In NbxFe0.5Al0.5–x alloys with x up to 0.005, the Nb atoms preferentially occupy the Fe sublattice sites and the normalized critical orderdisorder phase transformation temperature change by addition of Nb, ΔT/TO, varies from 47.35 for x = 0.00125 to 202.17 for x = 0.005. According to [2002Boz], however, the BozzoloFerrante-Smith method based on quantum approximate methods predicts that Nb will preferentially occupy the Al lattice sites in the B2 compound for both NbxFe0.5Al0.5–x and NbxFe0.5–xAl0.5. Several investigations have been focussed on the addition of Nb to alloys with a composition around Fe-25Al. Especially the effect of Nb on the ordered D03 and B2 and the disordered A2 structure has been studied. The temperature change for the D03 Ð B2 transition through addition of Nb has been measured by [1995Ant], [2003Ste], [2005Mor] and [2006Gol] with a maximum of about 100˚C for 1.5–2 at.% Nb. The B2 Ð A2 transition as determined by [2003Ste] and [2005Mor] shows a similar behavior. [2005Gol, 2006Gol] investigated the relaxation phenomena and atom diffusivity in Fe74Al26 with 0.005 to 0.04 at.% C and with and without 0.3 at.% Nb. Nb is supposed to occupy the FeI sites (4b) in the D03-order which gives the L21 structure. From the small change in 59Fe bulk diffusivity through the addition of 0.3 at.% Nb, it is deduced that Nb practically does not influence the vacancy concentration in the as-quenched B2 state. Nevertheless, the segregation of Nb to the grain boundaries results in a decrease of the Fe grain boundary self-diffusion by a factor of 2 to 3 at 1050 K. [1991Sik] describes the melting of Fe-25Al alloys with Nb and other additions by arc melting, air-induction melting, vacuum-induction melting, vacuum-arc remelting and electroslag remelting. The effect of solidification rate on the microstructure of Nb9.1Fe49.3Al41.2 and Nb9.7Fe67.5Al22.8 has been described in respectively [1999Mot] and [2005Mot]. XRD and Mo¨ssbauer studies of [1993Mah] on the FeAl phase and λ phase (Nb(Fe1–xAlx)2) within NbyFeAl1–y alloys with 0.1 ≤ y ≤ 0.3 show that neither phase is pure and stoichiometric. The lattice parameter a of FeAl is constant at 291 pm for all y, which is higher than the observed values for pure FeAl with less than 50 at.% Al. The lattice parameters a and c of λ decrease gradually from 496 to 491 pm, and from 812 to 800 pm as y increases from 0.1 to 0.3, thereby indicating the replacement of Fe by Al atoms. Investigations on the atomic structure, glass forming ability and crystallization behavior of partially or fully amorphous rapidly solidified Al-rich Al-Fe-Nb alloys have been performed for the composition Nb10Fe20Al70 by [1988Tsa], for Nb5Fe8Al87, Nb3Fe13Al84 and Nb10Fe10Al80 by [1997Aud2], for Nb3Fe10Al87 by [1997Aud1, 1998Aud1, 2004Aud]. In [1998Aud2] the mechanical and corrosion properties of (partially) amorphous Nb3Fe10Al87 alloys are discussed. These investigations show that through a proper choice of alloy composition and heat treatment, it is possible to prepare nanocomposites with an amorphous matrix containing metallic nanocrystals. [2001Rod] found that the high-energy ball milling of Nb5Fe5Al90 and Nb3Fe7Al90 also results in a composite microstructure with amorphous and nanocrystalline regions for which the effect of heat treatment on the microstructure is described in [2001Rod, 2003Rod].
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In the review of [2004Mor1] it is mentioned that the λ, Nb(Fe1–xAlx)2 phase has intermediate level of high temperature solubility and low temperature stability in comparison to possible carbides, oxides and nitrides in iron aluminides.
. Table 1 Investigations of the Al-Fe-Nb Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1966Ram]
X-ray investigation
1000˚C, 24 alloys in all composition range
[1970Bur]
X-ray and microstructural investigation
800˚C, 153 alloys in all composition range
[1971Dub]
Isothermal calorimetry using aluminothermic reduction reaction
(Nb0.38Fe0.62)100–xAlx with x = 1.3–6.3 (mass%) and (Nb0.68Fe0.32)100–xAlx with x = 1.9–9.5 (mass%)
[1985Tro]
X-ray powder diffraction, hardness
900–1200˚C, Nb0.8Fe2Al0.2 (at.%)
[1986Bla]
X-ray powder diffraction, hardness
800–1200˚C, NbFe2 to Nb0.5Fe2Al0.5 (at.%)
[1988Dim]
Transmission electron microscopy (SAED and CBED), electron probe microanalysis
600–800˚C, Nb2Fe73Al25 (at.%)
[1990Car]
Ab initio total energy and density of states calculation
Nb(Fe1–xAlx)3 with x = 0.167 and 0.334 (at.%)
[1993Mah]
X-ray diffraction, Mo¨ssbauer spectroscopy
830˚C, FeAl1–xNbx with 0.1≤ x ≤0.3 (at.%)
[1995Ant]
Differential scanning calorimetry
60˚C ≤ T ≤ 1030˚C, NbxFe74–xAl26 with 1 ≤ x ≤ 3 (at.%)
[1996Mac]
Optical and transmission electron T ≤ 1400˚C, Nb33.4Fe33.6Al33.0, microscopy, X-ray diffraction at room and Nb5.1Fe62.2Al32.7, Nb16.7Fe53.9Al29.4 (at.%) high temperature
[1999Mek]
Calculation using a quasi-chemical T = 1000–1250˚C, NbxFe0.5Al0.5–x with method combined with electronic theory x ≤ 0.005 (at.%) in the pseudopotential approximation
[1999Mot]
Directional solidification experiments, optical microscopy
Nb9.7Fe67.5Al22.8, Nb9.1Fe49.7Al41.2, Nb8.0Fe47.0Al45.0 (at.%)
[2001Rod]
X-ray diffraction, scanning electron microscopy, transmission electron microscopy, differential scanning calorimetry
T ≤ 600˚C, Nb5Fe5Al90, Nb3Fe7Al90 (at.%)
[2001Sud]
Thermodynamical calculation of enthalpy Al-Fe-Nb liquid of mixing
[2004Aud]
X-ray diffraction, transmission electron microscopy, differential scanning calorimetry
[2004Mor2] Transmission electron microscopy
DOI: 10.1007/978-3-540-88053-0_8 ß Springer 2009
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T = 550˚C, Nb3Fe7Al90, Nb3Fe10Al87 (at.%)
T = 900 and 1300˚C, Nb2Fe73Al25 (at.%) Landolt‐Bo¨rnstein New Series IV/11E1
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[2005Gol]
Internal friction using a vibrating-reed apparatus and an inverted torsion pendulum, positron annihilation spectroscopy, radiotracer diffusion measurements, transmission electron microscopy
1000˚C ≤ T ≤ 1250˚C, Nb0.3Fe73.7Al26 (at.%)
[2005Mor]
X-ray diffraction, transmission electron microscopy, scanning electron microscopy
as quenched from T = 1200˚C and annealed at T = 500 to 900˚C, Nb2Fe73Al25, Nb5Fe75Al20 (at.%)
[2005Mot]
Directional solidification experiments, T ≤ 1400˚C, Nb9.7Fe67.5Al22.8 (at.%) optical and scanning electron microscopy
[2006Gol]
Anelastic relaxation using a vibratingreed apparatus and an inverted torsion pendulum
[2006Mor1] Transmission electron microscopy, hardness
700˚C ≤ T ≤ 1100˚C, Nb0.3Fe73.7Al26 (at.%)
T ≤ 900˚C, Nb2Fe73Al25 (at.%)
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(Al) < 660.452
cF4 Fm 3m Cu
(αδFe) (αFe) < 912 (δFe) 1538 - 1394
cI2 Im 3m W
(γFe) 1394 - 912 (Nb) < 2469
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Lattice Parameters [pm]
Comments/References
a = 404.96
pure Al at 25˚C [Mas2]
a = 286.65
at 25˚C [Mas2]
a = 293.15
[Mas2] Strukturbericht designation: A2
cF4 Fm 3m Cu
a = 364.67
at 915˚C [V-C2, Mas2]
cI2 Im 3m W
a = 330.04
pure Nb at 25˚C [Mas2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Fe4Al13 < 1160
mC102 C2/m Fe4Al13
Fe2Al5 < 1169
oC24 Cmcm Fe2Al5
FeAl2 < 1156
aP18 P1 FeAl2
ε, Fe2Al3 1232 - 1102
cI16
α2, FeAl < 1310
cP2 Pm 3m CsCl
Fe3Al < 552
cF16 Fm 3m BiF3
NbAl3 < 1680
tI8 I4/mmm TiAl3
Lattice Parameters [pm]
a = 1549.2±0.2 76.0 at.% Al [1994Gri] b = 807.8 ± 0.2 c = 1247.1 ± 0.1 β = 107.69 ± 0.01˚ 70 to 73 at.% Al [1993Kat] at 71.5 at.% Al [1994Bur]
a = 767.5 b = 640.3 c = 420.3 a = 487.8 b = 646.1 c = 880.0 α = 91.75˚ β = 73.27˚ γ = 96.89˚
66 to 66.9 at.% Al [1993Kat] at 66.9 at.% Al [V-C2]
58 to 65 at.% Al [1993Kat] at 61 at.% Al [V-C2]
a = 598.0
DOI: 10.1007/978-3-540-88053-0_8 ß Springer 2009
Comments/References
a = 289.53 a = 289.66 a = 289.77 a = 290.17 a = 290.9
a = 579.30 a = 578.86
a = 384.1 c = 860.9
24 to 55 at.% Al [1993Kat] at 36.2 at.% Al at 38.3 at.% Al at 40.9 at.% Al at 43.8 at.% Al at 50 at.% Al [1958Tay] Strukturbericht designation: B2 25 to 37 at.% Al [2001Ike] at 23.1 at.% Al at 35.0 at.% Al [1958Tay] Strukturbericht designation: D03 Nb solubility: 1 at.% at 1300˚C [2004Mor2] 2 at.% at 1300˚C [1988Dim] 1.6 at.% at 1150˚C [2003Ste] 0.87 at.% at 1000˚C [2003Ste] 1 at.% at 1000˚C [1988Dim] 0.5 at.% at 700˚C [1988Dim] [V-C2] Strukturbericht designation: D022 5 at.% at 1000˚C [1966Ram]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] σ, Nb2Al < 1940
Pearson Symbol/ Space Group/ Prototype tP30 P42/mnm σCrFe
Lattice Parameters [pm]
Comments/References 30 to 42 at.% Al [Mas2] at 32 at.% Al
a = 995.2 c = 516.8 a = 989.6 c = 518.7
at 42 at.% Al [1980Jor] Fe solubility: 10 at.% at 1000˚C [1966Ram] 2 at.% at 800˚C [1970Bur]
Nb3Al < 2060
cP8 Pm 3n Cr3Si
μ, Nb19(Fe1–xAlx)21 Nb19Fe21
hR39 R 3m
< 1530
W6Fe7
μ’, Nb2FeAl
-
λ, Nb(Fe1–xAlx)2 NbFe2 < 1630
hP12 P63mmc MgZn2
τ Nb73–xFe10–yAl17+x+y -
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a = 518.6
18.6 to 25 at.% Al [Mas2]
0 ≤ x ≤0.4 [1966Ram] 48 to 52 at.% Nb (1100 to 1500˚C) [1993Bej1] a = 492.8 c = 268.3
for Nb6Fe7 [V-C2] Al solubility: 20 at.% at 1000˚C [1966Ram] μ type related crystal structure [1966Ram]
-
0 ≤x ≤ 0.75 [1966Ram] [1987Rag]
a = 482.6 c = 787.2 a = 483.2 c = 789.1 a = 483.2 c = 788.3 a = 482.8 c = 787.1 a = 483.8 c = 787.7 a = 484.6 c = 789.2 a = 496.7 c = 807.8
Nb0.9Fe2Al0.1 [1986Bla] Nb0.8Fe2Al0.2 [1985Tro] Nb0.7Fe2Al0.3 [1986Bla] Nb0.6Fe2Al0.4 [1986Bla] Nb0.5Fe2Al0.5 [1986Bla] Nb33.4Fe33.6Al33.0 [1996Mac] 0 ≤ x ≤ 5, 0 ≤ y ≤ 5 approximate concentration deduced from isothermal section at 1000˚C [1999Mot]
-
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. Table 3 Thermodynamic Data of Reaction or Transformation Reaction or Transformation
T [˚C]
Quantity, per mole of atoms [kJ, mol, K]
Comments
(100–x)Nb0.38Fe0.62(L) + xAl(L) Ð (Nb0.38Fe0.62)(100–x) Alx(L)
ΔsH = 105 ± 13 kJ·mol–1
x = 1.3–6.3 (mass%) Heat of solution by isothermal calorimetry using aluminothermic reaction [1971Dub]
(100–x)Nb0.68Fe0.32(L) + xAl(L) Ð (Nb0.68Fe0.32)(100–x) Alx(L)
ΔsH = 100 ± 13 kJ·mol–1
x = 1.9–9.5 (mass%) Heat of solution by isothermal calorimetry using aluminothermic reaction [1971Dub]
(100–x)NbyFe1–y(L) + xAl(L) Ð (NbyFe1–y)(100–x) Alx(L)
ΔmH = –(5–10) kJ·mol–1 ΔmH = –(20–25) kJ·mol–1 ΔmH = –(5–10) kJ·mol–1 ΔmH = –(15–20) kJ·mol–1
x = 0 at.%, y = 0.5 x = 50 at.%, y = 0 x = 50 at.%, y = 1 x = 33 at.%, y = 0.5 Calculated enthalpy of mixing [2001Sud]
. Table 4 Investigations of the Al-Fe-Nb Materials Properties Reference
Method / Experimental Technique
Type of Property
[1986Bla]
PMT-3 metallographic microscope with 100 and 200 gf load
Microhardness
[1991Ran]
Compression tests
Strength, superplastic behavior
[1993Kai]
Corrosion tests in H2/H2O/H2S atmospheres
Corrosion behavior
[1993Mah]
Mo¨ssbauer spectroscopy
Magnetic behavior
[1996Mac]
Vickers hardness, Compression and 4-point bending test
Vickers hardness, brittle-to-ductile transition temperature, stress-strain curve
[1997Cho]
Potentiostatic measurements
Corrosion behavior
[1997Par]
Yield strength measurements, Vickers hardness, charpy impact test
Strength, fracture behavior
[1998Aud2] Microhardness, nanoindentation, potentiodynamic polarization in a chloride containing solution
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Mechanical and corrosion behavior
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8
. Table 4 (continued) Reference
Method / Experimental Technique
Type of Property
[1999Den]
Positron lifetime measurements and calculations of density of valence electrons and grain boundary cohesion
Fracture behavior
[2002Ima]
Ab initio calculations (LAPW)
Magnetic properties
[2004Mor2] Compression tests
Strength
[2005Mor]
Strength
Microscopical investigation by transmission electron microscopy and X-ray diffractions
[2006Mor2] Vickers microhardness testing with 200 g load Hardness [2006Mor1] Vickers microhardness testing with 100 g load, Stress-strain curve, yield stress, compression testing hardness
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Al–Fe–Nb
. Fig. 1 Al-Fe-Nb. The Fe-Nb phase diagram
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. Fig. 2 Al-Fe-Nb. Tentative liquidus surface in the Fe corner
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Al–Fe–Nb
. Fig. 3 Al-Fe-Nb. Isothermal section at 1000˚C
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. Fig. 4 Al-Fe-Nb. Isothermal section at 800˚C
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Al–Fe–Nb
References [1958Tay] [1962Min]
[1966Lun] [1966Ram] [1970Bur]
[1971Dub]
[1980Jor] [1981Ell] [1982Kub] [1985Tro] [1986Bla] [1987Rag]
[1988Dim]
[1988Tsa] [1989Sub]
[1990Car] [1990Kum] [1991Bej]
[1991Ran]
[1991Sik] [1992Gam]
Taylor, A., Jones, R.M., “Constitution and Magnetic Properties of Iron Rich Iron-Aluminium Alloys”, J. Phys. Chem. Solids, 6, 16–37 (1958) (Crys. Structure, Experimental, 49) Mints, R.S., Samsonova, N.N., Malkov, Y.S., “The Effects of Elements of Group V in the Periodic System (V, Nb, Ta) on the Properties of Fe3Al” (in Russian), Dokl. Akad. Nauk SSSR, 144(6), 1324–1327 (1962) (Experimental, Electr. Prop., Mechan. Prop., 6) cited from abstract Lundin, C.E., Yamamoto, A.S., “The Equilibrium Phase Diagram, Niobium (Columbium)Aluminum”, Trans. Met. Soc. AIME, 236, 863–872 (1966) (Phase Diagram, Experimental, 20) Raman, A., “X-ray Investigations in Several T-T5-Al Systems” (in German), Z. Metallkd., 57(7), 535–540 (1966) (Phase Diagram, Phase Relations, Experimental, #, 5) Burnashova, V.V., Ryabov, V.R., Markiv, V.Y., “Investigation of the Nb-Fe-Al and Nb-Co-Al Systems”, Dopov. Akad. Nauk Ukr. RSR, Ser. A: Fiz.-Mat. Tekh. Nauki, A8, 747–750 (1970) (Phase Diagram, Experimental, Phase Relations, #, 11) Dubrovin, A.S., Gorelkin, O.S., Demidov, Yu.Ya., Chirkov, N.A., Kostylev, L.S., Kolesnikova, O.D., “Calorimetric Investigation of Heat Solution of Silicon and Aluminium in Aluminium Thermal Alloys” (in Russian), Metallotherm. Process. Khim. Met. Mater. 1971, 121–130 (1971) (Thermodyn., Experimental, 11) Jorda, J.L., Fluekiger, R., Mueller, J., “A New Metallurgical Investigation on the Niobium-Aluminium System”, J. Less-Common Met., 75, 227–239 (1980) (Phase Diagram, Experimental, #, 20) Elliott R.P., Shunk, F.A., “The Al-Nb System”, Bull. Alloy Phase Diagrams, 2, 75–81 (1981) (Phase Diagram, Crys. Structure, Review, 31) Kubaschewski, O., “Al-Fe” in “Iron Binary Phase Diagrams”, Springer Verlag, Berlin 5–9 (1982) (Phase Diagram, Review, 26) Trojko, R., Blazina, Z., “Metal-Metalloid Exchange in some Friauf-Laves Phases Containing Two Transition Metals”, J. Less-Common Met., 106, 293–300 (1985) (Crys. Structure, Experimental, 13) Blazina, Z., Trojko, R., “Structural Investigations of the Nb1–xSixT2 and Nb1–xAlxT2 (T = Cr, Mn, Fe, Co, Ni) Systems”, J. Less-Common Met., 119, 297–305 (1986) (Crys. Structure, Experimental, 6) Raghavan, V., “Section I. The Al-Fe-Nb (Aluminum-Iron-Niobium) System”, Phase Diagrams of Ternary Iron Alloys. Part 1, Ind. Inst. Techn. Delhi, 1, 3–8 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Review, 8) Dimiduk, D.M., Mendiratta, M.G., Banerjee, D., Lipsitt, H.A., “A Structural Study of Ordered Precipitates in an Ordered Matrix within the Fe-Al-Nb System”, Acta Metall., 36(11), 2947–2958 (1988) (Phase Relations, Morphology, Experimental, 17) Tsai, A.-P., Inoue, A., Masumoto, T., “A New Icosahedral Al-Fe-Ta Alloy Prepared by Rapid Solidification”, Jpn. J. Appl. Phys., 27(1), L5-L8 (1988) (Experimental, Morphology, 18) Subramanian, P.R., Simmons, J.P., Mendiratta, M.G., Dimiduk, D.M., “Effect of Solutes on Phase Stability in Al3Nb”, Mater. Res. Soc. Symp. Proc., 133(3), 51–56 (1989) (Phase Diagram, Experimental, Mechan. Prop., 12) Carlsson, A.E., Meschter, P.J., “Relative Stabilities of L12 and DO22 Structures in Ternary Mal3-Base Aluminides”, J. Mater. Res., 5(12), 2813–2818 (1990) (Crys. Structure, Thermodyn., Experimental, 15) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35(6), 293–327 (1990) (Crys. Structure, Phase Diagram, Review, 158) Bejarano, Z.J.M. Gama, S., Ribeiro, C.A., Effenberg, G., Santos, G., “On the Existence of the Fe2Nb3 Phase in the Fe-Nb System”, Z. Metallkd., 82, 615–620 (1991) (Phase Diagram, Phase Relations, Experimental, 7) Ranganath, S., Dutta, A., Subrahmanyam, J., “On the flow Behaviour of Combustion Synthesized FeAl-Nb System”, Scr. Metall. Mater., 25(7), 1593–1596 (1991) (Morphology, Experimental, Mechan. Prop., 6) (cited from abstract) Sikka, V.K., “Production of Fe3Al-Based Intermetallic Alloys”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 907–912 (1991) (Experimental, 2) Gama, S., “Aluminium - Iron - Niobium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.18061.1.20, (1992) (Crys. Structure, Phase Diagram, Assessment, 10)
DOI: 10.1007/978-3-540-88053-0_8 ß Springer 2009
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Al–Fe–Nb [1993Bej1] [1993Bej2] [1993Kai] [1993Kat]
[1993Mah]
[1994Bur]
[1994Gri]
[1995Ant] [1996Mac] [1997Aud1] [1997Aud2] [1997Cho]
[1997Par]
[1998Aud1]
[1998Aud2]
[1999Den]
[1999Mek]
[1999Mot]
[2001Ike]
[2001Rod] [2001Sud]
8
Bejarano, Z.J.M., Gama, S., Ribeiro, C.A., Effenberg, G., “The Iron-Niobium Phase Diagram”, Z. Metallkd., 84, 160–164 (1993) (Experimental, Phase Diagram, #, 6) Bejarano, Z.J.M, PhD Thesis, State Uni. Campinas, Brazil (1993) as quoted in [1999Mot] Kai, W., Douglass, D.I., “The Corrosion Behaviour of Fe-Nb-Al alloys in H2/H2O/H2S Atmospheres”, Oxidation of Metals, 39(5-6), 317–352 (1993) (Electrochem., Experimental, 12) (cited from abstract) Kattner, U.R., Burton, B.P., “Al-Fe (Aluminum-Iron)” in “Phase Diagrams of Binary Iron Alloys”, Okamoto, H. (Ed), ASM International, Materials Park, OH 44073–0002, 12–28 (1993) (Crys. Structure, Phase Diagram, Thermodyn., Review, Electr. Prop., Magn. Prop., 99) Mahmood, S.H., Awawdeh, M.A., Saleh, A.S., “Structural and Mossbauer Studies of the Alloy System FeAl1–xNbx”, J. Appl. Phys., 73(10), 5663–5665, part 2A (1993) (Phase Relations, Experimental, Magn. Prop., 12) Burkhardt, U., Grin, J., Ellner, M., Peters, K., “Structure Refinement of the Iron-Aluminum Phase with the Approximate Composition Fe2Al5”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., 50, 313–316 (1994) (Crys. Structure, Experimental, 9) Grin, J., Burkhardt, U., Ellner, M., Peters, K., “Refinement of the Fe4Al13 Structure and its Relationship to Quasihomological Homeotypical Structures”, Z. Kristallogr., 209, 479–487 (1994) (Crys. Structure, Experimental, 39) Anthony, L., Fultz, B., “Effect of Early Transition Metal Solutes on the D03 - B2 Critical Temperature of Fe3Al”, Acta Metall. Mater., 43(10), 3885–3891 (1995) (Phase Relations, Experimental, 35) Machon, L., Sauthoff, G., “Deformation Behaviour of Al-Containing C14 Laves Phase Alloys”, Intermetallics, 4, 469–481 (1996) (Phase Relations, Experimental, 41) Audebert, F, Sirkin, H., Escorial, A.G., “Approach to the Atomic Structure of Amorphous Al-Fe-Nb”, Philos. Mag. B, 76(4), 483–487 (1997) (Electronic Structure, Experimental, 18) (cited from abstract) Audebert, F, Sirkin, H., Escorial, A.G., “Aluminum-base Al-Fe-Nb Amorphous and Nanostructured Alloys”, Scr. Mater., 36(4), 405–410 (1997) (Morphology, Experimental, 19) Choe, H.-C., Choi, D.-C., “Effects of Cr, Mo, Nb and B on the Intergranular Corrosion Behaviour of Fe-25 at.% Al Intermetallic Compounds”, J. Korean Inst. Met. Mater., 35(4), 468–476 (1997) (Electrochem., Experimental, 16) (cited from abstract) Park, K.-I., Joo, S.-M., Choi, D.-C., “Effects of Cr, Mo, Nb and B Additions on the Microstructure and Mechanical Properties in Fe-25 at.% Al Alloy”, J. Korean Inst. Met. Mater., 35(3), 305–311 (1997) (Morphology, Experimental, Mechan. Prop., 17) (cited from abstract) Audebert, F., Sirkin, H., Escorial, A.G., “Study of the Crystallisation Process of the Al-Fe-Nb Amorphous Alloys”, Proceedings of the Fifth International Workshop on Non-Crystalline Solids. NonCrystalline and Nanoscale Materials. World Scientific. 1998, Singapore, Singapore, 367–372 (1998) (Morphology, Experimental, 19) (cited from abstract) Audebert, F., Vazquez, S., Gutierrez, A., Vergara, I., Alvarez, G., Garcia Escorial, A., Sirkin, H., “Mechanical and Corrosion Behaviour of Al-Fe-Nb Amorphous Alloys”, Mechanically Alloyed, Metastable and Nanocrystalline Materials, part 2, Mater. Sci. Forum, 269-272, 837–842 (1998) (Electrochemistry, Experimental, Mechan. Prop., 17) (cited from abstract) Deng, W., Zhong, X.P., Huang, Y.Y., Xiong, L.Y., Wang, S.H., Guo, J.T., Long, Q.W., “Effects of Nb and Si on Densities of Valence Electrons in Bulk and Defects of Fe3Al Alloys”, Sci. China, Ser. A, 42(1), 87–92 (1999) (Electronic Structure, Experimental, 15) (cited from abstract) Mekhrabov, A.O., Akdeniz, M.V., “Effect of Ternary Alloying Elements Addition on Atomic Ordering Characteristics of Fe-Al Intermetallics”, Acta Mater., 47(7), 2067–2075 (1999) (Thermodyn., Calculation, Theory, 63) Mota, M.A., Coelho, A.A., Bejarano, J.M.Z., Gama, S., Caram, R., “Directional Growth and Characterization of Fe-Al-Nb Eutectic Alloys”, J. Cryst. Growth, 198(part 1), 850–855 (1999) (Phase Relations, Experimental, 19) Ikeda, O., Ohnuma, I., Kainuma, R., Ishida, K., “Phase Equilibria and Stability of Ordered BCC Phases in the Fe-rich Portion of the Fe-Al System”, Intermetallics, 9, 755–761 (2001) (Phase Diagram, Experimental, Mechan. Prop., 18) Rodriguez, C.A.D., Botta, F.W.J., “High-Energy Ball Milling of Al-Based Alloys”, Key Eng. Mater., 189191, 573–578 (2001) (Crys. Structure, Experimental, 10) Sudavtsova, V.S., Vovkotrub, N.E., Kudin, V.G., “Thermodynamic Properties of the Fe-Nb(Ta,W), Fe-Nb(Ta,W)-Al Alloys” (in Russian), Metally, (2), 18–20 (2001) (Thermodyn., Experimental, 7)
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[2003Rod]
[2003Ste]
[2004Aud]
[2004Mor1]
[2004Mor2] [2005Gol]
[2005Mor] [2005Mot]
[2006Gol]
[2006Mor1]
[2006Mor2] [2006MSIT]
[Mas2] [V-C2]
Al–Fe–Nb Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2 Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Review, 27) Imaizumi, M., Las, B., Bejarano, J.M.Z., Gama, S., Coelho, A.A., Mota, M.A., Caram, R., “The Magnetic Properties of Fe-Al-Nb Intermetallic Compounds”, Intermag Europe 2002 Digest of Technical Papers, 2002 IEEE international Magnetics Conference, (Cat.No.02CH37323), Piscataway, NJ, USA, AR12 (2002) (Calculations, Magn. Prop., 5) (cited from abstract) Rodriguez C.A.D., Yavari, A.R., Kiminami, C.S, Botta F.W.J., “Milling and Hot Consolidation of Al-FeNb Alloy”, Mater. Sci. Forum, 415-418, 287–292 (2003) (Morphology, Experimental, 18) (cited from abstract) Stein, F., Schneider, A., Frommeyer, G., “Flow Stress Anomaly and Order-disorder Transition in Fe3Albased Fe-Al-Ti-X Alloys with X = V, Cr, Nb, or Mo”, Intermetallics, 11, 71–82 (2003) (Phase Relations, Experimental, 53) Audebert, F., Mendive, C., Vidal, A., “Structure and Mechanical Behaviour of Al-Fe-X and Al-Ni-X Rapidly Solidified Alloys”, Mater. Sci. Eng. A, 375-377, 1196–1200 (2004) (Electronic Structure, Experimental, 22) Morris, D.G., Munoz-Morris, M.A., Chao, J., “Development of High Strength, High Ductility and High Creep Resistant Iron Aluminide”, Intermetallics, 12(7-9), 821–826 (2004) (Interface Phenomena, Morphology, Phase Relations, Review, Mechan. Prop., 52) Morris, D.G., Munoz-Morris, M.A., Baudin, C., “The High-Temperature Strength of Some Fe3Al Alloys”, Acta Mater., 52(9), 2827–2836 (2004) (Morphology, Experimental, Mechan. Prop., 44) Golovin, I.S., Divinski, S.V., Cizek, J., Prochazka, I., Stein, F., “Study of Atom Diffusivity and Related Relaxation Phenomena in Fe3Al-(Ti,Nb)-C Alloys”, Acta Mater., 53(9), 2581–2594 (2005) (Thermodyn., Transport Phenomena, Experimental, 40) Morris, D.G., Requejo, L.M., Munoz-Morris, M.A., “A Study of Precipitation in D03 Ordered Fe-Al-Nb Alloy”, Intermetallics, 13(8), 862–871 (2005) (Morphology, Phase Relations, Experimental, 29) Mota, M.A., Coelho, A.A., Bejarano, J.M.Z., Gama, S., Caram, R., “Fe-Al-Nb Phase Diagram Investigation and Directional Growth of the (Fe,Al)2Nb-(Fe,Al,Nb)ss Eutectic System”, J. Alloys Compd., 399(1-2), 196–201 (2005) (Phase Relations, Experimental, 18) Golovin, I.S., Strahl, A., Neuhaeuser, H., “Anelastic Relaxation and Structure of Ternary Fe-Al-Me Alloys with Me = Co, Cr, Ge, Mn, Nb, Si, Ta, Ti, Zr”, Int. J. Mater. Res. (Z. Metallkd.), 97(8), 1078–1092 (2006) (Crys. Structure, Morphology, Phase Relations, Thermodyn., Experimental, Mechan. Prop., 79) Morris, D.G., Munoz-Morris, M.A., Requejo, L.M., Baudin, C., “Strengthening at High Temperatures by Precipitates in Fe-Al-Nb Alloys”, Intermetallics, 14, 1204–1207 (2006) (Morphology, Phase Relations, Experimental, Mechan. Prop., 24) Morris, D.G., Requejo, L.M., Munoz-Morris, M.A., “Age Hardening in some Fe-Al-Nb Alloys”, Scr. Mater., 54(3), 393–397 (2006) (Morphology, Phase Relations, Experimental, Mechan. Prop., 16) “Al-Fe (Aluminum-Iron)”, Diagrams as Published, in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services, GmbH, Stuttgart; Document ID: 30.10236.1.20, (2006) (Crys. Structure, Phase Diagram, Phase Relations, #, 11) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Molybdenum – Nickel Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Kostyantyn Korniyenko, Vasyl Kublii
Introduction Experimental investigations of the phase equilibria in the Al-Mo-Ni system were started by [1925Pfa] and [1933Roe] and, as summarized in [1976Mon], concerned the Ni-rich corner as well as the influence of additions of Mo and Ni on the Al solid solution, respectively. The investigations of the partial Ni-NiAl-Mo system were further developed by [1959Gua2, 1960Bag, 1965Ram, 1976Jac, 1977Aig, 1977Pea, 1978Gul, 1983Nas, 1983Wak, 1984Kov1, 1984Kov2, 1984Mir, 1985Nas, 1986Mas1, 1986Mas2, 1988Mas, 1989Mas, 1989Hon1, 1989Hon2, 1991Mis]. Results of phase equilibria studies for the Al-rich corner are presented by a series of isothermal sections [2002Gru]. The complete ternary system has been investigated experimentally at 600˚C [1971Pry], 800˚C [1969Mar] and 950˚C [1969Vir]. For preparation of the alloys most of the authors used arc melting, while [1971Pry] and [2002Gru] applied levitation induction melting, and [1984Mir] obtained specimens by both conventional arc-casting and powder metallurgy techniques. The traditional methods of investigations were X-ray diffraction (XRD), metallography, differential thermal analysis (DTA), electron microprobe analysis (EMPA). Some authors used scanning electron microscopy (SEM) [1989Hon1, 1989Hon2, 1991Mis, 2002Gru], transmission electron microscopy (TEM) [2002Gru], as well as energy-dispersive X-ray spectroscopy (EDXS) [1991Mis]. Calculations of phase equilibria were carried out by [1974Kau, 1999Kau] and [1999Lu]. A critical review of literature data on phase equilibria in the Al-Mo-Ni system was presented in the assessment of [1993Kub]. Further experimental studies are necessary in order to construct the liquidus surface and the reaction scheme of the complete ternary system as well as isothermal sections in the whole range of compositions.
Binary Systems The Al-Mo, Al-Ni and Mo-Ni systems are accepted from [2005Sch], [2004Sal] and [Mas2], respectively.
Solid Phases Crystallographic data on the known unary, binary and ternary phases are listed in Table 1. [1959Gua2] reported the existence of a ternary phase ψ of composition Mo7,5Ni58,0Al34,5 at Landolt‐Bo¨rnstein New Series IV/11E1
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1175˚C, but did not determine its crystal structure. However, the data of [1960Bag, 1969Mar, 1971Pry, 1983Nas] and [1984Mir] did not confirm its existence. A phase of similar composition was easily obtained by [1969Vir] in the alloys (at.%) Mo50Ni25Al25, Mo43Ni31Al26 and Mo9Ni53Al38, the latter being fairly close to the composition of the reported ψ phase [1959Gua2]. Thus, [1969Vir] concluded that the ψ phase in fact did not belong to the AlMo-Ni system, but was easily stabilized by small amounts of impurities (low purity 99.8 mass% nickel was used for preparation of the specimens!). The ψ phase was indexed as a MgZn2 type Laves phase (a = 474, c = 770 pm). Two ternary compounds have been identified in the Al-rich range of compositions, namely, τ1, Mo(NixAl1–x)3 [1969Mar, 1969Vir, 1971Pry] and τ2. The composition of τ2 was determined by [1969Mar] and [1971Pry] as Mo5Ni18Al77, but [2002Gru] corrected it and determined crystal system and lattice parameters (Table 1). According to the findings of [1969Mar] at 800˚C, [1971Pry] at 600˚C and [1965Ram, 1969Vir] at 950˚C, the τ1 phase is likely to exhibit a homogeneity range strongly dependent on temperature. The compound Mo(Al2,75Ni0,25), observed in the aluminothermic preparation of Al-Mo-Ni alloys from Mo- and Ni-oxides, is likely to be isotypic with the TiAl3 type, despite the fact that the c-axis corresponds to a twofold superstructure [1969Rec]. [1984Och1, 1984Och2, 1985Mis, 1988Mas, 1989Hon1, 1989Hon2] investigated the influence of the addition of a third element on the lattice parameter change of binary Ni solid solutions. The temperature dependence of the solid solubility is reflected in the isothermal sections (see below) and additional information on the γ/γ´ boundary may be obtained from [1989Gai] and [1989Hon1]. A model based on X-ray measurements to show the effect of Mo on the γ´ structure has been suggested by [1977Aig].
Quasibinary Systems On the basis of X-ray, DTA as well as optical microscopy data [1986Mas2] plotted the phase diagram of the partial quasibinary NiAl-Mo system. The temperature of the LÐβ+α equilibrium is equal to 1600±7˚C; the maximum solubility of molybdenum in the β phase is less than 4 at.%. The eutectic point is placed at 10 at.% Mo, and its co-ordinates were later confirmed by [1991Sas]. Part of the quasibinary section in the range of compositions 0 to 20 at.% Mo is presented in Fig. 1, with small changes according to the melting temperature of the β phase at 1651˚C [2004Sal], whereas 1638˚C was accepted by [1986Mas2]. Similar compositions of the eutectic point were reported earlier by [1970Cli] (9 at.% Mo) and [1971Pry] (10 at.% Mo), but considerably lower eutectic temperatures were presented (1427 and 1290˚C, respectively). According to the conclusion of [1993Kub], in view of the high melting temperature of the β phase and the reaction temperature of U1 (1340˚C), the higher quasibinary eutectic temperature (1600˚C) is recommended. The vertical section Mo-Ni3Al, according to the data of [1971Pry], demonstrates a peritectoid reaction β+γÐγ´+α, which is in contradictions to the observation of a eutectic solidification behavior in this area by [1976Spr] and [1983Nas].
Invariant Equilibria The reaction scheme of the partial Mo-NiAl-Ni system is presented in Fig. 2. One invariant three-phase equilibrium as well as six invariant four-phase reactions have hitherto been DOI: 10.1007/978-3-540-88053-0_9 ß Springer 2009
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observed in the system. [1977Pea] reported the equilibrium LÐα+β+γ at 1300˚C, but later it was established by [1977Aig, 1983Nas] and [1983Wak] that instead of the β phase the γ´ phase takes part in the eutectic reaction LÐα+γ+γ´, and the various authors merely agree on the temperature of this reaction at 1300˚C [1977Pea, 1983Wak, 1988Mas]. The reaction temperatures in Fig. 2 were measured and selected by [1986Mas1, 1986Mas2] and [1988Mas]. Table 2 presents the compositions of phases taking part in the invariant equilibria, estimated on the basis of isothermal sections as well as on the data calculated by [1987Sve] and experimentally determined by [1986Mas1] and [1986Mas2].
Liquidus Surface Liquidus surface projection of the Ni-rich region (the Ni-NiAl-Mo partial system) is presented in Fig. 3. It consists of five fields of primary crystallization corresponding to the α, β, γ, γ´ and δ phases. It has been constructed on the basis of constitution of the accepted binary phase diagrams and critically assessed experimental data of different authors. So, the position of the U2E monovariant curve is established using experimental data of [1977Pea, 1984Kov1] and [1984Kov2] on directionally solidified α+γ eutectic superalloys containing 8.58Al-27.22Mo (at.%) up to 18.66Al-15.50Mo (at.%) and 14.38Al-20.03Mo (at.%), as well as data for two alloys, crystallized by [1987Sve] using the Bridgeman method. The position of the U2 invariant point was accepted on the basis of data by [1977Pea] (Table 2), because data by [1987Sve] do not agree with the estimated compositions of the α, δ and γ phases participating in the equilibrium LU2+δÐα+γ at 1310˚C. [1974Tho, 1976Nes, 1976Hen] and [1976Spr] discovered by directional solidification studies the existence of eutectic reactions LÐα+γ´, LÐα+γ and LÐδ+γ. The Ni-NiAl-Mo liquidus surface projection was proposed by [1986Mas2] based on a rather schematic projection given by [1983Nas], but the constitution of the Al-Ni binary as used by [1983Nas] contradicts the assessment of [2004Sal]. A mathematical model was used to construct the isotherms at 1360 and 1340˚C, as well as the monovariant curves p2U2 and U2E [1987Sve, 1989Gai]. Results of thermodynamic calculations of the liquidus surface, carried out by [1999Lu] and based on the experimental data of [1987Sve], were also used in our assessment, except the position of U2, which is placed by [1999Lu] at a smaller Mo content. In Fig. 3 isotherms illustrating the shape of the surface, are added, in particular, the isotherms at 1415, 1425 and 1445˚C, using the experimental data of [1978Gul].
Isothermal Sections Partial isothermal section at 1260˚C is presented in Fig. 4 according to the data of [1984Mir]. Isothermal sections at 1200˚C constructed from the experimental results by [1983Nas] and [1988Mas] are in good agreement with each other and with the calculation performed by [1974Kau], see Fig. 5. The character of phase equilibria in the Ni-rich corner is similar to the character of the assessed equilibria, but the solubility of Mo in the γ´ phase is much smaller than experimental data, and also the position of the γ´+γ two-phase region is different. Phase equilibria at 1100˚C [1988Mas] and 1000˚C [2002Gru] are shown in Figs. 6 and 7, respectively. Figure 8 presents a combination of data at close temperatures: for 1050˚C in the Al-rich range by [2002Gru] and for 1038˚C in the Al-poor range [1984Mir]. Phase equilibria at 927˚C [1984Mir] for Al-poor range and at 950˚C [2002Gru] for Al-rich corner are merged Landolt‐Bo¨rnstein New Series IV/11E1
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in Fig. 9. Partial isothermal sections at 880 (Fig. 10) and 700˚C (Fig. 11) are accepted from [1988Mas]. In the assessed isothermal sections some minor modifications have been made taking into account the newly determined position of the γ/γ´ boundary according to SEM and DTA data by various groups (see “Solid Phases”) and according to the constitution of the boundary systems. In particular, Fig. 11 reflects the participation of the later determined Ni5Al3 phase in the equilibria at 700˚C. [1989Hon1, 1989Hon2] and [1991Mis] confirmed that the extent of the γ(Ni) solid solution area increases with rising temperature. The position of the nickel-rich boundary of the γ´ phase at 1200˚C, calculated by [1991Eno] using the cluster variation method, CVM (which utilizes the tetrahedron approximation and the phenomenological Lennard-Jones pair interaction potential), practically coincides with the data of [1983Nas].
Temperature – Composition Sections Figure 12 shows the partial isopleth at 14 at.% Al for a Ni content changing from 58 to 86 at.% according to the data of [1989Mas]. This isopleth crosses two volumes of primary crystallization, corresponding to the α and γ phases, and four planes of invariant four-phase equilibria, in one of which (at ≈ 1300˚C) the liquid phase takes part, and the others are with participation of only the solid phases (at ≈ 1130, ≈ 890 and ≈ 730˚C). The partial isopleth at 65 at.% Ni with Mo content changing from 15 to 35 at.%, as constructed by [1983Wak], and the isopleths Mo60Al40 - Ni, Mo45Al55 - Ni constructed by [1986Mas2] do not comply with the assessed liquidus surface.
Thermodynamics No experimental thermodynamic data concerning the Al-Mo-Ni system are published in literature. [1974Kau] calculated the isothermal sections at 1727, 1527 and 1200˚C using symmetrical functions for the excess free energies of mixing. There is a substantial disagreement between the calculated and the experimental data; moreover constitution of the calculated binaries Al-Ni, Al-Mo and Mo-Ni contradict the phase diagrams accepted in this assessment. [1999Kau] and [1999Lu] assessed the experimental phase equilibria data in order to evaluate the thermodynamic parameters of the ternary system by means of the CALPHAD method. A substitutional-solution model is used to describe liquid, face-centered cubic (fcc) and body-centered cubic (bcc) phases, while a sublattice model is used to describe the intermetallic phases. Two sets of thermodynamic descriptions have been obtained, and comparison has been made between them. There is satisfactory agreement between the calculations and experimental data. But phase diagrams of the boundary systems Al-Ni, Mo-Ni and Al-Mo, accepted by [1999Kau] and [1999Lu], disagree to some extent with the phase diagrams accepted in this assessment. [2000Bor] presented a general survey of the diffusion-controlled transformations (DICTRA) software as an engineering tool for diffusion simulations in multicomponent alloys. The model of coarsening of the γ´ phase particles in ternary Al-Mo-Ni alloys was used. In the calculation, the alloy composition was adjusted in
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order to have the same fraction of the γ´ phase as experimentally observed. This gave a small difference in composition compared with the experimental data.
Notes on Materials Properties and Applications The Al-Ni alloys with the addition of a refractory metal element (in particular, molybdenum) are interesting as materials for production of in situ composites of eutectic superalloys that can serve, in particular, as materials for specific hot section components of turbine engines, primarily blades or buckets and vanes as well as nozzles [1976Jac]. In spite of very complicated alloys compositions, commercial superalloys generally consist mainly of two phases, namely, γ and γ´. The β phase has potential applications such as hot sections of gas turbine engines for aircraft propulsion systems, coats under thermal barrier coating, electronic metallization compounds in advanced semiconductors [1998Mur] as well as surface catalysts [1971Nal, 1998Mur]. The influence of molybdenum additions on the structure and hardness of the γ´ phase based alloys has been studied by [1959Gua1], and three general effects have been observed, namely, solid-solution hardening, strain aging, and defect hardening arising from deviations from stoichiometry. A method for the determination of site preference of substitutional elements in intermetallic compounds was proposed by [2001Ter], and it was demonstrated that in the γ´ based alloys molybdenum substitutes preferentially for aluminium. The microstructure and chemical characteristics of the nanocrystalline β phase are studied by [2002Alb]. It was established that the addition of molybdenum tends to slightly refine the grain size of the β phase based alloy. The specimens with 2, 4 or 6 at.% Mo are polycrystalline containing, at the same time, the β phase, Ni, Al and Mo. The cast alloy NiAl-9Mo (at.%), prepared by [2002Guo], exhibited typical deformation characteristics shown in conventionally superplastic materials, and possessed finely grained structure. Properties of the directionally solidified eutectic superalloys were investigated by [1973Wal, 1974Tho, 1976Jac, 1976Nes, 1976Spr, 1981Sch, 1984Sch, 1985Nas, 1986Hor, 1986Kau] and [1987Sve]. It is shown that composites formed by directional eutectic solidification combined with a reinforcing α phase in the form of fibers have a considerable advantage over conventional superalloys [1973Wal, 1974Tho, 1976Jac, 1976Nes, 1976Spr]. Since the microstructure derives directly from the melt, the composites are extremely stable at elevated temperatures. In addition improved oxidation and creep resistances have been observed [1981Sch, 1985Nas]. The characteristic microstructure of the alloys consists of γ/γ´ matrix reinforced with faceted αMo fibres, which are primarily square in cross section and with the following orientation relationship: [001]γ´// [001]α, (010)γ//(010)γ´//(110)α, (100)γ´//(110) [1981Sch, 1984Sch, 1986Hor, 1986Kau, 1987Sve]. Precipitation in Ni-rich Al-Mo-Ni alloys has been investigated in the temperature range 600 to 1100˚C by transmission electron microscopy, selected-area electron diffraction and hardness measurements [1987San]. Various stable and metastable phases (α, γ´ and δ MoNi, MoNi2 (MoPt2 type), MoNi3 (TiAl3 type), MoNi4, MoNi8 and SRO) have been observed and the ranges of alloy composition and aging temperature for which each phase is formed have been determined and their strengthening influence on the mechanical properties of ternary Al-Mo-Ni alloys has been discussed [1987San]. Convergent-beam electron diffraction has been used by [1986Kau] to reveal local lattice distortions in directionally
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solidified γ´, Ni3Al type alloys with 12.8 at.% Al and 22.2 at.% Mo. TEM data [1990Yam] from a Mo20Ni75Al5 alloy annealed at 800˚C and quenched revealed the close-packed planes of the γ´´ and γ´ phases to be parallel: [100]γ´//[110]γ´, (010)γ´´//(111)γ´ and [001]γ//[112]γ´. [2001Kai] studied the effect of molybdenum on the morphological stability of the interface between the γ´ and β phases using Al-Mo-Ni ternary diffusion couples annealed at temperatures ranging from 900 to 1300˚C. Nonplanar interfaces with the Widmanstaetten-like structure were formed in the couples.
. Table 1 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] (Al) < 660.452
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] cF4 Fm 3m Cu
a = 404.88
cI2 Im 3m W
a = 314.7
pure Mo, at 25˚C [V-C]
a = 314.6 a = 314.5
x = 0, y = 0.004 [1980Fer] x = 0, y = 0.009 [1980Fer] y = 0, x = 0 to 0.035, at 1000˚C [1967Bel] y = 0, x = 0 to 0.055, at 1205˚C [1951Ham] y = 0, x = 0 to 0.068, at 1316˚C [1951Ham] y = 0, x = 0 to 0.077, at 1317˚C [1951Ham] y = 0, x = 0 to 0.096, at 1482˚C [1951Ham] y = 0, x = 0 to 0.11, at 1572˚C [1982Shi] y = 0, x = 0 to 0.108, at 1600˚C [1971Rex] y = 0, x = 0 to 0.114, at 1604˚C [1982Shi] y = 0, x = 0 to 0.14, at 1700˚C [1971Rex] y = 0, x = 0 to 0.138, at 1748˚C [1982Shi] y = 0, x = 0 to 0.195, at 2150˚C [1951Ham]
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pure Al at 24˚C [V-C]
x = 0, y = 0 to 0.004 [2004Sal] y = 0, x = 0 to 0.00028 at 400˚C [1960Vig] y = 0, x = 0 to 0.00062 at 640˚C [1960Vig] y = 0, x = 0 to 0.0007 at 640˚C, by extrapolation [1960Vig]
MoxNiyAl1–x–y
α, (Mo) < 2623 α,(Mo1–x–yNiyAlx)
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. Table 1 (continued) Phase/ Temperature Range [˚C] γ, (Ni) < 1455 γ, (MoxNi1–x)
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] cF4 Fm3m Cu
a = 352.40 a = 352.32 a = 355.8 a = 353.9 a = 355.2 a = 356.5 a = 356.3 a = 361.0
γ, (Ni1–xAlx)
a = 352.8 a = 353.2
Comments/References pure Ni at 25˚C [1984Och2, Mas2] pure Ni at 20˚C [V-C] quenched from 800˚C [V-C] x = 0.03, quenched from 1000˚C [1984Och1, 1984Och2, 1985Mis] x = 0.06, quenched from 1000˚C [1984Och1, 1984Och2, 1985Mis] x = 0.09, quenched from 1000˚C [1984Och1, 1984Och2, 1985Mis] x = 0.097 [1980Fer] x = 0.218 [1980Fer] x = 0 to 0.2 [2004Sal] x = 0.2 at 1372˚C [2004Sal] x = 0.025 Slowly cooled alloy [1952Tay] x = 0.05 Slowly cooled alloy [1952Tay]
θ, MoNi4 < 870
tI10 I4/m MoNi4
a = 572.0 c = 356.4
[V-C]
γ´´, MoNi3 < 910
oP8 Pmmn TiCu3
a = 506.4 b = 422.2 c = 444.8
[V-C]
δ, MoNi < 1362
oP112 Cmcm MoNi
(Mo,Ni,Al)1 (Ni,Mo,Al)1
a = 910.8 b = 910.8 c = 885.2 a = 455 b = 1663 c = 873
46 to 48 at.% Ni [Mas2] at 50.8 at.% Ni [1980Fer]
[1997Jin]
0 to 2 at.% Al [1969Vir] 0 to 1.6 at.% Al, T = 1260˚C [1984Mir] 0 to 1.2 at.% Al, T = 1200˚C [1988Mas] 0 to 1.1 at.% Al, T = 1171˚C [1984Mir] 0 to 1.1 at.% Al, T = 1100˚C [1988Mas] 0 to 0.6 at.% Al, T = 1093˚C [1984Mir] 0 to 0.5 at.% Al, T = 1038˚C [1984Mir] 0 to 0.3 at.% Al, T = 927˚C [1984Mir] MoAl12 < 712
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a = 757.3 a = 758.15
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. Table 1 (continued) Phase/ Temperature Range [˚C] MoAl5 (h2) 846 to 800 - 750
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] hP12 P63 WAl5
hP60 MoAl5 (h1) 3 800 - 750 to 648 P MoAl5 (h1)
Comments/References
a = 491.2 c = 886.0 a = 489 c = 880
83.8 at.% Al [1991Sch]
a = 493.3 c = 4398
at 83.3 at.% Al [1991Sch]
a = 495.1 c = 2623
at 83.3 at.% Al [1991Sch]
[1980Fer]
MoAl5 (r) ≲ 648
hP36 R3c MoAl5 (r)
Mo5Al22 964 to 831
oF216 Fdd2 Mo5Al22
Mo4Al17 < 1034
mC84 C2 Mo4Al17
MoAl4 1177 to 942
mC30 Cm WAl4
Mo1–xAl3+x 1260 to 1154
cP8 Pm 3n Cr3Si
MoAl3 1222 to 818
mC32 C2/m MoAl3
a = 1639.6 ± 0.1 b = 359.4 ± 0.1 c = 838.6 ± 0.4 β = 101.88 ± 0.07
Mo3Al8 < 1555 ± 10
mC22 Cm Mo3Al8
72 to 75 at.% Al [Mas2] a = 920.8 ± 0.3 [1962For] b = 363.78 ± 0.03 c = 1006.5 ± 0.3 β = 100.78 ± 0.05˚
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a = 7382 ± 3 b = 916.1 ± 0.3 c = 493.2 ± 0.2 a = 915.8 ± 0.1 b = 493.23 ± 0.08 c = 2893.5 ± 0.5 β = 96.71 ± 0.01
81.7 at.% Al [1991Sch] [1995Gri]
80.9 at.% Al [1991Sch] [1995Gri]
79 to 80 at.% Al [1991Sch] a = 525.5 ± 0.5 [1964Lea] b = 1776.8 ± 0.5 c = 522.5 ± 0.5 β = 100.88 ± 0.06˚ a = 525.5 [1991Sch] b = 1176.8 c = 522.5 β = 100.7˚ a = 494.5 ± 0.1
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. Table 1 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm]
Comments/References
Mo2Al3 1570 to 1490
-
-
MoAl 1750 to 1470
cP2 Pm 3m CsCl
46 to 51.7 at.% Al [Mas2] Called “ζ2” (h) [1971Rex] a = 309.8 [1971Rex] a = 309.8 to 309.9 [1980Fer]
Mo3Al ≲ 2150 (Mo,Ni,Al)3 (Al,Mo,Ni)1
cP8 Pm3n Cr3Si
ε, NiAl3 < 856
oP16 Pnma NiAl3
oP16 Pnma Fe3C Ni2Al3 < 1138
β´, Ni3Al4 < 702
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hP5 P3m1 Ni2Al3
cI112 Ia3d Ni3Ga4
Called “ζ1” (h) [1971Rex]
a = 495 a = 487.6 a = 661.15 b = 736.64 c = 481.18 a = 661.3 ± 0.1 b = 736.7 ± 0.1 c = 481.1 ± 0.1 a = 659.8 b = 735.1 c = 480.2 a = 403.63 c = 490.65 a = 402.8 c = 489.1 a = 1140.8 ± 0.1
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22 to 27 at.% Al [Mas2] [1958Woo] at 6 at.% Ni, 75 at.% Mo [1969Vir] [L-B]
[1996Vik]
[1997Bou, V-C]
36.8 to 40.5 at.% Al [Mas2] [L-B] [1997Bou, V-C] [1989Ell, V-C]
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. Table 1 (continued) Phase/ Temperature Range [˚C] β, NiAl < 1651
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] cP2 Pm 3m CsCl
a = 287.04 a = 287.26 a = 286.0 a = 287.0 a = 288.72 ± 0.02 a = 287.98 ± 0.02 a = 289.0 a = 289.7 a = 290.4 a = 291.2 a = 291.9 a = 293.2
(Ni,Mo,Al)1 (Al,Mo,Ni)1
Ni5Al3 < 723
oC16 Cmmm Pt5Ga3
DOI: 10.1007/978-3-540-88053-0_9 ß Springer 2009
Comments/References 42 to 69.2 at.% Ni [Mas2] 57.7 at.% Ni [L-B] 46.6 at.% Ni [L-B] [1987Kha] 63 at.% Ni [1993Kha] 50 at.% Ni [1996Pau] 54 at.% Ni [1996Pau] [1971Cli]: T = 0˚C T = 200˚C T = 400˚C T = 600˚C T = 800˚C T = 1000˚C 0 to 1.5 at.% Mo, T = 1200˚C [1983Nas] 0 to 0.3 at.% Mo, T = 1200˚C [1988Mas] 0 to 0.2 at.% Mo, T = 1100˚C [1988Mas] 0 to 4.0 at.% Mo, T = 1093˚C [1984Mir] 63 to 68 at.% Ni [1993Kha, Mas2] at 63 at.% Ni [1993Kha]
a = 753 b = 661 c = 376
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. Table 1 (continued) Phase/ Temperature Range [˚C] γ´, Ni3Al < 1372
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] cP4 P m3m Cu3Au
Ni3(Al1–xMox)
a = 356.6 a = 357.0 a = 356.77 a = 356.32 a = 357.92 a = 356.7
a = 357.0 a = 357.8 a = 356.8 a = 357.2
Ni2Al9
Landolt‐Bo¨rnstein New Series IV/11E1
mP22 P21/c Ni2Al9
a = 868.5 ± 0.6 b = 623.2 ± 0.4 c = 618.5 ± 0.4 β = 96.50 ± 0.05˚
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Comments/References 73 to 76 at.% Ni [Mas2] [1952Tay] [1984Och2, 1959Gua1] [1986Hua] disordered [1998Rav] ordered [1998Rav] at x = 0 [1963Arb] As scaled from diagram, linear da/dx, alloys quenched from 1000˚C [1984Och1, 1984Och2, 1985Mis]: at x = 0 at 4 at.% Mo, 75 at.% Ni at 1.5 at.% Mo, 75 at.% Ni [1963Arb] at 1.5 at.% Mo, 73.5 at.% Ni [1963Arb] 0 to 4 at.% Mo, T = 1260˚C [1984Mir] 0 to 4.6 at.% Mo, at 1200˚C [1988Mas] 0 to 4.8 at.% Mo, at 1171˚C [1984Mir] 0 to 4.9 at.% Mo, at 1100˚C [1988Mas] 0 to 5.7 at.% Mo, at 1038 - 1093˚C [1984Mir] 0 to 5 - 6 at.% Mo, at 1000˚C [1977Aig, 1983Och, 1983Nas, 1984Och1, 1984Och2, 1984Mir, 1985Nas, 1985Mis, 1988Mas, 1989Hon1, 1989Mas, 1993Kub] 0 to 5.9 at.% Mo, at 927˚C [1984Mir] Metastable; [1988Li, 1997Poh]
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. Table 1 (continued) Phase/ Temperature Range [˚C] NixAl1–x 0.60 < x < 0.68
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] tP4 P4/mmm AuCu
m**
Ni2Al
hP3 P 3m1 CdI2 aP126 P 1
a = 383.0 c = 320.5 a = 379.5 c = 325.6 a = 379.5 c = 325.6 a = 379.5 c = 325.6 a = 379.9 to 380.4 c = 322.6 to 323.3 a = 371.7 to 376.8 c = 335.3 to 339.9 a = 378.00 c = 328.00 a = 418 b = 271 c = 1448 β = 94.3˚ a = 407 c = 499 a ≈ 1252 b ≈ 802 c ≈ 1526 α ≈ 90˚ β ≈ 109.7˚ γ ≈ 90˚
Comments/References Martensite, metastable [1993Kha] 62.5 at.% Ni [1991Kim] 63.5 at.% Ni [1991Kim] 66.0 at.% Ni [1991Kim] 64 at.% Ni [1997Pot] 65 at.% Ni [1997Pot] [1998Sim] [1992Mur]
Metastable [1993Kha] [1994Mur]
D1 (Al-Ni)
decagonal
-
Metastable [1988Li]
D4 (Al-Ni)
decagonal
-
Metastable [1988Li]
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. Table 1 (continued) Phase/ Temperature Range [˚C] * τ1, Mo(NixAl1-x)3
Pearson Symbol/ Space Group/ Lattice Prototype Parameters [pm] tI8 I4/mmm TiAl3
a = 370.2 c = 836.1 a = 373.2 c = 843.0 a = 373 c = 1680 a = 376.1 ± 0.6 c = 841.2 ± 0.8 a = 373 c = 1680
superstructure? c = 2c0 * τ2, Mo11Ni14Al75
a = 1005.4 ± 0.4 b = 1528.8 ± 0.4 c = 851.9 ± 0.2
orthorhombic
Comments/References [1971Pry], at 3 to 8 at.% Ni 25 at.% Mo, 600˚C at 1.6 to 6.0 at.% Ni [2002Gru] Called “Mo2NiAl5” [1965Ram], from a threephase alloy Mo25Ni25Al50 at ≈4 to 12 at.% Ni, 25 at.% Mo, 900˚C [1969Vir] From a three-phase alloy Mo25Ni17Al58 [1969Vir] Called “N” [2002Gru] for Mo(Al2,75Ni0,25) [1969Rec] from aluminothermic synthesis Called “X” [2002Gru] Called “Mo5Ni18Al77” [1969Mar, 1971Pry]
. Table 2 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
LÐα+β
≈1600
e1 (max)
L α β
L + β Ð α + γ´
1340
U1
L α β γ´
L+δÐα+γ
L Ð α + γ + γ´
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≈1310
≈1300
U2
E
Phase
L α δ γ L α γ γ´
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Al 45 1 48.5
Ni
10 97.5 2.5
45 1.5 49
13 95.5 1.5 2
20.5 <4 <33 <25 8.58 <0.5 <2.5 <8 18 <2 <17 <20.5
Mo
66.5 >0.5 >65.5 >73
27.22 97.5 50.5 20
64.2 >2.0 >47 >72
16 96.0 10 5.5
66 >2 >73 >74
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Al–Mo–Ni
. Table 2 (continued) Composition (at.%) Reaction
T [˚C]
γ + α Ð γ´ + δ
≈1130
Type U3
Phase α δ γ γ´
Al
Mo
Ni
1 2 10 20
98.2 49 14 5
0.8 49 76 75
δ + γ Ð γ´+ γ´´
≈890
U4
δ γ γ´ γ´´
1 5.5 20.5 2.5
51 14 5.5 22.5
48 80.5 74 75
γ + γ´´ Ð γ´ + θ
≈730
U5
γ γ´ γ´´ θ
5.5 20 3 1
11.5 5 22 19
83 75 75 80
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. Fig. 1 Al-Mo-Ni. Partial quasibinary system Mo-NiAl
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. Fig. 2 Al-Mo-Ni. Reaction scheme of the partial Mo-NiAl-Ni system
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. Fig. 3 Al-Mo-Ni. Liquidus surface in the Ni-rich region
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Al–Mo–Ni
. Fig. 4 Al-Mo-Ni. Partial isothermal section at 1260˚C
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. Fig. 5 Al-Mo-Ni. Isothermal section at 1200˚C, calculated by [1974Kau]
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. Fig. 6 Al-Mo-Ni. Partial isothermal section at 1100˚C
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. Fig. 7 Al-Mo-Ni. Partial isothermal section at 1000˚C
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Al–Mo–Ni
. Fig. 8 Al-Mo-Ni. Partial isothermal section at 1038˚C in the Al-poor region and at 1050˚C in the Al-rich region
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. Fig. 9 Al-Mo-Ni. Partial isothermal section at 927˚C in the Al-poor region and at 950˚C in the Al-rich region
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. Fig. 10 Al-Mo-Ni. Partial isothermal section at 880˚C
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. Fig. 11 Al-Mo-Ni. Partial isothermal section at 700˚C
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. Fig. 12 Al-Mo-Ni. Partial isopleth at 14 at.% Al
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References [1925Pfa] [1933Roe] [1951Ham] [1952Tay] [1954Ada] [1958Woo]
[1959Gua1] [1959Gua2] [1960Bag] [1960Vig]
[1962For] [1963Arb]
[1964Lea] [1965Ram]
[1967Bel]
[1969Mar]
[1969Rec]
[1969Vir] [1970Cli] [1971Cli]
[1971Nal]
Pfautsch, H., “The System Aluminium-Molybdenum-Nickel” (in German), Z. Metallkd., 19(4), 125–127 (1925) (Experimental, Phase Diagram, 8) as quoted by [1993Kub] Roentgen, P., Koch, W., “Influence of Heavy Metals on Alloys of Aluminium” (in German), Z. Metallkd., 25, 182–185 (1933) (Experimental, 8) Ham, J.L., “An Introduction to Arc-Cast Molybdenum and its Alloys”, Trans. Amer. Soc. Mech. Eng. (ASME), 73, 723–732 (1951) (Experimental, 10) as quoted by [2005Sch] Taylor, A., Floyd, R.W., “The Constitution of Nickel-Rich Alloys of the Nickel-Chromium-Aluminium System”, J. Inst. Met., 81, 451–464 (1952–1953) (Experimental, Crys. Structure, Phase Diagram, 15) Adam, J., Rich, J.B., “The Crystal Structure of WAl5”, Acta Crystallogr., 7, 813–816 (1954) (Experimental, Crys. Structure, 14) Wood, E.A., Compton, V.B., Matthias, B.T., Corenzwit, E., “β-Wolfram Structure of Compounds Between Transition Elements and Aluminium, Gallium and Antimony”, Acta Crystallogr., 11, 604–606 (1958) (Experimental, Crys. Structure, 13) Guard, R.W., WestBrook, J.H., “Alloying Behavior of Ni3Al (γ´ phase)”, Trans. Met. Soc. AIME, 215, 807–814 (1959) (Phase Diagram, Experimental, 27) Guard, R.W., Smith, E.A., “Constitution of Nickel-Base Ternary Alloys”, J. Inst. Met., 88, 283–287 (1959–1960) (Phase Diagram, Experimental, 3) Bagaryatskiy, Y.A., Ivanovskaya, L.E., “Equilibrium Diagram for Ni-NiAl-Mo Alloys” (in Russian), Dokl. Akad. Nauk SSSR, 132, 339–342 (1960) (Experimental, Phase Diagram, 14) Vigdorovich, V.N., Glazov, V.M., Glagoleva, N.N., “Investigation of the Solubility of Cr, Mo and W in Al by the Microhardness Method” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvet. Met., 3(2), 143–146 (1960) (Experimental, Phase Diagram, 16) Forsyth, J.B., Gran, G., “The Structure of the Intermetallic Phase γ (Mo-Al)-Mo3Al8”, Acta Crystallogr., 15, 100–104 (1962) (Experimental, Crys. Structure, 13) Arbuzov, M.P., Zelenkov, I.A., “Structure of Ni3Al Alloys with Additions of a Third Element”, Phys. Met. Metallogr., 15(5), 71–73 (1963), translated from Fiz. Met. Metalloved., 15(5), 725–728 (1963) (Crys. Structure, Experimental, 6) Leake, J.A., “The Refinement of the Crystal Structure of the Intermetallic Phase Al4Mo”, Acta Crystallogr., 17, 918–924 (1964) (Experimental, Crys. Structure, 38) Raman, A., Schubert, K., “On the Constitution of Alloys Related to TiAl3, III. Investigations in Some T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56, 99–104 (1965) (Crys. Structure, Experimental, 14) Belyaeva, G.I., Ilyushchenko, N.G., Anfinogenov, A.I., “Thermodynamics of Solid Alloys of a Mo-Al System” (in Russian), Tr. Inst. Electrochim. Akad. Nauk SSSR, (10), 85–95 (1967) (Experimental, Thermodyn., 25) Markiv, V.Ya., Burnashova, V.V., Pryakhina, L.L., Myasnikova, K.P., “Phase Equilibria in the Mo-Ni-Al System”, Russ. Metall., (5), 117–119 (1969), translated from Izv. Akad. Nauk SSSR, Met., (5), 180 (1969) (Phase Diagram, Experimental, 14) Rechkin, V.N., Samsonova, T.I., “Production of Mo-Ni-Al Alloys by Aluminothermic Reaction”, Russ. Metall., (3), 61–63 (1969), translated from Izv. Akad. Nauk SSSR, Met., (3), 61–64 (1969) (Crys. Structure, Experimental, 7) Virkar, A.V., Raman, A., “Alloy Chemistry of τ (βU)-Related Phases”, Z. Metallkd., 60, 594–600 (1969) (Phase Diagram, Crys. Structure, Experimental, 25) Cline, H.E., Walter, J.L., “The Effect of Alloy Additions on the Rod-Plate Transition in the Eutectic NiAl-Cr”, Metall. Trans., 1, 2907–2917 (1970) (Phase Diagram, Experimental, 20) Cline, H.E., Walter, J.L., Koch, E.F., Osika, L.M., “The Variation of Interface Dislocation Networks with Lattice Mismatch in Eutectic Alloys”, Acta Metall., 19, 405–414 (1971) (Experimental, Crys. Structure, 14) Nalibaev, T.N., Fasman, A.B., Inayatov, N.S., “Structure of Multicomponent Foraminate Catalysts Based on Nickel”, Russian J. Phys. Chem., 45, 211–214 (1971), translated from Zh. Fiz. Khim., 45, 383–386 (1971) (Experimental, 8)
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9 [1971Pry]
[1971Rex] [1973Wal] [1974Kau] [1974Tho] [1976Hen] [1976Jac]
[1976Mon] [1976Nes]
[1976Spr] [1977Aig]
[1977Pea]
[1978Gul]
[1980Fer]
[1981Sch]
[1982Shi] [1983Nas] [1983Och] [1983Wak]
[1984Kov1]
[1984Kov2] [1984Mir]
Al–Mo–Ni Pryakhina, L.I., Myasnikova, K.P., Markiv, V.Ya., Burnasheva, V.V., “Investigation of the MolybdenumNickel-Aluminium Ternary System” in “Phase Diagrams of Metal Systems” (in Russian), Nauka, Moscow, 112–116 (1971) (Phase Diagram, Experimental, 4) Rexer, J., “Phase Equilibria in the System Al-Mo at Temperatures above 1400˚C” (in German), Z. Metallkd., 62, 844–848 (1971) (Experimental, Crys. Structure, Phase Diagram, 23) Walter, J.L., Cline, H.E., “Stability of the Directionally Solidified Eutectics NiAl-Cr and NiAl-Mo”, Metall. Trans., (4), 33–38 (1973) (Experimental, 10) Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams, Part II”, Metall. Trans., 5, 1623–1629 (1974) (Phase Diagram, Thermodyn., 20) Thompson, E.R., Lemkey, F.D., “Metallic Matrix Composites”, Kreider, K.G. (Ed.), Academic Press, New York, 101 (1974) as quoted by [1983Nas] Henry, M.F., “Precipitation of γ’ in γ-α (Ni-Al-Mo) Eutectics”, Scr. Metall., 10, 955–957 (1976) (Phase Diagram, Experimental, 3) Jackson, M.R., Walter, J.L., “Superalloy Eutectic Composites with the VIA Refractory Elements - Cr, Mo, W”, Superalloys-Metallurgy and Manufacture, AIME, N.Y., 341–350 (1976) (Review, Phase Diagram, 42) Mondolfo, L.F., “Aluminium Alloys: Structure and Properties”, Butterworths, London, 598–599 (1976) (Review, 5) as quoted by [1993Kub] Nesterovich, L.N., Kupchenko, G.V., Ivanov, N.P., Budnikov, V.T., “Structure and Properties of Some Directionally Crystallized Eutectics Based on Nickel”, Phys. Met. Metallogr., 42, 117–123 (1976), translated from Fiz. Metall. Metalloved., 42, 1034–1041 (1976) (Phase Diagram, Experimental, 11) Sprenger, H., Richter, H., Nickel, J.J., “Directional Solidification of Ni-Mo-Al Eutectic Alloys”, J. Mater. Sci., 11, 2075–2081 (1976) (Phase Diagram, Experimental, 17) Aigeltinger, E.R., Bates, S.R., Gould, R.W., Hren, J.J., Rhines, F.N., “Phase Equilibria in Rapidly Solidified Nickel-Rich Ni-Mo-Al Alloys”, Proc. Internat. Conf. Rapid Solidification Processing. Principles and Technologies, Reston, Virginia, Claitor’s Publishing Div., Baton Rouge, 291–305 (1977), (Publ. 1978) (Phase Diagram, Crys. Structure, Experimental, Review, 20) as quotet by [1993Kub] Pearson, D.D., Lemkey, F.D., “Solidification and Properties of γ/γ´-αMo Ductile/Ductile Eutectic Superalloy”, Proc. Conf. Solidification and Casting of Metals, Metals Soc. London, Sheffield, U.K., 526–532 (1977) (Publ. 1979) (Phase Diagram, Experimental, 18) as quoted by [1993Kub] Gulyaev, B.B., Grigorash, E. F., Efimova, M. N., “Solidification Range of Nickel Alloys”, Heat-Resistant Steels and Alloys, (11), 914–917 (1978) translated from Metallov. Term. Obrab. Met., 11, 34–37 (1978) (Experimental, Phase Diagram, 8) Ferro, R., Marazza, R., “Crystal Structure and Density Data, Molybdenum Alloys and Compounds other than Halides and Chalcogenides”, Atomic Energy Rev.: Spec. Iss. No.7, IAEA, Vienna (1980) (Crys. Structure, Review, 961) Schwam, D., Dirnfeld, S.F., “Influence of Solidification Parameters on Microstructure of γ/γ’-α(Mo) Eutectic Alloy”, Conf. Mater. Eng., Freund Publ. House Tel Aviv, 1981, 10–13 (1981) (Phase Diagram, Experimental, 7) as quoted by [1993Kub] Shilo, I., Franzen, H.F., “High Temperature Thermodynamic Study of the Molybdenum-Rich Regions of the Mo-Al System”, J. Electrochem. Soc., 129, 2613–2617 (1982) (Experimental, Thermodyn., 13) Nash, P., Fielding, S., West, D.R.F., “Phase Equilibria in Nickel-Rich Ni-Al-Mo and Ni-Al-W Alloys”, Met. Sci., 17(4), 192–194 (1983) (Phase Diagram, Experimental, #, 20) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull. P.M.E. (T.I.T.), 52, 1–16 (1983) (Phase Diagram, Review, Experimental, 4) Wakashima, K., Hoguchi, K., Suzuki, T., Umekawa, S., “Reinvestigation of Phase Equilibria in the System Ni-Al-Mo and its Implication to the Elevated Temperature Stability of γ/γ’ α-Mo Aligned Eutectics”, Acta Metall., (11), 1937–1944 (1983) (Phase Diagram, Experimental, 19) as quoted by [1993Kub] Kovacova, K., Kristin, J., “Morphological Properties of γ/γ’-αMo Eutectic Composite Material” (in Czech), Kovove Mater., 22(3), 347–356 (1984) (Phase Diagram, Experimental, 18) as quoted by [1993Kub] Kovacova, K., “Undirectional Solidification of Ni-Al-Mo Alloy”, J. Cryst. Growth, 66, 426–430 (1984) (Phase Diagram, Experimental, 9) Miracle, D.B., Lark, K.A., Srinivas, V., Lipsitt, H.A., “Nickel-Aluminium-Molybdenum Phase Equilibria”, Metall. Trans. A, 15A, 481–486 (1984) (Phase Diagram, Experimental, #, 12)
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[1985Nas]
[1986Hor]
[1986Hua] [1986Kau]
[1986Mas1]
[1986Mas2]
[1987Kha] [1987San]
[1987Sve]
[1988Li]
[1988Mas]
[1989Ell] [1989Gai]
[1989Hon1] [1989Hon2] [1989Mas] [1990Yam]
9
Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behavior of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32, 289–298 (1984) (Phase Diagram, Experimental, 90) Ochiai, S., Mishima, Y., Suzuki, T.S., “Lattice Parameter Data of Ni(γ), Ni3Al (γ’) and Ni3Ga (γ’) Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15–28 (1984) (Crys. Structure, Experimental, 66) Schwam, D., Dirnfeld, S.F., Nadiv, S., “Microstructural Instability of Ni-Mo-Al Unidirectionally Solidified Eutectics”, J. Mater. Sci. Lett., (3), 363–366 (1984) (Experimental, Crys. Structure, 6) Mishima, S., Ochiai, S., Suzuki, T.Y., “Lattice Parameters of Ni (γ), Ni3Al (γ’) and Ni3Ga (γ’) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33, 1161–1169 (1985) (Crys. Structure, Experimental, 64) Nash, P., “Ni-Base Intermetallics for High Temperature Alloy Design “High-Temperature Ordered Intermetallic Alloys”, Mater. Res. Soc. Symposia Proceedings, Kreh, C.C., Liu, C.T., Stoloff, N.S., (Eds.), MRS-Pennsylvania, Boston, Massachusetts (1984), 39, 423–427 (1985) (Phase Diagram, Review, 15) Horita, Z., Sano, T., Nemoto, M., “Identification of Fine Particles in Unidirectionally Solidified Ni-AlMo Eutectic Alloys by Means of EDX and SAD Analyses”, Acta Metall., 34(8), 1525–1531 (1986) (Experimental, Crys. Structure, *, 27) Huang, S.C., Briant, C.L., Chang, K.-M., Taub, A.I., Hall, E.L., “Carbon Effects in Rapidly Solidified Ni3Al”, J. Mater. Res., 1(1), 60–67 (1986) (Experimental, Mechan. Prop., 27) Kaufman, M.L., Pearson, D.D., Fraser, H.L., “The Use of Convergent beam Electron Diffraction to Determine Local Lattice Distortions in Nickel Base Superalloys”, Philos. Mag., A54, 79–92 (1986) (Crys. Structure, Experimental, 11) Maslenkov, S.B., Udovskii, A.L., Burova, N.N., Rodimkina, V.A., “Phase Diagram of the NickelAluminium-Molybdenum System at 1300–2000˚C”, Russ. Metall., (1), 203–209 (1986), translated from Izv. Akad. Nauk SSSR, Met., (1), 198–205 (1986) (Phase Diagram, Experimental, 9) Maslenkov, S.B., Rodimkina, V.A., “Phase Equilibrium of the System Ni-Al-Mo in the Composition Range Ni-NiAl-Mo”, Russ. Metall., (3), 215–220 (1986), translated from Izv. Akad. Nauk SSSR, Met., (3), 218–223 (1986) (Phase Diagram, Experimental, 7) Khadkikar, P.S., Vedula, K., “An Investigation of the Ni5Al3 Phase”, J. Mater. Res., 2(2), 163–167 (1987) (Experimental, Crys. Structure, 7) Sano, T., Nemoto, M., “Precipitates in Nickel-Rich Ni-Al-Mo Ternary Alloys”, Trans. Jpn. Inst. Met., 28, 8–19 (1987), translated from J. Jpn. Inst. Met., 49(8), 690–698 (1985) (Crys. Structure, Experimental, 52) Svetlov, I.L., Udovski, A.L., Monastyrskaya, E.V., Oldakovskii, I.V., Nazarova, M.P., “Calculation of the Monovariant Liq/(Liq + γ + α) Line in the Ni-Mo-Al System and Plane Front Solidification in γ/γ´ - α Alloys”, Russ. Metall., (6), 186–192 (1987), translated from Izv. Akad. Nauk SSSR, Met., (6), 183–189 (1987) (Phase Diagram, 14) Li, X.Z., Kuo, K.H., “Decagonal Quasicrystals with Different Peridicities Along the Tenfold Axis in Rapidly Solidified Al-Ni Alloys”, Phil. Mag. Lett., 58(3), 167–171 (1988) (Experimental, Crys. Structure, 14) Maslenkov, S.B., Burova, N.N., Rodimkina, V.A., “The Ni-NiAl-Mo State Diagram in the 1200–700˚C Temperature Range” (in Russian), Izv. Akad. Nauk SSSR, Met., (6), 183–190 (1988) (Phase Diagram, Experimental, #, 13) Ellner, M., Braun, J., Predel, B., “X-Ray Diffraction Investigation of Al-Cr Phases of the W Family” (in German), Z. Metallkd., 80, 374–383 (1989) (Experimental, Crys. Structure, 38) Gaidukov, A.M., Udovskii, A.L., Oldakovski, I.V., “Construction of Mathematical Models in the Liquid System Surface” (in Russian), Dokl. Akad. Nauk SSSR, 305, 643–648 (1989) (Review, Theory, 15) Hong, Y.M., Nakajima, H., Mishima, Y., Suzuki, T., “The γ Solvus Surface in Ni-Al-X (X = Cr, Mo and W) Ternary Systems”, ISIJ International, 29(1), 78–84 (1989) (Phase Diagram, Experimental, #, 25) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of γ’ Solvus in Ni-Al-X Ternary Systems”, Mater. Res. Soc. Symp. Proc., 733, 431–438 (1989) (Phase Diagram, Experimental, 35) Maslenkov, S.B., Rodimkina, V.A., “Phase Changes in Alloys of the System Ni-NiAl-Mo” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 194–198 (1989) (Phase Diagram, Crys. Structure, #, 12) Yamamoto, M., Iada, J., Nenno, S., “The Microstructure of a Two-Phase Mixture in a Ni75Mo20Al5 Alloy”, J. Mater. Sci. Lett., 9, 34–35 (1990) (Crys. Structure, Experimental, 5)
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[1991Kim]
[1991Mis] [1991Sas]
[1991Sch] [1992Mur]
[1993Kha] [1993Kub]
[1994Mur] [1995Gri] [1996Pau] [1996Vik] [1997Bou] [1997Jin] [1997Poh] [1997Pot]
[1998Mur] [1998Rav]
[1998Sim]
[1999Kau]
[1999Lu] [2000Bor]
Al–Mo–Ni Enomoto, M., Harada, H., Yamazaki, M., “Calculation of γ´/γ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143–158 (1991) (Assessment, Calculation, Phase Diagram, 34) Kim, Y.D., Wayman, C.M., “Transformation and Deformation Behavior of Thermoelastic Martensite Ni-Al Alloys Produced by Powder Metallurgy Method” (in Korean), J. Korean Inst. Met. Mater., 29(9), 960–966 (1991) (Mechan. Prop., Experimental, 15) as quoted by [2004Sal] Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the γ Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123–130 (1991) (Assessment, Phase Diagram, Experimental, 5) Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of NiAl”, Proc. Conf. Intermetal. Comp. - Struct. Mechan. Prop., 877–881 (1991) (Abstract, Phase Diagram, Experimental, Mechan. Prop., 10) Schuster, J.C., Ipser, H., “The Al-Al8Mo3 Section of the Binary System Aluminium- Molydenum”, Met. Trans., A22, 1729–1736 (1991) (Experimental, Crys. Structure, Phase Diagram, 20) Murakami, Y., Otsuka, K., Hanada, S., Watanabe, S., “Crystallography of Stress-Induced B2Ð7R Martensitic Transformation in a Ni-37.0 at.% Al Alloy”, Mater. Trans. JIM, 33(3), 282–288 (1992) (Crys. Structure, Experimental, 25) Khadkikar, P.S., Locci, I.E., Vedula, K., Michal, G.M., “Transformation to Ni5Al3 in a 63.0 At. Pct Ni-Al Alloy”, Metall. Trans. A, 24A, 83–94 (1993) (Phase Diagram, Crys. Structure, Experimental, 28) Kubaschewski, O., “Al-Mo-Ni (Aluminium - Molybdenum - Nickel),” MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12789.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, 54) Murthy, A.S., Goo, E., “Triclinic Ni2Al Phase in 63.1 at.% NiAl”, Met. Mater., A, 25A(1), 57–61 (1994) (Crys. Structure, Experimental, 10) Grin, Y.N., Ellner, M., Peters, K., Schuster, J.C., “The Crystal Structures of Mo4Al17 and Mo5Al22”, Z. Kristallogr., 210, 96–99 (1995) (Experimental, Crys. Structure, 11) Paufler, P., Faber, J., Zahn, G., “X-Ray Single Crystal Diffraction Investigation on Ni1+xAl1-x”, Acta Crystallogr., Sect. A: Found. Crystallogr., A52, C319 (1996) (Crys. Structure, Experimental, Abstract, 3) Viklund P., Haeussermann, U., Lidin, S., “NiAl3: a Structure Type of its Own?”, Acta Crystallogr., Sect. A: Found Crystallogr., A52, C321 (1996) (Crys. Structure, Experimental, Abstract) Bouche, K., Barbier, F., Coulet, A., “Phase Formation During Dissolution of Nickel in Liquid Aluminium”, Z. Metallkd., 88(6), 446–451 (1997) (Thermodyn., Experimental, 15) Jin, Y., Chaturvedi, M.C., Han, Y.F., Zhang, Y.G., “Crystal Structure of δ-NiMo Phase in a Ternary Ni-Mo-Al Alloy”, Mater. Sci. Eng. A, A225, 78–84 (1997) (Crys. Structure, Experimental, 13) Pohla, C., Ryder, P.L., “Crystalline and Quasicrystalline Phases in Rapidly Solidified Al-Ni Alloys”, Acta Mater., 45, 2155–2166 (1997) (Experimental, Crys. Structure, 48) Potapov, P.L., Song, S.Y., Udovenko, V.A., Prokoshkin, S.D., “X-Ray Study of Phase Transformations in Martensitic Ni-Al Alloys”, Metall. Mater. Trans. A, 28A, 1133–1142 (1997) (Crys. Structure, Experimental, 40) Murthy, B.S., Ranganathan, S., Int. Mater, Rev., 43(3), 101–141 (1998), as quoted by [2002Alb] Ravelo, R., Aguilar, J., Baskes, M., Angelo, J.E., Fultz, B., Holian, B.L., “Free Energy and Vibrational Entropy Difference between Ordered and Disordered Ni3Al”, Phys. Rev. B, 57(2), 862–869 (1998) (Thermodyn., Theory, Calculation, 43) Simonyan, A.V., Ponomarev, V.I., Khomenko, N.Yu., Vishnyakova, G.A., Gorshkov, V.A., Yukhvid, V.I., “Combustion Synthesis of Nickel Aluminides”, Inorg. Mater., 34(6), 558–561 (1998), translated from Neorgan. Mater., 34(6), 684–687 (1998) (Crys. Structure, Experimental, 12) Kaufman, L., Dinsdale, A.T., “Summary of the Proceedings of the CALPHAD XXVII Meeting, 17–22 May 1998, Beijing, China”, Calphad, 23(3-4), 265–303 (1999) (Assessment, Calculation, Phase Diagram, Thermodyn., #) Lu, X., Cui, Y., Jin, Z., “Experimental and Thermodynamic Investigation of the Ni-Al-Mo System”, Metall. Mater. Trans. A, 30A, 1785–1795 (1999) (Phase Diagram, Experimental, Thermodyn., #, 28) Borgenstam, A., Engstroem, A., Hoeglund, L., Agren, J., “DICTRA, a Tool for Simulation of Diffusional Transformations in Alloys”, J. Phase Equilib., 21(3), 269–280 (2000) (Calculation, Kinetics, Thermodyn.)
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Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8), 2314–2320 (2001) (Calculation, Crys. Structure, Experimental, Thermal Conduct., 63) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of γ’/β Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168–175 (2001) (Phase Diagram, Experimental, Thermodyn., 21) Albiter, A., Bedolla, E., Perez, R., “Microstructure Characterization of the NiAl Intermetallic Compound with Fe, Ga and Mo Additions Obtained by Mechanical Alloying”, Mater. Sci. Eng. A, 328, 80–86 (2002) (Crys. Structure, Experimental, 14) Grushko, B., Mi, S., Highfield, J.G., “A Study of the Al-Rich Region of the Al-Ni-Mo Alloy System”, J. Alloys Compd., 334, 187–191 (2002) (Crys. Structure, Phase Diagram, Experimental, 8) Guo, J.T., Du, X.H., Zhou, L.Z., Zhou, B.D., Qi, Y.H., Li, G.S., “Superplasticity in NiAl and NiAl-Based Alloys”, J. Mater. Res., 17(9), 2346–2356 (2002) (Experimental, Mechan. Prop., 17) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 20.10238.1.20, (2004) (Crys. Structure, Phase Diagram, Assessment, 164) Schuster, J.C., “Al-Mo (Aluminium - Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 30.12123.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 61) Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P., Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971); Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, Key Elements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, Subvol. a: Ac-Au … Au-Zr (1991); Springer-Verlag, Berlin. Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Molybdenum – Uranium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Zoya Alekseeva, updated by Kostyantyn Korniyenko
Introduction Phase relationships in the Al-Mo-U system are of great interest, above all because an alloy comprising Mo-U alloy particles dispersed in an aluminium matrix provides a candidate material for low-enriched uranium fuel with high loading. This is because a solid solution of molybdenum in (γU) possesses acceptable irradiation and mechanical properties and can be formed over a wide range of Mo concentrations. Below 560˚C, (γU) can exist only in a metastable state, as indicated by the constitution of the Mo-U binary system. Therefore, in fuel fabrication and service it is very important to retain the metastable (γU) below 560˚C [2002Lee]. Phase equilibria in the composition range U-UAl2-Mo3Al8-Mo have been studied by [1969Pet, 1971Pet] by means of physico-chemical analysis techniques. The liquidus surface, a projection of the four-phase equilibria, the UAl2-Mo3Al8 temperature-composition section and the reaction scheme involving the liquid phase were constructed; the phase compositions for ten four-phase equilibria were determined [1969Pet]. [1971Pet] presented isothermal sections at 1250, 1050 and 500˚C. Several temperature-composition sections were also reported. More recently, the attention of investigators of phase relationships in the ternary Al-Mo-U system has been concentrated mainly on the behavior of alloys that may serve as the basis of Mo-U alloy dispersion fuels [1979Sof, 1984Gom, 2002Lee, 2003Lee, 2003Ryu, 2004Lee, 2006Ryu2, 2006Wie]. Phase development in the Al rich Al-Mo-U alloys is presented in [2004Kei]. The crystal structures of phases of the Al-Mo-U system are reported in [1964Nic, 1979Sof, 1984Gom, 1994Nie, 1995Nie, 2002Lee, 2003Mir, 2003Ryu, 2004Kei, 2004Lee, 2006Ryu2, 2007Noe]. Thermodynamic properties are reported in [2003Ryu, 2006Kim, 2006Ryu1]. Publications relating to experimental studies of phase relationships, crystal structures and thermodynamics along with the techniques employed are listed in Table 1. A review of the crystal structures of the Al-Mo-U phases is presented in [1980Fer]. The character of the phase equilibria was assessed in the previous MSIT evaluation by [1993Ale]. The present report is supplemented by information from later publications, bringing together the data in accordance with the modern versions of the boundary binary phase diagrams.
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Binary Systems The Al-Mo boundary binary system is accepted from [2005Sch]. The boundary binary Al-U system is based on a critical assessment carried out by [1990Kas], reproduced in [Mas2], however, taking UAl4 as a fully ordered, strictly stoichiometric phase [2004Tou]. The Mo-U boundary binary system is accepted from [Mas2].
Solid Phases Crystallographic data on the solid unary, binary and ternary Al-Mo-U phases and their concentration and temperature ranges of stability are presented in Table 2. A noticeable solubility of Al and Mo in (U) and Al and U in (Mo) was reported by [1969Pet, 1971Pet]. There is negligible uranium solubility in the binary Al-Mo compounds. According to [1971Pet], aluminium can be substituted by molybdenum (up to 18.5 at.% of Mo at the temperatures of 1250, 1050 and 500˚C) in the UAl2 based λ2 phase. The ternary phases τ1, U9Mo16Al75 (with unknown structure - composition lies outside the studied U-U Al2-Mo3Al8-Mo range) and the τ2 Laves phase (with the MgZn2 type structure) were identified by [1969Pet, 1971Pet]. [1994Nie] and [1995Nie] found two new ternary phases (τ3, U6Mo4Al43 and τ4, UMo2Al20) from single crystal studies. Their existence was confirmed by [2007Noe], but phase equilibria involving these phases are in need of further study.
Invariant Equilibria Temperatures, reactions types and phase composition relating to invariant equilibria in the U-U UAl2-Mo3Al8-Mo region are listed in Table 3. The reaction scheme is shown in Figs. 1a and 1b. The presented data are based on [1969Pet, 1971Pet], with amendments to ensure compatibility with the accepted boundary binary systems. In Figs. 1a, 1b and Table 3, only invariant equilibria involving the liquid phase are presented. In addition to the presented reactions (Figs. 1a, 1b, Table 3), an invariant equilibrium between liquid, τ1, λ2 and UAl3 takes place at 1305˚C in the Al rich corner. This is deduced from the UAl2-Mo3Al8 temperaturecomposition section (see “Temperature-Composition Sections” chapter). A reaction scheme for the solid state was proposed in [1971Pet] but is omitted here as it is in need of revision in light of principal changes in the invariant temperatures of the Mo-U binary system. Further experimental verification of the solid state phase equilibria in the ternary Al-Mo-U system is also necessary.
Liquidus, Solidus and Solvus Surfaces Figure 2 presents the liquidus surface projection of the U-U Al2-Mo3Al8-Mo region on the basis of studies by [1969Pet], with slight corrections in accordance with the boundary binary systems, in particular in relation to the positions of the reactions p1 and p5. The isotherms positions on the (Mo) surface at the temperatures above 1600˚C contradict to the accepted Al-Mo and Mo-U binary diagrams and need further verification. They are omitted in Fig. 2. [1969Pet] also reported the compositions of solid phases taking part in the invariant DOI: 10.1007/978-3-540-88053-0_10 ß Springer 2009
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equilibria, and these compositions are listed in Table 3. But the solidus surface projection is not presented here because the compositions of the Al-Mo-based phases were not determined precisely making a correct representation difficult.
Isothermal Sections Figures 3, 4 and 5 show isothermal sections determined for 1250, 1050 and 500˚C, respectively, by [1971Pet]. The composition range studied was that within the region delineated by U, Mo, Mo3Al8 and UAl2. The sections are amended to ensure consistency with the boundary binary systems. A neutron diffraction study of an atomized uranium-based alloy with 10 mass% Mo (U78.4Mo21.6 in at.%) carried out by [2002Lee] showed that homogenizing annealing above 560˚C retards the decomposition of the γ phase at 400 and 500˚C, at which temperatures the phase is metastable. It was reported that in fuel with U78.4Mo21.6 alloy particles dispersed in an Al matrix, aluminium diffuses along the grain boundaries and reacts with the α phase formed from the decomposed γ phase at 400 and 500˚C. The formation of UAl3 and λ2 was observed. The formation of ternary compounds was not detected by neutron diffraction. The phase content of the U-Mo/Al dispersion fuels at temperatures of 500, 525 and 550˚C was determined by [2003Ryu, 2006Ryu2] using a diffusion couple technique. Diffusion couples based on atomized dispersion fuel with U-7 mass% Mo (U-15.7 at.% Mo) /Al were studied by [2003Mir, 2004Lee].
Temperature – Composition Sections A number of temperature-composition sections have been reported for the U-UAl2-Mo3Al8Mo region [1969Pet, 1971Pet]. Figure 6 shows the boundary (UAl2-Mo3Al8) section of the investigated part of the ternary system. It is not quasibinary (although declared by [1969Pet, 1971Pet]) because the τ1 phase, which is involved in equilibria within the section, lies outside the plane of the section. The remaining temperature-composition sections reported in [1971Pet] are presented in Figs. 7 to 13, amended to ensure agreement with the boundary binary systems. The character of phase equilibria at the Al-Mo side of the isopleth at 50 at.% Al proposed in [1971Pet] contradicts to the accepted Al-Mo diagram (the ξ1 phase does not take part in equilibria at 50 at.% Al) and therefore the corresponding part of the vertical section is plotted in Fig. 11 by dashed lines.
Thermodynamics [2003Ryu] presented results of studies on the kinetics of the growth of reaction layers. An as-cast U-10Mo (mass%) (U-21.6Mo (at.%)) ingot was heat-treated under vacuum at 900˚C for 100 h to ensure compositional homogeneity, followed by quenching to form the γ phase. This was used to make a diffusion couple with an Al sheet. Annealing of 10 vol% (U-21.6 at.% Mo)/Al dispersion fuels was carried out at temperatures from 500 to 550˚C for times between 10 min to 36 h with a view to determine the activation energy for the growth of the reaction Landolt‐Bo¨rnstein New Series IV/11E1
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layers. It was calculated using an Arrhenius plot according to two different models (277 kJ·mol–1 from the Jander’s model and 316 kJ·mol–1 from the Ginstling-Brounshtein model). The heat generation of U-Mo/Al dispersion fuels with 10–50 vol% of Mo-U fuel during a thermal cycle from room temperature to 700˚C was measured using DSC. The exothermic heat resulting from the reaction between Mo-U and the Al matrix was largest (235 kJ·kg–1) when the volume fraction of Mo-U fuel was about 30 vol%. Irradiation-enhanced interdiffusion in the diffusion zone of a U-Mo dispersion fuel in aluminium was studied in [2006Kim]. In order to analyze fuel performance accurately, a model to predict the diffusion kinetics was proposed. The authors developed a diffusion layer growth rate correlation for out-of-pile annealing tests and a similar correlation for in-reactor tests. The correlation for in-reactor tests is considerably different from that of out-of-pile tests because it contains factors that amplify diffusion kinetics by fission damage in the diffusion reaction zone. Using an appropriate computer code, fission damage factors were obtained as a function of diffusion reaction layer thickness and composition. The model correlation was established and fitted to the in-reactor data. As a result of the data fitting, the interaction layer growth rate was found to be proportional to the square root of the fission fragment damage rate and to have a temperature dependence characterized by the effective activation energy of 46 to 76 kJ·mol–1. The heats of formation of the (U1–xMox)Al3 intermetallic compound were obtained by [2006Ryu1] by measuring the reaction heats of U-Mo/Al dispersion samples using DSC. It was concluded that the heat of formation of (U1–xMox)Al3 becomes less negative as the molybdenum content increases (about –68 kJ·mol–1 and about –55 kJ·mol–1 at x = 0.15 and x = 0.225, respectively).
Notes on Materials Properties and Applications The Al-Mo-U alloys are of great practical interest as the basis of high-density U-Mo/Al dispersion fuel. Nuclear fuels can be placed into two main categories - for energy production and for production of neutrons. The alloys of the Al-Mo-U system are candidate materials for both of these fields of application [2007Noe]. Vickers microhardness of the γ phase in an U87.5Mo12.0Al0.5 (at.%) alloy that had been homogenized at 1050˚C and then reduced and quenched from 800˚C was reported in [1984Gom] as 1.78 GPa. The effects of fuel powder volume fraction and fuel particle shape on the green properties of pressed powder compacts produced from blended U-10Mo (mass%) (U-21.6Mo (at.%)) and Al powder were investigated by [2000Han]. The relative density of the compacts increased with decreasing volume fraction of fuel powder. The compressibility of compacts prepared from comminuted powder was larger than that of the compacts made from atomized powder owing to fragmentation of the comminuted particles. The green strength of the compacts made from comminuted powder was higher than that of the atomized powder compact. In the opinion of the authors, this seemed to have resulted from the smaller pore size and the larger contact area between the comminuted fuel powder and Al powder. It was suggested that an adjustment in the compacting conditions may be required in order to fabricate compacts from atomized powder having a comparable green strength. The effect of heat treatment on the thermal conductivity of U-Mo/Al alloy dispersion fuel was reported in [2003Lee]. Thermal conductivities were calculated from measured thermal diffusivities, specific heat capacities and densities, which were determined using DOI: 10.1007/978-3-540-88053-0_10 ß Springer 2009
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the laser flash, DSC, and Archimedes methods, respectively. The thermophysical properties were measured over a temperature range from room temperature up to 500˚C. The Mo-U alloy was annealed at between 525 and 550˚C for between 1 and 36 h. At high temperatures, the MoU particles reacted with the aluminium matrix forming reaction layers that adversely affected the thermal conductivity of the fuel core. The thermal conductivities of annealed samples appeared to decrease with increasing volume fraction of the reaction layers.
Miscellaneous The kinetics of the isothermal transformation of the γ phase in an alloy of the composition U87.5Mo12.0Al0.5 (in at.%) was investigated by [1984Gom] using dilatometry with a cooling rate from the γ phase range of 150˚C·min–1 between 700 and 500˚C. It was concluded that in comparison with the binary Mo-U system, the addition of a comparatively minor amount of the third component aluminium has a major effect on the shape of the isothermal transformation diagram, the kinetic parameters and the completeness of transformation. Precipitation of the second intermetallic λ2 phase takes place. The thermal compatibility of centrifugally atomized Mo-U alloys with aluminium was studied by [1997Lee]. The results of their investigations show that U-2 mass% Mo/Al dispersions increase in volume by 26% at 400˚C after 2000 h. This large volume change is due mainly to the formation of voids and cracks resulting from nearly complete interdiffusion of Mo-U and aluminium. No significant dimensional changes occur in the U-10 mass% Mo/Al dispersions. Interdiffusion between U-10Mo (mass%) and aluminium was found to be minimal. The different diffusion behavior is due primarily to the fact that U-10 mass% Mo (U-21.6Mo (at.%)) particles are much more supersaturated with substitutional molybdenum than U-2Mo (mass%) (U-4.8Mo (at.%)) particles. The aluminium diffuses into the U-2Mo (mass%) particles relatively rapidly along grain boundaries forming UAl3 almost fully throughout the 2000 h anneal, whereas the molybdenum supersaturation in the U-10Mo (mass%) particles inhibits the diffusion of aluminium atoms. U-10 mass% Mo displays superior thermal compatibility with aluminium compared to U-2Mo (mass%). The irradiation behavior of atomized U78.4Mo21.6 (at.%) Mo alloy aluminium matrix dispersion fuel meat at low temperature was reported by [2002Kim]. In order to examine the in-reactor behavior of very-high-density dispersion fuels for high flux performance research reactors, the U78.4Mo21.6 alloy dispersions in an aluminium matrix have been irradiated at low temperatures in the Advanced Test Reactor (ATR). The alloy fuel dispersant was produced by a centrifugal atomization process. The fuel particles had a fine and a relatively narrow fission gas bubble size distribution. In the author’s opinion, these appeared to be features in the microstructure that result from the segregation of the microstructure into molybdenum rich and depleted regions on solidification. Irradiation tests have been conducted by [2002Mey] with a view to evaluate the performance of a series of highdensity uranium-molybdenum (U-Mo) alloy aluminium matrix dispersion fuels. Fuel plates incorporating alloys with molybdenum contents in the range of 4–10 mass% were tested. As a whole, fuels with molybdenum contents of 6 mass% or higher showed stable in-reactor fission gas behavior, exhibiting a distribution of small, stable gas bubbles. The Bozzolo-FerranteSmith (BFS) method for alloys was applied by [2003Gar] to the analysis of aluminium interdiffusion in the Mo-U solid solution as a function of Mo concentration. The binary Al/U and Al/Mo system showed opposite behavior, which in the ternary case of Al/Mo-U translates into the role of regions rich in Mo acting as interdiffusion barriers to Al, in excellent Landolt‐Bo¨rnstein New Series IV/11E1
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agreement with experimental evidence. Later, an interdiffusional model for the prediction of interaction layer growth in the Al/Mo-U system using the computer codes PLACA and DPLACA simulating the behavior of a plate-type fuel containing in its core a foil of monolithic or dispersed fissile material, respectively, under normal operation conditions of a research reactor, was proposed by [2007Sob]. The PLACA code was used to simulate experimental results from planar Al/Mo-U diffusion couple studies published in the literature. DPLACA was used to simulate experiments performed with Mo-U particles dispersed in Al, with and without irradiation. Satisfactory prediction of the whole reaction layer thickness and of the individual fractions corresponding to alloy and matrix consumption was obtained. In [2006Wie], the similarity or even equivalence of the interdiffusion layer of U-Mo/Al fuel, that was found in-pile and alternatively generated by heavy ion irradiation, was shown. A similar failure under irradiation was reported.
. Table 1 Investigations of the Al-Mo-U Phase Relations, Structures and Thermodynamics Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1964Nic]
Arc melting; heat treatment; 900˚C with cooling to room temperature metallography; X-ray diffraction; electron at 4–5˚C·min–1; U rich corner, 0.75 to 2.8 microscopy at.% Mo, 2 to 8 at.% Al
[1969Pet]
Arc melting (the starting components Al - < 2623˚C; the range U-UAl2-Mo3Al8-Mo 99.99 mass%, Mo - 99.93, U - 99.9 mass%); DTA; chemical analysis; X-ray diffraction; optical microscopy
[1971Pet]
Arc melting (the starting components Al - 1250, 1050, 500˚C; U-UAl2-Mo3Al8-Mo 99.99 mass%, Mo - 99.93, U - 99.9 mass%); annealing; DTA; X-ray diffraction; optical microscopy
[1979Sof]
Arc melting; annealing; quenching; 1050–950˚C; U rich corner, 10 and 11.5 chemical analysis; EMPA; X-ray diffraction at.% Mo, up to 3 at.% Al
[1984Gom] Arc melting; annealing; reduction; quenching; X-ray diffraction; dilatometry
1050–400˚C; U87.5Mo12.0Al0.5 (at.%)
[1994Nie]
Cold pressing (the purities of starting U6Mo4Al43 components > 99.9 mass%); annealing at 800˚C for 3 weeks, single crystal growth, SEM; X-ray diffraction (Guinier camera)
[1995Nie]
Cold pressing (the purities of starting UMo2Al20 components > 99.9 mass%); annealing at 800˚C for 3 weeks, single crystal growth, SEM; X-ray diffraction (Guinier camera)
[2002Lee]
Blending, extrusion, annealing, atomization; neutron diffraction; SEM
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800–500˚C; U78.4Mo21.6 (at.%) with Al additions
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[2003Mir]
Arc melting; annealing and quenching; optical microscopy; SEM; EMPA; X-ray diffraction
580˚C; diffusion couples U-7 mass% Mo (U-15.7 at.% Mo) /Al
[2003Ryu]
Vacuum induction melting; annealing; SEM; X-ray diffraction; DSC
≤ 1400˚C; dispersion fuels U-10 mass% Mo (U-21.6 at.% Mo)/Al
[2004Kei]
Arc melting; SEM
As-cast and annealed at 500˚C alloys U0.8Mo0.2Al6 and U0.8Mo0.2Al7
[2004Lee]
Irradiation cycles (FUTURE irradiation rig of the BR2 reactor; optical microscopy; EMPA; wavelength dispersive X-ray analysis (WDX); SEM; X-ray diffraction
Low-enriched uranium fuel plates U7 mass% Mo (U-15.7 at.% Mo) atomized powder dispersed in Al matrix
[2006Kim]
SEM; EMPA
≤ 600˚C; dispersion fuels U-10 mass% Mo (U-21.6 at.% Mo)/Al
[2006Ryu1] DSC; SEM
≤ 677˚C; U-6 mass% Mo (U-13.7 at.% Mo)/ Al and U-10 mass% Mo (U-21.6 at.% Mo)/ Al dispersion samples
[2006Ryu2] Irradiation tests (HANARO reactor); X-ray diffraction; SEM; EDX
500–580˚C; U-6 to 10 mass% Mo (U-13.7 to 21.6 at.% Mo)/ (0.1 to 1.0 mass% Al) dispersion fuels
[2006Wie]
Heavy ion irradiation tests; optical microscopy; SEM; EDX; X-ray diffraction
U-6 mass% Mo (U-13.7 at.% Mo)/Al and U-10 mass% Mo (U-21.6 at.% Mo)/Al dispersion samples
[2007Noe]
EDX; X-ray diffraction
Phase equilibria with participation of the τ3 and τ4 phases
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] (Al) < 660.452 UxMoyAl1–x–y
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Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm] cF4 Im3m Cu
a = 404.96
Comments/References T = 25˚C [V-C2] x = 0, 0 < y . 10–4, T = 660˚C [2005Sch] y = 0, 0 < x ≤ 7·10–5, T = 641˚C [1990Kas]
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. Table 2 (continued) Phase/ Temperature Range [˚C] (Mo) < 2623
Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm] cI2 Im 3m W
a = 314.70
cI2 Im 3m W
a = 352.4
γ, U1–x–yMoxAly
a = 344.0
a = 343.9 β, (βU) (h1) 776 - 668 β, U1–x–yMoxAly
α, (αU) (r) < 668 α, U1–x–yMoxAly
λ2, U(MoxAl1–x)2
tP30 P42/mnm βU
a = 1075.9 c = 565.6
oC4 Cmcm αU
a = 285.37 b = 586.95 c = 495.48
cF24 Fd 3m Cu2Mg
UAl2 < 1620
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T = 25˚C [V-C2]
x = 0, 0 < y ≤ 0.195, T = 2150˚C [2005Sch] y = 0, 0 < x ≤ 0.04, T ≈ 1800˚C [Mas2]
UxMo1–x–yAly
γ, (γU) (h2) 1135 - 776
Comments/References
T > 776˚C [V-C2]
x = 0, 0 < y ≤ 0.047,T = 1105˚C [1990Kas] y = 0, 0 < x ≤ 0.42, T = 1284˚C [Mas2] in the U95.2Mo2.8Al2.0 and U93.9Mo2.1Al4.0 (at.%) alloys annealed at 900˚C 48 h and cooled to room temperature at 4–5˚C·min–1, together with α and λ2 phases [1964Nic] in the U87.5Mo12.0Al0.5 (at.%) alloy homogenized at 1050˚C, reduced and quenched from 800˚C [1984Gom] T > 668˚C [V-C2]
x = 0, 0 < y ≤ 0.0054, T = 758˚C [1990Kas] y = 0, 0 < x ≤ 0.02, T = 668˚C [Mas2] T = 25˚C [V-C2]
x = 0, 0 < y < 7·10–4, T = 665˚C [1990Kas] y = 0, 0 < x ≤ 0.02, T = 570˚C [Mas2] 0 ≤ x ≤ 0.27 [1971Pet]
a = 778
[V-C2]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm]
Comments/References
UAl3 < 1350
cP4 Pm3m AuCu3
a = 426.2
UAl4 < 731
oI20 Imma UAl4
a = 440.14 [2004Tou] b = 625.52 c = 1372.79
Mo3Al ≲ 2150
cP8 Pm3m Cr3Si
ξ2, MoAl 1750 - 1470
[V-C2]
a = 495
23 to 28.5 at.% Al at T ≈ 1720˚C [2005Sch] [1958Woo, 2005Sch]
a = 309.8
46 at.% Al at T ≈ 1720˚C to 52 at.% Al at T = 1570˚C [2005Sch] [1971Rex, 2005Sch]
cP2 Pm3m CsCl
ξ1, Mo2Al3 1570 - 1490
-
-
63 at.% Al [2005Sch]
Mo3Al8 < 1555
mC22 C2/m Mo3Al8
a = 916.4 b = 363.9 c = 1004 β = 100.5˚
[V-C2]
MoAl3 1222 - 818
mC32 C2/m MoAl3
a = 1639.6 [2005Sch] b = 359.4 c = 838.6 β = 101.88˚
Mo1–xAl3+x 1260 - 1154
cP8 Pm3m Cr3Si
76 at.% Al at T = 1222˚C to 78 at.% Al at T = 1154˚C [2005Sch] a = 494.5
[2005Sch]
MoAl4 1177 - 942
mC30 Cm WAl4
Mo4Al17 < 1034
mC84 C2 Mo4Al17
a = 915.8 b = 493.23 c = 2893.5 β = 96.71˚
[1995Gri, 2005Sch]
Mo5Al22 964 - 831
oF216 Fdd2 Mo5Al22
a = 7382 b = 916.1 c = 493.2
[1995Gri, 2005Sch]
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79 to 80 at.% Al at T = 1154˚C [2005Sch] a = 525.5 [V-C2] b = 1776.8 c = 522.5 β = 100.88˚
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm]
MoAl5 (h2) 750 - 800 < T < 846
hP12 P6322 WAl5
a = 491.2 c = 886.0 a = 493.7 c = 924.3
[2005Sch]
MoAl5 (h1) 648 < T < 750–800
hP60 P321
a = 493.3 c = 4398
[2005Sch]
Comments/References
Space group [P63V-C2]
MoAl5 (h1) MoAl5 (r) ≲ 648
hP36 R3c MoAl5 (r)
a = 495.1 c = 2623
[2005Sch]
MoAl12 < 712
cI26 Im3 WAl12
a = 757.3
[1954Ada, 2005Sch]
U2Mo ≲ 1252 (?)
tI6 I4/mmm MoSi2
a = 342.7 c = 983.4
* τ1, U9Mo16Al75
-
-
[1969Pet]
-
0.31 ≤ x ≤ 0.40 [1969Pet]
* τ2, U(MoxAl1–x)2 hP12 P63/mmc MgZn2
32.5 to 34 at.% Mo [Mas2] [V-C2]
* τ3, U6Mo4Al43
hP2 a = 1096.6 P63/mcm c = 1769.0 Ho6Mo4+xAl43–x (x = 0.11) or Yb6Cr4+xAl43–x (x = 1.15)
[1994Nie]
* τ4, UMo2Al20
cF8 Fd3m CeTi2Al20 or CeCr2Al20
[1995Nie]
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a = 1450.6
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. Table 3 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
L + ξ1 Ð Mo3Al8 + ξ2
1510
U1
L
L + ξ2 Ð Mo3Al8 + Mo3Al
L + λ2 Ð τ2
L Ð τ2 + Mo3Al
1480
1460
1410
U2
p3
e5
1380
U3
ξ1
63.0
37.0
0
Mo3Al8
72.7
27.3
0
ξ2
48.0
52.0
0
60.0
30.0
ξ2
48.0
52.0
0
Mo3Al8
72.7
27.3
0
Mo3Al
25.0
75.0
0
L
45.2
21.5
33.3
L
L Ð λ2 + Mo3Al8 + Mo3Al
1340
U4
E1
1220
U5
48.2
18.5
33.3
46.0
20.7
33.3
L
39.0
29.0
32.0
τ2
39.7
27.0
33.3
25.0
75.0
L
48.5
27.5
24.0
τ2
45.5
21.5
33.3
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0
49.0
18.0
25.0
75.0
66.4
18.3
15.3
τ1
75.0
16.0
9.0
λ2
61.0
6.0
33.0
Mo3Al8
72.7
27.3
0
L
58.0
26.0
16.0
58.5
8.5
33.0
Mo3Al8
72.7
27.3
0
Mo3Al
25.0
75.0
0
L
6.0
28.0
66.0
(Mo)
1.5
97.0
1.5
γ
1.0
39.0
60.0
25.0
75.0
L
Mo3Al
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λ2
λ2
L + (Mo) Ð γ + Mo3Al
4.0
τ2
Mo3Al 1352
U
29.5
λ2 L + τ1 Ð Ð λ2 + Mo3Al8
Mo 66.5
Mo3Al L + τ2 Ð λ2 + Mo3Al
Al
33.3 0
0
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. Table 3 (continued) Composition (at.%) Reaction
T [˚C]
Type
Phase
L + Mo3Al Ð γ + τ2
1140
U6
L
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U7
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Mo
U
9.0
22.0
25.0
75.0
γ
3.0
24.0
73.0
τ2
39.2
27.0
33.8
L
11.5
10.5
78.0
Mo3Al
L + τ2 Ð γ + λ 2
Al
69.0 0
τ2
44.8
21.5
33.7
γ
4.5
11.5
84.0
λ2
50.4
15.8
33.8
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. Fig. 1a Al-Mo-U. Partial reaction scheme, part 1
Al–Mo–U
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. Fig. 1b Al-Mo-U. Partial reaction scheme, part 2
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. Fig. 2 Al-Mo-U. Partial liquidus surface projection
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. Fig. 3 Al-Mo-U. Partial isothermal section at 1250˚C (U-UAl2-Mo3Al8-Mo region)
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. Fig. 4 Al-Mo-U. Partial isothermal section at 1050˚C (U-UAl2-Mo3Al8-Mo region)
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. Fig. 5 Al-Mo-U. Partial isothermal section at 500˚C (U-UAl2-Mo3Al8-Mo region)
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. Fig. 6 Al-Mo-U. Temperature - composition section UAl2-Mo3Al8
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. Fig. 7 Al-Mo-U. Isopleth at 33.3 at.% U
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. Fig. 8 Al-Mo-U. Temperature - composition section U52Al48-Mo3Al (plotted in at.%)
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. Fig. 9 Al-Mo-U. Isopleth at 40 at.% Mo
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. Fig. 10 Al-Mo-U. Isopleth at 20 at.% Al
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. Fig. 11 Al-Mo-U. Isopleth at 50 at.% Al
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. Fig. 12 Al-Mo-U. Isopleth at 20 at.% Mo
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. Fig. 13 Al-Mo-U. Isopleth at 80 at.% U
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References [1954Ada] [1958Woo]
[1964Nic]
[1969Pet] [1971Pet] [1971Rex]
[1979Sof]
[1980Fer] [1984Gom]
[1990Kas]
[1993Ale]
[1994Nie] [1995Gri] [1995Nie]
[1997Lee]
[2000Han]
[2002Kim]
[2002Lee]
[2002Mey]
Adam, J., Rich, J.B., “The Crystal Structure of WAl12, MoAl12 and (Mn, Cr)Al12”, Acta Cryst., (7), 813–816 (1954) (Crys. Structure, Experimental, 14) as quoted by [2005Sch] Wood, E.A., Compton, V.B., Matthias, B.T., Corenzwit, E., “β-Wolfram Structure of Compounds Between Transition Elements, Gallium and Antimony”, Acta Crystallogr., (11), 604–606 (1958) (Crys. Structure, Experimental, 13) as quoted by [2005Sch] Nicholson, S., Harris, D.G., Stobo, J.J., “The Effect of Ternary Additions on the Microstructures of Dilute U-Mo Alloys”, J. Nucl. Mater., 12(2), 173–183 (1964) (Crys. Structure, Morphology, Phase Relations, Experimental, 13) Petzow, G., Rexer, J., “Liquid Equilibria in the Uranium-UAl2-Al8Mo3-Mo System” (in German), Z. Metallkd., 60(5), 449–453 (1969) (Morphology, Phase Diagram, Phase Relations, Experimental, #, 14) Petzow, G., Rexer, J., “Phase Equilibria in the Uranium-UAl2-Al8Mo3-Mo System” (in German), Z. Metallkd., 62(1), 34–38 (1971) (Phase Diagram, Phase Relations, Experimental, #, 9) Rexer, J., “Phase Equilibria in the System Al-Mo at Temperatures above 1400˚C” (in German), Z. Metallkd., 62, 844–848 (1971) (Crys. Structure, Phase Diagram, Experimental, 23) as quoted by [2005Sch] Sofronova, R.M., Nikolaev, A.G., Lyutina, E.M., Voytekhova, E.A., “Influence of Al, Mn, Pd, Ir and Pt Additions on Martensitic Transformation γ → α’b in Uranium Alloys” (in Russian), Alloys for Atomic Energy, Ivanov, O.S., Alekseeva, Z.M. (Eds.), Nauka, Moscow, 131–134 (1979) (Crys. Structure, Morphology, Phase Relations, Experimental, 7) Ferro, R., Marazza, R., “Crystal Structure and Density Data”, Atomic Energy Rev.: Spec. Iss., (7), 359–507 (1980) (Crys. Structure, Phase Relations, Review, 961) Gomozov, L.I., Pokrovskii, A.A., “Influence of Additional Alloying on α-Phase Transformation Kinetics in a Uranium Alloy with 12 at.% Molybdenum”, Russ. Metall. (Engl. Transl.), (4), 142–146 (1984), translated from Izv. Akad. Nauk SSSR, Met., (4), 136–140 (1984) (Crys. Structure, Morphology, Phase Relations, Experimental, Kinetics, Mechan. Prop., 11) Kassner, M.E., Adamson, M.G., Adler, P.H., Peterson, D.E., “The Al-U (Aluminium-Uranium) System”, Bull. Alloy Phase Diagrams, 11(1), 82–89 (1990) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, 44) Alekseeva, Z.M., “Aluminium - Molybdenum - Uranium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.22243.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, #, 6) Niemann, S., Jeitschko, W., “Ternary Aluminides A6T4Al43 with A = Y, Nd, Sm, Gd-Lu, Th, U and T = Cr, Mo, W”, Z. Metallkd., 85(5), 345–349 (1994) (Crys. Structure, Experimental, 23) Grin, Y.N., Ellner, M., Peters, K., Schuster, J.C., “The Crystal Structure of Mo4Al17 and Mo5Al22”, Z. Krist., 210, 96–99 (1995) (Crys. Structure, Experimental, 11) as quoted by [2005Sch] Niemann, S., Jeitschko, W., “Ternary Aluminides AT2Al20 (A = Rare Earth Elements and Uranium; T = Ti, Nb, Ta, Mo, and W) with CeCr2Al20-Type Structure”, J. Solid State Chem., 114, 337–341 (1995) (Crys. Structure, Experimental, 18) Lee, D.B., Kim, K.H., Kim, C.K., “Thermal Compatibility Studies of Unirradiated U-Mo Alloys Dispersed in Aluminium, J. Nucl. Mater., 250(1), 79–82 (1997) (Morphology, Experimental, Interface Phenomena) cited from abstract Han, Y.-S., Park, J.-M., Kim, K.-H., Lee, Y.-S., Kim., C.-K., “An Investigation on the Green Properties of U-10 wt% Mo/Al and U3Si2/Al Powder Compacts”, Nucl. Eng. Design, 202(1), 1–9 (2000) (Morphology, Experimental, Mechan. Prop., 24) cited from abstract Kim, K.H., Park, J.M., Kim, C.K., Hofman, G.L., Meyer, M.K., “Irradiation Behavior of Atomized U-10 wt.% Mo Alloy Aluminum Matrix Dispersion Fuel Meat at Low Temperature”, Nucl. Eng. Design, 211 (2-3), 229–235 (2002) (Morphology, Experimental, Kinetics) cited from abstract Lee, J.-S., Lee, Ch.-H., Kim, K.H., Em, V., “Study of Decomposition and Reactions with Aluminum Matrix of Dispersed Atomized U-10 wt% Mo Alloy”, J. Nucl. Mater., 306(2-3), 147–152 (2002) (Crys. Structure, Morphology, Phase Relations, Experimental, Kinetics, 6) Meyer, M.K., Hofman, G.L., Hayes, S.L., Clark, C.R., Wiencek, T.C., Snelgrove, J.L., Strain, R.V., Kim, K.H., “Low-Temperature Irradiation Behavior of Uranium-Molybdenum Alloy Dispersion Fuel, J. Nucl. Mater., 304(2-3), 221–236 (2002) (Morphology, Experimental, Kinetics) cited from abstract
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[2004Kei]
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Al–Mo–U Garces, J.E., Marino, A.C., Bozzolo, G., “Theoretical Description of the Interdiffusion of Al in the U-Mo Solid Solution”, Appl. Surf. Sci., 219, 47–55 (2003) (Morphology, Calculation, Theory, Interface Phenomena, 11) Lee, S.H., Kim, J.C., Park, J.M., Kim, C.K., Kim, S.W., “Effect of Heat Treatment on Thermal Conductivity of U-Mo/Al Alloy Dispersion Fuel”, Intern. J. Thermophys., 24(5 Special Issue SI), 1355–1371 (2003) (Morphology, Calculation, Experimental, Phys. Prop., 10) cited from abstract Mirandou, M.I., Balart, S.N., Ortiz, M., Granovsky, M.S., “Characterization of the Reaction Layer in U-7 wt% Mo/Al Diffusion Couples”, J. Nucl. Mater., 323(1), 29–35 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, Transport Phenomena, 18) Ryu, H.J., Han, Y.S., Park, J.M., Park, S.D., Kim, C.K., “Reaction Layer Growth and Reaction Heat of U-Mo/Al Dispersion Fuels Using Centrifugally Atomized Powders”, J. Nucl. Mater., 321, 210–220 (2003) (Crys. Structure, Morphology, Phase Relations, Thermodyn., Experimental, Interface Phenomena, Kinetics, 31) Keiser, D.D., Jr., Clark, C.R., Meyer, M.K., “Phase Development in Al-Rich U-Mo-Al Alloys”, Scr. Mater., 51, 893–898 (2004) (Crys Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, 11) Leenaers, A., Van den Berghe, S., Koonen, E., Jarousse, C., Huet, F., Trotabas, M., Boyard, M., Guillot, S., Sannen, L., Verwerft, M., “Post-irradiation Examination of Uranium-7 wt% Molybdenum Atomized Dispersion Fuel”, J. Nucl. Mater., 335 (1), 39–47 (2004) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, 18) Tougait, O., Noel, H., “Stoichiometry of UAl4”, Intermetallics, 12, 219–223 (2004) (Crys. Structure, Experimental, 24) Schuster, J.C., “Al-Mo (Aluminium-Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 30.12123.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 61) Kim, Y.S., Hofman, G.L., Ryu, H.J., Hayes, S.L., “Irradiation-Enhanced Interdiffusion in the Diffusion Zone of U-Mo Dispersion Fuel in Al”, J. Phase Equilib. Diffus., 27(6), 614–621 (2006) (Morphology, Thermodyn., Calculation, Experimental, Interface Phenomena, Kinetics, 16) Ryu, H.J., Kim, Y.S., Hofman, G.L., Park, J.M., Kim, C.K., “Heats of Formation of (U,Mo)Al3 and U(Al, Si)3”, J. Nucl. Mater., 358, 52–56 (2006) (Morphology, Thermodyn., Experimental, 10) Ryu, H.J., Park, J.M., Kim, C.K., Kim, Y.S., Hofman, G.L., “Diffusion Reaction Behaviors of U-Mo/Al Dispersion Fuel”, J. Phase Equilib. Diffus., 27(6), 651–658 (2006) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, 21) Wieschalla, N., Bergmaier, A., Boeni, P., Boening, K., Dollinger, G., Grossmann, R., Petry, W., Roehrmoser, A., Schneider, J., “Heavy Ion Irradiation of U-Mo/Al Dispersion Fuel”, J. Nucl. Mater., 357, 191–197 (2006) (Morphology, Phase Relations, Experimental, Interface Phenomena, 9) Noel, H., Tougait, O., Potel, M., “Crystal Structures and Phase Stoichiometry of Nuclear Materials in the U-Si-C and U-Mo-Al Systems”, Collected Abstracts of the X International Conference on Crystal Chemistry of Intermetallic Compounds, Lviv, Ukraine, 17–20 September, 2007, Ivan Franko National University of Lviv, 5 (2007) (Crys. Structure, Experimental, 4) Soba, A., Denis, A., “An Interdiffusional Model for Prediction of the Interaction Layer Growth in the System Uranium-Molybdenum/Aluminum”, J. Nucl. Mater., 360, 231–241 (2007) (Morphology, Calculation, Theory, Interface Phenomena, Kinetics, 22) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Niobium – Nickel Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Lesley Cornish, Damian M. Cupid, Joachim Gro¨bner, Annelies Malfliet
Introduction Knowledge of the phase equilibria and thermodynamic properties of the Al-Nb-Ni system is relevant for many applications, especially for nickel-based superalloys. Pertinent to this is the order-disorder transformation between the disordered (Ni) and the ordered L12-Ni3Al [2003Du]. The excellent high temperature strength and high corrosion resistance of the binary intermetallic compound Ni3Al can be improved by alloying Nb to this compound, and this element has a large ternary solubility [2003Cer1]. The first experimental work on the Al-Nb-Ni system was by [1962Min1, 1962Min2] and [1965Kor], who both constructed the quasibinary section between Ni3Al and NbNi3. [1966Mar] determined the isothermal sections of most of the phase diagram at 900 and 1140˚C. Other investigations are mostly concentrated on the Ni rich corner [1970Cis, 1970Duv, 1979Oma, 1983Och, 1969Gus, 1989Hon]. The phase relations and homogeneity range of Ni3Al are studied by [1980Nas, 1994Jia]. An evaluation of the ternary Al-Nb-Ni system is made by [1993Sau]. [1992Lee, 2003Du, 2006Rag] reviewed the literature published up to this date. A Calphad type calculation of the entire phase diagram is reported by [2003Du], which is in good agreement with most of the experimental data. Although this calculation did not agree with all the experimental data, it has the best overall fit for describing the whole system. In cases where the agreement was less good, this was deduced to be because the samples had been annealed for relatively short times, considering the high melting points of most of the phases. Typically, at least 1200 h would be a reasonable annealing time for high temperature phases in bulk alloys with high melting compounds. Additionally, in all the isothermal sections calculated by [2003Du], the NbNi3 phase had an improbable shape in that there was some extension into the ternary on the Ni rich side, but not on the Nb rich side, and this has been smoothed in the present evaluation. An overview of the investigations considering phase equilibria, solid phases and thermodynamics of the system is given in Table 1.
Binary Systems The phase diagram of the binary Nb-Ni system is adopted from the calculations of [2006Che], which is in good agreement with most available experimental data. It is an extension of [2003Du], but used more recent experimental data. This diagram is redrawn in Fig. 1, but to be consistent with the other binary and ternary data the melting point of Nb is lowered from 2477 to 2469˚C. The binary Al-Ni phase diagram is taken from the MSIT evaluation by
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[2004Sal]. For the Al-Nb phase diagram, [Mas2] was accepted, with additional information on the Al rich liquidus Al-Nb was taken from [1973Wil].
Solid Phases The binary and ternary phases of the Al-Nb-Ni system are listed in Table 2. [1966Ben, 1968Hun] reported that the NbNi phase takes up a considerable quantity of Al. The binary phases, NbNi and Ni3Al have considerable extension into the ternary system [1966Ben]. [1968Hun] confirmed the large solubility of Al in the NbNi phase to be about 35 at.% Al. [1969Gus] measured a solubility of about 10 at.% Nb in Ni3Al, which only varied slightly over the temperature range 800 - 1200˚C. [1980Nas] confirmed this solubility at 1200˚C. The variation of the lattice parameter of Ni3Al along the line Ni3Al - NbNi3 was quite well established by [1962Min1, 1963Arb, 1969Gus, 1970Cis, 1980Nas, 1984Och1]. [1980Nas] reported a solubility of at least 5.5 at.% Nb in NiAl in equilibrium with τ1 and NbNi3 at 1200˚C. There are at least three true ternary phases τ1, τ2 and τ3 (denoted as H, L and M in the original literature) which exist only in the ternary diagram. The first data on these phases arose from [1964Sch, 1965Mar, 1965Ram]. They basically agree on the existence of the τ1, NbNi2Al, a Heusler type phase and τ2, Nb(Al,Ni)2 with a MgZn2 type structure. [1965Ram] reported that a phase, isostructural with Ti2Ni, exists in equilibrium with the τ1 phase, but no other studies report this phase. The τ2 phase has extensive solubility of Ni and Al confirming the formula Nb(AlxNi1–x)2, where x varies from 0.19 to 0.83 [1966Ben]. This result agrees with the data of [1966Mar]. In addition, above 900˚C an orthorhombic phase τ3 structurally related to the NbNi phase forms [1966Ben, 1967Sho, 1968Hun]. According to the phase diagram of [1966Ben], the τ3 phase extends from about 3 to 30 at.% Al. However, [1968Hun] observed a smaller homogeneity range for the τ3 phase. In addition, he found an additional ternary phase with higher Al content and like τ3, with a structure similar to NbNi. The discrepancy between the two authors could result from the difference in annealing time as [1966Ben] annealed only 20 h, while [1968Hun] used 168 h. However, [1968Hun] used only a limited number of samples, and the phase relations in his isothermal section are unusual. As no other experimental investigations report on this phase region, this part of the phase diagram is questionable. [1966Mar] did not observe the ternary τ3 phase at 900˚C and only a small extension of NbNi into the ternary. Thermal conductivity contours were used to determine the site preference of Nb in Ni3Al [2001Ter], and ab-initio calculations were used to investigate the site preference of Nb substitution in NiAl [2000Boz, 2001Son, 2002Boz]. Calculations based on the d-orbital level of the transition metals [1985Mor] as well as ab-initio methods [1991Eno] were used to predict the phase boundaries between NiAl and Ni3Al with dissolved Nb. First principles calculations were also used to calculate the relative stability of L12-Cu3Au and D0a-βCu3Ti structures of Ni3Al and NbNi3, each with additional Nb and Al respectively [1996Rav]. [2001Sav] used electrical resistivity measurements to examine the long range ordering of Ni3Al with Nb additions as a function of temperature. Order-disorder transition temperatures and the kinetics of the transition were studied [2001Sav].
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Quasibinary Systems [1965Kor] established that a simple quasibinary eutectic exists between NiAl and τ2 with stoichiometry NbNiAl, with the eutectic temperature and composition being respectively 1437˚C and 16.5 at.% Nb (Fig. 2). However, it should be noted that the thermal analysis data reported in tabular form give a 40˚C temperature range for the eutectic which would be very close to the maximal experimental error at that time. Thus, it is not fully conclusive with these data that the section is a true quasibinary. The position of the eutectic point does not coincide with that of the calculation of [2003Du], and in Fig. 6, the surface of solidification for τ2 has been moved to accommodate the experimental data. In addition, the solubility of Nb in the NiAl phase is less than might be expected from the measurements for NiAl in equilibrium with the τ1 phase [1980Nas], although the two measurements were from slightly different NiAl compositions.
Invariant Equilibria A ternary eutectic reaction (L Ð (Ni) + NiAl3 + NbNi3) was identified by [1973Kra] and [1973Lem] at 1270˚C and 13.6 at.% Nb and 79.1 at.% Ni. From this and the work of [1962Min1, 1962Min2, 1965Kor, 1969Nov] invariant reactions involving the liquid were deduced by [1992Lee]. [1992Lee] also included a possible ternary transition type solid state reaction NiAl + NbNi3 Ð NiAl3 + τ1. However, the evidence for this is considered too speculative here, and the reaction has not been included in Fig. 3, which is from the calculation of [2003Du]. [2003Rio] reported an invariant reaction L Ð NbAl3 + Nb2Al + τ2 at 1553˚C by DTA, with the ternary eutectic composition very near to the binary eutectic L Ð NbAl3 + Nb2Al. This is 150˚C higher than calculated by [2003Du], and the micrographs of [2003Rio] convincingly show three eutectic phases. Therefore, the calculated temperature is taken to be less reliable. The compositions of the phases participating in the invariant reactions are listed in Tables 3 and 4 as digitized from Fig. 4.
Liquidus, Solidus and Solvus Surfaces The liquidus of Ni rich alloys was determined by [1969Nov, 1973Kra, 1973Lem]. Based on these studies [1969Nov, 1973Kra, 1973Lem], [2003Du] undertook thermodynamic calculations and constructed a liquidus projection, which is shown in Fig. 4 [2003Du]. [2001Miu] investigated liquidus and solidus lines in the Ni rich corner between 1325 and 1450˚C (Figs. 5 and 6), which are consistent with the calculated results of [2003Du]. Both Fig. 5 and Fig. 6 are not adjusted to the accepted binary diagrams. The solvus line of (Ni) was reported by [1989Hon] and [2001Miu]. The solvus line of (Ni) [2001Miu] is shown in Fig. 7, and is qualitatively in agreement with [1989Hon].
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Isothermal Sections There have been a number of detailed studies of solid state equilibria in the Ni-rich region of the phase diagram: at 1300˚C by [1994Jia], at 1200˚C by [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1980Nas, 1975Bok, 1994Jia, 1997Uey], at 1100˚C [1994Jia], at 1080˚C by [1975Bok], at 1000˚C by [1970Cis, 1983Och] at 800˚C by [1970Cis, 1969Gus, 1975Bok, 1994Jia], at 750˚C by [1970Duv] and at 600˚C by [1979Oma]. [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1983Och] used both X-ray diffraction and metallography, while [1969Gus] used X-ray diffraction. A study by [1980Nas] used a variety of techniques which included metallography, X-ray diffraction and electron microprobe analysis. [1966Mar] using X-ray diffraction and [1966Ben] using both X-ray diffraction and metallography determined isothermal sections of almost the complete phase diagram at 900 and 1140˚C, respectively. In all the calculated sections, the NbNi3 phase showed some extension on the Ni rich side but not on the Nb rich side, and all of these have been smoothed here. The isothermal section at 1300˚C is given in Fig. 8. It is taken from the calculation by [2003Du] and in good agreement with the tie lines between (Ni) and Ni3Al measured by [1994Jia], although the agreement between NiAl3 and NiAl is less good, but acceptable. At 1200˚C (Fig. 9), [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1975Bok, 1980Nas, 1994Jia, 1997Uey] show good agreement with [2003Du], although the (Ni) boundaries of [1970Cis] were slightly reduced, but within experimental error, especially considering that the samples have been annealed for only 65 h. [1980Nas] agrees with [2003Du], except for the minor phase analyses of two alloys which were probably compromised by the major phase (since the alloy composition was nearer to that of the major phase). In [1966Ben], the isothermal section at 1140˚C shows a three-phase field between τ2, NbAl3 and NiAl. This section is characterized by extensive solubility of the Ni3Al and NbNi phases in the ternary and the existence of three ternary compounds τ1, τ2 and τ3. In the calculated section [2003Du] (Fig. 10), the data point of NbAl3 [1966Ben] falls in a different phase field. However, this could be explained by the retention of NbAl3 due to a short annealing time (20 h), and so does not compromise the isothermal section of [2003Du]. At 1080˚C, the calculated isothermal section of [2003Du] (Fig. 11) is in good agreement with the experimental isothermal section of [1994Jia]. For the 1027˚C isothermal section, as shown in Fig. 12, the experimental data of [1983Och] show very good agreement with [2003Du], although Ni3Al boundary has a slight discrepancy. However, the experimental results of [1989Hon] for the three-phase region (Ni) + NbNi3 + Ni3Al do not fall in the calculated three-phase region of [2003Du], this could be because the alloys were only annealed for 555 h. The isothermal section at 900˚C of [2003Du] is given in Fig. 13. Most of the data at 900˚C [1966Mar] agree with [2003Du], except for the extension for NbNi for which [1966Mar] shows a very limited extent. The greater extent of NbNi might not have been observed due to the lack of equilibrium after annealing the alloys for only the relatively short time of 700 h (Nb having a high melting point would be slow to diffuse), and the high temperature Nb2Al phase might not have disappeared in this time. The extent of NbNi would fall in the NbNi-Nb2Al two-phase field of [1966Mar], and the identified three-phase fields could be due to a third phase which is still disappearing on annealing. Only partial isothermal sections were calculated at 800˚C, and these are shown in Figs. 14a and 14b. In the Ni rich corner at 800˚C (Fig. 14a), the experimental results for the position of the two-phase (Ni) + Ni3Al region [1969Gus, 1970Cis, 1994Jia] are not in agreement with the DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
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calculation of [2003Du], but they were not used for the optimization. However, it must be pointed out that these three sets of experimental data agree well with each other, suggesting that the (Ni) actually has a lower solubility at 800˚C than is accepted by [2003Du]. Additionally, [1969Gus] shows that the (Ni) + Ni3Al two-phase field penetrates further into the ternary than that calculated by [2003Du]. The isothermal section at 800˚C in Fig. 14a was redrawn ensuring the NbNi3 phase is consistent with the accepted binary diagram. The Al rich corner at 800˚C (Fig. 14b) was investigated experimentally by [1965Mar] and the results are in fairly good agreement with the calculated work of [2003Du], even though they were not included in the optimization [2003Du]. Figure 15 shows the isothermal section at 750˚C. The experimental results of [1970Duv] in the Ni rich corner are in good agreement with each other, although (Ni) shows less solubility of Al and more for Nb than calculated by [2003Du]. Similarly, the isothermal section at 600˚C (given in Fig. 16) has some discrepancy between the experimental [1979Oma] and calculated results [2003Du], with the position of the two-phase Ni2Al3+NbAl3 region being slightly different. In general, the studies of [1962Min1, 1962Min2, 1969Gus, 1970Duv, 1980Nas, 1983Och] agree to within ±2 at.% Al and 1 at.% Nb on the composition of the three-phase equilibrium (Ni), Ni3Al and NbNi3, except [1980Nas] who reported up to 2 at.% Al solubility in NbNi3, while the other studies suggest that the solubility is very small. The results of [1970Cis] are quite different from these studies and have not been accepted here.
Temperature – Composition Sections A temperature-composition section at a constant 2.5 mass% Al showing the (Ni), Ni3Al and NbNi3 phases was produced by [1975Bok] experimentally, but is not consistent with the ternary diagrams accepted here (that of [2003Du] with some modifications in the liquidus surface) because of the shape of the liquidus near (Ni). [1962Min1, 1962Min2, 1965Kor] examined liquid/solid as well as solid state equilibria in two separate isopleths. [1962Min1, 1962Min2] examined alloys formed between Ni3Al and NbNi3 using thermal analysis, metallography, X-ray diffraction, hardness and resistivity measurement and established that a eutectic reaction exists between Ni3Al and NbNi3 at 1280˚C and 16 at.% Nb. This temperature-composition section is not a completely true quasibinary, since Ni3Al does not melt congruently. The temperatures of [1962Min1, 1962Min2] are close, although not in exact agreement, to the accepted Al-Ni binary, but are acceptable within experimental limits of that time. The composition of Ni3Al at the eutectic temperature was also determined by [1963Arb], and later [1969Tho], using metallography, reported that the eutectic composition is slightly lower in Nb at 15.4 at.% Nb. [1963Arb] carefully determined the solubility of Nb in Ni3Al at the quasibinary eutectic temperature using diffraction methods. Both of these additional studies are taken into account in Fig. 17 which has also been slightly altered from [1962Min1, 1962Min2] to be consistent with the AlNi binary system accepted here [2004Sal]. The calculated temperature-composition Ni3Al NbNi3 by [2003Du] is in good agreement with the experimental data of [1962Min1, 1963Arb, 1969Tho], except that the experimental results show less temperature dependence for Ni3Al.
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Thermodynamics A detailed thermodynamic modelling after the Calphad method was published by [2003Du]. Most of the available experimental information for the Al-Nb-Ni system was used for this calculation, except for those alluded to above [1965Mar, 1969Gus, 1970Cis, 1994Jia]. No fundamental discrepancies between the different experimental information taken from literature were observed. Experimental thermodynamic data were not available for this work. More recently [2006Hu] reported enthalpies of formation for two ternary phases τ1 and τ2 as well as other phases, and selected values are given in Table 5.
Notes on Materials Properties and Applications The most important application of the Al-Nb-Ni system is in the Ni-based superalloys, which can have up to 14 components. The matrix of this alloys comprised (Ni) and the major precipitates are based on Ni3Al. However, with increasing addition, the precipitates can be more complex. Nb is a useful addition for increasing the melting points. The Al-Nb-Ni system is also of interest as a metallic glass, with Al particles embedded in the glass phase, for structural materials. The glass formation and crystallization temperatures of amorphous (Ni60Nb40)100–xAlx alloys were investigated using differential scanning calorimetry [1987Akh, 2007Yu] and differential scanning calorimetry with differential thermal analysis [2004Lee]. [2007Yu] found that although the amorphous phase was resistant to attack by Al, the crystallized phase was not. Al87Ni10Nb3 was examined by [2004Aud], and He2+ ion bombardment [1991Ska], along with rapid quenching, and mechanical milling [2004Dia] were used to generate amorphous metal phases from NbNiAl and Nb45Ni50Al5 alloys. The Young’s modulus, yield strength, and ultimate tensile strength of an Al-particle reinforced Ni70Nb30 metallic glass was studied by [2006Yu]. [1998Far] produced NiAl matrix composites reinforced with Nb particles by reactive hot pressing of the elemental powders. Microhardness and compressive yield strength of the alloys were measured. [1995Net] used combustion synthesis (of elemental powders) to prepare selected Al-Nb-Ni alloys. [2000DeL] found that excess Al not participating in the aluminothermic reduction reaction for the formation of a Ni-65Nb (mass%) alloy dissolved into the NbNi phase, which is in agreement with the solubility of Al in NbNi. The effect on creep strength of tungsten and molybdenum layers applied to NiAl + Nb powders prior to pressing was studied by [2006Bel]. The ternary eutectic reactions in the Al-Nb-Ni system were used to produce directionally solidified (DS) eutectic alloys with in-situ, aligned composite microstructures without grain boundaries as potential candidate materials for replacement of Ni-based superalloys. After confirming the NbAl3-NbNiAl-Nb2Al ternary eutectic [2000Rio], the influence of growth rate on microstructure [2002Rio, 2004Rio, 2005Cos] and mechanical properties [2002Rio] were examined. [1992Rev] directionally solidified a NiAl-NbNiAl alloy near the eutectic maximum in the ternary. [1975Kra] directionally solidified an alloy near to the (Ni)-Ni3Al-NbNi3 eutectic, and high-temperature strengths of similarly solidified alloys in the same region were investigated by [1984Tor]. Dislocation structures and mechanical properties of a directionally solidified eutectic alloy with (Ni) / Ni3Al matrix with strengthening NbNi3 phase were investigated by [1978Sve]. DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
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Investigations have been made on NiAl for potential shape memory effect [2002Kim], although the beneficial transformation was blocked by the transformation to Ni5Al3. However, Nb additions were found to facilitate the desired transformation. [1991Gra] investigated the oxidation behavior of two- and three-phase alloys in the Al-Nb-Ni system with thermogravimetric analysis. Oxidation of τ2 (NbNiAl) displayed quasilinear oxidation kinetics, whereas oxidation of ternary alloys of τ2 (NbNiAl) with additions of NiAl and (NiAl + Nb3Al) displayed parabolic oxidation rates. The corrosion behavior of Al-Nb-Ni alloys in mixed gas atmospheres of H2/H2O/H2S were generally parabolic, although linear behavior was also observed [1993He]. [1980Bhe] investigated the effect of different Al-Nb-Ni coating compositions on the oxidation of a selected Al-Nb-Ni alloy. Three types of oxidation behaviors were observed that depended on coating microstructure: formation of protective Al2O3 coatings, formation of non-protective Nb-oxide scales, and formation of non-protective NiO scales.
Miscellaneous Ni76Al22Nb22 and Ni3Al0.75Nb0.25 alloys were shown to obey an empirical relationship between bulk modulus B, molar volume V, and electron density η determined by [2004Li] as B = η2V. [2003Cer1] studied the effect of Nb concentration on interdiffusion coefficients in Ni3Al alloys with Nb using diffusion couples. Similarly, [2003Cer2] measured the interdiffusion coefficients of Nb and Al in the temperature interval 1173–1533 K with a Ni3Al + Ni-0.15Al0.075Nb couple to ascertain the effect of site preference on interdiffusion coefficients. The elastic modulus of a single crystal of composition Ni-0.19Al-0.06Nb was investigated using the velocity of elastic waves of frequencies 5.5, 15, and 50 MHz [2006Rin]. The coefficient of linear expansion was examined in several alloys based on the Ni3Al-NbNi3 composition line [1967Arb]. The results showed an inverse relationship between amount of NbNi3 and the linear expansion coefficient. The low-temperature (3.2 K-10.3 K) specific heat of NbNi2Al was determined using adiabatic calorimetry [1999Roc] and used to calculate the electronic specific heat, Debye temperature, and Einstein temperature. The experimental results were compared to theoretical calculations using the linear “muffin-tin” orbital-tight-binding (LMTO-TB) method. Other experimental work includes the microalloying Nb to NiAl [1991Sas], the characterization of the unstable growth interface with Widmansta¨tten-like precipitates between ternary Ni3Al and ternary NiAl [2001Kai], and Ni6NbAl metastable phase formation [1984Lia] through melt spinning and splat quenching experiments. The reversible hydrogen absorption properties at room temperature of NbNi and Nb10Ni9Al3 were also explored [2003Jou].
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. Table 1 Investigations of the Al-Nb-Ni Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1962Arb]
XRD
Ni3Al-NbNi3 alloys
[1962Min1, 1962Min2]
Microscopy, thermal analysis, hardness, XRD
Vertical section Ni3Al-Ni3Nb
[1963Arb]
XRD, thermal analysis
Ni3Al-NbNi3 alloys at 1300˚C
[1964Sch]
XRD
Three ternary phases NbNi2Al, NbNiAl, NbNiAl2
[1965Gol]
Microchemical analysis, XRD
Ni3Al-NbNi3
[1965Kor]
Microscopy, thermal analysis, hardness, XRD
Vertical section NiAl to 40 at.% Nb
[1965Mar]
XRD, microstructure investigations NbNi2Al, NbNiAl annealed at 800˚C
[1965Ram]
XRD
NbNi2Al at 800˚C
[1966Ben]
Metallography, XRD
NbNiAl at 1140˚C
[1966Mar]
XRD
Complete composition range at 800 and 900˚C
[1967Sho]
XRD
Nb48Ni39Al13
[1968Hun]
XRD
Nb2Al region at 1000˚C
[1969Gus]
XRD
800, 1080 and 1200˚C in Ni rich range
[1969Nov]
Thermal analysis
Liquidus surface of Ni-Ni3Al-NbNi3 system
[1969Tho]
Metallography, chemical analysis
Ni3Al-NbNi3 alloys
[1970Cis]
Metallography, XRD and EPMA
Nickel rich corner at 800, 1000 and 1200˚C
[1970Duv]
Spectroscopy, light metallography, Ni rich corner (> 75 at.% Ni) at 750 and XRD 1200˚C
[1971Min]
XRD, microstructural analysis
Ni-Ni3Al-NbNi3 from room temperature to liquidus temperature
[1972Duv]
Optical metallography, TEM
750˚C Ni rich corner
[1973Kra, 1975Kra]
Directional solidification study, metallography, DTA
Ni-NbNi3-Ni3Al subsystem
[1973Lem]
Metallography, thermal analysis, XRD
Liquidus, Ni rich corner
[1973Wil]
Electronic structure calculation
Ni3Al at 1000˚C
[1975Bok]
Electron microscopy, solidification Ni-Ni3Al-NbNi3 at 800, 1080 and 1200˚C study
[1975Kau]
Calphad calculation
Complete composition range
[1978Gul]
Thermal analysis
Ni rich corner
[1979Oma]
Metallography, XRD
Complete system
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1980Nas]
XRD, EPMA
Ni rich corner at 1200˚C
[1983Och]
Metallography, XRD
Solubility in Ni3Al at 1000˚C
[1984Arg]
XRD
Selected phases
[1984Bre]
Metallography, optical microscopy, Ni-Ni3Al-Nb Ni3 at 800, 1000 and 1200˚C SEM
[1984Och1]
XRD
(Ni), Ni3Al lattice parameters
[1984Och2]
Miedema calculation
Bond energies for Ni3Al with Nb
[1985Mis]
XRD
Ni3Al
[1985Tro]
XRD
τ2, NbNiAl (Laves phase)
[1986Bla]
XRD, metallography
NbNi2 from 800 to 1200˚C
[1988Mor]
Solidification study, XRD
Ni3Al at 1000˚C
[1989Hon]
DTA, optical microscopy, SEM
Solvus of Ni between 827 and 1227˚C
[1989Sub]
DTA, SEM, EPMA, XRD
Nb Al3 at 1300˚C
[1991Eno]
Cluster variation calculation
Ni and Ni3Al compositions below 1000˚C
[1991Ino]
Band structure calculations
NbAl3 with Ni substitutions
[1991Mis]
Metallography, XRD
(Ni) solvus surface
[1992Rev]
Metallography, SEM, EPMA
Al-16.5Nb-41.75Ni microstructures
[1994Jia]
EPMA, SEM
Ni rich corner at 800 and 1300˚C
[1997Uey]
SEM, EPMA, XRD
NbNi3 at 1200˚C
[2001Miu]
DTA
Ni rich corner from 900 to 1450˚C
[2002Kri]
XRD, DSC
Al-10Nb-50Ni and Al-10Nb-62Ni from room temperature to 700˚C
[2003Du]
Calphad calculation
Complete composition range above 600˚C
[2003Rio]
SEM, WDS, STA
Alloys: Al: 57.3–54.4; Nb: 41.8–33.3; Ni:0.9–12.3 (at.%)
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. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Al) < 660.452
cF4 Fm 3m Cu
a = 404.96
pure Al at 25˚C [V-C2]
(Nb) < 2469
cI2 Im 3m W
a = 330.04
pure Nb at 25˚C [Mas2]
(Ni) < 1455
cF4 Fm 3m Cu
a = 352.40
pure Ni at 25˚C [Mas2]
NbAl3 < 1680
tI8 I4/mnm TiAl3
a = 384.1 c = 860.9
[1980Jor]
Nb2Al < 1940
tP30 P42/mnm σCrFe
a = 994.4 c = 517.2
66.7 at.% Nb [1980Jor]
Nb3Al < 2060
cP8 Pm 3n Cr3Si
a = 518.0
75 at.% Nb [1980Jor]
NiAl3 < 856
oP16 Pnma Fe3C
a = 661.3 b = 736.7 c = 481.1
[2004Sal]
Ni2Al3 < 1138
hP5 P3m1 Ni2Al3
a = 402.8 c = 489.1
36.8 to 40.5 at.% Ni [2004Sal]
β, NiAl < 1651
cP2 Pm 3m CsCl
a = 287
42 to 69.2 at.% Ni [2004Sal]
Ni5Al3 < 723
oC16 Cmmm Pt5Ga3
a = 753 b = 661 c = 376
63 to 68 at.% Ni [2004Sal]
γ´, Ni3Al < 1372
cP4 Pm 3m AuCu3
a = 356.77 a = 358.9 a = 356.32
73 to 76 at.% Ni [2004Sal]
NbNi < 1290
hR39 R3m W6Fe7
a = 489.4 c = 2674.0 a = 499.3 c = 2710.0
50 at.% Ni, 0 at.% Al [1968Hun] 24 at.% Ni, 30 at.% Al [1968Hun]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C]
Lattice Parameters [pm]
Comments/References
δ, NbNi3 < 1399
oP8 Pmmn TiCu3
a = 509.65 b = 423.23 c = 453.92
[1997Uey]
NbNi8
-
-
-
* τ1, NbNi2Al
cF16 Fm 3m BiF3
a = 597.0
25 at.% Al, 25 at.% Nb [1966Ben] ("H" phase)
* τ2, NbNiAl
hP12 P63/mmc MgZn2
a = 487.0 to 505.0 for formula Nb(AlxNi1–x)2 where x = 0.19 to c = 793.0 to 837.0 0.83 [1966Ben] ("L" phase)
* τ3, Nb10Ni9Al3
oP52 Pnma Nb10Ni9Al3
a = 930.3 b = 1626.6 c = 493.3
13 at.% Al, 48 at.% Nb [1967Sho], ("M" phase)
Metastable NbNi8 < 515
-
-
-
. Table 3 Invariant Four-Phase Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
Al
L + Nb3Al Ð NbNi + (Nb)
1634
U1
L
Nb
Ni
29.00
51.00
20.00
L + Nb3Al Ð Nb2Al + NbNi
1633
U2
L
31.06
49.97
18.97
L+NbNi Ð NbNiAl + Nb10Ni9Al3
1547
U3
L
27.82
42.09
30.09
L + NbNi Ð Nb2Al + NbNiAl
1488
U4
L
49.18
38.11
12.71
L Ð NbAl3 +Nb2Al +NbNiAl
1408
E1
L
55.66
34.87
9.47
L Ð NbNi3 + Ni3Al + (Ni)
1267
E2
L
7.86
14.07
78.07
L + NiAl Ð NbNiAl + Ni3Al
1206
U5
L
16.28
21.56
62.16
L + NbNi3 Ð NbNiAl + Ni3Al
1205
U5
L
15.94
21.63
62.43
L + NbNiAl Ð Nb10Ni9Al3 + NbNi3
1203
U7
L
1.04
38.88
60.08
Ni3Al + NbNiAl Ð NiAl + NbNi3
1200
U8
-
-
-
-
L + Nb10Ni9Al3 Ð NbNi + NbNi3
1191
U9
L
0.58
40.21
59.21
NbNiAl + NiAl Ð NbNi2Al + NbNi3
1174
U10
-
-
-
-
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. Table 3 (continued) Composition (at.%) Reaction
T [˚C]
Type
Phase
NiAl + NbNi3 ÐNbNi2Al + Ni3Al
1117
U11
-
L + NbNiAl Ð NiAl + NbAl3
1082
U12
L + NiAl Ð NbAl3 + Ni2Al3
1051
Nb10Ni9Al3 + NbNi3 Ð NbNiAl + NbNi Nb2Al + NbNi Ð NbNiAl + Nb3Al
Nb
Ni
-
-
-
L
63.62
11.59
24.79
U13
L
69.86
6.47
23.67
1019
U14
-
-
-
-
947
U15
-
-
-
-
L
-
-
-
-
-
-
L + Ni2Al3 Ð NbAl3 + NiAl3
849
U16
Nb3Al + NbNi Ð NbNiAl + (Nb)
797
U17
NbNi + (Nb) Ð NbNi2Al + NbNiAl
779
U18
NbNiAl + (Nb) Ð NbNi2Al + Nb3Al
728
U19
L + NbAl3 Ð NiAl3 + (Al)
647
U20
Al
-
L
. Table 4 Invariant Three-Phase Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
Al
L + Nb3Al Ð NbNi
1635
p3
L
L + NbNi Ð Nb10Ni9Al3
1615
p4
L Ð NbNi + NbNiAl
1571
L Ð NiAl + NbNiAl
Nb
Ni
29.68
50.56
19.76
L
24.6
45.10
30.30
e1
L
37.42
39.09
23.49
1460
e3
L
37.06
18.87
44.07
L Ð NbAl3 + NbNiAl
1419
e4
L
57.5
31.25
11.25
L Ð NbNi3 + NbNiAl
1276
e7
L
7.28
29.96
62.76
L Ð NbNi3 + Ni3Al
1275
e8
L
11.32
15.44
73.24
L Ð NbNiAl + Ni3Al
1204
e9
-
-
-
-
NiAl + NbNiAl Ð NbNi2Al
1178
p7
-
-
-
-
Nb10Ni9Al3 Ð NbNiAl +NbNi
910
e11
-
-
-
-
NbNi Ð NbNi2Al + (Nb)
815
e12
-
-
-
-
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. Table 5 Thermodynamic Properties of Single Phases Temperature Range [˚C]
Phase
Property, per mole of atoms [J, mol, K]
Comments
τ1, Al25Ni50Nb25
25
–38600 ± 1100
Enthalpy of formation [2006Hu]
τ2, Al27Ni40Nb33:
25
–39600 ± 1600
Enthalpy of formation [2006Hu]
τ2, Al33.3Ni33.3Nb33.3
25
–40900 ± 1500
Enthalpy of formation [2006Hu]
τ2, Al40Ni20Nb33: (N.B. not 100%)
25
–46200 ± 1300
Enthalpy of formation [2006Hu]
NiAl3 of composition: Al75Ni17Nb8
25
–40300 ± 1100
Enthalpy of formation [2006Hu]
Ni3Al of composition: Al20Ni76Nb4
25
–28400 ± 800
Enthalpy of formation [2006Hu]
. Table 6 Investigations of the Al-Nb-Ni Materials Properties Reference
Method / Experimental Technique
Type of Property
[1967Arb]
Hardness, mechanical testing
Vickers Hardness, Young’s modulus, thermal expansion
[1978Sve, 2006Gre]
TEM
Dislocation structures
[1980Bhe, 1991Gra]
Thermogravimetric analysis
Oxidation behavior, oxidation rates
[1984Tor]
Metallography, directional solidification, mechanical testing
Phases, high temperature strength, fracture toughness
[1991Laa]
HIP poweder compacts, mechanical testing
Young’s modulus, Poisson’s ratio, and fracture toughness at RT
[1991Och]
Compression testing
Stress-strain behavior
[1992Gra]
Thermogravimetric analysis
Oxidation resistance of intermetallic phases
[1992Rev]
Directional solidification, microhardness
Microhardness and fracture strength
[1993He]
Thermogravimetric analysis
Corrosion behavior
[1993Rei]
Compression testing
High temperature strength, fracture toughness
[1996Mac]
Mechanical testing
Deformation behaviour of τ2, NbNiAl
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. Table 6 (continued) Reference
Method / Experimental Technique
[1998Far]
Solid state sintering of elemental powders; Hardness testing, compression testing
NiAl-Nb composites; Hardness testing, compression testing
[1999Roc]
Calorimetry
Low temperature specific heat
[2002Kim]
Metallography, DTA
Phase transformations
[2002Rio]
Hardness testing
Hardness, fracture toughness
[2003Cer1, 2003Cer2]
Diffusion couples
Interdiffusion coefficients
[2006Bel]
Creep testing
Creep strength
[2006Hag]
Three-point bending
Fracture toughness
[2006Rin]
Velocity of elastic waves
Elastic modulus
[2006Yu]
Compression testing
Fabrication and characterization of Ni-Nb particle-reinforced Al-based composite
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. Fig. 1 Al-Nb-Ni. Adopted binary Nb-Ni system
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Al–Nb–Ni
. Fig. 2 Al-Nb-Ni. Experimental quasibinary system between NiAl and τ2 ( NbNiAl)
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. Fig. 3a Al-Nb-Ni. Reaction scheme, part 1
Al–Nb–Ni
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. Fig. 3b Al-Nb-Ni. Reaction scheme, part 2
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. Fig. 3c Al-Nb-Ni. Reaction scheme, part 3
Al–Nb–Ni
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Al–Nb–Ni
. Fig. 4 Al-Nb-Ni. Liquidus surface projection
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. Fig. 5 Al-Nb-Ni. Experimental liquidus lines in the Ni-rich corner between 1325 and 1450˚C
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Al–Nb–Ni
. Fig. 6 Al-Nb-Ni. Experimental solidus lines in the Ni-rich corner between 1325 and 1450˚C
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. Fig. 7 Al-Nb-Ni. Dependence of the solvus line of (Ni) on Al concentration
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Al–Nb–Ni
. Fig. 8 Al-Nb-Ni. Isothermal section at 1300˚C
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. Fig. 9 Al-Nb-Ni. Isothermal section at 1200˚C
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. Fig. 10 Al-Nb-Ni. Isothermal section at 1140˚C
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. Fig. 11 Al-Nb-Ni. Isothermal section at 1080˚C
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Al–Nb–Ni
. Fig. 12 Al-Nb-Ni. Isothermal section at 1027˚C
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. Fig. 13 Al-Nb-Ni. Isothermal section at 900˚C
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Al–Nb–Ni
. Fig. 14a Al-Nb-Ni. Isothermal section at 800˚C, Ni rich corner
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. Fig. 14b Al-Nb-Ni. Isothermal section at 800˚C, Al rich corner
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Al–Nb–Ni
. Fig. 15 Al-Nb-Ni. Isothermal section at 750˚C
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. Fig. 16 Al-Nb-Ni. Isothermal section at 600˚C
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Al–Nb–Ni
. Fig. 17 Al-Nb-Ni. Experimental temperature - composition section Ni3Al - NbNi3
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References [1962Arb]
[1962Min1]
[1962Min2]
[1963Arb] [1964Sch] [1965Gol]
[1965Kor]
[1965Mar]
[1965Ram]
[1966Ben] [1966Mar]
[1967Arb]
[1967Sho] [1968Hun]
[1969Gus]
[1969Nov]
[1969Tho]
[1970Cis] [1970Duv]
Arbuzov, M.P., Chuprina, V.G., “X-Ray Determination of the Crystal Structures of Alloys in the Ni3AlNi3Nb System” (in Russian), Akad. Nauk SSSR, Inst. Metall. Im A.A. Baikova, (8), 85–87 (1962) (Crys. Structure, Experimental, 13) Mints, R.S., Belyaeva, G.F., Malkov, Yu.S., “The Reaction Between the Metallic Compounds Ni3Al and Ni3Nb” (in Russian), Dokl. Akad. Nauk SSSR, 143(4), 871–874 (1962) (Phase Relations, Phase Diagram, *, 26) Mints, R.S., Belyaeva, G.F., Malkov, Y.S., “Equlibrium Diagram of the Ni3Al-Ni3Nb System”, Russ. J. Inorg. Chem. (Engl. Transl.), 7, 1236–1239 (1962), translated from Zh. Neorg. Khim., 7(10), 2382–2385 (1962) (Experimental, Morphology, Phase Relations, *, 32) Arbuzov, M.P., Chuprina, V.G., “Ageing Alloys of the System Ni3Al-Ni3Nb” (in Russian), Izv. V. U. Z., Fiz., 5, 82–85 (1963) (Crys. Structure, Phase Diagram, Experimental, *, 2) Schubert, K., Meissner, H.G., Raman, A., Rossteutscher, W., “Structural Data of Some Metallic Phases”, Naturwissenschaften, 51, 287 (1964) (Crystal Structure, 0) Golubtsova, R.B., “Selective Isolation of Metallic Compounds from an Alloy of the Ni3Al-Ni3Nb System” (in Russian), Dokl. Akad. Nauk SSSR, 160, 1311–1314 (1965) (Experimental, Phys. Prop., Kinetics, 13) Kornilov, I.I., Mints R.S., Guseva, L.N., Malkov, Y.S., “Interaction Between the Compound NiAl and Niobium”, Russ. Metall. (Engl. Transl.), (6), 93–96 (1965), translated from Izv. Akad. Nauk SSSR, Met., (6), 132–136 (1965) (Experimental, Morphology, Phase Diagram, Phase Relations, #, 5) Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of the MnCu2Al and MgZn2 Types Containing Aluminum and Gallium”, Sov. Phys. Crystallogr., 9, 619–620 (1965) translated from Kristallografiya, 9, 737–738 (1964) (Crys. Structure, Experimental, 4) Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3. III. Investigations in Several T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56, 99–104 (1965) (Crys. Structure, Experimental, Phase Diagram, 14) Benjamin, J.S., Giessen, B.C., Graut, N.J., “Intermediate Phases in the Ternary System Nb(Cb)-Ni-Al at 1140˚C”, Trans. Metall. Soc. AIME, 236, 224–226 (1966) (Experimental, Phase Relations, #, 8) Markiv, V.Ya., Matushevskaya, N.F., Kuz’ma, Yu.B., “X-Ray Diffraction of the Nb-Ni-Al System”, Russ. Metall. (Engl. Transl.), (6), 72–74 (1966), translated from Izv. Akad. Nauk SSSR, Met., (6), 127–129 (1966) (Experimental, Phase Relations, *, 2) Arbuzov, M.P., Chuprina, V.G., “Determination of Some Physical Properties of the Ni3Al-Ni3Nb Alloys”, Russ. Metall. (Engl. Transl.), (3), 78–80 (1967), translated from Izv. Akad. Nauk SSSR, Met., (3), 174–179 (1967) (Experimental, Phase Relations, Phys. Prop., 7) Shoemaker, C.B., Shoemaker, D.P., “The Crystal Structure of the M Phase, NbNiAl”, Acta Crystallogr., 23, 231–238 (1967) (Crys. Structure, Experimental, 18) Hunt, C.R. Jr., Raman, A., “Alloy Chemistry of σ (βαffl U)-Related Phases. I. Extens Ion of μ- and Occurrence of μ’-Phases in the Ternary Systems Niobium(Tantalum)-X-Aluminum X = Iron, Cobalt, Nickel, Copper, Chromium, Molybdenum)”, Z. Metallkd., 59(9), 701–707 (1968) (Crys. Structure, Phase Diagram, *, 14) Guseva, L.N., Mints, R.S., Malhov, Y.S., “Phase Equilibria in the Ni-Ni3Al-Ni3Nb System at 800–1200˚C”, Russ. Metall. (Engl. Transl.), (5), 120–122 (1969), translated from Izv. Akad. Nauk SSSR, Met., (5), 186–188 (1969) (Experimental, Phase Relations, *, 12) Novic, F.S., Mints, R.S., Malkov, Yu.S., “Plotting the Liquidus Surface of a Ni-Ni3Al-Ni3Nb Ternary System” (in Russian) in “Teor. Eksp. Metody Issled. Diagramm Sostoyaniya Metal. Sistem”, Ageev, N.V. (Ed.), Nauka, Moscow, 145–150 (1969) (Phase Diagram, Experimental, *, 10) Thompson, E.R., Lemkey, F.D., “Structure and Properties of Ni3Al Eutectic Alloys Produced by Unidirectional Solidification”, Trans. Quart., A.S.M., 62, 140–154 (1969) (Phase Diagram, Experimental, *, 35) Cisse, J., Davies, R.G., “Nickel-Rich Portion of the Nickel-Aluminum-Niobium Phase Diagram”, Met. Trans., 1, 2003–2006 (1970) (Experimental, Mechan. Prop., Morphology, Phase Diagram, 9) Duvall, D.S., Donachie, M.J., “Phase Equilibria in Nickel-Rich Ni-Al-Nb Alloys”, J. Inst. Met., 96, 182–187 (1970) (Experimental, Morphology, Phase Diagram, *, 16)
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DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
35
36
11 [1971Min]
[1972Duv] [1973Kra]
[1973Lem]
[1973Wil] [1975Bok]
[1975Kau] [1975Kra]
[1978Gul] [1978Sve]
[1979Oma]
[1980Bhe]
[1980Jor]
[1980Nas] [1983Och] [1984Arg] [1984Bre] [1984Lia] [1984Och1]
[1984Och2] [1984Tor]
Al–Nb–Ni Mints, R.S., Malkov, Y.S., Davydor, N.I., Kofanova, G.N., “Polythermal Sections of the Ni-Ni3Al-Ni3Nb System at Nb 5 and 10 wt.-%” (in Russian), Diagrammy Sost. Metall. Sistem, 158–160 (1971) (Phase Diagram, Phase Relations, Experimental, Mechan. Prop., 8) Duvall, D.S., Donachie, M.J. Jr., “Precipitation Characteristics of Nickel-Rich Ni-Al-Nb Alloys”, J. Inst. Met., 100, 6–12 (1972) (Experimental, Morphology, Phase Relations, 22) Kraft, E.H., Thompson, E.R., “Directional Solidification of Ni-Nb, Ni-Al-Nb, Ni-Cr-Nb and Ni-Cr-AlNb Alloys”, Proc. Conf. “In Situ Composites”, National Academy of Science, Washington, 297–303 (1973) (Phase Relations, Experimental, *, 18) Lemkey, F.D., Thompson, E.R., “Eutectic Superalloys Strengthened by Aligned S-Ni3Nb Lamellas, γ’(Ni3Al) Precipitates and Reduced Interlamellar Spacing”, Proc. Conf. “In Situ Composites”, National Academy of Science, Washington, 161 (1973) (Phase Relations, Experimental) Willey, L.A., “The Al-Nb System” in “Metals Handbook”, 8th Ed., Vol. 8, ASM, Metals Park Ohio, 342 (1973) (Phase Relations, Review, #, 3) Bokshteyn, S.Z., Vasilenok, L.B., Kishkin, S.T., Nazarova, M.P., Svetlov, I.L., Sorokina, L.P., Khusnetdinov, F.M., “Formation of the Eutectic γ/γ’ - δ Microstructure During Directional Crystallization and After Heat Treatment”, Phys. Met. Metallogr., 79–85 (1975) (Experimental, Morphology, Phase Diagram, 17) Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams. Part III”, Metall. Trans. A, 6(11), 2115–2122 (1975) (Calculation, Crys. Structure, Phase Diagram, Thermodyn., 35) Kraft, E.H., Thompson, E.R., “Directional Solidification of Ni-Nb, Ni-Al-Nb, Ni-Cr-Nb and Ni-Cr-AlNb Alloys”, 2nd Conf. on in Situ Composites, Boston, 297–307 (1975) (Experimental, Thermodyn., Phase Relations, 18) Gulyaev, B.B., Grigorash, E.F., Efimova, M.N., “Solidification Range of Ni Alloys”, Met. Sci. Heat Treat., 20, 914–917 (1978) (Experimental, Phase Relations, 8) Svetlov, I.L., Vasilenok, L.B., Khusnetdinov, F.M., Sidorov, V.V., Nazarova, M.P., “Plastic Deformation of Directionally Crystallized Eutectic Ni/Ni3Al-Ni3Nb”, Phys. Met. Metallogr., 45(1), 124–129 (1978) (Experimental, Mechan. Prop., 12) Omarov, A.K., Smagulov, S.U., Askarov, M.S., “Study of Phase Equilibriums in Alloys of Al-Ni-Zn and Al-Nb-Ni Ternary Systems” (in Russian), Metall. Obogashch., 133–137 (1979) (Abstract, Experimental, Phase Diagram, Phase Relations, 5) Bhedwar, H.C., Heckel, R.W., Laughlin, D.E., “The Oxidation Behavior of Aluminide-Coated γ/δ Directional Eutectics”, Metall. Trans. A, 11, 1303–1314 (1980) (Experimental, Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Phys. Prop., 25) Jorda, J.L., Fluekiger, R., Mueller, J., “A New Metallurgical Investigation of the NiobiumAluminium System”, J. Less-Common Met., 75, 227–239 (1980) (Crys. Structure, Phase Diagram, Experimental, #, 20) Nash, P., Kavishe, F.P.L., West, D.R.F., “Nickel-Rich Region of Ni-Al-Nb System at 1473 K”, Met. Sci., 14(4), 147–149 (1980) (Experimental, Phase Relations, Crys. Structure, *, 21) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull. P.M.E. (T.I.T.), 52, 1–17 (1983) (Phase Diagram, Experimental, 45) Argent, B.B., “Phase Diagrams of Alloys Based on Niobium”, Metals Society/AIME Acc. No. 84(7), 72–486, 325–415 (1984) (Crys. Structure, Phase Diagram) Brezovsky, M., “Morphology of Eutectic Composite Material γ-γ´”, Kovove Mater., 1 (22), 104–112 (1984) (in Slovakija) (Experimental, 17) Liang, W.W., Standley, R., Nash, P., Skowron, M., “The Relative Stabilities of the Ni6AlX (X = V, Nb, Ta) Phases”, J. Mater. Sci. Lett., 3(3), 259–261 (1984) (Experimental, 5) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni (γ), Ni3Al (γ’) and Ni3Ga (γ’) Solid Solutions”, Bull. Res. Lab. Precis. Machin. Electron., Tokyo Inst. Technol., 53, 15–28 (1984) (Crys. Structure, Experimental, *, 66) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289–298 (1984) (Experimental, Phase Diagram, 90) Toropov, V.M., Bondarenko, Yu.A., “Properties of High-Temperature Alloys of the System Ni-Al-Nb with a Unidirectional Eutectic Structure”, Met. Sci. Heat Treat., 26(9-10), 660–664 (1984), translated from Metalloved. Term. Obrab. Met., (9), 11–15, 1984 (Experimental, Mechan. Prop., Phase Diagram, 8)
DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
MSIT1
Landolt‐Bo¨rnstein New Series IV/11E1
Al–Nb–Ni [1985Mis]
[1985Mor] [1985Tro] [1986Bla] [1987Akh] [1988Mor]
[1989Hon] [1989Sub]
[1991Eno]
[1991Gra]
[1991Ino] [1991Laa]
[1991Mis] [1991Och]
[1991Sas]
[1991Ska]
[1992Gra]
[1992Lee] [1992Rev]
[1993He]
[1993Rei]
11
Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161–1169 (1985) (Experimental, 64) Morinaga, M., Yukawa, N., Ezaki, H., Adachi, H., “Solid Solubilities in Nickel-Based F.C.C. Alloys”, Philos. Mag. A, 51(2), 247–252 (1985) (Phase Diagram, Phase Relations, Experimental, 26) Trojko, R., Blazina, Z., “Metal-Metalloid Exchange in Some Friauf-Laves Phases Containing Two Transition Metals”, J. Less-Common Met., 106, 293–300 (1985) (Crys. Structure, Experimental, 13) Blazina, Z., Trojko, R., “Structural Investigations of the Nb(1–x)SixT2 and Nb(1–x)AlxT2 (T = Cr, Mn, Fe, Co, Ni) Systems”, J. Less-Common Met., 119, 297–305 (1986) (Crys. Structure, Experimental, 6) Akhtar, D., Misra, R.D.K., “Formation and Stability of Ni60Nb40-xAlx”, J. Mater. Sci. Lett., 6(1), 29–30 (1987) (Experimental, 19) Morinaga, M., Sone, K., Kamimura, T., Ohtaka, K., Yukawa, N., “X-Ray Determination of Static Displacements of Atoms in Alloyed Ni3Al”, J. Appl. Crystallogr., 21, 41–46 (1988) (Crys. Structure, Experimental, 18) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of γSolvus in Ni-Al-X Ternary Systems”, Mater. Res. Soc. Symp. Proc., 133, 429–440 (1989) (Experimental, Phase Diagram, Phase Relations, 35) Subramanian, P.R., Simmons, J.P., Mendiratta, M.G., Dimiduk, D.M., “Effect of Solutes on Phase Stability in Al3Nb”, Mater. Res. Soc. Symp. Proc., 133(3), 51–56 (1989) (Experimental, Mechan. Prop., Phase Diagram, 12) Enomoto, M., Harada, H., Yamazaki, M., “Calculation of γ’/γ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143–158 (1991) (Assessment, Calculation, Phase Diagram, 34) Grabke, H.J., Steinhorst, M., Brumm, M., Wiemer, D., “Oxidation and Intergranular Disintegration of the Aluminides NiAl and NbAl3 and Phases in the System Nb-Ni-Al”, Oxid. Met., 35(3-4), 199–222 (1991) (Electronic Structure, Experimental, Kinetics, Thermodyn., 30) Inoue, H.R.P., Kitamura, M., Wayman, C.M., Chen, H., “Phase Stability of Al3Nb as a Function of Nickel Additions”, Philos. Mag. Lett., 63(6), 345–353 (1991) (Crys. Structure, Experimental, 19) Laag, R., Kaysser, W.A., Petzow, G., “A Comparative Study on the Influence of Nb and Ti Additions to Different Processed Atomized NiAl Powders”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 821–826 (1991) (Experimental, Mechan. Prop., 9) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the γ Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123–130 (1991) (Assessment, Experimental, Phase Diagram, 5) Ochiai, S., Shirokura, T., Doi, Y., Kojima, Y., “Microstrucutres and Mechanical Properties of Ni-Nb Aluminides Produced by MA Process”, ISIJ Int., 31(10), 1106–1112 (1991) (Experimental, Mechan. Prop., 27) Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on thr Solidified Structure of NiAl”, Proc. Conf. Intermetal. Comp. - Struct. Mechan. Prop., 877–881 (1991) (Abstract, Experimental, Mechan. Prop., Phase Diagram, 10) Skakov, Yu.A., Djakonova, N.P., Edneral, N.V., Koknaeva, M.R., Semina, V.K., “Some Peculiarities of the Atomic Structure of Metallic Phase Formed During Liguid Quenching and Solid State Reactions”, Mater. Sci. Eng. A, 133, 560–564 (1991) (Experimental, Phase Relations, 5) Grabke, H.J., Brumm, M., Steinhorst, M., “Development of Oxidation Resistant High Temperature Intermetallics”, Mater. Sci. Technol., 8, 339–344 (1992) (Experimental, Interface Phenomena, Phase Diagram, Thermodyn., 21) Lee, K.J., Nash, P., “The Al-Nb-Ni System”, to be published in J. Phase Equilib., (Crys. Structure, Phase Diagram, Review, #, 27) a copy is available at MSI Reviere, R.D., Noebe, R.D., Oliver, B.F., “Processing, Microstructure and Low-Temperature Properties of Directionally Solidified NiAl/NiAlNb Alloys”, Mater. Lett., 14(2-3), 149–155 (1992) (Morphology, Phase Relations, Mechan. Prop., 19) He, Y.R., Douglass, D.L., “The Corrosion Behavior of Ni-Al Alloys and Ni-Nb-Al Alloys in a H2/H2O/ H2S Gas Mixture”, Oxid. Met., 40(3-4), 337–371 (1993) (Experimental, Kinetics, Morphology, Phase Relations, 28) Reip, C.-P., Sauthoff, G., “Deformation Behaviour of the Intermetallic Phase Al3Nb with DO22 Structure and of Al3Nb-base Alloys: Part 1. Physical Properties and Short-Term Behaviour”, Intermetallics, 1, 159–169 (1993) (Experimental, Phase Diagram, Phys. Prop., 30)
Landolt‐Bo¨rnstein New Series IV/11E1
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DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
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38
11 [1993Sau]
[1994Jia]
[1995Net] [1996Mac] [1996Rav]
[1997Uey]
[1998Far]
[1999Roc]
[2000Boz] [2000DeL]
[2000Rio] [2001Kai]
[2001Miu]
[2001Sav]
[2001Son] [2001Ter]
[2002Boz] [2002Kim]
[2002Kri]
[2002Rio]
Al–Nb–Ni Saunders, N., “Aluminium - Niobium - Nickel”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.10206.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, XX) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between γ (A1), γ’ (L12) and β (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473–485 (1994) (Crys. Structure, Experimental, Phase Diagram, 25) Neto, R.M.L., Ferreira, P.I., “Reaction Sintering of Nb-Ni-Al Intermetallic Alloys”, Mater. Sci. Eng. A, 193, 549–555 (1995) (Experimental, Kinetics, Mechan. Prop., Phase Relations, Thermodyn., 18) Machon, L., Sauthoff, G., “Deformation Behaviour of Al-Containing C14 Laves Phase Alloys”, Intermetallics, 4, 469–481 (1996) (Experimental, Phase Relations, 41) Ravindran, P., Subramoniam, G., Asokamani, R., “Ground-State Properties and Relative Stability Between the L12 and Doa Phases of Ni3Al by Nb Substitution”, Phys. Rev. B, 53(3), 1129–1137 (1996) (Calculation, Crys. Structure, Thermodyn., 44) Ueyama, T., Ghanem, M.M., Miura, N., Takeyama, M., Matsuo, T., “Phase Stability of Ni3Nb-δ Phase in Ni-Nb-M Systems at Elevated Temperatures”, THERMEC´97, Intern. Conf. Thermomechan. Proc. Steels Other Mater., TMS, Warrendale, USA, 2, 1753–1760 (1997) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 16) Farber, L., Gotman, I., Gutmanas, E.Y., Lawley, A., “Solid State Synthesis of NiAl-Nb Composites from Fine Elemental Powders”, Mater. Sci. Eng. A, 244(1), 97–102 (1998) (Experimental, Morphology, Phase Relations, 29) da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., da Silva, C.M., Gomes, A.A., “Specific Heat and Electronic Structure of Heusler Compounds Ni2Tal (T = Ti, Zr, Hf, V, Nb, Ta)”, Physica B (Amsterdam), 269, 154–162 (1999) (Crys. Structure, Experimental, 20) Bozzolo, G., Noebe, R.D., Honecy, F., “Modeling of Ternary Element Site Substitution in NiAl”, Intermetallics, 8, 7–18 (2000) (Crys. Structure, Review, 34) De Lima, B.B., Ramos, A.S., Nunes, C.A., Conte, R.A., “Ni-65 wt.% Nb Alloy by Aluminothermic Reduction Process”, Int. J. Refract. Met. Hard Mater., 18, 267–271 (2000) (Experimental, Phase Diagram, 6) Rios, C.T., Milenkovic, S., Caram, R., “Directional Growth of Al-Nb-X Eutectic Alloys”, J. Cryst. Growth, 211, 466–470 (2000) (Experimental, Phase Relations, 8) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of γ´/β Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168–175 (2001) (Experimental, Phase Relations, Thermodyn., 21) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: V, Nb and Ta) Ternary Systems”, J. Phase Equilib., 22, 345–351 (2001) (Experimental, Phase Relations, 9) Savin, O.V., Stepanova, N.N., Akshentsev, Yu.N., Rodionov, D.P., “Ordering Kinetics in Ternary Ni3Al-X Alloys”, Scr. Mater., 45(8), 883–888 (2001) (Crys. Structure, Electr. Prop., Experimental, Kinetics, Thermodyn., 18) Song, Y., Guo, Z.X., Yang, R., Li, D., “First Principles Study of Site Substitution of Ternary Elements in NiAl”, Acta Mater., 49, 1647–1654 (2001) (Calculation, Electronic Structure, 17) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8), 2314–2320 (2001) (Calculation, Crys. Structure, Experimental, Transport Phenomena, 63) Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2 Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Review, 27) Kim, S.-H., Oh, M.-H., Wee, D.-M., “Phase Transformation Behavior and Critical Temperature for Operation of Ni-Al Shape Memory Alloys Including Ternary Elements” (in Korean), J. Korean Inst. Met., 40(6), 621–627 (2002) (Experimental, Morphology, Phase Relations, 18) Krivoroutchko, K.A., Kulik, T., Fadeeva, V.I., Portnoy V.K., “Formation of Stable and Metastable Phases in Ni-Al-Nb and Ni-Al-Me-C (Me = Ti, Nb or V) Powder Systems during Mechanical Alloying and Thermal Treatment”, J. Alloys Compd., 333, 225–230 (2002) (Crys. Structure, Experimental, 13) Rios, C.T., Milenkovic, S., Gama, S., Caram, R., “Influence of the Growth Rate on the Microstructure of a Nb-Al-Ni Ternary Eutectic”, J. Cryst. Growth, 237-239, 90–94 (2002) (Experimental, Phase Relations, 8)
DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
MSIT1
Landolt‐Bo¨rnstein New Series IV/11E1
Al–Nb–Ni [2003Cer1]
[2003Cer2]
[2003Du]
[2003Jou]
[2003Rio] [2004Aud]
[2004Dia]
[2004Lee]
[2004Li]
[2004Rio]
[2004Sal]
[2005Cos]
[2006Bel]
[2006Che] [2006Gre]
[2006Hag]
[2006Hu] [2006Rag]
11
Cermak, J., Rothova, V., “Concentration Dependence of Ternary Interdiffusion Coefficirnts in Ni3Al/ Ni3Al-X Couples with X = Cr, Fe, Nb and Ti”, Acta Mater., 51(15), 4411–4421 (2003) (Electronic Structure, Experimental, Transport Phenomena, 15) Cermak, J., Gazda, A., Rothova, V., “Interdiffusion in Ternary Ni3Al/Ni3Al-X Diffusion Couples with X = Cr, Fe, Nb and Ti”, Intermetallics, 11(9), 939–946 (2003) (Experimental, Kinetics, Transport Phenomena, 24) Du, Y., Chang, Y.A., Gong, W., Huang, B., Xu, H., Jin, Zh., Zhang, F., Chen, S.-L., “Thermodynamic Properties of the Al-Nb-Ni System”, Intermetallics, 11(1-2), 995–1013 (2003) (Assessment, Phase Diagram, Thermodyn., 45) Joubert, J.-M., Pommier, C., Leroy, E., Percheron-Guegan, A., “Hydrogen Absorption Properties of Topologically Close-Packed Phases of the Nb-Ni-Al System”, J. Alloys Compd., 356-357, 442–446 (2003) (Crys. Structure, Experimental, 17) Rios, C.T., Milenkovic, S., Caram, R., “A Novel ternary Eutectic in the Nb-Al-Ni System”, Scr. Mater., 48(10), 1495–1500 (2003) (Crys. Structure, Experimental, Kinetics, Morphology, Phase Relations, 10) Audebert, F., Mendive, C., Vidal, A., “Structure and Mechanical Behaviour of Al-Fe-X and Al-Ni-X Rapidly Solidified Alloys”, Mater. Sci. Eng. A, 375-377, 1196–1200 (2004) (Electronic Structure, Experimental, 22) Diakonova, N.P., Sviridova, T.A., Semina, V.K., Skakov, Yu.A., “Intermetallic Phase Stability on High Energy Treatments (Rapid Quenching, Ion Irradiation and Mechanical Milling)”, J. Alloys Compd., 367, 199–204 (2004) (Crys. Structure, Experimental, Phase Relations, 18) Lee, M.H., Kim, W.T., Kim, D.H., Kim, Y.B., “The Effect of Al Addition on the Thermal Properties and Crystallization Behavior of Ni60Nb40 Metallic Glass”, Mater. Sci. Eng. A, 375-377, 336–340 (2004) (Crys. Structure, Experimental, Interface Phenomena, 12) Li, C., Chin, Y.L., Wu, P., “Correlation between Bulk Modulus of Ternary Intermetallic Compounds and Atomic Properties of their Constituent Elements”, Intermetallics, 12, 103–109 (2004) (Electronic Structure, Thermodyn., 24) Rios, C.T., Oliveira, M.F., Caram, R., Botta F.W.J., Bolfarini, C., Kiminami, C.S., “Directional and Rapid Solidification of Al-Nb-Ni Ternary Eutectic Alloy”, Mater. Sci. Eng. A, 375-377, 565–570 (2004) (Experimental, Interface Phenomena, 15) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.10238.1.20, (2004) (Phase Diagram, Phase Relations, Review, 164) Costa, C.A.R., Batista, W.W., Rios, C.T., Milenkovic, S., Goncalves, M.C., Caram, R., “Eutectic Alloy Microstructure: Atomic Force Microscopy Analysis”, Appl. Surf. Sci., 240(1-4), 414–423 (2005) (Crys. Structure, Experimental, Morphology, 17) Belomyttsev, M.U., Laptev, A.I., Ezhov, I.P., Chertov, S.S., “Strength and Creep of Structural Materials Based on Intermetallic Compound NiAl”, Phys. Met. Metallogr. (Engl. Transl.), 101(5), 397–403 (2006), translated from Fiz. Metal. Metallov., 101(5), 429–435 (2006) (Experimental, Mechan. Prop., Morphology, 6) Chen, H., Du, Y., “Refinement of the Thermodynamic Odelling of the Nb-Ni System”, Calphad, 30, 308–315 (2006) (Calculation, Phase Diagram, #, *, 37) Greenberg, B.A., Antonova, O.V., Ivanov, M.A., Patselov, A.M., Plotnikov, A.V., “Some Features of the Formation and Destruction of Dislocation Barriers in Intermetallic Compounds: II. Observation of Blocked Superidislocations Upon Heating Without Stress”, Phys. Met. Metallogr. (Engl. Transl.), 102(1), 69–75 (2006), translated from Fiz. Metal. Metallov., 102(1), 749–755 (2006) (Crys. Structure, Experimental, Morphology, 5) Hagihara, K., Yokotani, N., Umakoshi, Y., “Temperature and Orientation Dependence of Fracture Behavior of Directionally Solidified Duplex-Phase Crystals Composed of Ni3X-Type Intermetallic Compounds”, Mater. Sci. Forum, 512, 67–72 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 12) Hu, R., Nash, P., “Review: Experimental Enthalpies of Formation of Compounds in Al-Ni-X Systems”, J. Mater. Sci., 41(3), 631–641 (2006) (Experimental, Thermodyn., 101) Raghavan, V., “Al-Nb-Ni (Aluminum-Niobium-Nickel)”, J. Phase Equilib. Diffus., 27(4), 397–402 (2006) (Crys. Structure, Phase Diagram, Review, 28)
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[2007Yu]
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Al–Nb–Ni Rinkevich, A.B., Stepanova, N.N., Burkhanov, A.M., “Acoustical Properties of Ni3Al Single Crystals Alloyed With Cobalt and Niobium”, Phys. Met. Metallogr. (Engl. Transl.), 102(6), 632–636 (2006), translated from Fiz. Metal. Metallov., 102(6), 678–682 (2006) (Experimental, Mechan. Prop., Phys. Prop., Thermodyn., 13) Yu, P., Kim, K.B., Das, J., Baier, F., Xu, W., Eckert, J., “Fabrication and Mechanical Properties of Ni-Nb Metallic Glass Particle-Reinforced Al-Based Metal Matrix Composite”, Scr. Mater., 54(8), 1445–1450 (2006) (Experimental, Mechan. Prop., Morphology, Phase Relations, 30) Yu, P., Zhang, L.C., Zhang, W.Y., Das, J., Kim, K.B., Eckert, J., “Interfacial Reaction During the Fabrication of Ni60Nb40 Metallic Glass Particles-Reinforced al Based MMCs”, Mater. Sci. Eng. A, 444(1-2), 206–213 (2007) (Experimental, Morphology, 31) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009
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Aluminium – Niobium – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Christian Baetzner, Joerg Beuers, Michael Hoch, updated by Kostyantyn Korniyenko
Introduction Phase relations in the ternary Al-Nb-Si system are of great interest above all because alloys of this system are the splendid superconducting materials and also the perspective materials for application in the niobium silicide-based composites with the excellent practical properties similar to the nickel-based superalloys. Experimental data on phase equilibria in the Al-Nb-Si system are represented mainly by the series of isothermal sections at various temperatures [1961Bru, 1971Mue, 1972Pan, 1973All, 1977Mue, 1984Pan, 2001Mur1, 2003Zha]. Crystal structures of the phases taking part in equilibria in the Al-Nb-Si system are reported in [1961Bru, 1961Now, 1971Mue, 1974Joh, 1975Kha, 1977Ale, 1977Gur, 1977Mue, 1978Cat, 1978Dew, 1978Gol, 1987Ves, 2001Mur1, 2003Man, 2003Mur, 2003Zha, 2006Mat]. Experimental study of thermodynamic properties has been done in [1971Dub]. A thermodynamic assessment of the Al-Nb-Si system was carried out in [2004Sha] using the CALPHAD method. Two isothermal sections and the liquidus surface were calculated. Publications concerned with experimental studies of phase relations, crystal structures and thermodynamics, and applied techniques are listed in Table 1. Reviews on phase relations in the Al-Nb-Si system are presented in [1963Eng, 2005Rag, 2006Rag]. The character of phase equilibria was assessed in the previous MSIT evaluation by [1993Bae]. In comparison with that, the present report is supplemented by the information from later publications.
Binary Systems The Al-Si boundary binary system is accepted from the MSIT evaluation by [2004Luk]. The binary boundary Al-Nb and Nb-Si systems are based on [Mas2].
Solid Phases Crystallographic data on the known solid unary, binary and ternary Al-Nb-Si phases and their concentration and temperature ranges of stability are presented in Table 2. The τ1, Nb3Al2Si5, ternary phase has been detected by X-ray diffraction in pressed powder samples annealed at 1400˚C [1961Bru, 1961Now] and the structure and an X-ray pattern are given by [1961Now]. The phase equilibria above 25 at.% Nb were also studied by [1961Bru] and another phase, τ2, appears in the ternary which is suggested to be the Al-stabilized high temperature form of the βNb5Si3 [1961Now]. This view is supported by [1971Mue] who Landolt‐Bo¨rnstein New Series IV/11E1
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concludes the extension of the Nb5Si3 phase up to 21 at.% Al at 1840˚C from metallography and X-ray diffraction in agreement with the binary βNb5Si3 Ð αNb5Si3 transition at 1650˚C [Mas2]. This view is also supported by [1984Pan] who found the βNb5Si3 phase in the ternary Nb5(AlxSi1–x)3 samples with x ranging from 0.6 to 0.75 at 1500˚C [1984Pan] and also at 1400˚C [1961Bru]. It was noted in the assessment of [1993Bae] that the solubility of 11 at.% Al in the αNb5Si3 phase directing towards NbAl3 at 1500˚C [1984Pan] is inconsistent with the solubility of less than 2 at.% Al at 1400˚C given by [1961Bru]. The ternary τ1 phase was also detected by [1973All] in a study of the liquid-solid equilibria at 1500 and 1300˚C (below 33 at.% Nb) by electromagnetic separation of phases and chemical analysis, metallography, electron microprobe, X-ray diffraction and DTA. At 1145 ± 10˚C the invariant transition type reaction L + τ1 Ð NbAl3 + NbSi2 was proposed by [1973All]. But it was concluded in [2003Mur] on the basis of X-ray diffraction patterns of the Nb3Al2Si5 matrix compacts annealed at 1000 and 750˚C that the above-mentioned reaction does not occur at 1145 ± 10˚C and that the τ1 phase is stable below this temperature. This conclusion is confirmed by the participation of the τ1 phase in the phase equilibria at 1000˚C established by [2003Zha]. The homogeneity range of the superconducting Nb3Al phase has been studied and is given in the ternary diagrams at 1840˚C [1971Mue], 1820˚C [1984Pan], 1700˚C [1972Pan], 1500˚C [1984Pan] and 1400˚C [1961Bru]. The solubility of Si was found to range between 3 to 9 at.% Si. Only [1974Joh] concluded a large solubility of about 70 mol% Nb3Si in Nb3Al from X-ray data and Tc (critical temperature) data in sputtered thin film samples annealed at 750˚C for 3 h. The lattice parameter of Nb3Al decreases with Si content [1971Mue, 1974Joh, 1975Kha, 1977Ale, 1977Mue, 1978Cat].
Invariant Equilibria Temperatures, reaction types and phase compositions relating to the invariant equilibria of the system are listed in Table 3. The reaction scheme is shown in Figs. 1a, 1b. The presented scheme is mainly based on the calculated liquidus surface [2004Sha]. It is amended in accordance with the accepted binary boundary systems. The existence of the quasibinary eutectic L Ð Nb5Si3 + NbAl3 proposed in [1984Pan] contradicts to the constitution of the isothermal section at 1500˚C [1973All] and to the calculated liquidus surface projection [2004Sha] and not accepted in the present evaluation.
Liquidus, Solidus and Solvus Surfaces The liquidus surface projection of the Al-Nb-Si is presented in Fig. 2 according to the calculations of [2004Sha] and the constitution of the accepted boundary binary systems. It was predicted that neither the τ1 nor the τ2 phases can melt congruently. Melting of alloys containing the τ2 phase will not occur until above 1475˚C, being consistent with the experimental data [1961Bru, 1984Pan, 2001Mur1]. It was noted by [2004Sha] that for the development of refractory composites based on Nb5Si3 and NbSi2, one should avoid phase fields such as the Nb5Si3 + τ1, so as to avoid problems arising from partial melting and hence liquidinduced fracture at high service temperatures. DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009
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The alloy Nb42.7Al55.7Si1.6 (at.%) containing the binary eutectic Nb2Al-NbAl3 exhibits a regular microstructure characterized by the eutectic cells in the as-cast condition. Ternary element (Si) additions modify the as-cast eutectic microstructure from cellular to dendritic [2000Rio].
Isothermal Sections The Nb rich part of the isothermal section at 1840˚C is presented in Fig. 3 on the basis of data reported by [1971Mue] with amendments according to the accepted here binary edge diagrams. The amendments concerned the extension of the homogeneity ranges of the intermediate phases. It was communicated by [1984Pan] that at 1820˚C the homogeneity range of the Nb3Al phase exists inside the ternary system without contiguity to the boundary binary Al-Nb system, but this assumption was not confirmed by the relevant experimental data. Figure 4 shows the isothermal section at 1500˚C compiled on the basis of [1973All] data for the Al-NbAl3-NbSi2-Si phase region, results of [1984Pan] for the Nb-Nb5Si3-NbAl3 range and the data on the homogeneity ranges discussed in the section “Solid Phases”. The isothermal section at 1400˚C was constructed by [1961Bru] (reproduced in the review [1963Eng]) but it contradicts to the Al-Si binary phase diagram because at this temperature aluminium should be liquid. However, only the solid phases were shown in equilibria. Later it was calculated by [2004Sha]. The calculated section agrees well with the experimental data of [1961Bru] in the Nb rich corner and with the later data of [2001Mur1] concerning phase equilibria with the participation of the τ1 and τ2 phases. It is presented in Fig. 5. A tentative isothermal section at 1300˚C is shown in Fig. 6. It was adopted from the data of [1973All] for the Al-NbAl3-NbSi2-Si phase region and the data on the homogeneity ranges discussed above. The isothermal section at 1000˚C was constructed from the experimental data obtained by [2003Zha] using a high-efficiency diffusion-multiple approach. The phase fields positions in the Nb-NbSi2-NbAl3 region were obtained from the tri-junction area of the diffusion multiple annealed at 1000˚C for 2000 h, at the same time three-phase equilibria at the Al-Si side were estimated from the high-temperature data of [1973All]. The of calculation of the isothermal section at 1000˚C by [2004Sha] agrees well with the experimental data of [2003Zha]. This section is presented in Fig. 7.
Thermodynamics A calorimetric investigation of heats of solution of silicon and aluminium in aluminothermal alloys was carried by [1971Dub]. The heat of dissolution was reported as 100.3 ± 12.5 kJ·mol–1 when the mass ratio Si:Nb was 0.25 and the content of Al varied in the interval from 1.9 to 7.7 mass% and when the mass ratio Si:Nb was 0.60 and the content of Al varied from 1.3 to 7.1 mass%. Thermodynamic assessment of the Al-Nb-Si system was carried out in [2004Sha] by evaluating the available equilibrium data, using the CALPHAD method. The Nb-Si diagram was reassessed to ensure good description of phase equilibria in the ternary system. The assessed binary phase diagram is nearly identical with that of [Mas2] accepted in the present evaluation. They differ only in the invariant temperatures by a few degrees. Thermodynamic Landolt‐Bo¨rnstein New Series IV/11E1
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models for the intermediate phases are based on crystal structures, as well as orientation and solubility ranges of single-phase fields in the experimental isothermal sections at different temperatures. The assessed isothermal sections at the temperatures of 1400 and 1000˚C agree well with the experimental data obtained on the ternary alloys by [1961Bru, 2001Mur1] and [2003Zha], respectively. The liquidus surface projection also was calculated.
Notes on Materials Properties and Applications The Al-Nb-Si alloys are of great practical interest due to the various aspects. In the first place, they possess excellent superconductive properties, with high values of the superconducting transition temperatures (Tc). Also these alloys are the constituents of the niobium silicidebased composites that show great promise for application as the next generation turbine airfoil materials with significantly higher operating temperatures than the current generation advanced nickel-based superalloys. The Nb-Si binary composites possess excellent creep strength but poor oxidation resistance and poor room temperature fracture toughness, at the same time alloying with Al can improve the oxidation resistance [2003Zha]. The applied experimental techniques and studied types of properties are listed in Table 4. It was established by [1971Mue] that the alloy Nb75.3Al21.5Si3.2 (in at.%) possesses a value of the superconducting transition temperature of 19.2 K. A similar value (18.3 K) was measured by [1977Ale] on the Nb75Al19.7Si5.3 specimen homogenized at 1650˚C for 5 h and then annealed at 700˚C for 250 h. For the alloys Nb3Si-Nb3Al alloys (with the aluminium content of 5 at.% and higher) annealed at 750˚C this parameter attains about 14.5 K [1974Joh]. It was shown that additions of Si decrease the superconducting transition temperature of the Nb3Al-based alloys [1974Joh, 1978Cat]. According to [1978Gol], the character of the concentration dependence of the Tc for the films of the Nb3Al-based alloys with silicon additions is smooth with a maximal value of 18.8 K corresponding to the composition of Nb3Al0.8Si0.2. Aiming at further improvement of high-field critical current density, Jc, in the Nb3Al conductors, Si addition was attempted by [2004Ban1]. The value of microhardness for the single-crystal NbAl3.3Si0.01 was reported in [1977Gur] as 4.51 ± 0.39 GPa. Its electrical resistivity was 6.4·10–11 Ω·m or 3.2·10–10 Ω·m at the temperatures of 77 K and 298 K (25˚C), respectively. The effect of mechanical properties of the Nb-based solid solution on toughness and strength of multiphase alloys in the Al-Nb-Si ternary system was studied by [1999Mur]. The toughness of (Nb) single-phase alloys and multiphase alloys estimated from SP energy increases with decreasing of the total (Si + Al) content in (Nb). It was reported by [2001Mur1, 2001Mur2] that the matrix compact corresponding to the composition of the τ1, Nb3Al2Si5 phase showed extremely good oxidation resistance at 1300˚C, although the compact showed poor oxidation resistance at 750˚C.
Miscellaneous Analysis of the ternary Al-Nb-Si alloys by atomic absorption spectrometry was carried out by [1973Mol]. For each element the influence of the acid solvent (HF or HNO3), the influence of two other components and the flame conditions were tested. The formation of the high Tc Nb3Al compound was unsuccessfully attempted by [1978Dew] using solid state diffusion from DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009
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ternary bronzes. The formation of the Nb(AlxSi1–x)2 phase on niobium dipped into the Al-Si liquid saturated by Si was investigated by [1998Nan]. The Nb(AlxSi1–x)2 consisted of submicron-order fine grains. A liquid phase consisting of Al-Si was observed at the grain boundary of the intermetallic layers and at the interface between the product and the refractory metals. This implied the formation of the intermetallics by solution-reprecipitation process. A new method for surface modification based on the arc surface alloying was proposed in [2003Mat]. Its feasibility was investigated performing NbAl3 coated on niobium. When tungsten arc was used to melt an aluminium plate placed on a Nb block, the niobium surface was also melted and a melt pool of Al-Nb binary alloy was formed on a niobium block. The melt pool solidified into niobium aluminides on the surface of the Nb block, forming a thick NbAl3 layer on the top surface of the coating layer. When an Al-Si alloy plate was used instead of the aluminium plate, a niobium alumino-silicide layer was formed on the niobium block. A new fabrication method for multifilamentary Nb/Al-Si precursors was developed by [2004Ban2]. The drawability of the composite wires has been significantly improved by using intermediately a reel-to-reel rapid heating and quenching (RHQ) technique. Unlike the wellknown RHQ process for the Nb3Al processing, this new RHQ process was specially performed at an early stage of the precursor processing to yield a fine microstructure of the core by solidifying rapidly only the Al alloy core from the molten state, without any reaction with the Nb matrix. In authors opinion, using this technique, not only the workability of the core but also the hardness balance between the matrix and the core can be improved, thereby making the subsequent restack and draw process easier. Based on the optimal process parameters, fabrication of multifilamentary Nb/Al-Si precursors with a piece length of 30 m was achieved as laboratory samples, reducing the matrix ratio.
. Table 1 Investigations of the Al-Nb-Si Phase Relations, Structures and Thermodynamics Reference [1961Bru]
Method / Experimental Technique Pressing; sintering; X-ray diffraction
Temperature / Composition / Phase Range Studied 1400˚C; the whole range of compositions
[1961Now] X-ray diffraction
The NbSi2-based and τ1 phases
[1971Dub]
Isothermal water calorimetry; aluminothermy; sintering
1.3 to 9.5 mass% Al
[1971Mue]
High-frequency melting; anode oxidation; X-ray diffraction (Debye-Scherrer technique)
1840˚C; 60–100 at.% Nb
[1972Pan]
Arc melting; annealing; metallography; X-ray 1700˚C; the Nb rich corner diffraction
[1973All]
High-frequency melting; electromagnetic separation; optical microscopy; EMPA
[1974Joh]
Sputtering simultaneously from two different 750˚C; 75 at.% Nb binary alloy electrodes; arc melting; pressing; vacuum annealing; anodization; X-ray diffraction; SEM; optical microscopy
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1975Kha]
Arc melting; annealing; X-ray diffraction; metallography
Nb3Al-based alloys with additions up to 10 at.% Si
[1977Ale]
Arc melting; chemical analysis; annealing; X-ray diffraction
1650, 700˚C; Nb3Al-based alloy with addition of 5.3 at.% Si
[1977Gur]
Spontaneous crystallization; chemical analysis; X-ray diffraction
NbAl3.3Si0.01
[1977Mue]
Arc melting; annealing; metallography; X-ray 1750˚C, as-cast state; 75 at.% Nb diffraction; EMPA
[1978Cat]
Arc melting; rapid quenching; X-ray diffraction;
Rapid quenching from 2000˚C; Nb3Albased alloys with additions up to 15 at.% Si
[1978Dew]
Annealing; EMPA; X-ray diffraction
Nb3Al-based alloys with Si additions
[1978Gol]
Films precipitation; EMPA; X-ray diffraction
600–800˚C; 75 at.% Nb
[1984Pan]
Arc melting; annealing; metallography; X-ray 1820 and 1500˚C; the range diffraction; DTA Nb-Nb5Si3- NbAl3
[1987Ves]
Heating of the pressed workpieces; X-ray diffraction
The Nb3Al-Nb3Si section (up to 70 at.% Si)
[2000Rio]
Arc melting; ingots growth; optical microscopy; SEM; EDS
Nb42.7Al55.7Si1.6 (at.%)
[2001Mur1] Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA)
1200–1600˚C
[2001Mur2] Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA)
1200–1700˚C; Nb3Al2Si5
[2003Man]
Mechanical alloying; ball milling; X-ray diffraction; HRTEM; SAD
Nb40Al40Si20, Nb40Al30Si30
[2003Mat]
Arc surface alloying; SEM; EMPA
Aluminide coating on Nb
[2003Mur]
Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA); EDS
750 and 1000˚C; the τ1, Nb3Al2Si5 phase- containing matrix compacts
[2003Zha]
Diffusion-multiple approach; hot isostatic 1000˚C pressing (HIP); SEM; EMPA; electro- discharge machining (EDM)
[2006Mat]
Self-propagating high temperature synthesis 1100˚C; NbAl3 and Nb5Si3 phases containing alloys (SHS); spark plasma sintering (SPS); X-ray diffraction; SEM; EMPA; optical microscopy
[2007Qu]
Arc melting; EDM; back-scattered electron image (BEI); X-ray diffraction
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. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C] (Al) (I) < 660.452
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
cF4 Fm3m Cu
a = 404.96
hP2 P63/mmc Mg
a = 269.3 c = 439.8
cI2 Im3m W
a = 330.04
T = 25˚C [V-C2]
x = 0, 0 < y ≤ 0.015, T = 577˚C [2004Luk] y = 0, 0 < x · 6·10–4, T = 661.4˚C [Mas2]
NbxAl1–x–ySiy
(Al) (II) NbxAl1–x–ySiy (Nb) < 2469
T = 25˚C, p = 20.5 GPa [V-C2] x = 0, 0 < y ≲ 0.048, T ≈ 605˚C, p = 2.1 GPa [2004Luk] T = 25˚C [V-C2]
x = 0, 0 < y ≤ 0.035, T = 1920˚C [Mas2] x = 0, 0 < y ≤ 0.005, T = 1770˚C [Mas2] y = 0, 0 < x ≲ 0.215, T = 2060˚C [Mas2]
NbxAl1–x–ySiy
(Si) < 1414
cF8 a = 543.06 Fd3m C (diamond)
T = 25˚C [V-C2]
x = 0, 0 < y ≤ 1.8·10–5, T = 577˚C [2004Luk]
NbxAl1–x–ySiy
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. Table 2 (continued)
Phase/ Temperature Range [˚C] Nb3Al < 2060
Pearson Symbol/ Space Group/ Prototype cP8 Pm3m Cr3Si
Nb3(Al1–xSix)
Nb2Al < 1940
tP30 P42/mnm CrFe
Lattice Parameters [pm]
18.6 to 25 at.% Al [Mas2] Strukturbericht designation: A15 a = 518.6 [V-C2] a = 517.3 x = 0, T = 1840˚C [1971Mue] a = 518 x = 0.2, in the alloy sintered at 1400˚C [1961Bru] a = 517.3 x = 0.12, T = 1840˚C [1971Mue] a = 518.8 to 517 0 ≤ x ≤ 0.4, in the as-cast alloys [1975Kha] a = 518.9 to 517.3 0.08 ≤ x ≤ 0.4, in the alloys annealed at 900˚C for 14 d [1975Kha] a = 516.1 x = 0.212, subsequent annealing at 1650 and 700˚C [1977Ale] a = 517 0 ≤ x ≤ 0.5, in the alloys annealed at 1750˚C for 1 h [1977Mue] a = 518.4 to 516.25 0 ≤ x ≤ 0.6, in the alloys quenched from 2000˚C [1978Cat] a = 519 x = 0.2, in the film precipitated and annealed at 600–800˚C [1978Gol] a = 518.4 to 517.4 0 ≤ x ≤ 0.7 [1987Ves]
a = 994.3 c = 518.6 a = 988.2 to 994.2 c = 517.6 to 515.3
NbAl3 < 1680
tI8 I4/mmm TiAl3
a = 384 c = 857
Nb3Si 1980 - 1770
tP32 P42/n Ti3P
a = 1022.4 c = 518.9
βNb5Si3 (h) 2520 - 1650
tI32 I4/mcm W5Si3
αNb5Si3 (r) < 1940
tI32 I4/mcm Cr5B3
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Comments/References
30 to 42 at.% Al [Mas2]; labeled elsewhere as σ phase [V-C2] in the binary alloys with 25 to 35 at.% Al sintered at 1400˚C [1961Bru] 75 at.% Al [Mas2] [V-C2] 25 at.% Si [Mas2] [V-C2] 37.5 to 40.5 at.% Si [Mas2] [V-C2]
a = 1004.0 c = 508.1
37.5 to 38.5 at.% Si [Mas2] [V-C2]
a = 657.0 c = 1188.4
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. Table 2 (continued)
Phase/ Temperature Range [˚C] NbSi2 < 1940
Pearson Symbol/ Space Group/ Prototype hP9 P6222 CrSi2
Nb(AlxSi1–x)2 * τ1, Nb3Al2Si5 Nb(AlxSi1–x)2
oF24 Fddd TiSi2
* τ2, Nb10Al3Si3 Nb5(AlxSi1–x)3
tI32 I4/mcm W5Si3
Lattice Parameters [pm]
Comments/References 66.7 at.% Si [Mas2] Strukturbericht designation: C40 [V-C2]
a = 481.9 c = 659.2 a = 478.7 to 480 c = 658 to 663
0 ≤ x ≤ 0.15, T = 1400˚C [1961Bru, 1961Now] 0.3 ≤ x ≤ 0.375 [1973All] Strukturbericht designation: C54 in the Nb33Al20Si47 alloy [1961Now]
a = 838.6 b = 489.1 c = 877.6
Strukturbericht designation: D8m 0.45 ≤ x ≤ 0.60, T = 1400˚C [1961Bru]
a = 1014 to 1019 c = 507
. Table 3 Invariant Equilibria Composition* (at.%) Reaction
T [˚C]
Type
Phase
Al
Nb
Si
L Ð βNb5Si3 + Nb3Al
-
e1
L
21
73
6
L + Nb3Al Ð βNb5Si3 + (Nb)
2000
U1
L
16
76
8
L + βNb5Si3 Ð Nb3Si + (Nb)
-
U2
L
4
81
15
L + Nb3Al Ð Nb2Al + βNb5Si3
-
U3
L
35
63
2
L + βNb5Si3 Ð αNb5Si3
1800
p4
L
44
34 .5
21 .5
L + βNb5Si3 + NbSi2 Ð τ1
1750
P
L
28
32
40
L + βNb5Si3 Ð αNb5Si3 + τ1
1700
U4
L
37
30
33
L + Nb2Al + NbAl3 + βNb5Si3
-
D
L
57
42
1
L + βNb5Si3 Ð αNb5Si3 + NbAl3
-
U5
L
60
32
8
L + αNb5Si3 Ð τ1 + NbAl3
1490
U6
L
69
18
13
Note: * - values are estimated from the diagram
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. Table 4 Investigations of the Al-Nb-Si Materials Properties Reference
Method / Experimental Technique
Type of Property
[1971Mue]
Anode oxidation, superconductivity tests Superconducting transition temperature, critical current density
[1972Pan]
Superconductivity tests
Superconducting transition temperature, curves of the superconductive transition
[1974Joh]
Anode oxidation, sputtering electrodes superconductivity tests
Superconducting transition temperature
[1975Kha]
Superconductivity tests
Superconducting transition temperature, curves of the superconductive transition
[1977Ale]
Superconductivity tests, string magnetometer tests
Superconducting transition temperature, paramagnetic susceptibility
[1977Gur]
Microdurometry, electrical resistivity measurements
Microhardness, electrical resistivity
[1977Mue]
Superconductivity tests
Superconducting transition temperature
[1978Cat]
Superconductivity tests
Superconducting transition temperature
[1978Dew]
Superconductivity tests (inductive technique)
Superconducting transition temperature
[1978Gol]
Superconductivity tests (resistive potentiometry)
Superconducting transition temperature, curves of the superconductive transition
[1980Mat]
Superconductivity tests (inductive NQR spectra; superconducting transition technique); magnetic and electrical tests; temperature; critical magnetic field; nuclear quadrupole resonance (NQR) electrical resistivity
[1987Ves]
Superconductivity tests (inductive technique); magnetic measurements
Superconducting transition temperature, curves of the superconductive transition; critical magnetic field dependences on temperature
[1989Kri]
Superconductivity tests (four-probe technique)
Volt-ampere characteristics
[1999Mur]
Small Punch (SP) toughness tests; tensile tests at room temperature; compression tests at high temperatures
Toughness; strength
[2001Mur1] Vickers microhardness tests; compression Microhardness; yield stress; oxidation tests; oxidation tests resistance
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. Table 4 (continued) Reference
Method / Experimental Technique
Type of Property
[2001Mur2] Vickers microhardness tests; compression Microhardness; yield stress; oxidation tests; oxidation tests; Knoop indentation resistance; elastic modulus; fracture method; Vickers indentation toughness; thermal expansion coefficient microfracture method; thermal expansion tests [2004Ban1] Superconductivity tests [2006Mat]
Critical current density
Vickers hardness tests; ultrasonic method; Hardness; Young’s modulus; shear four-point bending tests; Archimedes modulus; Poisson’s ratio; density; method bending strength; fracture toughness
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. Fig. 1a Al-Nb-Si. Reaction scheme, part 1
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. Fig. 1b Al-Nb-Si. Reaction scheme, part 2
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. Fig. 2 Al-Nb-Si. Calculated liquidus surface projection
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. Fig. 3 Al-Nb-Si. Nb-rich part of the isothermal section at 1840˚C
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. Fig. 4 Al-Nb-Si. Isothermal section at 1500˚C
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. Fig. 5 Al-Nb-Si. Isothermal section at 1400˚C
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. Fig. 6 Al-Nb-Si. Tentative isothermal section at 1300˚C
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. Fig. 7 Al-Nb-Si. Isothermal section at 1000˚C
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References [1961Bru]
[1961Now] [1963Eng]
[1971Dub]
[1971Mue] [1972Pan]
[1973All]
[1973Mol]
[1974Joh] [1975Kha]
[1977Ale]
[1977Gur]
[1977Mue]
[1978Cat]
[1978Dew]
[1978Gol]
[1980Mat]
Brukl, C., Nowotny, H., Benesovsky, F., “Investigations in the Ternary Systems: V-Al-Si, Nb-Al-Si, CrAl-Si, Mo-Al-Si or Cr(Mo)-Al-Si” (in German), Monatsh. Chem., 92, 967–980 (1961) (Crys. Structure, Phase Diagram, Experimental, *, 20) Nowotny, H., Benesovsky, F., Brukl, C., “The Ternary: Niobium-Aluminium-Silicon” (in German), Monatsh. Chem., 92(1), 193–196 (1961) (Crys. Structure, Experimental, 14) English, J.J., “Binary and Ternary Phase Diagrams of Columbium, Molybdenum, Tantalum and Tungsten”, Defense Metals Information Center, Batelle Memorial Institute, Columbus 1, Ohio, 183, (99–2)-63 (1963) (Phase Diagram, Review, *, 1) Dubrovin, A.S., Gorelkin, O.S., Demidov, Yu.Ya., Chirkov, N.A., Kostylev, L.S., Kolesnikova, O.D., “Calorimetric Investigation of Heats Solution of Silicon and Aluminium in Aliminothermal Alloys” (in Russian), Metalloterm. Process. Khim. Met. Mater. Conference (1971), 121–130 (1971) (Thermodyn., Experimental, 11) Mueller, A., “Superconductivity in the A15-Phase in the Nb-Al-Si System” (in German), Z. Naturforsch., 26(A), 1035–1039 (1971) (Crys. Structure, Phase Relations, Experimental, Superconduct., *, 7) Pan, V.M., Latysheva, V.I., Sudovtsov, A.I., “Superconductivity of Niobium-Aluminium-Silicon Alloys”, Phys. Met. Metallogr. (Engl. Transl.), 33, 180–183 (1972), translated from Fiz. Met. Metalloved., 33(6), 1311–1313 (1972) (Phase Diagram, Phase Relations, Experimental, Superconduct., *, 3) Allibert, C., Wicker, A., Driole, J., Bonnier, E., “Study of the System Niobium-Aluminium-Silicon. I. Partial Isothermal Sections at 1500 and 1300˚C and Behaviour of the Phase Nb(Si, Al)2” (in French), J. Less-Common Met., 31(2), 221–228 (1973) (Morphology, Phase Diagram, Phase Relations, Experimental, *, 4) Molins, R., Garden, J., Bozon, H., Driole, J., “Study of the System Niobium-Aluminium-Silicon. II. Analysis of Ternary Alloys Niobium-Aluminum-Silicon by Atomic Absorption Spectrometry” (in French), J. Less-Common Met., 31(2), 229–237 (1973) (Morphology, Experimental, Kinetics, 7) Johnson, G.R., Douglass, D.H., “Superconductivity in New A-15 Niobium Alloys”, J. Low Temp. Phys., 14(5), 575–595 (1974) (Crys. Structure, Morphology, Experimental, Superconduct., 27) Khan, H.R., Raub, Ch.J., “Structure and Superconductivity of Ternary and Quaternary A15 Phases Based on Nb3Al” (in German), Metall, 29(7), 673–677 (1975) (Crys. Structure, Morphology, Experimental, Electronic Structure, Superconduct., 10) Alekseyevskiy, N.Ye., Ageyev, N.V., Shamray, V.F., “Superconductivity of Some Three-Component Solid Solutions Based on the Compound Nb3Al”, Fiz. Met. Metalloved., 43(1), 29–35 (1977), translated from Fiz. Met. Metalloved. (USSR), 43(1), 38–44 (1977), (Crys. Structure, Phase Relations, Experimental, Electronic Structure, Magn. Prop., Superconduct., 14) Gurin, V.N., Korsukova, M.M., Popov, V.E., Elizarova, O.V., Belousov, N.N., Kuz’ma, Yu.B., “Solid Solutions of B, C, Si in Aliminides of Transition Metals” in “Single-Crystals of Refractory and Rare Metals, Alloys and Compounds” (in Russian), Akad. Nauk SSSR, Nauka, Moscow, 39–42 (1977) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., 6) Mueller, P., “Superconductivity in Quasibinary Alloys of the A3B-Nb3Si Type with A15 Structure” (in German), Z. Metallkd., 68(6), 421–427 (1977) (Crys. Structure, Morphology, Phase Relations, Experimental, Superconduct., 26) Caton, R., Sweedler, A.R., “The Dependence of the Superconducting Transition Temperature on Silicon Concentration in the NbAlSi Ternary System”, J. Less-Common Met., 60, 91–100 (1978) (Crys. Structure, Morphology, Experimental, Superconduct., 10) Dew-Hughes, D., Luhmann, T.S., “The Thermodynamics of A15 Compound Formation by Diffusion from Ternary Bronzes”, J. Mater. Sci., 13, 1868–1876 (1978) (Crys. Structure, Phase Relations, Calculation, Experimental, Interface Phenomena, Superconduct., 41) Golovashkin, A.I., Levchenko, I.S., Lobanov, N.N., Motulevich, G.P., “Properties of Films of Ternary Superconducting Nb3(AlSi) Alloy” (in Russian), Fiz. Met. Metalloved., 46(1), 45–49 (1978) (Crys. Structure, Morphology, Experimental, Superconduct., 7) Matukhin, V.L., Safin, I.A., Shamray, V.F., “Nuclear Quadrupole Resonance 93Nb in the Ternary Nb3AlBased Phases” (in Russian), Fiz. Met. Metalloved., 50(3), 526–532 (1980) (Morphology, Experimental, Electronic Structure, Magn. Prop., Superconduct., 12)
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[1987Ves]
[1989Kri] [1993Bae]
[1998Nan]
[1999Mur]
[2000Rio] [2001Mur1]
[2001Mur2]
[2003Man]
[2003Mat]
[2003Mur]
[2003Zha]
[2004Ban1]
[2004Ban2]
[2004Luk]
[2004Sha]
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Pan, V.M., Latysheva, V.I., Kulik, O.G., Popov, A.G., Litvinenko, E.N., “The Nb-NbAl3-Nb5Si3 Phase Diagram”, Russ. Metall. (Engl. Transl.), (4), 233–235 (1984), translated from Izv. Akad. Nauk SSSR, Met., (4), 225–226 (1984) (Phase Diagram, Experimental, *, 6) Vesnin, Yu.I., Starikov, M.A., “On the Features of the Superconductivity of the Nb-Al-X Solid Solutions with the A15 Structure” (in Russian), Doklady Akademii Nauk SSSR, USSR, 296(1-3), 98–100 (1987) (Crys. Structure, Experimental, Magn. Prop., Superconduct., 11) Krivko, N.I., “Investigation of a Superconductor-Silicon Interface” (in Russian), Fizika Tverdogo Tela (USSR), 31(6), 225–230 (1989) (Morphology, Experimental, Interface Phenomena, Superconduct., 7) Baetzner, C., Beuers, J., Hoch, M., “Aluminium - Niobium - Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, Document ID: 10.16068.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, 13) Nanko, M., Takahashi, A., Ogura, T., Kitahara, A., Yanagihara, K., Maruyama, T., “Formation of Intermetallic Compounds on Refractory Metals in Aluminum-Silicon Liquid”, Materials Science and Engineering Serving Society. Proceedings of the Third Okinaga Symposium on Materials Science and Engineering Serving Society, Elsevier Science, Amsterdam, Netherlands, 299–302 (1998) (Morphology, Experimental, Interface Phenomena, 7) cited from abstract Murayama, Y., Hanada, S., “Effect of (Si + Al) Content in Nb Solid Solution on Mechanical Properties of Multiphase Nb-Si-Al Alloys” (in Japanese), J. Jpn. Inst. Met., 63(12), 1519–1526 (1999) (Morphology, Experimental, Mechan. Prop., 18) cited from abstract Rios, C.T., Milenkovic, S., Caram, R., “Directional Growth of Al-Nb-X Eutectic Alloys”, J. Cryst. Growth, 211, 466–470 (2000) (Phase Relations, Experimental, 8) Murakami, T., Sasaki, S., Ichikawa, K., Kitahara, A., “Microstructure, Mechanical Properties and Oxidation Behavior of Nb-Si-Al and Nb-Si-N Powder Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 9(7), 621–627 (2001) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *, 17) Murakami, T., Sasaki, S., Ichikawa, K., Kitahara, A., “Oxidation Resistance of Powder Compacts of the Nb-Si-Cr System and Nb3Si5Al2 Matrix Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 9(7), 629–635 (2001) (Morphology, Phase Relations, Experimental, Kinetics, Mechan. Prop., Phys. Prop., *, 17) Manna, I., Chattopadhyay, P.P., Banhart, F., Fecht, H.-J., “Solid State Synthesis of Al-Based Amorphous and Nanocrystalline Al-Nb-Si and Al-Zr-Si Alloys”, Z. Metallkd., 94(7), 835–841 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, 24) Matsuura, K., Koyanagi, T., Ohmi, T., Kudoh, M., “Aluminide Coating on Niobium by Arc Surface Alloying”, Mater. Trans., JIM, 44(5), 861–865 (2003) (Morphology, Phase Relations, Experimental, Interface Phenomena, 14) Murakami, T., Sasaki, S., Ito, K., “Oxidation Behavior and Thermal Stability of Cr-Doped Nb(Si, Al)2 and Nb3Si5Al2 Matrix Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 11(3), 269–278 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, *, 28) Zhao, J.-C., Peluso, L.A., Jackson, M.R., Tan, L., “Phase Diagram of the Nb-Al-Si Ternary System”, J. Alloys Compd., 360, 183–188 (2003) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Interface Phenomena, #, 36) Banno, N., Takeuchi, T., Kikuchi, A., Iijima, Y., Inoue, K., Yuyama, M., Wada, H., “First Trial to Fabricate Nb3(Al, Si) Multifilamentary Superconductors by Rapid-Heating and Quenching (RHQ) Process”, AIP Conference Proceedings (USA), 711(2), 515–522 (2004) (Morphology, Experimental, Magn. Prop., Superconduct., 7) cited from abstract Banno, N., Takeuchi, T., Kikuchi, A., Iijima, Y., Inoue, K., Yuyama, M., Wada, H., “Multifilamentary Nb/Al-Ge and Nb/Al-Si Precursor Fabrication Using the Intermediately Rapid Heating and Quenching Technique”, Superconduct. Sci. Techn., 17(3): 320–326 (2004) (Morphology, Experimental, 5) cited from abstract Lukas, H.L., Lebrun, N., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 20.14887.1.20, (2004) (Crys. Structure, Phase Diagram, Assessment, 29) Shao, G., “Thermodynamic Assessment of the Nb-Si-Al System”, Intermetallics, 12(6), 655–664 (2004) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, #, 38)
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[2006Rag] [2007Qu]
[Mas2] [V-C2]
Al–Nb–Si Raghavan, V., “Al-Nb-Si-Ti (Aluminum-Niobium-Silicon-Titanium)”, J. Phase Equilib. Diffus., 26(6), 638 (2005) (Phase Relations, Review, 6) Matsuura, K., Kata, D.B., Lis, J.T., Kudoh, M., “Grain Refinement and Improvement in Mechanical Properties of Nb-Al-Si Intermetallic Alloys”, ISIJ Int., 46(6), 875–879 (2006) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 7) Raghavan, V., “Al-Nb-Si (Aluminum-Niobium-Silicon)”, J. Phase Equilib. Diffus., 27(2), 163–165 (2006) (Crys. Structure, Phase Diagram, Phase Relations, Review, 9) Qu, S., Han, Y., Song, L., “Effects of Alloying Elements on Phase Stability in Nb-Si System Intermetallics Materials”, Intermetallics, 15(5-6), 810–813 (2007) (Morphology, Phase Relations, Experimental, 10) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Nickel – Vanadium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Frederick H. Hayes, Peter Rogl, Eberhard Schmid, updated by Viktor Kuznetsov
Introduction A large number of studies of the system was devoted to alloying behavior of V in the Ni aluminides, mainly γ’Ni3Al and βNiAl [1959Gua, 1977Pel, 1983Och, 1984Och1, 1984Och2, 1988Mor, 1989Hon, 1994Jia, 1997Pri, 2001Zap, 2004Ish]. Some studies of the system were also motivated by the possible influence of V on catalytic action of Al rich nickel aluminides [1971Mya, 1977Mya]. Data on the formation of intermetallics in the system are briefly reviewed in [1990Kum]. [2005Rag] summarized recent publications on phase equilibria. Phase equilibria are studied fairly well. There exist data for primary fields of crystallization of phases [1977Mya], for the liquidus, solidus [2001Miu] and solvus [1989Hon, 1991Mis] surfaces in the Ni rich region, as well as several complete [1977Mya] and partial [1991Cot, 1997Pri] isothermal sections. Two vertical sections NiAl3-VAl3 and Ni2Al3-V5Al8 were studied by [1971Mya], and the section between Ni3Al and VNi3 by [1984Gup]. The quasibinary eutectic that exists on the NiAl-V section, was studied by [1977Pel, 1991Kim, 2000Mil, 2002Mil]. Thermodynamic data are restricted to the low-temperature heat capacity data [1999Dar] and enthalpy of formation of liquid on the section V/Ni = 3/7, presented by [2005Sud]. Some mechanical properties of alloyed γ’Ni3Al and βNiAl were studied by [1985Ino, 1988Dar, 1991Hay, 1991Wu]. The kinetic simulation of ordering process in the alloy Ni-6.6Al-15.1V (at.%) at 800˚C was performed by [2001Par]. A series of works of chinese authors [2003Zha, 2004Zha, 2005Hou, 2005Li1, 2005Li2, 2005Li3, 2005Zha, 2006Li, 2007Li1, 2007Li2] who used phase field model is devoted to theoretical studies of kinetics of transformations between γ, γ’Ni3Al and VNi3 phases and the resulting morphology. Experimental investigations of phase relations, structures and thermodynamics of phases are summarized in Table 1.
Binary Systems The binary phase diagram for Al-V was taken from [2000Ric] and for Ni-V from [Mas2, 1982Smi]. However, it shall be noted that a recent reinvestigation of the Al-V binary system by [2000Ric] revealed a significantly lower peritectic formation temperatures of 1408˚C for V5Al8 and 1270˚C for VAl3 than those accepted by [Mas2] (1670˚C and 1360˚C, respectively). For the Al-Ni binary, the latest version [2004Sal] evaluated within the MSIT Binary Evaluation Program is accepted. Landolt‐Bo¨rnstein New Series IV/11E1
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Solid Phases V substitutes for Al to a considerable degree both in NiAl [1965Ram, 1977Mya, 1991Cot, 1997Pri] and in Ni3Al [1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas, 1989Hon]. The solubility on Ni in the vanadium aluminides is very small [1971Mya, 1981Yin]. No ternary phases have been definitely reported for this system. In particular, [1975Mar] found for the alloy of VNi2Al composition, annealed at 500˚C, a two-phase structure, though with the possible participation of a phase with the MnCu2Al (cF16) structure (X-ray data were inconclusive). A phase with this structure, but with a totally different lattice space (580.31 pm as compared with 633.0 pm after [1975Mar]) was reported by [1999Dar] in an alloy annealed at 800˚C. Moreover, [2004Ish] reports a second-order transition with the tricritical point between the phases with CsCl (cP2) and MnCu2Al (cF16) structures in the Al-Ni-Ti-V quaternary system. Nevertheless, for the ternary system the present data seem not to be conclusive, and the existence of this phase needs more definitive confirmation. [1984Lia] did not manage to prepare a VNi6Al phase (neither stable or metastable), although similar phases MNi6Al exist in the systems with Nb and Ta (metastable NbNi6Al and stable TaNi6Al). Data on the solid phases are listed in Table 2.
Invariant Equilibria A partial reaction scheme given in Fig. 1, which is complete for the Ni rich corner, is based largely on the fields of primary crystallization determined by [1977Mya] and on the findings of [1977Pel] who observed a quasibinary eutectic at 1360˚C and 40 at.% V on the NiAl to V section. The existence of the latter was confirmed also by [1991Cot, 2000Mil]. Between this maximum and the Ni-V binary edge there are three invariant reactions: two transition reactions, U1 and U2, and a ternary eutectic reaction at 1180˚C in which the Ni rich liquid containing 15 at.% Al and 32 at.% V is, on the basis of the solubility data of [1959Gua, 1984Gup], in equilibrium with the Ni rich solid solution containing 7 at.% Al, 34 at.% V, NiAl containing 21 at.% Al, 28 at.% V, and σ phase containing 5 at.% Al, 53 at.% V. At the Al rich side of the saddle point there is a sequence of transition reactions towards the Al rich corner. Two of them, U3 and U4 can be identified from [1977Mya]. In both Fig. 1 and Fig. 2 the σ/σ’transition of the Ni-V σ phase [Mas2] has been disregarded. Unfortunately, the data for the compositions of the phases are nearly completely absent, so the invariant equilibria cannot be tabulated.
Liquidus, Solidus and Solvus Surfaces Figure 2 gives the primary crystallization field boundaries based on [1977Mya, 1977Pel]. Figures 3 and 4 present the data of [2001Miu] showing the dependence of the γ liquidus and solidus temperatures on the variation of Al at the parametric V content, and on the variation of V at the parametric Al content, respectively. Figure 5 gives the (Ni)-solvus curves for the Ni rich corner from [1989Hon, 1991Mis]. Beyond the solubility limit the two-phase field γ+γ’ is entered (see also Fig. 8 and Fig. 9 below). In the V rich end of Fig. 5 the precipitation of the congruently transforming VNi3 is indicated. DOI: 10.1007/978-3-540-88053-0_13 ß Springer 2009
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Isothermal Sections A partial isothermal section at 1300˚C is given in Fig. 6 [1991Cot, 2005Rag]. A partial isothermal section at 1200˚C is presented in Fig. 7 [1997Pri, 2005Rag]. The isothermal sections at 1100 and 800˚C shown in Fig. 8 and Fig. 9, respectively, are based primarily on the results of [1977Mya] modified to be consistent with the binaries, to exhibit close agreement with the X-ray results of [1965Ram] concerning equilibria among (V), NiAl and the σ phase at 800˚C, and to be consistent with the results of [1989Hon] for the Ni rich corner at both temperatures. In addition, the boundaries and extension limit of γ’Ni3Al are accepted from [1983Och, 1984Gup]. The results of [1994Jia] for the equilibria between γ, γ’Ni3Al and βNiAl phases, which are presented in a tabular form, are reproduced in Tables 3 and 4.
Temperature – Composition Sections The section βNiAl-V contains quasibinary eutectic at 40 at.% V, 1360˚C [1977Pel, 1991Cot, 2000Mil]. However at lower temperatures the composition of phases is shifting away from the join, as may be seen from the isothermal section above, so the section cannot be quasibinary. The vertical sections VAl3-NiAl3 and V5Al8-Ni2Al3 are presented in Figs. 10 and 11, respectively. Both are taken from [1971Mya]. Figure 12 gives the 75 at.% Ni isopleth from [1989Hon] which contains a saddle point for the eutectoid decomposition of the (Ni) solid solution into VNi3 and Ni3Al on cooling.
Thermodynamics [1999Dar] performed low-temperature (3.2 to 10.3 K) measurements of the heat capacity of the VNi2Al phase. The results, when treated in the standard way (Cp(T ) = γelT + CD(θ/T)), give γel = 14.17±0.10 mJ·mol–1·K2, θD = 359±1.9 K. This equation is valid only below 7 K. It should be noted that the existence of this phase is questionable (see Solid Phases), so the result may really correspond to the βNi(Al0.5V0.5) phase. The enthalpies of formation of liquid along the section xV/xNi = 3/7 are presented on Fig. 13 taken from [2005Sud].
Notes on Materials Properties and Applications The possibility of plastifying of γ’Ni3Al and βNiAl by V additions was considered by [1985Ino, 1988Dar, 1991Wu] from both theoretical and experimental points of view. [1991Cot, 1991Kim, 1994Ben, 2000Mil, 2002Mil] studied morphology and strength of the directionally solidified quasibinary eutectic βNiAl-V. Creep behavior of off-stoichiometric binary and ternary γ’Ni3Al was studied by [1991Hay]. [1995Ino] measured the strength of an intermetallic phase with nanoinclusions of amorphous phase.
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Miscellaneous [1985Mor] used the system among the tests for their method of predicting solubility limits in the γ phase. [1988Mor] investigated static displacements of atoms in the alloyed γ’Ni3Al compound. [1994Ben] studied the evolution of morphology of phases in the near-eutectoid alloy Ni5Al20V (at.%) at 650 to 900˚C. [1997Kai, 2001Kai] discovered the formation of abrupt compositional changes in the βNiAl phase grown in diffusion pairs Ni0.8Al0.2+V (as well as several other third components). [2000Boz, 2002Boz] suggested a simple theoretical scheme for prediction of alloying behavior of third elements in βNiAl. [2001Sav] investigated experimentally the ordering kinetics in the alloy Ni75Al21V4 (at.%). [2001Son] performed ab initio analysis of site preference of V in βNiAl. [2001Ter] suggested usage of thermal conductivity measurements for determination of site preferences in γ’ Ni3Al phase. [2003Oza] tested hydrogen permeability of fcc Al-Ni-V alloys. [2005Cos] suggested using atomic force microscopy as a tool for studying the morphology of eutectics and used it for study of NiAl-V one. [2006Tan] investigated role of antiphase boundaries in kinetics of the L12→DO22 transformation.
. Table 1 Investigations of the Al-Ni-V Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/ Phase Range Studied
[1959Gua]
Microhardness measurement, XRD
Ni3Al + up to 8 at.% V, 1150˚C
[1965Ram]
XRD
25 to 70 at.% Al, 25 and 50 at.%V (6 compositions), 800˚C
[1971Mya]
Metallography, DTA, XRD, hardness and microhardness measurements
Sections NiAl3-VAl3 and Ni2Al3V5Al8, 600 to 1400˚C
[1975Mar]
XRD
VNi2Al composition, 500˚C
[1977Mya]
DTA, XRD, metallography, density, hardness and 800 to 1100˚C electrical resistivity measurements
[1977Pel]
DTA, SEM, directional solidification
Section NiAl-V (3 compositions)
[1981Yin]
XRD
Section VAl3-NiAl3, room temperature
[1983Och]
XRD
Ni3AlxV1–x, x = 1 to 0.4, 1000˚C
[1984Gup]
DSC, XRD
Ni3AlxV1–x, x = 0 to 1, 1000˚C
[1984Lia]
Melt quenching, XRD
VNi6Al composition Ni3Al+ up to 8 at.% V, 1000˚C
[1984Och1] XRD
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. Table 1 (continued) Reference
Temperature/Composition/ Phase Range Studied
Method/Experimental Technique
[1989Hon]
DTA, metallography and SEM-EDX analysis
Up to 25 at.% V + Al, 1000 to 1320˚C
[1991Cot]
Metallography, DTA, TEM, SEM, EPMA
NiAl + up to 20 at.% V, 1300 to 1370˚C
[1991Mis]
DTA, metallography and SEM-EDX analysis
Up to 25 at.% V + Al, 1000 to 1320˚C
[1994Jia]
Diffusion pairs technique, metallography, EPMA 2 to 6, 13 and 31 mass% Al, 1 to 5 mass% V, 800 to 1300˚C
[1997Pri]
Metallography, microhardness measurement, EPMA
20 to 30 at.% Ni, 9.5 to 22.5 at.% Al, 1200˚C
[1999Dar]
Adiabatic calorimetry
VNi2Al composition, 1.2 to 10 K
[2001Miu]
DTA
0 to 16 at.% Al, 0 to 17 at.% V
[2005Sud]
Calorimetry
Liquid phase, V/Ni = 3/7, up to 16 at.% V
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Al) < 660.452
cF4 Fm 3m Cu
a = 404.96 a = 408.4
pure Al at 25˚C [Mas2] at 0.17 at.% V [V-C]
γ, (Ni) < 1455
cF4 Fm 3m Cu
a = 352.40
at 25˚C [Mas2]
(V) < 1910
cI2 Im 3m W
a = 302.40 a = 307.55
at 25˚C [Mas2] at 50 at.% Al [V-C2]
NiAl3 < 856
oP16 Pnma Fe3C
a = 661.3 ± 0.1 b = 736.7 ± 0.1 c = 481.1 ± 0.1
[2004Sal]
Ni2Al3 < 1138
hP5 P3ml Ni2Al3
a = 402.8 c = 489.1
[2004Sal]
βNiAl < 1651
cP2 Pm 3m CsCl
a = 288.72 ± 0.02 a = 287.98 ± 0.02
at 50 at.% Ni [2004Sal] at 54 at.% Ni [2004Sal]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Ni5Al3 < 723
oC16 Cmmm Pt5Ga3
a = 753 b = 661 c = 376
at 63 at.% Ni [2004Sal]
γ’Ni3Al < 1372
cP4 Pm 3m AuCu3
a = 357.18 ± 0.02
[V-C2]
VAl3 < 1360
tI8 I4/mmm TiAl3
a = 378.0 c = 832.2
[V-C2]
V5Al8 < 1673
cI52 I 43m Cu5Zn8
a = 923.4 ± 0.5
[V-C2]
V4Al23 < 736
hP54 P63/mmc V4Al23
a = 769.28 c = 1704
[1989Mur]
VNi3 < 1045
tI8 I4/mmm TiAl3
a = 354.3 c = 720.2 a = 354.1 c = 721.8
at 23.44 at.% V [P]
VNi2 < 922
oI6 Immm MoPt2
-
[Mas2]
σ < 1280
tP30 P42/mnm σCrFe
a = 895.4 c = 463.5 a = 899.6 c = 465.3
at 57.5 at.% V [1982Smi]
cP8 Pm 3n Cr3Si
a = 471.2
Phase/ Temperature Range [˚C]
V3Ni ≲900
Comments/References
at 25.60 at.% V [P]
at 63.2 at.% V [1982Smi] [1982Smi]
. Table 3 Equilibrium Compositions of the γ and γ’Ni3Al Phases and V Partition Coefficients [1994Jia] γ phase
Temperature [˚C]
γ’Ni3Al phase
V, at.%
Al, at.%
V, at.%
Al, at.%
Partition coefficient kVγ’/γ
1300
3.05
16.6
3.09
19.6
1.01
1200
0.45
17.7
0.89
22.3
1.98
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. Table 3 (continued) γ phase
Temperature [˚C]
γ’Ni3Al phase
V, at.%
Al, at.%
V, at.%
Al, at.%
Partition coefficient kVγ’/γ
0.54
15.3
0.83
21.1
1.53
3.38
12.6
4.27
18.5
1.26
1000
1.93
12.3
3.85
17.6
1.99
900
3.15
9.9
4.73
17.0
1.50
3.64
9.0
5.50
16.4
1.51
1.95
10.4
3.43
19.3
1.76
1100
800
. Table 4 Equilibrium Compositions of the γ’Ni3Al and βNiAl Phases and V Partition Coefficients [1994Jia] γ’Ni3Al phase
βNiAl phase
V, at.%
Al, at.%
V, at.%
Al, at.%
Partition coefficient kVβ/γ’
1300
2.17
24.0
0.98
32.4
2.21
1100
2.81
24.1
0.74
35.4
3.80
5.29
21.2
2.56
35.0
2.07
Temperature [˚C]
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. Fig. 1 Al-Ni-V. Partial reaction scheme
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. Fig. 2 Al-Ni-V. Liquidus surface projection
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. Fig. 3 Al-Ni-V. Dependence of the γ liquidus and solidus temperatures on Al variation at the parametric V content
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. Fig. 4 Al-Ni-V. Dependence of the γ liquidus and solidus temperatures on V variation at the parametric Al content
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. Fig. 5 Al-Ni-V. The γ (Ni) solvus surface
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. Fig. 6 Al-Ni-V. Partial isothermal section at 1300˚C
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. Fig. 7 Al-Ni-V. Partial isothermal section at 1200˚C
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. Fig. 8 Al-Ni-V. Isothermal section at 1100˚C
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. Fig. 9 Al-Ni-V. Isothermal section at 800˚C
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. Fig. 10 Al-Ni-V. Isopleth at 75 at.% Al
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. Fig. 11 Al-Ni-V. Vertical section V5Al8-Ni2Al3
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. Fig. 12 Al-Ni-V. Isopleth at 75 at.% Ni
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. Fig. 13 Al-Ni-V. Enthalpy of formation of liquid along the section xV/xNi = 3/7
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References [1959Gua] [1965Ram]
[1971Mya]
[1975Mar]
[1977Mya]
[1977Pel]
[1981Yin]
[1982Smi] [1983Och] [1984Gup]
[1984Lia] [1984Och1]
[1984Och2] [1985Ino]
[1985Mis]
[1985Mor] [1985Nas] [1988Dar]
[1988Mor]
Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al (γ’ Phase)”, Trans. Metall. Soc. AIME, 215, 807–814 (1959) (Experimental, Phase Diagram, Phase Relations, 27) Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3. III. Investigations in Several T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56, 99–104 (1965) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 14) Myasnikova, K.P., Ponomareva, L.F., Pryakhina, L.I., Marshakov, I.K., “Examination of Alloys of the Systems NiAl3-VAl3 and Ni2Al3-V5Al8”, Russ. Metall. (Engl. Transl.), (1), 126–128 (1971), translated from Izv. Akad. Nauk SSSR, Met., (1), 186–188 (1971) (Experimental, Phase Relations, #, 5) Marazza, R., Ferro, R., Rambaldi, G., “Some Phases in Ternary Alloys of Titanium, Zirconium and Hafnium, with a MgAgAs or AlCu2 Type Structure”, J. Less-Common Met., 39(2), 341–345 (1975) (Crys. Structure, Experimental, Morphology, 11) Myasnikova, K.P., Markiv, V.Ya., Pryakhina, L.I., Motrychuk, G.Yu., “Phase Equilibria in the V-Ni-Al System and Some Alloy Properties”, Russ. Metall. (Engl. Transl.), (3), 192–199 (1977), translated from Izv. Akad. Nauk SSSR, Met., (3), 222–229 (1977) (Experimental, Phase Diagram, Phase Relations, Phys. Prop., #, 12) Pellegrini, P.W., Hutta, J.J., “Investigations of Phase Relations and Eutectic Directional Solidification in NiAl-V Join”, J. Cryst. Growth, 42, 536–539 (1977) (Experimental, Morphology, Phase Relations, Phys. Prop., *, 3) Ying-Hong, Z., Jing-Qi, L., Jiang-Xuang, Z., Cheng, C.S., “A Room-Temperature Section of the Phase Diagram of TiAl3- VAl3-MAl3 of the System Alloys of Al-Ti-V-M (M = Ni, Fe)”, Acta Phys. Sin. (Chin. J. Phys.), 30, 972–975 (1981) (Phase Relations, Experimental, 9) Smith, J.F., Carlson, O.N., Nash, P.G., “The Ni-V (Nickel-Vanadium) System”, Bull. Alloy Phase Diagrams, 3, 342–348 (1982) (Phase Diagram, Review, Phase Relations, Crys. Structure, #, 35) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull. P.M.E. (T.I.T.), 52, 1–17 (1983) Phase Relations, 7) Gupta, A., Horton, J.A., Liu, C.T., “Phase Formation and Stability in the Pseudobinary Ni3Al-Ni3VAlloy System”, Metals Soc. AIME, Conf: High-Temp. All., Bethesda, Maryland, USA, 115–123 (1984) (Phase Relations, Experimental, Crys. Structure, Morphology, 10) Liang, W.W., Standley, R., Nash, P., Skowron, M., “The Relative Stabilities of the Ni6AlX (X = V, Nb, Ta) Phases”, J. Mater. Sci. Lett., 3(3), 259–261 (1984) (Experimental, Phase Relations, 5) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions”, Bull. Res. Lab. Precis. Machin. Electron., Tokyo Inst. Technol., 53, 15–28 (1984) (Crys. Structure, Experimental, 66) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289–298 (1984) (Experimental, Phase Diagram, Phase Relations, 90) Inoue, A., Masumoto, T., Tomioka, H., Yano, N., “Microstructure and Mechanical Properties of Ductile Intermetallic Compounds Produced by Melt Quenching”, Int. J. Rapid Solidification, 1, 115–142 (1985) (Morphology, Mechan. Prop., Review, 28) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33, 1161–1169 (1985) (Experimental, Crys. Structure, 64) Morinaga, M., Yukawa, N., Ezaki, H., Adachi, H., “Solid Solubilities in Nickel-Based F.C.C. Alloys”, Philos. Mag. A, 51(2), 247–252 (1985) (Phase Relations, Theory, 26) Nash, P., “Nickel-Base Intermetallics for High Temperature Alloy Design”, Mater. Res. Soc. Conf.: HighTemp. Ordered Intermet. Alloys, Boston, 423–427 (1985) (Review, Phase Diagram, Phase Relations, 15) Darolia, R., Lahrman, D.F., Field, R.D., Freeman, A.J., “Alloy Modeling and Experimental Correlation for Ductility Enhancement in NiAl”, High-Temperature Ordered Intermetallic Alloys III, Mater. Res. Soc. Symp. Proc., 133, 113–118 (1989) (Mechan. Prop., Experimental, Theory, 14) Morinaga, M., Sone, K., Kamimura, T., Ohtaka, K., Yukawa, N., “X-Ray Determination of Static Displacements of Atoms in Alloyed Ni3Al”, J. Appl. Crystallogr., 21, 41–46 (1988) (Crys. Structure, Experimental, 18)
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[1989Mur] [1990Kum] [1991Cot] [1991Hay]
[1991Kim]
[1991Mis]
[1991Wu]
[1994Ben]
[1994Jia]
[1995Ino]
[1997Kai]
[1997Pri] [1999Dar]
[2000Boz] [2000Mil]
[2000Ric]
[2001Kai]
[2001Miu]
[2001Par] [2001Sav]
Al–Ni–V Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of γ’ Solvus in Ni-Al-X Ternary Systems”, Mater. Res. Soc. Symp. Proc., 133, 429–440 (1989) (Experimental, Phase Diagram, Phase Relations, #, 35) Murray, J.L., “The Al-V (Aluminum-Vanadium) System”, Bull. Alloy Phase Diagrams, 10, 351–357 (1989) (Phase Diagram, Crys. Structure, Review, Phase Relations, #, 34) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35(6), 293–327 (1990) (Crys. Structure, Phase Diagram, Review, 158) Cotton, J.D., Kaufman, M.J., Noebe, R.D., “Constitution of Pseudobinary Hypoeutectic β-NiAl+α-V Alloys”, Scr. Metall. Mater., 25, 1827–1832 (1991) (Phase Relations, Mechan. Prop., #, 9) Hayashi, T., Shinoda, T., Mishima, Y., Suzuki, T., “Effect of Off-Stoichiometry on the Creep Behavior of Binary and Ternary Ni3Al”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 617–622 (1991) (Experimental, Mechan. Prop., 7) Kim, Y.D., Wayman, C.M., “Effect of Vanadium Transformation Behavior and Martensite Morphology in a Ni-Al Shape Memory Alloy”, Scr. Metall. Mater., 25(8), 1863–1868 (1991) (Abstract, Crys. Structure, 0) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the γ Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123–130 (1991) (Assessment, Experimental, Phase Diagram, Phase Relations, #, 5) Wu, Y.P., Sanchez, J.M., Tien, J.K., “Effect of APB Microsegregation on the Strength of Ni3Al with Ternary Additions”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 87–94 (1991) (Mechan. Prop., Calculation, 22) Bendersky, L.A., Biancaniello, F.S., Williams, M.E., “Evolution of the Two-Phase Microstructure L12 + D022 in Near-Eutectoid Ni3(Al,V) Alloy”, J. Mater. Res., 9(12), 3068–3082 (1994) (Morphology, Experimental, 16) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (γ (A1), γ’ (L12) and β (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473–485 (1994) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 25) Inoue, A., Kimura, H., Sasamori, K., Masumoto, T., “High Strength Al-V-M (M = Fe, Co or Ni) Alloys Containing High Volume Fraction of Nanoscale Amorphous Precipitates”, Mater. Trans. JIM, 36, 1219–1228 (1995) (Mechan. Prop., Experimental, 21) Kainuma, R., Ikenoya, H., Ohnuma, I., Ishida, K., “Pseudo-Interface Formation and Diffusion Behaviour in the B2 Phase Region of NiAl-Base Diffusion Couples”, Def. Diffus. Forum, 143-147, 425–430 (1997) (Crys. Structure, Experimental, Phase Relations, Phys. Prop., 10) Prima, S.B., Morozova, E.A., Bega, N.D., “Phase Equilibria in Vanadium-rich Alloys of the V-N-Al System”, Powder Metall. Met. Cer., 36(7-8), 390–393 (1997) (Experimental, Phase Relations, #, 6) da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., da Silva, C.M., Gomes, A.A., “Specific Heat and Electronic Structure of Heusler Compounds Ni2TAl (T = Ti, Zr, Hf, V, Nb, Ta)”, Physica B (Amsterdam), 269, 154–162 (1999) (Crys. Structure, Experimental, 20) Bozzolo, G., Noebe, R.D., Honecy, F., “Modeling of Ternary Element Site Substitution in NiAl”, Intermetallics, 8, 7–18 (2000) (Crys. Structure, Review, 34) Milenkovic, S., Coelho, A.A., Caram, R., “Directional Solidification Processing of Eutectic Alloys in the Ni-Al-V System”, J. Cryst. Growth, 211(1-4), 485–490 (2000) (Crys. Structure, Experimental, Morphology, 13) Richter, K.W., Ipser, H., “The Al-V Phase Diagram between 0 and 50 Atomic Percent Vanadium”, Z. Metallkd., 91(5), 383–388 (2000) (Crystal Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 13) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of γ’/β Interface Morphologies in NiAl-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168–175 (2001) (Experimental, Phase Relations, Kinetics, 21) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: V, Nb and Ta) Ternary Systems”, J. Phase Equilib., 22, 345–351 (2001) (Experimental, Phase Diagram, Phase Relations, #, 9) Pareige, C., Blavette, D., “Simulation of the FCC -> FCC+L12+DO22 Kinetic Reaction”, Scr. Mater., 44 (2), 243–247 (2001) (Experimental, Kinetics, Phase Relations, 8) Savin, O.V., Stepanova, N.N., Akshentsev, Yu.N., Rodionov, D.P., “Ordering Kinetics in Ternary Ni3AlX Alloys”, Scr. Mater., 45(8), 883–888 (2001) (Crys. Structure, Electr. Prop., Experimental, Kinetics, 18)
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[2004Sal]
[2004Zha] [2005Cos]
[2005Hou]
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[2005Li2] [2005Li3]
[2005Rag] [2005Sud]
[2005Zha]
[2006Li]
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Song, Y., Guo, Z.X., Yang, R., Li, D., “First Principles Study of Site Substitution of Ternary Elements in NiAl”, Acta Mater., 49, 1647–1654 (2001) (Calculation, Crys. Structure, Electronic Structure, 17) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8), 2314–2320 (2001) (Calculation, Crys. Structure, Experimental, Transport Phenomena, 63) Zapolsky, H., Pareige, C., Marteau, L., Blavette, D., Chen, L.Q., “Atom Probe Analyses and Numerical Calculation of Ternary Phase Diagram in Ni-Al-V System”, Calphad, 25(1), 125–134 (2001) (Calculation, Phase Relations, Thermodyn., 14) Milenkovic, S., Caram, R., “Microstructure of the Microalloyed NiAl-V Eutectics”, Mater. Lett., 55(12), 126–131 (2002) (Experimental, Kinetics, Morphology, 9) Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2 Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Review, 27) Ozaki, T., Zhang, Y., Komaki, M., Nishimura, Ch., “Hydrogen Permeation Characteristics of V-Ni-Al Alloys”, Int. J. Hydrogen Energy, 28, 1229–1235 (2003) (Crys. Structure, Electrochemistry, Experimental, 8) Zhao, Y.H., Hou, H., Xu, H., Wang, Y.X., Chen, Z, Sun, X.D., “Atomic Scale Computer Simulation for Early Precipitation Process of Ni75Al6V19 alloy”, J. Mater. Sci. Techn., 19(Suppl. 1), 17–19 (2003) (Kinetics, Morphology, Calculation, 6) Ishikawa, K., Ohnuma, I., Kainuma, R., Aoki, K., Ishida, K., “Phase Equilibria and Stability of HeuslerType Aluminides in the NiAl-Ni2AlTi-Ni2AlY (Y: V, Cr or Mn) Systems”, J. Alloys Compd., 367(1-2), 2–9 (2004) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 20) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.10238.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 164) Zhao, Y.H., Chen, Z., Wang, Y.X., Lu, Y.L., “Atomic-Scale Computer Simulation for Early Precipitation Process of Ni75Al10V15 Alloy”, Progr. in Natural Sci., 14, 241–246 (2004) (Kinetics, Calculation, 16) Costa, C.A.R., Batista, W.W., Rios, C.T., Milenkovic, S., Goncalves, M.C., Caram, R., “Eutectic Alloy Microstructure: Atomic Force Microscopy Analysis”, Appl. Surf. Sci., 240, 414–423 (2005) (Morphology, Experimental, 17) Hou, H., Zhao, Y.H., Chen, Z., Xu, H., “Prediction for the Early Precipitation Process of Ni75AlxV25–x System with Lower Al Concentration by the Phase-Field Model”, Acta Metallurgica Sinica, 41, 695–702 (2005) (Kinetics, Calculation, 18) Li, Y.-S., Chen, Z., Wang, Y.-X., Lu, Y.-L., “Computer Simulation of γ’ and θ Phase Precipitation of NiAl-V Alloy Using Microscopic Phase-Field Method”, Trans. Nonferrous Met. Soc. China, 15, 57–63 (2005) (Kinetics, Morphology, Calculation, 20) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., “Computer Simulation of the Interphase Boundary Evolution in Ni75AlxV25–x Alloy”, J. Mater. Sci. Techn., 21, 395–398 (2005) (Kinetics, Calculation, 20) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., Zhang, J.J, “Microscopic Phase-Field Simulation for Nucleation Incubation Time of Ni75AlxV25–x Alloy”, J. Cent.-South Univ. Technol., 12, 635–640 (2005) (Kinetics, Calculation, 16) Raghavan, V., “Al-Ni-V (Aluminum-Nickel-Vanadium)”, J. Phase Equilib. Diffus., 26, 273–275 (2005) (Crys. Structure, Phase Diagram, Phase Relations, Review, 15) Sudavtsova, V.S., Makara, V.A., Kudin, V.G., “Part 3 (Alloys of Nickel and Tin, Methods of Modeling and Prognosis of Thermodynamic Properties), Ch. 6. Thermodynamic Properties and Phase Equilibria of Nickel Alloys” (in Ukrainian) in “Thermodynamics of Metallurgical and Welding Melts”, “Logos” Publ., Kiev, 2005 (Thermodyn., Review, 192) Zhao, Y.H., Ju, D.Y., Hou, H., “Atomic-Scale Computer Simulation of Mixture Precipitation Mechanism for Ni75AlxV25–x Alloy”, Mater. Sci. Forum, 475-479, 3115–3118 (2005) (Calculation, Kinetics, Phase Relations, 4) Li, Y., Chen, Z., Lu, Y., Wang, Y., Chu, Z., “Computer Simulation of Ordered Interphase Boundary Structure of Ni-Al-V Alloy Using Microscopic Phase-Field Method”, Rare Metal Mater. Eng., 35, 200–204 (2006) (Kinetics, Morphology, Calculation, 14) Tanimura, M., Koyama, Y., “The Role of Antiphase Boundaries in the Kinetic Process of the L12→D022 Structural Change of an Ni3Al0.45V0.50 Alloy”, Acta Mater., 54, 4385–4391 (2006) (Crys. Structure, Experimental, Kinetics, Morphology, 26)
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[Mas2] [V-C] [V-C2]
Al–Ni–V Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., Lai, Q.B., “Microscopic Phase-Field Simulation of Atomic Migration Characteristics in Ni75AlxV25–x Alloys”, Mater. Lett., 61, 974–978 (2007) (Calculation, Experimental, Phase Relations, 17) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., “Coarsening Kinetics of Intermetallic Precipitates in Ni75AlxV25–x Alloys”, J. Mater. Res., 22, 61–67 (2007) (Experimental, Kinetics, Morphology, Phase Relations, Thermodyn., 30) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Aluminium – Oxygen – Zirconium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Mireille Harmelin, updated by Olga Fabrichnaya
Introduction The ZrO2-Al2O3 system is of interest because of the excellent mechanical properties of these composites widely used as structural materials [2002Boc], implant materials and cutting tools [2004Bas, 2004Ker, 2004Lee, 2006San1, 2006San2], protective and thermal resistant coatings [2006Por, 2008Che]. Laminated Al2O3/(Al2O3-ZrO2) composites exhibited good tribological behavior [2000Tar, 2003Tos]. Nanocomposite materials obtained in this system find widespread applications in structural field because of their excellent mechanical properties [2006Jia, 2006Ran, 2007Kon]. Mesoporous nanocrystalline composites Al2O3-(Y-stabilized) ZrO2 also used in catalysts, as well as in adsorption, separation and photoelectric devices [2005Che]. Fully stabilized ZrO2 (6.5–9 mol% Y2O3) in its cubic form is used as electrolyte in solid oxide fuel cell (SOFC). The Al2O3 is used to improve thermo-mechanical properties of this material [1995Nav, 2005Cho]. The Al2O3-ZrO2 eutectic composites are prospective materials for SOFC and oxygen sensors working at high temperatures due to the combination of ionic conductivity with improved mechanical properties and corrosion resistance. The ZrO2-Al2O3 system is a part of the high-order system CaO-Al2O3-SiO2-UO2-ZrO2 important for modeling of chemical interactions between nuclear reactor core debris and concrete [1990Rel, 1993Bal]. The Al2O3-ZrO2 thin films deposited on Si substrate are promising candidates for gate dielectric in metal-metal oxide semiconductors [2001Mor, 2003Zhu, 2005Biz, 2006Biz]. Reinforcement of Al matrix alloys by Al2O3 and Al3Zr by direct melt reaction improves mechanical properties of these materials [2001Zha, 2003Zha2, 2007Zha]. Most of the experimental studies on phase equilibria, microstructure development and properties are devoted to the ZrO2-Al2O3 system. Isothermal sections for the Al-O-Zr system were constructed by [1977Guk, 1978Guk1, 1978Guk2, 1978Guk3] based on diffusion couple studies. No ternary phases have been observed. The equilibrium quasibinary section ZrO2-Al2O3 has been determined by [1932War, 1964Alp, 1967Alp, 1994Lak, 1997Lak, 2000Jer, 2005Kam] and from 20 mass% Al2O3 by [1968Cev]. The quasibinary section ZrO2-Al2O3 was calculated by [1979Doe, 1980Wei, 1990Rel, 1992Wu, 1993Bal, 2004Fab, 2006Lak] using the CALPHAD approach. The previous MSIT evaluation of the experimental studies up to 1986 was presented by [1987Har]. Table 1 contains summary of experimental and theoretical studies from 1986 up to the present time.
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Binary Systems The binary system Al-Zr is accepted according to the critical evaluation of [2004Sch]. The binary system O-Zr is accepted from [2006Wan]; information about high pressure phases is from [V-C2, 2005Oht, 1999Des]. The binary phase diagram of the Al-O system is accepted from [1992Tay]. Crystallographic data for solid phases are from [1985Wri] and from [1998Lev] and [V-C2] for metastable Al2O3 phases.
Solid Phases There are no ternary phases in the Al-O-Zr system. Monoclinic αZrO2 practically does not dissolve Al2O3. According to [1964Alp] the solubility of ZrO2 in Al2O3 is 0.83 mol%. [2000Jer] found that it is even smaller (0.008 mol% ZrO2). There is uncertainty in the solubility of Al2O3 in the tetragonal (βZrO2) and fluorite (γZrO2) structures of ZrO2. Maximal solubility of Al2O3 in the tetragonal phase βZrO2 is ranging from 1.3 to 7 mol% according to different experimental data: 7 mol% [1967Alp], 5 mol% [1997Lak] or 1.3 mol% [2000Jer]. [1997Lak] determined maximal solubility of Al2O3 in the fluorite phase γZrO2 to be 6 mol% based on thermal analysis data. [2006Lak] derived maximal solubility of Al2O3 in the tetragonal phase βZrO2 as 4.5 mol% and in the fluorite phase γZrO2 as 3 mol% by a Calphad type assessment. According to [1988Tan, 1994Bal, 1994Ish] non-equilibrium monoclinic, tetragonal and cubic phases could contain up to 20–40 mol% Al2O3. The tetragonal phase could have different ratio of a/c lattice parameters [1986Ady, 1988Tan, 1994Bal, 1994Ish]. A metastable tetragonal phase formed by diffusionless transformation on quick cooling is called t’ and its c/a ratio slightly exceed unity. In several works [1990Jay, 1991And, 1994Bal, 1995Nar, 2003Kin, 2007Pod], a formation of metastable Al2O3 phases, such as γ, δ or θ Al2O3, was observed during crystallization from amorphous and metastable solid solution of Zr(1–x)AlxO(2–x/2) or by rapid solidification. The γAl2O3 phase has a structure of spinel, the other phases δ and θ are the superstructures of γAl2O3. Investigation of structures of Al2O3 metastable phases was performed in [1998Lev] using TEM. A high temperature hexagonal phase referred to as the ε phase, was found by [1968Cev] above 1930˚C, at the Al2O3 side of the ZrO2-Al2O3 section, with a = 784.9 and c = 1618.3 pm. Crystallographic data for the solid phases are presented in Table 2.
Quasibinary Systems The ZrO2-Al2O3 quasibinary system was experimentally studied in several works [1932War, 1961Suz, 1964Alp, 1967Alp, 1968Cev, 1987Vol, 1997Lak, 2000Jer, 2005Kam]. Samples were melted in an Ar atmosphere by [1964Alp, 1967Alp] and in vacuum and tungsten crucibles by [1968Cev]. In all cases, temperature was measured with an optical pyrometer. [1994Lak, 1997Lak] studied liquidus in the ZrO2-Al2O3 system using derivative thermal analysis in air in solar furnace at the temperatures up to 3000˚C. Jerebtsov et al. [2000Jer] and Kamaev et al. [2005Kam] studied phase transformations in the ZrO2-Al2O3 system using DTA and XRD. There is a remarkable scatter in the eutectic temperature and composition obtained in the above mentioned experimental studies. The eutectic reaction was observed at 1920˚C by DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009
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[1932War], 1880˚C by [1964Alp, 1967Alp] and 1710±10˚C by [1968Cev]. The composition of the eutectic point was found to be 61.95 mol% Al2O3 by [1968Cev] and 63.93 mol% Al2O3 by [1964Alp, 1967Alp]. [1968Cev] observed no solid solution at low temperature between Al2O3 and ZrO2 on the Al2O3 side and [1964Alp, 1967Alp] found that less than 1 mass% ZrO2 goes into the solid solution in αAl2O3. According to [1964Alp, 1967Alp] about 7 mass% Al2O3 dissolves in solid ZrO2. Lakiza et al. [1994Lak, 1997Lak] determined a metatectic reaction between fluorite and tetragonal structures of ZrO2 and liquid at 2260˚C with the liquid composition of 20 mol% Al2O3 and the maximum solubility in ZrO2 of 5 mol%. The eutectic reaction was determined at 1860˚C and 63 mol% Al2O3 by [1994Lak, 1997Lak]. The results of [2000Jer] and [2005Kam] for the eutectic reaction are close to each other. The composition was obtained the same 64.47 mol% ZrO2 in both works and the temperature of 1866˚C [2000Jer] and 1861˚C [2005Kam]. The recommended phase diagram was derived by thermodynamic calculations taking into account phase equilibrium data obtained in the ZrO2-Al2O3 system together with the high-order system ZrO2-Al2O3-Y2O3 by [2006Lak]. The calculated phase diagram of the quasibinary system based on the thermodynamic description of [2006Lak] is presented in Fig. 1. It should be mentioned that the shape of the liquidus obtained in the works of [1964Alp, 2000Jer, 2005Kam] show a tendency to phase separation in the liquid state at temperatures slightly exceeding the liquidus line and compositions 55-45 mol% Al2O3.
Invariant Equilibria The ternary invariant equilibria of the quasibinary section ZrO2-Al2O3 are given in Table 3.
Isothermal Sections [1977Guk, 1978Guk1, 1978Guk2, 1978Guk3] have studied the reaction diffusion within the ternary system with metal/oxide diffusion couples annealed at the temperatures between 1000 and 1300˚C by means of electron microprobe analysis. Two isothermal sections at 1300 and 1130˚C have been deduced from the phases developing at the interface of the (Zr,Al)/ZrO2 and (Zr,Al)/Al2O3 samples. Figures 2 and 3 show the isothermal sections at 1300 and 1130˚C after [1978Guk2] with a minor correction due to the absence of Zr4Al3 above 1030˚C. In the diffusion couples Zr4Al3 has been observed and not the high temperature phase Zr5Al4 because of the transformations during cooling. It is interesting to note that a very small Al solubility in (αZr) was observed in the diffusion couples at the αZr/βZr interface [1978Guk1]. Since the two-phase field (αZr)+(βZr) extends from the O-Zr to the Al-Zr system (metastable above 940˚C), the tie lines of this field must shift eventually to end up with a higher Al content in (αZr) than in (βZr). From the binary O-Zr system it is presumable that αZrO2 occurs in the isothermal section at 1130˚C and βZrO2 at 1300˚C.
Thermodynamics A CALPHAD type assessment of thermodynamic parameters in the ZrO2-Al2O3 system was performed by [1979Doe, 1980Wei, 1990Rel, 1992Wu, 1993Bal, 2004Fab, 2006Lak]. Different Landolt‐Bo¨rnstein New Series IV/11E1
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models of liquid solution were used in the calculations. [1979Doe, 1980Wei] used the ideal solution model for the ZrO2-AlO1.5 liquid, [1992Wu] applied the quasi-chemical model. [1990Rel, 1993Bal] used the simple substitutional model with the excess Gibbs energy term expressed by the Redlich-Kister polynomial. No solubility in the ZrO2 phases was taken into account by [1979Doe, 1980Wei, 1992Wu, 1990Rel, 1993Bal]. Partially the ionic liquid model was used for the liquid phase and the compound energy formalism for the ZrO2-based solid solutions by [2004Fab, 2006Lak]. Excellent agreement with the experimental data of [1964Alp, 1967Alp] was obtained.
Notes on Materials Properties and Applications The ZrO2-Al2O3 composite materials are attractive due to their excellent mechanical properties combining properties of ZrO2 and Al2O3 [1997He, 2004Bas]. Introduction of ZrO2 in the Al2O3 matrix leads to composite materials with increased toughness due to the stress induced transformation of tetragonal to monoclinic ZrO2 phase [1997He, 2004Bas, 2006Ran]. Introducing Al2O3 in the ZrO2 matrix also improves thermal shock resistance and hardness in comparison with ZrO2 ceramics [1997He]. [1997He] studied mechanical properties (wear and friction test) of 20 mass% Al2O3 dispersed in the Y2O3 stabilized tetragonal ZrO2 matrix (Y-TZP) and Al2O3-15 mass% ZrO2 materials. [2004Bas] investigated mechanical properties of ZrO2(Y-TZP)-28 vol% Al2O3 composite. Influence of other metal oxides on mechanical properties of the Y-TZP-Al2O3 composite was studied by [2004Ker]. [2004Lee, 2005Lee, 2005Kim] produced Al2O3-m-ZrO2/t-ZrO2 fibrous composites using extrusion process and measured their mechanical properties. [2004Llo] studied influence of the Y2O3 on mechanical properties of directionally solidified Al2O3-ZrO2 eutectic. [2006San1, 2006San2] studied effect of composition and sintering condition on mechanical properties of ZrO2-Al2O3 biocompatible composites. Influence of sintering method (conventional one or in the arc plasma reactor) on density, mechanical and dielectric properties of the ZrO2-Al2O3 composites was studied by [2005Sah, 2006Sah]. [2000Tar, 2003Tos] prepared laminated composites by superimposing alternated layer of Al2O3/ZrO2 composite and ZrO2 and studied mechanical properties of these materials. The studied composites demonstrated improvement of surface toughness, reduced friction coefficient and increased wear resistance. [2002Cic] studied conductivity of the Al2O3-ZrO2 directionally solidified eutectic composites. These composites have good combination of mechanical and electrical properties making them prospective materials for SOFC and oxygen sensors working at high temperatures. [2005Cho] studied mechanical properties of Y2O3-stabilized cubic ZrO2 (10YSZ) reinforced with 0–30 mol% Al2O3. [2001Zha, 2003Zha2, 2007Zha] studied mechanical properties of Al matrix alloys reinforced by Al2O3 and Al3Zr produced by direct melt reaction. Tensile tests [2003Zha2] indicated that (Al3Zr+Al2O3)/Al-alloy composite exhibited high strength both at room and elevated temperatures. [2001Mor] studied stability of ZrAlxOy thin film obtained by sputtering on a Si substrate for a possible replacement of SiO2 as gate dielectric material in Si-based complimentary metaloxide semiconductors. [2003Zhu, 2005Biz, 2006Biz] measured dielectric constant of Al-O-Zr thin film deposited on a Si substrate as an alternative for gate dielectric application. Also Zr-doped Al2O3 is promising candidate for gate dielectric materials [2003Jun]. DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009
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Experimental studies of materials properties in the Al-O-Zr system are summarized in Table 4.
Miscellaneous [1984Kli] studied the rate of oxygen absorption by Al-Zr alloys (16.5, 16 mass% Al) at 100–500˚C, because they are used as gas absorber to protect various devises in electronic industries. [1993Pal] studied oxidation kinetics of Zr2Al3 in the stream of dry oxygen at atmospheric pressure with constant heating rate or isothermally at 602–692˚C. Activation energy was estimated from the obtained results to be 270 kJ·mol–1. [1994Dyb] suggested different mechanisms for oxidation process studied by [1993Pal]. [1994Pal] studied oxidation rate for αZr-1mass%Al, Zr3Al, Zr2Al, ZrAl, Zr2Al3, ZrAl2 and ZrAl3 with the increasing temperature and isothermally. ZrO2 in the tetragonal and monoclinic forms was found as main oxidation product, while Al escapes by diffusion in the bulk of alloy. [1985Gar1] and [1985Gar2] applied end-point thermodynamic analysis to tetragonal/ monoclinic transformation in the Al2O3-ZrO2 composites in the presence and absence of applied stress field. [1985Gar1] took into account the contribution of chemical, dilatational, residual shear strain and interfacial energy to total Gibbs energy and calculated dependence of transformation temperature from inversed critical size of particles. [1985Gar2] modelled microcrystal of tetragonal ZrO2 constrained in a matrix subjected to hydrostatic tensile stress field. [1995Nar] studied thin films grown on cubic YSZ by liquid precursor route. Phase transformation and epitaxy of these films were studied as a function of the heat treatment temperature and time. [1999Raj] developed a process of in-situ reduction of ZrO2 with excess aluminium for preparation of an Al-Zr master alloy. [2001Zha, 2001She, 2003Deq, 2003Zha1, 2007Zha] investigated the process of in-situ fabrication of Al matrix composite reinforced by Al3Zr and Al2O3. [2003Deq] used differential scanning calorimetry to study temperature of reduction of the ZrO2 by molten Al. [2002Pen] studied influence of physicochemical treatment and Al2O3 additions to transformation of the metastable tetragonal ZrO2 phase to monoclinic modification and showed that pressing at elevated temperature accelerates the formation of the equilibrium phase, whereas the Al2O3 additions stabilize a metastable phase. [2005Cal] studied effect of the solidification rate on the microstructure of the directionally solidified eutectic. [2004Muo] studied wetting and spreading behavior of the Ag-Cu-Ti alloys on aluminazirconia ceramics important for production of metal-ceramic joints.
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. Table 1 Investigations of the Al-O-Zr Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1986Ady]
DTA, XRD, immersion method (refractive index determination)
300–1800˚C, ZrO2-Al2O3 (15–85 mol% ZrO2)
[1987Vol]
Melting, XRD, vibromilling, annealing
797–1747˚C, ZrO2-30 mass% Al2O3
[1988Tan]
Co-precipitation, heat treatment, cooling with different rates, XRD
1000, 1250˚C, ZrO2-Al2O3 0–21 mol% Al2O3
[1990Jay]
Co-precipitation, TEM, SEM/EDX, metallography, XRD
1797˚C sintering, ZrO2-Al2O3 (0–90 mass% ZrO2)
[1990Rel]
CALPHAD
1727–2727˚C, ZrO2-Al2O3
[1991And]
Rapid solidification by melt-extraction technique, TEM, SEM, XRD
1000–2000˚C, Al2O3-ZrO2 (25, 42 mol% ZrO2)
[1993Bal]
CALPHAD
1000–2800˚C, ZrO2-Al2O3
[1993Pal]
Arc-melting, XRD
602–692˚C, Zr2Al3 oxidation
[1994Bal]
Synthesis from pecursors, spray pyrolyzed 25–1200˚C (DTA), heat treatment and heat treated; 1000˚C, ZrO2-Al2O3 (10, 40 mol% XRD, DTA/TGA, electron diffraction, Raman Al2O3) spectroscopy, TEM, SEM
[1994Ish]
Co-precipitation, high isostatic pressing, heat treatment, XRD, TEM, SEM, DTA, IR spectroscopy
[1994Lak, 1997Lak]
DTA in He and TA in air in solar furnace, XRD 1100–3000˚C, ZrO2-Al2O3
[1994Pal]
XRD, EPMA, thermobalance
[1995Nar]
Thin film growth by liquid precursor route, 1100–1400˚C, ZrO2-20 mol% Al2O3 XRD, SEM, TEM
[1995Nav]
Co-precipitation, surface area 25–1000˚C, ZrO2-3–8 mol% Y2O3-5–20 measurements, XRD, TG/DTA, IR, SEM, TEM mass% Al2O3
[2000Jer]
DTA, X-ray fluorescence microanalysis
1800–2130˚C, ZrO2-Al2O3
[2001Mor]
Film sputtering on Si upon thermal annealing in vacuum or in O2, XRD, NRA, X-ray photoelectron spectroscopy, ion scattering spectroscopy
600˚C, Zr80Al20 in oxygen-containing plasma
[2001She]
TG/DTA, XRD
25–1100˚C, 55 vol% Al-25 vol% α-Al2O3 -20 vol% β-ZrO2
[2001Zha]
XRD, SEM, EPMA, TEM
800˚C, in-situ reaction ZrOCl2 with liquid Al (composite: Al matrix with Al3Zr and Al2O3)
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1250˚C, ZrO2-Al2O3 (5–25 mol% Al2O3)
374–830˚C, Zr-1 mass% Al, Zr3Al, Zr2Al, ZrAl, Zr2Al3, ZrAl2, ZrAl3 oxidation
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[2002Cic]
XRD, light microscopy
1860˚C, 54–59 mol% Al2O3- 39–44 mol % ZrO2 (1–2 mol% Y2O3)
[2003Mpe]
Sintering, SEM
1500–1600˚C, 3YTZ (3 mass% Y2O3 ZrO2) - 0–75 vol% Al2O3
[2003Deq]
Uniaxial pressing, heating in Ar, DSC, XRD, SEM, thermodynamic calculations
150–1200˚C; Al-10–20 mass% ZrO2
[2003Jun]
X-ray emission and absorption spectroscopy
Zr-doped Al2O3
[2003Kin]
Plasma Al and Zr interaction with O2 gas, XRD, X-ray photoelectron spectroscopy (XPS), TEM/EDS
1000–1200˚C annealing, ZrO2-Al2O3
[2003Zha1] [2003Zha2]
Direct melt reaction, SEM, TEM, EPMA
800˚C, ZrOCl2+Al
[2003Zho]
Plasma spray method, sintering, annealing, 1400˚C, ZrO2(3 mol%Y2O3)-20, XRD, SEM 57 mass% Al2O3
[2003Zhu]
Laser deposition on Pt coated Si and Si substrate at 20 Pa O2, XRD, DTA
800–950˚C deposition, 25–1300˚C DTA; ZrO2-50 mol% Al2O3
[2004Fab]
CALPHAD
227–2727˚C, ZrO2-Al2O3
[2005Cal]
Arc-melting, XRD, SEM/EDX, micro-Raman
ZrO2-Al2O3 eutectic, rapid solidification
[2005Che]
XRD, TG/DSC, IR spectroscopy, TEM/EDS, nitrogen adsorption
100–1400˚C, ZrO2-Al2O3
[2005Kam]
DTA, XRD, SEM/EDX, calculations
1800–2100˚C; ZrO2-Al2O3
[2006Jia]
Coating (layer-by-layer method), cold and hot isostatic pressing, sintering, XRD, SEM, TEM
1350–1550˚C, Al2O3-12 mass% ZrO2
[2006Lak]
CALPHAD
1727–2727˚C, ZrO2-Al2O3
[2006Por]
Coating (magneton sputtering) in Ar and O2 gas mixture from metallic target, XRD
ZrO2-2–9 mass% Al2O3
[2006Ran]
Gel-precipitation, XRD, DSC/TG, Archimedes 25–1200˚C (DSC),1400–1600˚C method, IR-spectroscopy (sintering); Al2O3-ZrO2
[2006Sah] [2005Sah]
Conventional sintering and arc melting, XRD, SEM
1400–1600˚C, ZrO2-Al2O3
[2006San1] [2006San2]
Cold pressing, sintering, XRD, SEM
1500–1600˚C, ZrO2-0–30 mass% Al2O3
[2007Kon]
Pechini sol-gell process, HIP, DTA/TGA, XRD, 600–1500˚C, ZrO2-20 mass% Al2O3 TEM/EDX, SEM
[2007Pod]
Co-precipitation, heat treatment, XRD, DTA 700–1400˚C, ZrO2-50 mol% Al2O3
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. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[2007Sch]
Microwave plasma process, XRD, TEM, Perturbed Angular Correlation (PAC) measurement
≤627˚C, ZrO2 nanocrystalline particles covered by Al2O3
[2007Zha]
In-situ magneto-chemistry reaction, XRD, SEM, TEM
900˚C, Al-x mass% Zr(CO3)2; x = 5–25
[2008Car]
XRD, Neutron diffraction, stress 550 MPa
500˚C, 20 mass% ZrO2 - Al2O3
[2008Che]
Co-precipitation, plasma spray deposition, heat-treatment, XRD, SEM, chemical analysis (XPS – X-ray photoelectron spectroscopy)
1100–1500˚C, Al2O3-40 mass% ZrO2 (7 mass% Y2O3)
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αAl) < 660.452
cF4 Fm 3m Cu
a = 404.96
at 25˚C [Mas2]
(βAl)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
25˚C, 20.5 GPa [Mas2]
(βZr) 1855 - 863
cI2 Im 3m W
a = 360.90
[Mas2]
(αZr) < 863
hP2 P63/mmc Mg
a = 323.16 c = 514.75
at 25˚C [Mas2]
Zr3Al < 1019
cP4 Pm 3m AuCu3
a = 437.2
[2004Sch]
Zr2Al(r) < 1250
hP6 P63/mmc Ni2In
a = 489.39 c = 592.83
[2004Sch]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Zr5Al3(h) 1395 - 1000
tI32 I4/mcm W5Si3
a = 1104.4 c = 539.1
[2004Sch]
Zr5Al3(r)?
hP16 P6/mcm Mn5Si3
a = 817.0 c = 569.8
[2004Sch]
Zr3Al2 < 1480
tP20 P42/mnm Zr3Al2
a = 763.0 c = 699.8
[2004Sch]
Zr4Al3 ≤ 1030
hP7 P6/mmm Zr4Al3
a = 543.3 c = 539.0
[2004Sch]
Zr5Al4 1550 - 1000
hP18 P63/mcm Ti5Ga4
a = 844.7 c = 580.5
[2004Sch]
ZrAl < 1275
oC8 Cmcm CrB
a = 335.9 b = 1088.7 c = 427.4
[2004Sch]
Zr2Al3 < 1590
oF40 Fdd2 Zr2Al3
a = 960.1 b = 1390.6 c = 557.4
[2004Sch]
ZrAl2 < 1660
hP12 P63/mmc MgZn2
a = 528.24 c = 874.82
[2004Sch]
ZrAl3 < 1580
tI16 I4/mmm ZrAl3
a = 399.93 c = 1728.3
[2004Sch]
ZrAl3 (m)
cP4 Pm 3m AuCu3
a = 408
[2004Sch], metastable
αAl2O3 < 2054
hR30 R3c Al2O3
a = 475.4 c = 1299
[V-C2], corundum, congruent melting 2054˚C [Mas2]
γAl2O3
cF56 Fd 3m MgAl2O4
a = 794.7
[V-C2] metastable phase [Mas2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
δAl2O3
oP64 P21212
a = 1589.4 b = 794.7 c = 1192.05
[1998Lev]
θAl2O3
mC8 C2/m Ga2O3
a = 1192.05 b = 280.97 c = 561.94 β = 104
[1998Lev]
λAl2O3
mP64 P21/c
a = 1685.81 b = 1589.4 c = 1192.05 β = 115
[1998Lev]
κAl2O3
hP44 P63mc Al2O3
a = 554.4 c = 902.4
[V-C2]
αZrO2 ≤ 1094˚C
mP12 P21/c ZrO2
a = 514.15 b = 520.56 c = 531.28 β = 99.30 a = 516.2 b = 519.4 c = 532.5 β = 99.11
baddeleyite [2006Wan]
a = 360.55 c = 517.97
up to 4.5 mol% Al2O3 for x = 0, at 1393˚C [V-C2] for 16.7 mol% Al2O3, Metastable [1988Tan] 5–40 mol% Al2O3 metastable t’ [1994Ish]
βZrO2 2311 - 1094
γZrO2 2710 - 2311
tP6 P42/nmc ZrO2
cF12 Fm3m CaF2
a = 359.5 c = 519.1 a = 508.5 c = 512.7 to 518.2 a = 513.2 a = 507 to 510.5
δZrO2
oP16 Pbcm ZrO2
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a = 503.64 b = 525.46 c = 508.55
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for 16.3 mol% Al2O3, Metastable [1988Tan]
fluorite up to 3 mol% Al2O3 in equilibrium conditions, 5–40 mol% Al2O3 metastable [1994Ish] high pressure phase 3 < p < 12.5 GPa [2005Oht] for x = 0 [V-C2]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C] εZrO2
oP12 Pnma PbCl2
ωZrO2
tP12 P4m2
Lattice Parameters [pm]
Comments/References high pressure phase [1999Des] p > 12.5 GPa [2005Oht]
a = 562 b = 334.7 c = 650.3 a = 504.6 c = 521.9
high pressure phase 3.5 < p < 15 GPa [V-C2]
. Table 3 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
γZrO2 Ð L + βZrO2
2314
e1
γZrO2 βZrO2
L Ð βZrO2 + Al2O3
1856
e2
e3
Zr
1.9815
66.3356
31.6828
1.6832
66.3861
31.9307
L
14.2800
64.2867
21.4334
29.2985
61.7836
8.9179
2.9521
66.1746
30.8732
Al2O3 1130
O
L βZrO2
γZrO2 Ð αZrO2 + Al2O3
Al
40
60
0.0
γZrO2
0.878
66.5205
32.6015
αffl ZrO2
0.082
66.6530
33.2646
Al2O3
40
60
0.0
. Table 4 Investigations of the Al-O-Zr Materials Properties Reference
Method / Experimental Technique
Type of Property
[1997He]
Friction and wear test (ball-on-plate wear Mechanical properties test)
[2000Tar]
Vickers hardness, indentation technique, wear test (pin-on-disc apparatus), XRD, SEM, EDAX
Mechanical properties (hardness, residual stress, Young’s modulus, wear resistance)
[2001Zha]
Tensile test
Mechanical properties
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. Table 4 (continued) Reference [2002Cic]
Method / Experimental Technique Impedance analysis for conductivity measurements
Type of Property Electric properties
[2003Mpe] Conductivity, dielectric constant
Electric properties
[2003Gan]
Compression test, SEM, optical microscopy
Mechanical properties of amorphous ZrO2 - 40 mol% Al2O3
[2003Tos]
Indentation technique, wear test (pin-on- Mechanical properties (residual stress, disc apparatus) surface toughness, friction coefficient, wear resistance)
[2003Zha2] Tensile test
Mechanical properties of (Al-7Si-0.3Mg)/ (Al3Zr+Al2O3)
[2004Bas]
Archimedes method, resonance frequency method, hardness test, indentation XRD, SEM, EPMA/EDS
Density, mechanical properties (elastic modulus, Vickers hardness, fracture toughness)
[2004Ker]
Wear friction test (pin-on-disc apparatus), Mechanical properties (friction XRD, SEM/EDX coefficient, wear resistance)
[2004Lee] [2005Lee] [2005Kim]
Archimedes method, hardness test, indentation, four point bending test, XRD, SEM, TEM
Density, mechanical properties (Vickers hardness, fracture toughness, bending strength)
[2004Llo]
Three point bending test, hardness test, indentation fracture method, XRD, SEM, TEM, Raman spectroscopy
Mechanical properties of directionally solidified eutectic Al2O3-ZrO2
[2005Cho]
XRD, SEM, TEM/EDS, mass-volume method, flexure strength test, fracture toughness test, impulse excitation of vibration method, microhardness test, fatigue test
Density and mechanical properties (fracture toughness, flexure strength, elastic properties, Vickers hardness, fatigue) ZrO2 (10 mol% Y2O3)-0–30 mol% Al2O3
[2006Biz] [2005Biz]
XRD, IR-spectroscopy, X-ray spectroscopy ZrO2-Al2O3 thin films, dielectric properties (XPS), capacitance measurements
[2006Sah] [2005Sah]
Archimedes method, Vickers indentation Density, mechanical properties method, dielectric bridge, XDT, SEM (hardness), dielectric properties ZrO2Al2O3 composites
[2006San1] Vickers indentation method [2006San2]
Mechanical properties (Vickers hardness, fracture toughness) of ZrO2-0–30 mass% Al2O3
[2007Zha]
Mechanical properties (tensile and yield strength, elongations) of Al/(Al3Zr+Al2O3) composite
Tensile test
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. Fig. 1 Al-O-Zr. Quasibinary system ZrO2-AlO1.5
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Al–O–Zr
. Fig. 2 Al-O-Zr. Isothermal section at 1300˚C
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. Fig. 3 Al-O-Zr. Isothermal section at 1130˚C
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Al–O–Zr
References [1932War]
[1961Suz]
[1964Alp] [1967Alp]
[1968Cev]
[1977Guk] [1978Guk1]
[1978Guk2]
[1978Guk3]
[1979Doe]
[1980Wei]
[1984Kli]
[1985Gar1]
[1985Gar2]
[1985Wri] [1986Ady]
[1987Har]
[1987Vol]
von Wartenberg, H., Reusch, H.J., “The Melting Diagrams of Some High-Refractory Oxides. IV (Aluminium Oxides)” (in German), Z. Anorg. All. Chem., 207, 1–20 (1932) (Phase Relations, Experimental, 2) Suzuki, H., Kimura, S., Yamada, H., Yamauchi, T., “Studies of the Systems Al2O3-ZrO2 and Na2O-ZrO2 Studies on the Refractories of the System Na2O-Al2O3-ZrO2, I”, J. Ceram. Soc. Jpn., 69(2), 72–79 (1961) (Experimental, Morphology, Phase Relations, 10) Alper, A.M., McNally, R.N., Doman, R.C., “Phase Equilibria in the Al2O3-ZrO2 System”, Amer. Ceram. Soc. Bull., 43, 642 (1964) (Phase Relations, Abstract, 0) Alper, A.M., “Inter-Relationship of Phase Equilibria Microstructure and Properties in Fusion-Cast Ceramics” in “Science of Ceramics”, Vol. 3, Stewart, G.H. (Ed.), Academic Press Inc. (London) Ltd., 337–344 (1967) (Phase Relations, Experimental, 10) Cevales, G., “Phase-Equilibrium Diagram of Al2O3-ZrO2 and Examinations of a New HighTemperature Phase (ε-Al2O3)” (in German), Ber. Deut. Keram. Ges., 45, 216–219 (1968) (Phase Relations, Experimental, 5) Gukelberger, A., Steeb, S., “Chemical Diffusion in the Zr-Al-O Ternary System” (in German), Mikrochim. Acta, Supp., 7, 373–388 (1977) (Phase Relations, Experimental, 9) Gukelberger, A., Steeb, S., “Diffusion Weldings Within the Zr-Al-O Ternay System at Temperatures Between 1000 and 1300oC. Part II. Zr-Al2O3 Couple” (in German), Z. Metallkd., 69, 385–393 (1978) (Experimental, 14) Gukelberger, A., Steeb, S., “Diffusion Weldings Within the Zr-Al-O Ternary System at Temperatures Between 1000 and 1300˚C. Part III. Study of the Al2O3-(Zr, Al) Alloys, ZrO2-(Zr, Al) Alloys and Al2O3ZrO2” (in German), Z. Metallkd., 69, 462–469 (1978) (Phase Relations, Experimental, 11) Gukelberger, A., Steeb, S., “Diffusion Weldings Within the Zr-Al-O Ternary System at Temperatures Between 1000 and 1300˚C Part 1. Zr-Zr2Al3 Couple” (in German), Z. Metallkd., 69(4), 255–260 (1978) (Experimental, Morphology, Phase Relations, 13) Do¨rner, P., Gauckler, L.I., Kreig, H., Lukas, H.L., Petzow, G., Weiss, J., “On the Calculation and Representation of Multicomponent Systems”, Calphad, 3, 241–257 (1979) (Phase Diagram, Phase Relations, Thermodyn., 24) Weiss, J., “Constitutional Investigations and Thermodynamic Calculations in the Si-Al-Zr/ N-O System” (in German), Thesis, Univ. Stuttgart, (1980) (Phase Diagram, Phase Relations, Thermodyn., 92) Klimentenko, O.P., Kozik, V.V., Serebrennikov, V.V., Khvesevich, Yu.G., “Absorption of Oxygen at Zirconium-Aluminum Alloys”, J. Appl. Chem. USSR (Engl. Transl.), 57(5), 1051–1052 (1984), translated from Zh. Priklad. Khimii, 57(5), 1136–1137 (1984) (Experimental, 3) Garvie, R.C., Swain, M.V., “Thermodynamics of the Tetragonal to Monoclinic Phase Transformation in Constrained Zirconia Microcrystals. 1. In the Absence of an Applied Stress Field”, J. Mater. Sci., 20 (4), 1193–1200 (1985) (Experimental, Thermodyn., 32) Garvie, R.C., “Thermodynamic Analysis of the Tetragonal to Monoclinic Transformation in a Constrained Zirconia Microcrystal. II. In the Presence of an Applied Stress”, J. Mater. Sci., 20(10), 3479–3486 (1985) (Experimental, Thermodyn., 22) Wriedt, H.A., “The Al-O (Aluminium-Oxygen) System”, Bull. Alloy Phase Diagrams, 6(6), 548–553 (1985) (Phase Relations, Phase Diagram, Review, 46) Adylov, G.T., Urazaeva, E.M., Mansurova, E.P., “Phase Relations in the Al2O3-ZrO2 System under Various Melt Cooling Conditions”, Inorg. Mater. (Engl. Trans.), 22(2), 1474–1477 (1986), translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 22(10), 1683–1686 (1986) (Experimental, Morphology, Phase Relations, 7) Harmelin, M., “Aluminium - Oxygen - Zirconium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12755.1.20, (1992) (Crys. Structure, Phase Diagram, Assessment, 11) Volkova, I.Yu, Semenov, S.S., Kravchik, A.E., Ordanyan, S.S., Kozlovskii, L.V., “Influence of Rate of Cooling of Eutectic of System Al2O3-ZrO2 on Stability of Phase Components”, Inorg. Mater. (Engl. Trans.), 23(3), 394–398 (1987), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 23(3), 448–451 (1987) (Experimental, Phase Diagram, Phase Relations, 7)
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[1990Rel]
[1991And]
[1990Jay]
[1992Tay]
[1992Wu] [1993Bal] [1993Pal] [1994Bal]
[1994Dyb] [1994Ish]
[1994Lak]
[1994Pal] [1995Nav]
[1995Nar]
[1997He] [1997Lak] [1998Lev]
[1999Des] [1999Raj] [2000Jer]
14
Tananaev, I.V., Khodzhamberdiev, M.S., Salyn, A.L., Beresnev, E.N., Pakhomov, V.I., Orlovskii, V.P., “Solubility of Al2O3 in ZrO2”, Inorg. Mater. (Engl. Trans.), 24, 1651–1653 (1988), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 24(11), 1925–1927 (1988) (Crys. Structure, Experimental, 10) Relave, O., Chevalier, P.Y., Cheynet, B., Cenerino, G., “Thermodynamic Calculation of Phase Equilibria in a Quinary Oxide System of First Interest in Nuclear Energy Field UO2-ZrO2-SiO2-CaO-Al2O3”, User Aspects of Phase Diagrams, Proc. Int. Conf., 25–27 June 1990, Petten, Hayes, F.H. (Ed.), Int. Met., 55–63 (1991) (Calculation, Phase Diagram, Phase Relations, 19) Ando, T., Shiohara, Y., “Metastable Alumina Structures in Melt-Extracted Alumina-25% Zirconia and Alumina-42% Zirconia Ceramics”, J. Am. Ceram. Soc., 74(1), 410–417 (1991) (Experimental, Morphology, Phase Diagram, Phys. Prop., Thermodyn., 23) Jayaram, V., Levi, C.G., Witney, T., Mehrabian, R., “Characterization of Al2O3-ZrO2 Powders Produced by Electrohydrodynamic Atomization”, Mater. Sci. Eng., 124(A), 65–81 (1990) (Experimental, Phase Relations, 15) Taylor, J.R., Dinsdale, A.T., Hillert, M., Selleby, M., “A Critical Assessment of Thermodynamic and Phase Diagram Data for the Al-O System”, Calphad, 16(2), 173–179 (1992) (Calculation, Phase Diagram, Thermodyn., 22) Wu, P., “Optimization and Calculations of Thermodynamic Properties and Phase Diagrams of Multicomponent Oxide Systems”, Ph. D. Thesis, Ecole Polytechnique, Montreal, Canada (1992) Ball, R.G.J., Mignanelli, M.A., Barry, T.I., Gisby, J.A., “The Calculation of Phase Equilibria of Oxide Core-Concrete Systems”, J. Nucl. Mater., 201, 238–249 (1993) (Thermodyn., Calculation, 66) Paljevic, M., “Selectiv Oxidation of Zr2Al3”, J. Alloys Compd., 191, 27–29 (1993) (Experimental, Interface Phenomena, Phase Relations, 13) Balmer, M.L., Lange, F.F., Levi, C.G., “Metastable Phase Selection and Partitioning for Zr(1–x) AlxO(2–x/2) Materials Synthesized with Liquid Precursors”, J. Am. Ceram. Soc., 77(8) 2069–2075 (1994) (Experimental, Phase Relations, 20) Dybkov, V.I., “Comment on the Paper (Selective Oxidation of Zr2Al3) by M. Paljevic”, J. Alloys Compd., 215, L1-L1 (1994) (Experimental, 3) Ishida, K., Hirota, K., Yamaguchi, O., “Formation of Zirconia Solid Solutions Containing Alumina Prepared by New Preparation Method”, J. Am. Ceram. Soc., 77(5), 1391–1395 (1994) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 20) Lakiza, S.N., Lopato, L.M., Shevchenko, A.V., “Interaction in the Al2O3-ZrO2-Y2O3 System”, Powder Metall. Met. Ceram., 33(9-10), 486–490 (1994), translated from Poroshk. Metall., (9–10), 46–51 (1994) (Experimental, Morphology, Phase Diagram, Phase Relations, 22) Paljevic, M., “High-Temperature Oxidation Behavior in the Zr-Al System”, J. Alloys Compd., 204, 119–126 (1994) (Experimental, 57) Navarro, L.M., Recio, P., Duran, P., “Preparation and Properties Evaluation of Zirconia-Based Al2O3 Composites as Electrolytes for Solid Oxide Fuel Cell Systems. 1. Powder Preparation and Characterization”, J. Mater. Sci., 30(8), 1931–1938 (1995) (Experimental, Mechan. Prop., Morphology, Thermodyn., 19) Narwankar, P.K., Speck, J.S., Lange, F.F., “Phase Partitioning and Epitaxy of Zr(Al)O2 Thin Films on Cubic Zirconia Substrates”, J. Mater. Res., 10(7), 1756–1763 (1995) (Experimental, Morphology, Phase Relations, Phys. Prop., 22) He, Y.J., Winnubst, A.J.A., Burggraaf, A.J., Verweij, H., van der Varst, P.G.T., de With, G., “Sliding Wear of ZrO2-Al2O3 Composite Ceramics”, J. Eur. Ceram. Soc., 1371–1380 (1997) (Experimental, 25) Lakiza, S.M., Lopato, L.M., “Stable and Metastable Phase Relations in the System Alumina-ZircomiaYttria”, J. Am. Ceram. Soc., 80(4), 893–902 (1997) (Experimental, Phase Relations, 26) Levin, I., Gemming, Th., Brandon, D.G., “Some Metastable Polymorphs and Transient Stages of Transformation in Alumina”, Phys. Status Solidi A, 166(1), 197–218 (1998) (Crys. Structure, Experimental, Phase Relations, 24) Desgreniers, S., Lagarec, K., “High-Density ZrO2 and HfO2: Crystalline Structures and Equations of State”, Phys. Rev. B, 59(13), 8467–8472 (1999) (Crys. Structure, Experimental, 25) Rajagopalan, P.K., Sharma, I.G., Krishnan, T.S., “Production of Al-Zr Master Alloy Starting from ZrO2”, J. Alloys Compd., 285(1-2), 212–215 (1999) (Experimental, Thermodyn., 8) Jerebtsov, D.A., Mikhailov, G.G., Sverdina, S.V., “Phase Diagram of the System: Al2O3-ZrO2”, Ceram. Int., 26, 821–823 (2000) (Experimental, Phase Diagram, 10)
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[2001Mor]
[2001She] [2001Zha] [2002Boc]
[2002Cic] [2002Pen]
[2003Deq]
[2003Gan] [2003Jun]
[2003Kin]
[2003Mpe]
[2003Tos]
[2003Zha1]
[2003Zha2]
[2003Zho] [2003Zhu] [2004Bas] [2004Fab] [2004Ker]
[2004Lee]
Al–O–Zr Tarlazzi, A., Roncari, E., Pinasco, P., Guicciardi, S., Melandri, C., de Portu, G., “Tribological Behaviour of Al2O3/ZrO2-ZrO2 Laminated Composites”, Wear, 244(1-2), 29–40 (2000) (Experimental, Mechan. Prop., Morphology, Phys. Prop., 23) Morais, J., da Rosa, E.B.O., Pezzi, R.P., Miotti, L., Baumvol, I.J.R., “Composition, Atomic Transport, and Chemical Stability of ZrAlxOy Ultrathin Films Deposited on Si(001)”, Appl. Phys. Lett., 79(13), 1998–2000 (2001) (Experimental, 18) Sheedy, P.M., Caram, H.S., Chan, H.M., Harmer, M.P., “Effects of Zirconium Oxide on the Reaction Bonding of Aluminum Oxide”, J. Am. Ceram. Soc., 84(5), 986–990 (2001) (Experimental, 16) Zhao, Y. T., Sun, G. X., “In-Situ Synthesis of Novel Composites in the System Al-Zr-O”, J. Mater. Sci. Lett., 20(20), 1859–1861 (2001) (Crys. Structure, Experimental, Mechan. Prop.,6) Bocanegra-Bernal, M.H., Diaz De La Torre, S., “Review. Phase Transitions in Zirconium Dioxide and Related Materials for High Performance Engineering Ceramics”, J. Mater. Sci., 37(23), 4947–4971 (2002) (Crys. Structure, Mechan. Prop., Phase Relations, Review, Thermodyn., 175) Cicka, R., Trnovcova, V., Starostin, M.Y., “Electrical Properties of Alumina-Zirconia Eutectic Composites”, Solid State Ionics, 148(3-4), 425–429 (2002) (Experimental, Morphology, 14) Pentin, I.V., Oleinikov, N.N., Murav’eva, G.P., Eliseev, A.A., Tret’yakov, Yu.D., “Stability of Tetragonal ZrO2 Toward External Influences”, Inorg. Mater. (Engl. Trans.), 38(10), 1012–1014 (2002) (Crys. Structure, Experimental, 8) Deqing, W., Neumann, J., Lopez, H.F., “Reactive Synthesis of in-Situ ZrAl3-Al2O3-Al Composites”, Metall. Mater. Trans. A, 34(6), 1357–1360 (2003) (Crys. Structure, Experimental, Morphology, Thermodyn., 21) Gandhi, A.S., Jayaram, V., “Plastically Deforming Amorphous ZrO2-Al2O3”, Acta Mater., 51(6), 1641–1649 (2003) (Experimental, Mechan. Prop., 47) Jung, R., Lee, J.-Ch., So, Y.-W., Noh, T.-W., Oh., S.-J., Lee, J.-Ch., Shin, H.-J., “Bandgap States in Transition-Metal (Sc, Y, Zr, and Nb)-Doped Al2O3”, Appl. Phys. Lett., 83(25), 5226–5228 (2003) (Electronic Structure, Experimental, 7) Kinemuchi, Y., Mouri, H., Suzuki, T., Suematsu, H., Jiang, W., Yatsui, K., “Increase in Phase Transition Temperature of Activated Alumina with Nano-Zirconia Synthesized by Pulsed Wire Discharge”, J. Am. Ceram. Soc., 86(9), 1522–1526 (2003) (Crys. Structure, Experimental, Phase Relations, 16) M’Peko, J.-C., Spavieri Jr., D.L., da Silva, Ch.L., Fortulan, C.A., de Souza, D.P.F., de Souza, M.F., “Electrical Properties of Zirconia-Alumina Composites”, Solid State Ionics, 156(1-2), 59–69 (2002) (Electr. Prop., Experimental, 23) Toschi, F., Melandri, C., Pinasco, P., Roncari, E., Guicciardi, S., de Portu, G., “Influence of Residual Stresses on the Wear Behavior of Alumina/Alumina-Zirconia Laminated Composites”, J. Am. Ceram. Soc., 86(9), 1547–1553 (2003) (Calculation, Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 33) Zhao, Y.T., Cheng, X.N., Dai, Q.X., Cai, L., Sun, G.X., “Crystal Morphology and Growth Mechanism of Reinforcements Synthesized by Direct Melt Reaction in the System Al-Zr-O”, Mater. Sci. Eng. A, 360(1-2), 315–318 (2003) (Experimental, Morphology, Phys. Prop., 10) Zhao, Y.T., Dai, Q.X., Cheng, X.N., Sun, S.C., “Microstructure and Properties of In-Situ Synthesized (Al3Zr + Al2O3)p/A356 Composites”, Internat. J. Mod. Phys. B, 17(8-9), 1292–1296 (2003) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 10) Zhou, X., Shukla, V., Cannon, W.R., Kear, B.H., “Metastable Phase Formation in Plasma-Sprayed ZrO2 (Y2O3)-Al2O3”, J. Am. Ceram. Soc., 86(8), 1415–1420 (2003) (Crys. Structure, Phase Diagram, 13) Zhu, J., Liu, Z.G., “Structure and Dielectric Properties of Zr-Al-O Thin Films Prepared by Pulsed Laser Deposition”, Microelectr. Eng., 66(1-4), 849–854 (2003) (Crys. Structure, Experimental, Phys. Prop., 10) Basu, B., Vleugels, J., van der Biest, O., “ZrO2-Al2O3 Composites With Tailored Toughness”, J. Alloys Compd., 372, 278–284 (2004) (Experimental, Mechan. Prop., Morphology, 23) Fabrichnaya, O., Aldinger, F., “Assessment of Thermodynamic Parameters in the System ZrO2-Y2O3Al2O3”, Z. Metallkd., 95(1), 27–39 (2004) (Assessment, Phase Diagram, Thermodyn., 61) Kerkwijk, B., Garcia, M., van Zyl, W.E., Winnubst, L., Mulder, E.J., Schipper, D.J., Verweij, H., “Friction Behaviour of Solid Oxide Lubricants as Second Phase in α-Al2O3 and Stabilised ZrO2 Composites”, Wear, 256(1-2), 182–189 (2004) (Experimental, Morphology, Phys. Prop., 38) Lee, B.-T., Kim, K.-H., Han, J.-K., “Microstructures and Aterial Properties of Fibrous Al2O3-(mZrO2)/t-ZrO2 Composites Fabricated by a Fibrous Monolithic Process”, J. Mater. Res., 19(11), 3234–3241 (2004) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phys. Prop., 17)
DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009
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[2004Muo] [2004Sch]
[2005Biz]
[2005Cal]
[2005Che]
[2005Cho] [2005Kam]
[2005Kim]
[2005Lee]
[2005Oht]
[2005Sah]
[2006Biz]
[2006Jia]
[2006Lak]
[2006Por]
[2006Ran]
[2006Sah]
[2006San1]
14
LLorca, J., Pastor, J.Y., Poza, P., Pena, J.I., de Francisco, I., Larrea, A., Orera, V.M., “Influence of the Y2O3 Content and Temperature on the Mechanical Properties of Melt-Grown Al2O3-ZrO2 Eutectics”, J. Am. Ceram. Soc., 87(4), 633–639 (2004) (Experimental, Mechan. Prop., Morphology, Phase Relations, 27) Muolo, M.L., Ferrera, E., Morbelli, L., Passerone, A., “Wetting, Spreading and Joining in the AluminaZirconi-Inconel 738 System”, Scr. Mater., 50, 325–330 (2004) (Experimental, Kinetics, Morphology, 31) Schuster, J.C., “Al-Zr (Aluminium - Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13524.1.20, (2004) (Phase Diagram, Phase Relations, Assessment, 138) Bizarro, M., Alonso, J.C., Ortiz, A., “The Effect of the Process Conditions on the Synthesis of Zirconium-Aluminum Oxide Thin Films Prepared by Ultrasonic Spray Pyrolysis”, J. Electrochem. Soc., 152(11), F179-F184 (2005) (Experimental, Thermodyn., 29) Calderon-Moreno, J.M., Yoshimura, M., “Stabilization of Zirconia Lamellae in Rapidly Solidified Alumina-Zirconia Eutectic Composites”, J. Eur. Ceram. Soc., 25(8), 1369–1372 (2005) (Experimental, Morphology, 22) Chen, H., Wang, X.M., Shia, J., Xiao, P., Yan, D., “A Novel Structural Mesoporous Alumina/Yttrium Doped Zirconia Nanocrystalline Composite Derived by Solvothermal Approach”, J. Mater. Res., 20(1), 42–47 (2005) (Crys. Structure, Experimental, Morphology, Phase Relations, Thermodyn., 19) Choi, S.R., Bansal, N.P., “Mechanical Behavior of Zirconia/Alumina Composites”, Cer. I., 31(1), 39–46 (2005) (Mechan. Prop., 29) Kamaev, D.N., Archugov, S.A., Mikhailov, G.G., “Behavior of the Al2O3-ZrO2 System at High Temperatures”, Russ. J. Appl. Chem., 78(3), 347–350 (2005) (Experimental, Phase Diagram, Phase Relations, 16) Kim, T.-S., Goto, T., Lee, B.-T., “Microstructural Control and Mechanical Properties of Fibrous Al2O3/ ZrO2 Composites Fabricated by Extrusion Process”, Scr. Mater., 52(8), 725–729 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, Phys. Prop., 10) Lee, B.-T., Jang, D.-H., Kang, I.-Ch., Lee, Ch.-W., “Relationship Between Microstructures and Material Properties of Novel Fibrous Al2O3-(m-ZrO2)/t-ZrO2 Composites”, J. Am. Ceram. Soc., 88(10), 2874–2878 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, 21) Ohtaka, O., Andrault, D., Bouvier, P., Schultz, E., Mezouar, M., “Phase Relations and Equation of State of ZrO2 to 100 GPa”, J. Appl. Crystallogr., 38(5), 727–733 (2005) (Crys. Structure, Experimental, Phys. Prop., Thermodyn., 38) Sahu, D.R., Mishra, D.K., Swain, D., Ray, M., Roul, B.K., “Effect of Compositional Variation in Sintering Behaviour of Al-Zr Oxide Composites”, Mater. Sci. Eng. B, 119(1), 29–35 (2005) (Experimental, Morphology, Phase Relations, Phys. Prop., 32) Bizarro, M., Alonso, J.C., Fandino, J., Ortiz, A., “Effect of Water Addition in the Spray Solution on the Synthesis of Zr-Al Oxide Films Prepared by the Pyrosol Process”, J. Electrochem. Soc., 153(7), F153F159 (2006) (Experimental, 44) Jia, Y., Hotta, Y., Sato, K., Watari, K., “Homogeneous ZrO2-Al2O3 Composite Prepared by Nano-ZrO2 Particle Multilayer-Coated Al2O3 Particles”, J. Am. Ceram. Soc., 89(3), 1103–1106 (2006) (Experimental, Morphology, Phase Relations, 26) Lakiza, S., Fabrichnaya, O., Zinkevich, M., Aldinger, F., “On the Phase Relations in the ZrO2-YO1.5AlO1.5 System”, J. Alloys Compd., 420(1-2), 237–245 (2006) (Assessment, Crys. Structure, Experimental, Phase Diagram, Phase Relations, Thermodyn., 23) Portinha, A., Teixeira, V., Carneiro, J., Newton, R., Fonseca, H., “Structural Characterization of Sputtered Composite Stabilized ZrO2 Thin Films”, Mater. Sci. Forum, 514-516, 1150–1154 (2006) (Crys. Structure, Experimental, Nano, 18) Rana, R.P., Pratihar, S.K., Bhattacharyya, S., “Effect of Powder Treatment on the Crystallization Behaviour and Phase Evolution of Al2O3-High ZrO2 Nanocomposites”, J. Mater. Sci., 41(21), 7025–7032 (2006) (Crys. Structure, Experimental, 14) Sahu, D.R., Roul, B.K., Singh, S.K., Choudhury, R.N.P., “Studies on Sintering Behaviour of Al2O3ZrO2 Oxide Composites Processed by Extended Arc Thermal Plasma and Conventional Heating”, J. Mater. Sci., 41(17), 5480–5489 (2006) (Crys. Structure, Electr. Prop., Experimental, Morphology, 41) Santos, C., Teixeira, L.H.P., Daguano, J.K.M.F., Strecker, K., Elias, C.N., “Effect of Isothermal Sintering Time on the Properties of the Ceramic Composite ZrO2-Al2O3”, Mater. Sci. Forum, 530-531, 526–531 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 8)
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[2006Wan]
[2007Kon]
[2007Pod]
[2007Sch]
[2007Zha]
[2008Che]
[2008Car]
[Mas2] [V-C2]
Al–O–Zr Santos, C., Teixeira, L.H.P., Strecker, K., Elias, C.N., “Effect of Al2O3 Addition on the Mechanical Properties of Biocompatible ZrO2-Al2O3 Composites”, Mater. Sci. Forum, 530-531, 575–580 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 8) Wang C., Zinkevich M., Aldinger F., “The Zirconia - Hafnia: DTA measurements and Thermodynamic Calculations”, J. Am. Ceram. Soc., 89(12) 3751–3758 (2006) (Experimental, Calculation, Phase Relations, Thermodyn., 78) Kong, Y.-M., Kim, H.-E., Kim, H.-W., “Production of Aluminum-Zirconium Oxide Hybridized Nanopowder and Its Nanocomposite”, J. Am. Ceram. Soc., 90(1), 298–302 (2007) (Experimental, Mechan. Prop., Morphology, Nano, 17) Podzorova, L.I., Ilicheva, A.A., Shvorneva, L.I., “Effect of the Precipitation Sequence on Phase Formation in the ZrO2-CeO2-Al2O3 System”, Inorg. Mater. (Engl. Trans.), 43(9), 972–975 (2007), translated from Neorg. Mater., 43(9), 1086–1089 (2007) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 14) Schlabach, S., Szabo, D.V., Vollath, D., de la Presa, P., Forker, M., “Zirconia and Titania Nanoparticles Studied by Electric Hyperfine Interactions, XRD and TEM”, J. Alloys Compd., 434-435, 590–593 (2007) (Crys. Structure, Experimental, Nano, 13) Zhang, S.-L., Zhao, Y.-T., Chen, G., Cheng, X.-N., Dai, Q.-X., “Microstructures and Mechanical Properties of Aluminum Matrix Composites Fabricated from Al-x wt.% Zr(CO3)2 (x = 5, 10, 15, 20, 25) Systems”, J. Alloys Compd., 429(1-2), 198–203 (2007) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 25) Chen, D., Jordan, E.H., Gell, M., Ma, X., “Dense Alumina-Zirconia Coatings Using the Solution Precursor Plasma Spray Process”, J. Am. Ceram. Soc., 91(1), 359–365 (2008) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 20) Carter, G.A., van Riessen, A., “Neutron Diffraction of Zirconia-Dispersed Alumina with Increasing Stress and Temperature”, J. Am. Ceram. Soc., 91(2), 559–562 (2008) (Crys. Structure, Experimental, Mechan. Prop., 10) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009
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Landolt‐Bo¨rnstein New Series IV/11E1
Al–Ta–Ti
15
Aluminium – Tantalum – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Tamara Velikanova, Mikhail Turchanin, Svitlana Ilyenko, Guenter Effenberg, Vasyl Tomashik, Dmytro Pavlyuchkov
Introduction Alloys based on the Al-Ta-Ti system are promising candidate materials for various structural applications. Alloys based on the γαffl phase (TiAl) have an excellent potential to become an important next generation aerospace materials because of their low density, high melting temperature, good elevated-temperature strength, high resistance to oxidation and hydrogen absorption, and excellent creep properties. The alloys of the Al-Ta-Ti system with the composition in the range 45–50 at.% Al are potential high temperature engineering materials. Additions of Ta to the γαffl phase (TiAl) are useful to improve the fracture toughness, ductility and corrosion and oxidation resistance at high temperatures. However, there is still a scope for alloy development to optimize microstructures of these alloys. Several experimental investigations of phase relations in the system were undertaken and, as a result, a set of isothermal sections at 1000 [1966Ram, 2000Kai], 1100 [1983Sri, 1991McC2, 1991Das, 1992Per, 1993Das, 1993Jew], 1200 [2000Kai], 1300 [2000Kai], 1330 [1991Boe], 1350 [1995Wea], 1440 [1991McC2, 1993Das], and 1450˚C [1991Wea, 1995Wea] is reported. Based on the solidification behavior of Al-Ta-Ti alloys [1991McC1, 1991McC2, 1992McC] constructed a partial liquidus surface. Table 1 summarizes the experimental studies on phase equilibria. The above mentioned experimental results were reviewed in [2005Das, 2005Rag].
Binary Systems The Al-Ta phase diagram from [1996Du] (Fig. 1), Ta-Ti from [Mas2] are accepted. For the AlTi system, the MSIT evaluation from [2004Sch] is available. More recently, a new review by [2006Sch] and a thermodynamic assessment by [2007Wit] appeared which show differences in two regions. [2006Sch] and [2007Wit] show two peritectoid reactions involving (αTi), (βTi) and Ti3Al and a small two-phase field between (βTi) and Ti3Al in the 1150–1200˚C temperature range. Evaluating the same experimental literature information, [2004Sch] concluded that Ti3Al transforms congruently and that there is no invariant equilibrium involving these three phases. According to [2004Sch], the different phase equilibria may be related to very small differences in Gibbs energy, i.e. very small driving forces, leading to phase transitions in that range. Thus, it is very difficult to decide which interpretation of the experimental data is actually correct.
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The second difference considers the equilibria between γ(TiAl), TiAl2 and the onedimensional antiphase domain structures or “long period structures (LP)” which is stable at high temperature in the range of 65 to 75 at.% Al. The assessments by [2006Sch] and [2007Wit] are based on more and recent experimental data and give a more complete discussion of the stable equilibria in this region. Thus the phase diagram calculated by [2007Wit] is accepted in the present evaluation because of its larger base of experimental data. The majority of these were assessed in [2006Sch] and complemented by own key experiments. The Al-Ti phase diagram from [2007Wit] is included in the present volume in the evaluation of the Al-C-Ti system.
Solid Phases No ternary compounds exist in the system. Above 882˚C a continuous solid solutions β, (βTi, Ta) is found and above 724˚C a continuous solid solutions ε, (Ti,Ta)Al3 forms [1966Ram, 1983Sri, 1990Abd, 1993Das, 1993Jew]. Single crystals of the (Ti,Ta)Al3 compound were produced by spontaneous crystallization during slow cooling of Al-0.38Ti-0.10Ta (at.%) melt [1990Abd]. The chemical composition of the aluminide crystals separated from melt was Al-75Ti-20Ta (at.%). The crystals had the same crystal structure as ε(h) (TiAl3) and were formed by partial substituting in the lattice Ti atoms by Ta atoms. At 1100˚C the β phase (including the ordered β0 phase) extends up to about 40 at.% Al into the ternary system. In the range of 1350 to 1450˚C it extends to 42–50 at.% Al after [1995Wea, 1993Das, 1993Jew]. The ordering of the bcc phase was established based on the existence of the thermal antiphase boundaries in ternary alloys quenched from 1350 and 1450˚C in [1995Wea]. Such a structure could indicate that these alloys are disordered at high temperature and undergo ordering during cooling. The ordered phase was experimentally revealed at 1100˚C by [1993Jew, 1993Das]. [1993Das] reported that the ordering takes place in the bcc phase (coexisting with other phases) within a composition range of about 40 to 60 at.% Ti and 15 to 60 at.% Al. Near the composition Ti-25Al-25Ta (at.%) the existence of the β0 phase in the ternary region is established. The phase reactions in a Ti-33Al-17Ta (at.%) alloy were examined in detail following solidification and solid-state processing treatments. Differential thermal analysis (DTA) proved that β0-phase in this alloy is stable up to 1205˚C, where it experiences a solid-state order to disorder transformation. The ordering reaction is too fast that it could not be arrested by rapid solidification processing. It is the presence of thermal antiphase boundaries in the microstructure that confirms the solid-state ordering of the β0 phase from the disordered β phase. Taking into account that the β0 phase exists in the binary Al-Ti system at about 1100 to 1430˚C, its range in the ternary system has to adjoin to the Al-Ti side of the Gibbs triangle in the above mentioned temperature interval. Thus, its composition limits depend considerably on temperature. The solubility of the β stabilizer Ta in (αTi) was determined in [1963Luz]. A single-phase region around Ti4Al3Ta, denoted α*, is assumed by [1993Jew] to be a new ternary phase based on α2 (of the close crystal structure). It was observed both in bulk alloys at 1100˚C and also in some of the diffusion couples. There is, however, no confirmation of the existence of this phase by the other authors. Crystallographic data of the Al-Ta-Ti phases and their temperature ranges of stability are listed in Table 2. DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009
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Invariant Equilibria No experimentally established invariant equilibria in the ternary system are available. A tentative reaction scheme is published by [2005Rag] based on the limited data on monovariant lines on the liquidus, namely β/α, α/γ and γ/η [1992McC, 1991McC1, 1991McC2], the data for the isothermal sections at 1440˚C to 1000˚C [1993Jew, 1995Wea, 1991Boe, 2000Kai] and the boundary binary phase diagrams accepted from [Mas2]. A partial reaction scheme is presented in Fig. 2 for liquid/solid equilibria based only on the accepted tentative liquidus surface projection shown in Fig. 3. Since no data on the temperatures of invariant equilibria in the ternary system exist, in the reaction scheme they were approximated by [2005Rag] based on the binary systems and on the conclusion of [1992McC, 1991McC1, 1991McC2] concerning the direction of the liquidus surface decrease in the ternary system. All the five invariant reactions of Fig. 2 are shown in Fig. 3 without providing temperatures.
Liquidus, Solidus and Solvus Surfaces Limited information on the liquidus surface, obtained by [1991McC1, 1991McC2, 1992McC], concerns only the monovariant l Ð β + α, l Ð α + γ lines and partly l Ð γ + ε. A steep slope of the liquidus β- and α primary crystallization surfaces towards the Al-Ti binary side was concluded. This was concluded from microstructure observations during solidification of ternary alloys in the vicinity of the 50 at.% Al isoconcentrate above 20 at.% Ti. Below 20 at.% Ti an estimated monovariant line l Ð β + σ has been assumed, however not supported by experimental points. No data are available on the solidification of ternary alloys in other parts of the system. From the above mentioned information together with the data on the phase equilibria at 1440˚C by [1993Jew], the existence of two four-phase peritectic reactions was proposed by [2005Rag] who published a schematic liquidus projection and the above tentative reaction scheme including solid state equilibria. The tentative liquidus projection taken from [2005Rag] is presented here in Fig. 3 with amendments made according to the accepted binary diagrams. There is no explicit experimental work describing the equilibria on the solidus surface. In the present work it is assumed that it must be similar to the highest isothermal sections studied in literature by [1995Wea] at 1450˚C and by [1993Jew, 1993Das] at 1440˚C.
Isothermal Sections The isothermal sections at 1450, 1440, 1350, 1330, 1300, 1200, 1100 and 1000˚C were studied experimentally [2000Kai, 1995Wea, 1993Das, 1993Jew, 1991Boe, 1983Sri, 1966Ram]. The experimental results of the above works are in satisfactory agreement with each other. [2005Das] discussed and presented isothermal sections at 1450, 1440, 1350, 1100 and 1000˚C which he derived from the works of [1995Wea, 1993Jew, 1993Das, 1983Sri, 1966Ram]. [2005Rag] considered the above sections at 1440, 1350 and 1100˚C, an isothermal section at 1330˚C constructed from the data of [1991Boe] and partial sections at 1300 and 1200˚C after [2000Kai]. Both review papers, [2005Das] and [2005Rag], apply corrections to Landolt‐Bo¨rnstein New Series IV/11E1
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add missing phase fields or to ensure consistency with the accepted binary diagrams and the phase relations proposed are not in conflict. In this work the isothermal sections given in Figs. 4 to 11 are drawn from the experimental data. They take into account the here accepted edge binary diagrams and the acceptable corrections made in [2005Das] and [2005Rag]. The phase diagrams at 1450 and 1440˚C, Figs. 4 and 5, are constructed after [1995Wea] and [1993Jew, 1993Das]. Both show the three-phase fields σ + β + α, σ + α + δ, α + δ + η and α + η + L and, in addition, at 1440˚C the σ + TaAl + δ field appears because of the existence of the equiatomic tantalum aluminide at this temperature. The solid solution of Al in the α phase reaches up to 85–87 at.% Al at 1450, 1440˚C and then sharply decreases with decreasing temperature after [1995Wea] and [1993Jew, 1993Das], as one can see at 1350 and 1100˚C, Figs. 6 and Fig. 10 respectively. The γ solid solution region at 1450˚ should reach the Al-Ti side unlike the interpretations given by [1995Wea] in his original figure and by [2005Das] in his review. The γ phase is shown to permeate into the ternary system up to about 80 at.% Ta at 1450, 1440 and 1350˚C. Its extension varies only slightly with temperature as one can see from comparison with the data at 1100˚C. The ε phase extension at 1440˚C given by [1993Das, 1993Jew] had to be corrected to agree with that at 1450 and 1350˚C by the same authors. At 1440˚C and below, equilibria with the participation of tantalum monoaluminide had to be added. Below 1345˚C the Ta5Al7 phase exists in the isothermal section and below 1226˚C the Ta2Al3 phase appears. But the Al69Ta39 phase (δ) disappears at 1183˚C from the phase equilibria because of the eutectoid decomposition of this phase. For the equilibria involving α and δ a principal change takes place in the temperature interval of 1440 to 1330˚C. The α + δ equilibria change for σ + ε and σ + γ and these equilibria are retained down to the lowest temperature under investigation (1000˚C [1966Ram]). The existence of the γ + σ equilibrium and a significant temperature dependence of the tantalum solubility in γ were confirmed by [1992McC] who observed the decomposition of γ into a two-phase (γ + σ) lamellar structure. A partial isothermal section at 1330˚C after [1991Boe] is given in Fig. 7. At 1350˚C the ε phase reaches up to the Al-Ti binary. At the high Ti border it is in equilibrium with γ and at the high Al border with L, contrary to [1995Wea], who assumed that L is in equilibrium with γ. Other equilibria missed out by 1995Wea] concern the equilibria with the TaAl and δ phases are shown in the isothermal section at 1350˚C (Fig. 6). That means the two three-phase fields σ + δ + ε and σ + δ + TaAl have been added tentatively by dashed lines. Also the β0, the ordered bcc phase, which exists at this temperature in the binary Al-Ti is added to Fig. 6. The β0 ternary solid solution was found by [1993Das, 1993Jew] at 1100˚C, just below the temperature where it disappears in the binary Al-Ti system. The 1100˚C isothermal section in Fig. 10 is based on the experimental data of [1993Jew, 1993Das] with the following corrections: the position of the β (β0)-phase fields is adjusted to match the accepted Al-Ti diagram; equilibria with the η and ζ phases at the Al-Ti side have been added as well as the equilibria with the Ta5Al7 and Ta2Al3 phases near the Al-Ta side. The main changes in the phase relationships at 1100˚C - with respect to higher temperatures - are (i) that the α phase field became very small (supported by [1983Sri]), (ii) an extended α2 field exists and (iii) an equilibrium exists between γ and β0 (the γ + β0 field in Fig. 10) instead of α + σ at higher temperatures. If one compares the isothermal sections at 1300 and 1200˚C, Figs. 8 and 9, published by [2000Kai] a transformation α + σ Ð γ + β0 must have taken place in this temperature interval. DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009
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The isothermal section at 1000˚C given in Fig. 11 is based on the preliminary data published by [1966Ram], again amended to match the accepted binary diagrams. A remarkable difference of the phase relationships from that at 1100˚C is the existence of the σ + α2 equilibrium instead of β0 + γ and the absence of the ordered bcc (β0) phase, which is in agreement with its absence in the binary Al-Ti. The extension of the γ filed into the ternary system seems to be too large by [1966Ram] and has been reduced slightly in Fig. 11 to nearly the same extension as at high temperatures. The phase equilibria at 1100˚C, reported by [1983Sri] and later accepted by [1993Kub], look like those at 1000˚C published by [1966Ram]. They are significantly different from those established by [1993Jew, 1993Das]. The reason of this contradiction is not quite clear, although it is suspected that contamination during the sample preparation could be the reason. The initial materials were of about the same purity and the methods of preparation and investigations were almost the same. The alloys in [1983Sri] were prepared by sintering the powder mixtures of the components as well as by arc melting and subsequently annealed in quartz ampoules. The duration of thermal treatment was at least 170 h. As stated in [1983Sri], the alloys could be contaminated by silicon during the preparation, because the impurity phases Ti5Si3 and Ta5Si3 were observed. In contrast to [1983Sri], [1966Ram] prepared specimens by arc-melting and heat treated them for 7 d in evacuated quartz containers, encapsulated by Ta-foil with subsequent quenching. Since the results of [1983Sri] do not follow the general trend with changing temperature, these data are not accepted in the present evaluation.
Notes on Materials Properties and Applications The works devoted to the study of various mechanical, structural and physical properties [1993Has, 1995Muk, 1999Hao, 2001Gar, 2002Red, 2005Leg, 2006Joh, 2006How] are summarized in Table 3. Electrical resistance measurements and hardness tests as well as X-ray and metallographic analyses were carried out by [1963Luz] to determine the solubility of the β stabilizer Ta in (αTi). Mechanical properties of Ti-46.4Al-2.5Ta (at.%) alloy were investigated in [1993Has]. With increasing temperature from 800 to 1100˚C the elongation of sample increased from 16 to 104%. At the same time yield stress decreased from 394 to 28.4 MPa. The X-ray technique was used together with micro-indentation to measure average residual stresses and local hardness of the alloy of Ti-47Al-2Ta (at.%) composition [2001Gar]. The measured values of residual stresses were compared to a composite cylinder model, which incorporates the effect of creep. In general, the residual stresses decrease as heat treatment temperature increases, reaching a minimum and then increases as the temperature and/or time of treatment increases. The fiber-matrix interface shows no degradation during the heat treating process. Such results compare well with the composite cylinder model once the effect of creep is incorporated. The analysis of the micro-hardness results indicate an increase in hardness in regions close to the fiber-matrix interface at 593˚C but a softening occurs at 815˚C. This behavior is reversed once the heat treatment temperature is increased to 982˚C. At this temperature, the hardness at the interface increases for prolonged heat treatment duration. Landolt‐Bo¨rnstein New Series IV/11E1
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Miscellaneous Mechanical alloying of blended elemental Ti-25Al-25Ta (at.%) was investigated in [1995Muk]. Mechanical alloying led initially to the formation of a disordered bcc phase with a = 332 pm, and at still longer times a fully amorphous phase. The disordered bcc phase transformed to the B2 (the ordered bcc) structure after heat treatment at 600˚C and to an orthorhombic phase (with a = 608.9 pm, b = 857.0 pm, c = 498.0 pm) after a heat treatment at 800˚C. The site occupancies of Ta in γαffl (TiAl) alloys with different compositions, and in Ti3Al with the compositions of Ti - 26 at.% Al - (1 to 2) at.% Ta, were measured by the atom location channeling enhanced microanalysis (ALCHEMI) method [1999Hao]. For TiAl alloys, the results show that Ta atoms invariably occupy Ti sites. For Ti3Al alloys Ta atoms occupy Ti sites too. The experimental results were interpreted in terms of a Bragg-Williams-type model and bond-order data were obtained from electronic structure calculation. Qualitative agreement between the model and measurements was reached. A theoretical model that relates the (γ + α2) two-phase equilibrium in ternary Al-Ta-Ti alloys to the substitution behavior of alloying elements in the two ordered phases has been suggested in [1999Yan]. The method used experimentally measured occupation probabilities of the alloying elements on different sublattices as input, and allowed estimates of the mole fractions of the two phases for a given alloy composition. The review summarized theoretical and experimental investigations of sublattice substitution of alloying elements in γαffl phase (TiAl) presented in [2000Yan]. The isothermal oxidation behavior of a ternary Ti-25Al-18Ta (at.%) Ta intermetallic alloy has been investigated in pure oxygen over the temperature range of 850 to 1100˚C [2002Red]. The oxidation kinetics was found to follow a parabolic rate. Effective activation energy of 259 kJ·mol–1 was deduced from the oxidation data. The oxidation products were a mixture of TiO2, the main component, Al2O3 (alumina), and small amounts of tantalum oxide. The addition of Ta to Ti3Al alloy decreased the oxidation rate of the alloy. However, the oxidation scale was not compact and exhibited significant spallation especially at high temperatures. The microchemistry and morphology of the oxide layer produced on a single γαffl phase (TiAl) alloy (Ti-52.1Al-2Ta (at.%)) following anneals at 550 and 900˚C in low-pressure oxygen and hydrogen environments were studied in [2005Leg]. The oxide structure on the electropolished surface consisted of two layers, an outer one of Al2O3 and an inner one of TiO2. Annealing in low partial pressures of oxygen retained the same oxide structure but increased the total thickness. After the low-pressure oxygen anneals at 550˚C, the oxide surface was smooth, whereas after anneals at 800˚C and above the surface consisted of small, round particles, whose size and density increased with increasing annealing temperature and oxygen partial pressure. In contrast, the oxide structure produced by annealing in a hydrogen environment after the low-pressure oxygen treatment consisted of an outer layer of TiO2 and a sub-oxide of Al2O3. The morphology of this oxide consisted of elongated, rod-like particles, whose size increased with annealing temperature. Microstructural selection during the directional solidification of binary γαffl (TiAl) alloys grown from Ti-52Al-8Ta (at.%) seeds was examined [2006Joh]. By using a seed crystal, the high-temperature hcp α phase can be correctly oriented so that an aligned lamellar γαffl (TiAl)/ Ti3Al microstructure results from the subsequent solid-state transformations upon cooling. From the equilibrium phase diagram, primary bcc β phase solidification is expected in the compositional range for which the binary γ (TiAl) alloys were successfully seeded. Thus successful crystal growth of γ (TiAl) was found to be dependent upon the undercooling DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009
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necessary for β nucleation. From these data, a microstructural selection map for the seeded growth of the α phase was constructed. The atomic-level structural, compositional, kinetic and mechanistic aspects of phase transformations that occur during growth of product phases in fcc to hcp reactions were investigated in [2006How] for Ti-48Al- 2Ta (at.%) alloy.
. Table 1 Investigations of the Al-Ta-Ti Phase Relations and Structures Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1963Luz]
Electrical resistance measurements, hardness tests, X-ray diffraction (XRD), metallographic analysis
solubility of Ta in (αTi).
[1966Ram]
XRD
isothermal section at 1000˚C
[1983Sri]
XRD
isothermal section at 1100˚C
[1990Abd]
XRD, scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS)
crystal structure of ε phase
[1991Boe]
metallographic analysis, transmission alloy Al-25Ti-25Ta (at.%) at 1200 to electron microscopy (TEM), SEM, electron 1550˚C, isothermal section at 1330˚C probe microanalysis (EPMA), differential thermal analysis (DTA)
[1991Das]
TEM, EPMA, DTA, diffusion couple technique
isothermal section at 1100˚C
[1991McC1] SEM, TEM, EDS
undercooled and splat quenched alloy Ti-15Ta-48Al (at.%), solidification path
[1991McC2] XRD, SEM, TEM, EDS
alloys at 50 at.% Al, isothermal section at 1100˚C, partial liquidus surface
[1991Wea]
metallographic analysis, XRD, TEM, DTA
alloys at 26 to 70 at.% Al, 8 to 54 at.% Ti, 5 to 50 at.% Ta, isothermal section at 1450˚C, morphology
[1992Kim]
metallographic analysis, XRD, EPMA
alloy Ti-46.4Al-2.5Ta (at.%) at 1000 to 1300˚C, phase constitution, morphology
[1992McC]
XRD, SEM, TEM, EDS
alloys Ti-0 to 27 at.% Ta-45 to 70 at.% Al, partial liquidus surface, morphology and phase constitution at 1200˚C
[1992Per]
SEM, EPMA, DTA, XRD
isothermal section at 1100˚C
[1992Wea]
metallographic analysis, SEM, EPMA, TEM, alloys at 64 to 71 at.% Al, 7 to 13 at.% Ti, XRD 16 to 32 at.% Ta, isothermal section at 1450˚C
[1993Das]
TEM, SEM, EPMA, DTA, XRD study of diffusion couple and bulk alloy samples
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1993Jew]
metallographic analysis, TEM, SEM, EPMA, isothermal sections at 1100 and 1440˚C DTA, XRD
[1993Has]
metallographic analysis, XRD
[1995Wea]
metallographic analysis, SEM, EPMA, TEM alloys at 38 to 55 at.% Al, 19 to 30 at.% Ti, 20 to 38 at.% Ta, solid state phase transformations, isothermal sections at 1350 and 1450˚C, morphology
[2000Kai]
SEM, EPMA
alloy Ti-47Al-6Ta (at.%) at 1000 to 1300˚C, isothermal sections at 1000, 1200 and 1300˚C
[2003Das]
DTA, TEM, splat-quenching technique
ordering of the bcc-phase
alloy Ti-46.4Al-2.5Ta (at.%) at 1100˚C, phase constitution, morphology
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(Al) < 660.5
cF4 Fm3m Cu
β, (βTi,Ta) Ti1–x–yTaxAly < 3020
cI2 Im3m W
(Ta) < 3020 (βTi) 1670 - 882 β0, Ti1–x–yTaxAly cP2 < 1427 Pm3m CsCl
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Lattice Parameters [pm] a = 404.96 a = 405.2
Comments/References pure Al at 25˚C [Mas2] Solubility limits: at 0% Ti, 0.04 at.% Ta [1972Fer], at 0% Ta, 0.8 at.% Ti [2007Wit] or 0.6 at.% Ti [1992Kat]
a = 330.30
complete solid solution [Mas2], 0 < x < 1 at T = 1440˚C, x = 0 to 1 and y = 0 to 0.5 [1993Das]; at T = 1350˚C, x = 0 to 1 and y = 0 to 0.42 [1995Wea]; at T = 1100˚C, x = 0 to 1 and y = 0 to 0.37 [1993Das]; pure Ta at 25˚C [Mas2]
a = 330.65
pure Ti at 900˚C [Mas2]
-
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. Table 2 (continued) Phase/ Temperature Range [˚C] α, (αTi) Ti1–x–yTaxAly 1491 - 1119 and < 1159
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
hP2 P63/mmc Mg
a = 295.08 c = 468.35
α2, Ti3Al Ti1–x–yTaxAly < 1189
hP8 P63/mmc Ni3Sn
α*, Ti1–x–yTaxAly crystal structure close to α2
γ, TiAl Ti1–x–yTaxAly < 1463
ζ, Ti2Al5 Ti1–x–yTaxAly 1432 - 976
1d-APD
tP28 P4/mmm “Ti2Al5”
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at x = 0: 17 to 38.5 at.% Al [2007Wit]; at 22 at.% Al [2004Sch] at 38 at.% Al [2004Sch] at T = 1100˚C, y = 0.33 to 0.48 and x = 0.07 to 0.09 [1993Das] Ternary phase? The existence needs confirmation
-
a = 400.0 c = 407.5 a = 398.4 c = 406.0 set of tetragonal superstructures* based on AuCu-lattice;
at T = 1440˚C, x = 0.28 and y = 0.60 [1993Das]; at T = 1350˚C, x = 0.22 and y = 0.51 [1995Wea]; at x = 0.02 and T = 1100˚C, y = 0.2 [1993Das]; pure Ti, at 25˚C [V-C2] at x = 0, y = 0 to 0.28 [2007Wit]; at x = 0 and T = 1189˚C, y = 0.32 [2007Wit]; at x = 0 and T = 1456˚C, y = 0.5 [2007Wit]; 0.27 < y < 0.48 and 0 < x < 0.09 at 1100˚C (including α2*)
tP4 P4/mmm AuCu
TiAl
Comments/References
at T = 1100˚C, 0<x<0.15 and 0.50
a* = 395.3 c* = 410.4 a* = 391.8 c* = 415.4 a = 390.53 c = 2919.63
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at 66 at.% Al for AuCu subcell only [2001Bra]; at 71 at.% Al for AuCu subcell only [2001Bra]; 63.7 to 70.4 at.% Al [2007Wit]; “Ti2Al5”, 1215 - 985˚C [1990Sch]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
η, TiAl2 Ti1–x–yTaxAly < 1224
tI24 I41/amd HfGa2
a = 397.0 c = 2497.0
Ti3Al5 < 810
tP32 P4/mbm Ti3Al5
a = 1129.3 c = 403.8
ε, (Ti,Ta)Al3(h) Ti1–x–yTaxAly < 1608
tI8 I4/mmm TiAl3(h)
a = 384.9 c = 860.9
TiAl3(h) < 1396 TaAl3 < 1608
a = 384.88 c = 859.82 a = 383.9 c = 853.5
ε(l), TiAl3(l) < 932
tI32 I4/mmm TiAl3(l)
σ, Ta2Al Ti1–x–yTaxAly < 2061
tP30 P42/mnm σCrFe
a = 387.7 c = 3382.8
65.1 to 67.3 at.% Al [2007Wit]; stable structure at T < 1216˚C [2001Bra]; denoted TiAl2(r) by [1990Sch] [2007Wit]; [2001Bra] 0 < x < 1 at T = 1000˚C [1966Ram], at T = 1100˚C [1993Das], at T = 1350˚C [1995Wea]; at 1396 to 724˚C: 0 < x ≲ 0.25 and y = 0.75; at 1440˚C: 0.05 < x ≲ 0.25 and y ≲ 0.75 [1993Das]; at x = 0, 74.5 to 76.5 at.% Al [2007Wit]; at x = 0, 74.5 to 75 at.% Al [2001Bra]; at y = 0 [1996Du]; 75 at.% Al [1990Sch]; 75 at.% Al [V-C2, 1985Sch] 74.6 to 75.7 at.%Al [2007Wit]; 74.5 to 75 at.%Al [2001Bra] 0.20 < y < 0.44 at x ≲ 0.40 and 0.20 < x+y < 0.26 for 1100 to 1450˚C [1993Das, 1995Wea]; y = 0.40 and x = 0.40 at 1350˚C [1995Wea]; y = 0.35 and x ≲ 0.40 at 1100˚C [1993Das]
Ta2Al a = 982.0 c = 523.0 a = 998.0 c = 516.0 a = 986.4 c = 521.5
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Comments/References
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65 to 79 at.% Ta [1996Du]; ≲ 50 to 70 at.% Ta at <2000˚C [1985Sch]; at 67 at.% Ta [1972Fer]; at 75 at.% Ta [1972Fer]; Ta2Al [1972Fer]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
TaAl < 1446
mP*
Ta5Al7 < 1345
hP*
Ta2Al3 < 1225
aP*
δ, Ta39Al69 1548 - 1183
cF444 F43m Ta39Al69
Lattice Parameters [pm] a = 1485.2 b = 988.1 c = 986.4 β = 99.99˚
Comments/References [1996Du]; [1993Mah]
[1996Du]; [1995Mah]; above 700 atoms in the unit cell
a = 3214 c = 1341 a = 1074.8 b = 1107.8 c = 1043.3 α = 90.45˚ β = 97.60˚ γ = 63.19˚ a = 1915.3
[1996Du]; [1995Mah]
[1996Du]; [1994Mah]
. Table 3 Investigations of the Al-Ta-Ti Materials Properties Reference
Method / Experimental Technique
Type of Property
[1963Luz]
Electrical resistance measurements
electrical resistivity of (αTi)-phase
[1993Has]
Tensile tests, SEM
elongation and yield stress of Ti- 46.4Al2.5Ta (at.%) alloy at 800 to 1100˚C
[1995Muk] Ball milling, XRD
crystal structure of metastable mechanically alloyed alloys of Ti-25Al- 25Ta (at.%) composition
[1999Hao] atom location channeling enhanced microanalysis (ALCHEMI) method
site occupancies of Ta in,TiAl alloys with different compositions and in Ti3Al with the compositions of Ti - 26 at.% Al - (1 to 2) at.% Ta
[2001Gar]
XRD, micro-indentation measurements
micro-hardness and residual stresses of Ti-47Al-2Ta (at.%) alloy
[2002Red]
XRD
oxidation behavior of Ti-25Al-18Ta (at.%) alloy at 850 to 1100˚C
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. Table 3 (continued) Reference [2005Leg]
Method / Experimental Technique
Type of Property
SEM, Auger electron spectrometry, X-ray oxidation behavior of Ti-52.1-Al-2Ta (at.%) photoelectron spectrometry, secondary alloy at 550 and 900˚C ion mass spectrometry
[2006How] high resolution TEM
atomic-level structural, compositional, kinetic and mechanistic aspects of phase transformations of fcc to hcp
[2006Joh]
microstructural selection during the directional solidification of γ alloys (TiAl)
SEM, optical microscopy
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. Fig. 1 Al-Ta-Ti. The Al-Ta binary system
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. Fig. 2 Al-Ta-Ti. Partial reaction scheme
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. Fig. 3 Al-Ta-Ti. Tentative liquidus projection
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. Fig. 4 Al-Ta-Ti. Isothermal section at 1450˚C
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. Fig. 5 Al-Ta-Ti. Isothermal section at 1440˚C
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. Fig. 6 Al-Ta-Ti. Isothermal section at 1350˚C
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. Fig. 7 Al-Ta-Ti. Partial isothermal section at 1330˚C
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. Fig. 8 Al-Ta-Ti. Partial isothermal section at 1300˚C
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. Fig. 9 Al-Ta-Ti. Partial isothermal section at 1200˚C
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. Fig. 10 Al-Ta-Ti. Isothermal section at 1100˚C
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. Fig. 11 Al-Ta-Ti. Isothermal section at 1000˚C
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References [1963Luz]
[1966Ram] [1972Fer] [1983Sri]
[1985Sch] [1990Abd]
[1990Sch] [1991Boe]
[1991Das] [1991McC1]
[1991McC2]
[1991Wea]
[1992Kat]
[1992Kim]
[1992McC]
[1992Per]
[1992Wea]
[1993Das]
Luzhnikov, L.P., Novikova, V.M., Mareev, A.P., “Solubility of β Stabilisers in α-Ti”, Met. Sci. Heat Treat., 5(2), 78–81 (1963), translated from Metallov. Term. Obrab. Met., (2), 13–16 (1963) (Experimental, Phase Diagram, Phase Relations, 4) Raman, A., “X-Ray Studies of Some T-T5-Al-Systems” (in German), Z. Metallkd., 57, 535–540 (1966) (Crys. Structure, Phase Relations, Phase Diagram, 5) Ferro, R., “Tantalum Crystal Structure and Density Data”, Atomic Energy Review, Special Issue No.3, IAEA, Vienna, 67–129 (1972) (Crys. Structure, Review) Sridharan, S., Nowotny, H., “Studies in the Ternary System Ti-Ta-Al and in the Quaternary System Ti-Ta-Al-C”, Z. Metallkd., 74, 468–472 (1983) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, 8) Schuster, J.C., “Phases and Phase Relations in the System Ta-Al”, Z. Metallkd., 76, 724–727 (1985) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, 13) Abdel-Hamid, A.A., “Crysatllization of Complex Aluminide Compounds fron Dilute Al-Ti Melts Containing One or Two Other Transition Metals of IVB to VIB Groups”, Z. Metallkd., 81(8), 601–605 (1990) (Crys. Structure, Experimental, 16) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81, 389–396 (1990) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, Review, #, 33) Boettinger, W.J., Shapiro, A.J., Cline, J.P., Gayle, F.W., Bendersky, L.A., Biancaniello, F.S., “Investigation of the Phase Constitution of Al2TiTa”, Scr. Metall. Mater., 25, 1993–1998 (1991) (Crys. Structure, Experimental, Morphology, Phase Diagram, 8) Das, S., Perepezko, J.H., “Ternary Phase Development in the Ti-Al-Ta System”, Scr. Metall. Mater., 25(5), 1193–1198 (1991) (Experimental, Phase Relations, Phase Diagram, 20) McCullough, C., Valencia, J.J., Levi, C.G., Mehrabian, R., “Microstructural Benefits of RSP in a γ-(TiTa-Al) Alloy”, Alloy Phase Stability and Design Symp., Mater. Res. Soc. 1991, Pittsburgh, PA, USA, 186, 155–160 (1991) (Experimental, Morphology, Phase Diagram, Phase Relations, 8) McCullough, C., Valencia J.J., Levi, G.G., Mehrabian, R., Maloney, M., Hecht, R., “Solidification Paths of Ti-Ta-Al Alloys”, Acta Metall. Mater., 39(11), 2745–2758 (1991) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 19) Weaver, M.L., Guy, S.L., Stone, R.K., Kaufman, M.J., “An Investigation of Phase Equilibria in the Ternary Al-Ti-Ta System”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 163–168 (1991) (Experimental, Phase Diagram, 6) Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the Ti-Al System”, Metall. Trans. A, 23(8), 2081–2090 (1992) (Assessment, Calculation, Phase Relations, Phase Diagram, Thermodyn., 51) Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems at High Temperatures”, Mater. Sci. Eng. A, 152A(1–2), 54–59 (1992) (Abstract, Crys. Structure, Experimental, Phase Diagram, 12) McCullough, C., Levi, C.G., Valencia, J.J., Mehrabian, R., “Peritectic Solidification of Ti-Al-Ta Alloys in the Region of γ-TiAl”, Mater. Sci. Eng. A, 156(2), 153–166 (1992) (Experimental, Morphology, Phase Diagram, 32) Perepezko, J.H., Jewett, T.J., Das, S., Mishurda, J.C., “High Temperature Phase Stability in Ternary Titanium Aluminides”, Society for the Advancement of Material and Process Engineering. P.O. Box 2459, Covina, California 91722, USA. Conference: Advancements in Synthesis and Processes, Toronto, Canada, 20–22 Oct. 1992, M357-M365 (1992) (Experimental, Phase Diagram, 15) Weaver, M.L, Kaufman, M.J., “An Investigation of Al2Ta and Related Phases in the Ternary Al-Ta-Ti System”, Scr. Metall. Mater., 26(3), 411–416 (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 19) Das, S., Jewett, T.J., Perepezko, J.H., “High Temperature Phase Equilibria of Some Ternary Titanium Aluminides” in “Structural Intermetallics”, Proc. 1st Int. Symp. Struct. Intermetallics, Champion, Pa., Sept. 1993, Daralia, R., Lewandovski, J.J., Liu, C.T., Martin, P.L., Miracle, D.B., Nathal, M.V. (Eds.), TMS, Warrendale, Pa., 35–43 (1993) (Experimental, Phase Diagram, Phase Relations, 46)
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[1993Jew]
[1993Kub]
[1993Mah]
[1994Mah] [1995Mah]
[1995Muk]
[1995Wea]
[1996Du] [1999Hao] [1999Yan]
[2000Kai]
[2000Yan]
[2001Bra]
[2001Gar]
[2002Red] [2003Das]
[2004Sch]
[2005Leg]
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Hashimoto, K., Masao, K., “Effects of Third Element Addition on Mechanical Properties of TiAl” in “Structural Intermetallics”, Proc. 1st Int. Symp. Struct. Intermetallics, Champion, Pa., Sept. 1993, Daralia, R., Lewandovski, J.J., Liu, C.T., Martin, P.L., Miracle, D.B., Nathal, M.V. (Eds.), TMS, Warrendale, Pa., 309–318 (1993) (Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 18) Jewett, T.J., Das, S., Perepezko, J.H., “High Temperature Phase Equilibria in the Al-Ta-Ti Ternary System” in “Titanium 92: Sci. Technol.”, Froes, F.H., Caplan, I. (Eds.), Proc. Symp. 1992 (Pub. 1993), TMS, 1, 713–719 (1993) (Experimental, Phase Diagram, Phase Relations, 13) Kubaschewski, O., “Al-Ta-Ti (Aluminum-Tantalum-Titanium)”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.20331.1.20 (1993) (Phase Diagram, Phase Relations, Review, 6) Mahne, S., Krumeich, F., Harbrecht, B., “Phase Relations in the Al-Ta System: On the Translational Symmetries of Al3Ta2 and AlTa”, J. Alloys Compd., 201, 167–174 (1993) (Crys. Structure, Experimental, Phase Diagram, 20) Mahne, S., Harbrecht, B., “Al69Ta39 - a New Variant of a Face-Centred Cubic Giant Cell Structure”, J. Alloys Compd., 203, 271–279 (1994) (Crys. Structure, Experimental, 28) Mahne, S., Harbrecht, B., Krumeich, F., “Phase Relations in the Al-Ta System: on the Translational Symmetrics of a Triclinic Structure and a New Hexagonal Giant Cell Structure”, J. Alloys Compd., 218(2), 177–182 (1995) (Crys. Structure, Experimental, Phase Diagram, 9) Mukhopadhyay, D.K., Suryanarayana, C., Froes, F.H., “Synthesis of Metastable Phases in Blended Elemental Ti-Al-Ta Powders by Mechanical Alloying”, Synthesis/Processing Light Weight Metallic Mater., Proc. Symp. TMS Annual meetings, TMS., Warrendale, PA, USA, 65–73 (1995) (Crys. Structure, Experimental, Phase Relations, 26) Weaver, M.I., Kaufman, M.J., “Phase Relationships and Transformations in the Ternary AluminumTitanium-Tantalum System”, Acta Metall. Mater., 43(7), 2625–2640 (1995) (Experimental, Kinetics, Mechan. Prop., Morphology, Phase Diagram, Phase Relations, 24) Du, Y., Schmid-Fetzer, R., “Thermodynamic Modeling of the Al-Ta System”, J. Phase Equilib., 17(4), 311–324 (1996) (Assessment, Calculation, Phase Relations, Thermodyn., 50) Hao, Y.L., Xu, D.S., Cui, Y.Y., Yang, R., Li, D., “The Site Occupancies of Alloying Elements in TiAl and Ti3Al Alloys”, Acta Mater., 47(4), 1129–1139 (1999) (Crys. Structure, Experimental, 41) Yang, R., Hao, Y.L., “Estimation of (γ+α2) Equilibrium in Two-Phase Ti-Al-X Alloys by Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41(3), 341–346 (1999) (Calculation, Phase Relations, 13) Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among α (hcp), β (bcc) and Gama (L1(0)) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855–867 (2000) (Crys. Structure, Experimental, Phase Relations, 29) Yang, R., Hao, Y., Song, Y., Guo, Z.X., “Site Occupancy of Alloying Additions in Titanium Aluminides and its Application to Phase Equilibrium Evaluation”, Z. Metallkd., 91(4), 296–301 (2000) (Crys. Structure, Phase Relations, Review, 38) Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the Binary System Ti-Al”, Metall. Mater. Trans. A, 32A, 1037–1048 (2001) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, 34) Garmestani, H., Al-Haik, M., Townsley, T.A., Sabinash, C., “Residual Stress Development During Fabrication and Processing of γ-Titanium Based Composites”, Scr. Mater., 44(1), 179–185 (2001) (Experimental, Morphology, 19) Reddy, R.G., Li, Y., Arenas, M.F., “Oxidation of a Ternary Ti3Al-Ta Alloy”, High Temp. Mater. Proces., 21(4), 195–205 (2002) (Experimental, Phys. Prop., 24) Das, K., Das, S., “Order-Disorder Transformation of the Body Centered Cubic Phase in the Ti-Al-X (X = Ta, Nb, or Mo) System”, J. Mater. Sci., 38(19), 3995–4002 (2003) (Crys. Structure, Experimental, Morphology, Phase Relations, 26) Schmid-Fetzer, R, “Al-Ti (Aluminum-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.15634.1.20 (2004) (Phase Diagram, Phase Relations, Review, 85) Legzdina, D., Robertson, I.M., Birnbaum, H.K., “Oxidation Behavior of a Single Phase γ-TiAl Alloy in Low-Pressure Oxygen and Hydrogen”, Acta Mater., 53(3), 601–608 (2005) (Calculation, Experimental, Morphology, Thermodyn., 24)
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Al–Ta–Ti Das, K., Das, S., “A Review of the Ti-Al-Ta (Titanium-Aluminum-Tantalum) System”, J. Phase Equilib. Diffus., 26(4), 322–329 (2005) (Experimental, Kinetics, Phase Diagram, Phase Relations, Thermodyn., 20) Raghavan, V., “Al-Ta-Ti (Aluminum-Tantalum-Titanium)”, J. Phase Equilib. Diffus., 26(6), 629–634 (2005) (Phase Diagram, Phase Relations, Review, 12) Johnson, D.R., Inui, H., Muto, S., Omiya, Y., Yamanaka, T., “Microstructural Development During Directional Solidification of α-Seeded TiAl Alloys”, Acta Mater., 54(4), 1077–1085 (2006) (Experimental, Morphology, 22) Howe, J.M., “Comparison of the Atomic Structure, Composition, Kinetics and Mechanisms of Interfacial Motion in Martensitic, Bainitic, Massive and Precipitation Face-Centered Cubic-Hexagonal Close-Packed Phase Transformations”, Mater. Sci. Eng. A, 438-440, 35–42 (2006) (Experimental, 67) Schuster, J.C., Palm, M., “Reassessment of the Binary Aluminum-Titanium Phase Diagram”, J. Phase Equilib. Diffus., 27(3), 255–277 (2006) (Assessment, Calculation, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 272) Witusiewicz, V.T., Bondar, A.A., Hecht, U., Rex, S., Velikanova, T.Ya., “The Al-B-Nb-Ti System. III. Thermodynamic Re-Evaluation of the Constituent Binary System Al-Ti”, in press, J. Alloys Compd., DOI:10.1016/j.jallcom.2007.10.061 (2007) (Phase Relations, Phase Diagram, Experimental, Thermodyn., Calculation, Review, 89) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Chromium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction As chromium is one of the most important alloying species in steels and borided steels, many research groups have dealt with the phase relations in the ternary system B-C-Cr [1952Gla, 1959Por, 1965Lev, 1966Mar, 1969Mar, 1973Pap, 1974Pra, 1978Pra, 1985Lan, 1989Ord]. Early studies focused on the interaction between the various chromium borides, chromium carbides, boron carbide and carbon. From samples hot-pressed in graphite dies above 1350˚C all chromium borides, Cr2B, Cr5B3, CrB, Cr3B4, CrB2, were said to be stable in combination with carbon [1952Gla] and accordingly CrB2 + C [1965Lev] and CrB2 + ‘B4C’ [1952Gla, 1959Por, 1989Ord] were observed to be quasibinary eutectic systems. The great stability of CrB and CrB2 with respect to carbon was confirmed by [1966Mar, 1969Mar], whereas the lower chromium borides, in particular Cr2B, are unstable in combination with carbon at higher temperatures and form chromium carbides which contain up to 10 at.% B [1966Mar, 1969Mar]. Based on a cursory study of the isothermal section of the B-C-Cr system at 1000˚C by [1973Pap], a detailed investigation was performed by [1974Pra, 1978Pra] of the liquidus surface projection and of the isothermal section at 1450˚C in the chromium rich region (> 50 at.% Cr). Although no signs for the existence of ternary compounds were found in the experiments of [1952Gla, 1965Lev, 1966Mar, 1973Pap, 1974Pra, 1978Pra], agreement exists about the formation of a ternary boron carbide, first labelled as “Cr7BC4” [1966Mar, 1967Kon, 1967Nes, 1971Nes] and later described as Cr3(B0.44C0.56)C0.85 with the filled Re3B-type structure on the basis of a neutron diffraction study [1982Rog]. The ternary compound was only obtained in the temperature range from 1550 to 1710˚C in a small field of primary crystallization [1966Mar, 1982Rog]. From X-ray and metallographic observations a phase field distribution at 1100˚C was assigned by [1985Lan] for the Cr poor region thereby complementing data established by [1974Pra, 1978Pra] i.e. revealing no compatibility between Cr and ‘B4C’. A compilation of the most relevant data to the topology of the B-C-Cr system is due to [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information in literature for the B-C-Cr system up to 1996 was assessed in a general review of phase relations for metal-boroncarbon systems [1998Rog]. Experimental details for all investigations in the B-C-Cr system are summarized in Table 1.
Binary Systems The binary system C-Cr is taken from [2008Bon], Fig. 1, generally based on [1987Ere]. It shall, however, be noted that a recent reinvestigation of the Cr rich part of the system by [1987Ere] revealed a significantly higher melting point of chromium (1878±20˚C) as well as significantly Landolt‐Bo¨rnstein New Series IV/11E1
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higher reaction isotherms than those accepted by [Mas2]. A recent study of the chromium carbides by reaction diffusion couples in the temperature range 1100˚C to 1400˚C [1999May], confirmed the homogeneity region for Cr3C2 and Cr7C3, whilst a slightly higher carbon concentration was revealed for Cr23C6 (21 to 21.5 at.% C). A review on the thermodynamic stabilities of the Cr-carbides is due to [1992Lee]. The binary system B-Cr is adopted from [1992Rog] including the compound Cr2B3 and the low temperature modification of CrB with the αMoB type [1973Pap], Fig. 2. "Cr6B", mentioned by [1974Pra, 1978Pra], is rather oxygen stabilized than a binary chromium boride (see [1992Rog]). The B-C system, Fig. 3, is taken from a recent assessment and thermodynamic calculation by [1998Kas, 1996Kas]. The peritectic l + ‘B4C’ Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan].
Solid Phases Mutual solid solubilities among chromium borides, chromium carbides, boron carbide and graphite were said to be insignificant [1952Gla, 1959Por, 1965Lev, 1966Mar, 1973Pap, 1985Lan, 1989Ord]. This statement generally holds for a small carbon solubility in all the chromium borides and was particularly observed for CrB2-‘B4C’ [1959Por, 1985Lan, 1989Ord], CrB - C [1966Mar, 1973Pap, 1974Pra, 1985Lan] and CrB2 - C [1965Lev, 1966Mar, 1985Lan]. Whereas the authors of [1973Pap, 1974Pra, 1985Lan] agree on a small carbon solubility in Cr2B and Cr5B3, the boron solubility in the chromium carbides is controversial: a rather extended boron solubility in the chromium carbides was observed at 1450˚C by [1966Mar, 1974Pra], whilst boron solubility in chromium carbides at 1000˚C to 1100˚C was claimed to be insignificant by [1973Pap, 1985Lan]. Maximal boron solubilities at 1450˚C were given as follows [1974Pra]: Cr23(B0.3C0.7)6, Cr7(B0.35C0.65)3 and Cr3(B0.05C0.95)2. Ideal C/B-substitution is inferred from the linear variation of the lattice parameters vs boron content (see Table 2). The ternary compound Cr3(B,C)C1–x, earlier denoted as "Cr7BC4" [1966Mar, 1967Kon, 1967Nes, 1971Nes] was only observed [1982Rog] in a small temperature region from 1550 to 1710˚C in good agreement with the reported melting point of Cr7BC4 at 1700˚C [1966Mar, 1971Nes]. Practically homogeneous samples were produced by [1982Rog] near the nominal composition Cr59B8C33 (in at.%) containing small amounts of Cr7(C,B)3. Neutron powder diffraction [1982Rog] confirmed the structure type of Re3B, earlier assigned by [1967Kon] from X-ray single crystal photographs, and revealed full occupancy of the triangular metal prisms by a statistical distribution of B, C atoms. Chromium octahedra are only partially occupied by carbon atoms, confirming the chemical formula Cr3(B0.44C0.56)C0.85 = Cr7B1.03C3.3. Lattice parameter data in Table 2 suggest a small homogeneous range. Crystallographic data of all solid phases pertinent to the B-C-Cr ternary system are listed in Table 2.
Quasibinary Systems Two sections CrB2-C [1965Lev] and CrB2-‘B4C’ [1959Por, 1989Ord] were mentioned as quasibinary of the eutectic type without significant mutual solid solubilities. For CrB2-C only the eutectic temperature at 1875˚C was reported by [1965Lev] from pyrometric measurements on CrB2+C powder mixtures within a graphite heater. Figure 4 presents the CrB2 - ‘B4C’ DOI: 10.1007/978-3-540-88053-0_16 ß Springer 2009
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quasibinary system revealing the eutectic at 2150±30˚C and at 72 mol% CrB2, as established by [1989Ord]. To comply with the accepted B-C binary system, congruently melting ‘B4C’ is to be taken at B0.817C0.183. For lattice parameters see Table 2.
Invariant Equilibria The partial reaction scheme including six experimentally determined isothermal reactions is presented in Figs. 5a and 5b. Whilst the compositions of the liquid phase were reported in the original papers for e2 [1989Ord], and for U2, U3, U4, U7, E3 by [1978Pra], the composition of the corresponding solid equilibrium phases were estimated on the basis of the solid solubility data supplied by [1974Pra] for the 1450˚C partial isothermal section. In order to match the invariant equilibria, as measured by [1978Pra] in the course of a determination of the liquidus surface, with the phase relations established in the isothermal section for the Cr rich part of the system at 1450˚C, five additional transition reactions U1, U5, U6, U8, U9 were included. These reactions relate the observed equilibria of the chromium rich borides with graphite to the observed equilibria between chromium rich borides and chromium carbides at 1450˚C. With respect to the reported quasibinary eutectics, e2 (CrB2-‘B4C’) and e4 (CrB2-C), two further invariant equilibria E1 and E2 were formulated in agreement with earlier observations on the compatibility of CrB, CrB2 with graphite, and of CrB2 with ‘B4C’ [1952Gla, 1966Mar, 1973Pap, 1989Ord]. Accordingly, a peritectoid ternary decomposition of CrB4 is likely at P1 and at a temperature slightly higher than the B-Cr binary peritectoid reaction. The reaction scheme presented in Figs. 5a and 5b as well as the listing of the compositions of the corresponding equilibrium phases in Table 3 is essentially based on the experimental findings of [1965Lev, 1974Pra, 1978Pra, 1989Ord], but does not include information on the ternary compound Cr3(B,C)B1–x and is still incomplete as far as the boron rich area is concerned, awaiting further experimental evidence.
Liquidus, Solidus and Solvus Surfaces A partial liquidus surface for the Cr rich area of the B-C-Cr ternary system was established by [1978Pra] and is presented in Fig. 6 with small changes to comply with the accepted binaries. The small field of primary crystallization of the ternary compound Cr3(B,C)C1–x (1550 to 1710˚C) is not shown in Fig. 6, as it appears at compositions richer in carbon. The dashed lines refer to the extended solid solutions Cr23(BxC1–x)6, Cr7(BxC1–x)3 and Cr3(BxC1–x)2 at 1450˚C.
Isothermal Sections Figure 7 summarizes the available information on the isothermal section at 1450˚C. The Cr rich region is based on the results derived by [1974Pra] with small changes for consistency with the accepted binary systems. “Cr6B”, which very likely is impurity stabilized, was removed from the triangulation proposed by [1974Pra]. The tie lines CrB-C, CrB2-C and CrB2-‘B4C’ are due to the results obtained by [1966Mar, 1973Pap, 1965Lev, 1989Ord]. The boron rich area is suggested and thus shown with dashed tie lines. The position of the ternary compound, Landolt‐Bo¨rnstein New Series IV/11E1
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Cr3(B,C)C0.85 only stable in the range between 1550˚C and 1710˚C [1982Rog] is indicated by a filled circle. Figure 8 represents the partial isothermal section at 1100˚C for concentrations less than 60 at.% Cr based on the findings of [1973Pap, 1985Lan].
Thermodynamics Partial and integral enthalpies of mixing of liquid components in the Cr rich part of the system (xCr > 0.6) have been derived from isoperibolic high temperature calorimetry [1993Vit, 1994Wit]. Integral enthalpies of mixing are represented in Figs. 9a and 9b (in Fig. 9a the reference states are liquid Cr, crystalline B, C; in Fig. 9b the reference states are the liquid elements, B, C in hypothetical supercooled liquid state HSL). Among 3d metal-B-C systems B-C-Cr alloys revealed the largest exothermic effects. The concentration dependence of the partial enthalpies of mixing for Cr, HSL-B and HSL-C are shown in Fig. 10 along three sections with the ratio xB/xC = 0/1, 0.1/0.9, 0.2/0.8 and 0.25/0.75 (curves 1 to 4). From model calculations employing the concentration dependencies of the enthalpies of mixing and the stoichiometry of associates (model of Zielinski and Matyja), the glass forming tendency GFT was estimated for the Cr rich liquids in the ternary system. On the GFT surface (see Fig. 11) a maximum appears at Cr0.75B0.15C0.1 as the most favorable concentration for the formation of amorphous alloys [1994Wit]. A thermodynamic CALPHAD calculation of the phase equilibria as part of the quaternary system C-B-Cr-Fe is due to [1995Rab].
Notes on Materials Properties and Applications [1968Alp] discussed phase assemblages, microstructures, and properties of fused carbides and/or borides from refractory M-C-Cr systems (also for B-C-Cr) containing free graphite in terms of compatibility and phase diagrams. All these bodies have excellent thermal-shock resistance. Other properties (such as electrical, thermal, mechanical, chemical) can be modified by choosing different phase assemblages. Some of these materials have been cast into large shapes more than 45 cm long, and they can be machined into articles. Sinterability and physical properties of hot-pressed CrB2/Cr3C2 sinterbodies were investigated revealing Vickers hardness at RT of about 14 and 15 GPa for 10 and 15 mass% Cr3C2. The Vickers hardness decreased with increasing measurement temperature to 6 GPa at 1273 K for the 10 mass% Cr3C2 composite (from XPD all composites contained CrB and B4C!) [2007Mat]. Addition of up to 25 mol% CrB2 to B4C increased the fracture toughness of the hot-pressed ceramic body from 2.5 to 3.5 MPa·m0.5 and reached a flexural strength of 630 MPa; a maximal value of 684 MPa was reached for a liquid phase sintered ceramic (B4C + 22.5 mol% CrB2) with fine microstructure [2002Yam1, 2002Yam2, 2002Yam3, 2003Yam1, 2003Yam2]. For bodies sintered at 2030˚C a density of 98.1% was reached. Between 2002 and 2350˚C abnormal grain growth was observed on B4C [2002Yam3]. Improvement of fracture toughness was thought to result from micro-crack formation and propagating crack deflection arising from the thermal expansion mismatch of CrB2 and B4C [2002Yam2, 2003Yam2]. An increase of electrical conductivity in specimens with more than 15 mol% CrB2 was attributed to the formation of a CrB2 network [2003Yam1]. The oxidation behavior of CrB2-‘B4C’ composites was studied by [1993Rad]. Preparation of O-free exothermic mixtures of ‘B4C’-Cr cermets (thermites) was DOI: 10.1007/978-3-540-88053-0_16 ß Springer 2009
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reported by [1975Kry]. Physical properties of ‘B4C’-Cr cermets with 0.3, 0.5, 3.0, 5.0, 10.0 mass% Cr were studied: microhardness [1975Mar, 1976Mar], thermal expansion [1975Mar, 1976Mar] and electrical resistance [1975Mar]. Temperature dependence of the electrical resistance indicated that all alloys were semiconducting. Thermal expansion coefficients were all similar and slightly higher than that of pure B4C. Maximum hardness was shown for 3 and 5 mass% Cr alloys superior to that of B4C. [1980Vla] tried to explain the rise of hardness on Cr incorporation into the B4C lattice from EPR radio-spectroscopy. Physical properties of Cr3(B,C)C1–x were reported by [1967Nes, 1971Nes] for a hotpressed sample with 8% porosity prepared with a nominal composition (in at.%) Cr58.2B8.1C33.7: electrical resistivity, ρ = 103 μΩcm; thermal conductivity, κ = 9.6 W·m–1·K–1; thermo-emf = –6.8 μV·K–1; Hall coefficient, R = 0.072 m3·K–1; microhardness 25.5 GPa. Paramagnetism was said to be higher for Cr3(B,C)C1–x than for elemental Cr, Cr-borides or Cr-carbides. Cr3(B,C)C1–x was said to be more resistant against HCl, HNO3 and NaOH, KOH than the binary borides and carbides of chromium but is easily soluble in boiling sulfuric acid [1966Mar]. Oxidation in dry oxygen was reported to be low below 900˚C, however, raises rapidly at higher temperatures and at 1200˚C combustion is almost complete [1966Mar].
Miscellaneous From the contact reaction of materials (B4C + up to 3 mass% Cr) with liquid nickel the formation of Cr2NiB4 at the interface was anticipated [1979Pan].
. Table 1 Investigations of the B-C-Cr Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1952Gla]
Hot-pressing of powder mixtures XPD on seven hot-pressed samples. (powders of Cr, Cr3C2, C, CrB and CrB2) in Compatibility of CrB2 with ‘B4C’. graphite dies at 1500 to 2800˚C. Electrical resistivity and densities were measured. Chemical analyses, XPD
[1959Por]
Specimens sintered from elemental powders to form CrB2 and B4C followed by further reaction sintering. Study of microstructure (LOM), microhardness, solidus points, crystal structure (XRD), heat resistance, tensile strength on samples from the system CrB2-‘B4C’
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. Table 1 (continued) Reference [1965Lev]
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
Interaction of CrB2 with C was claimed to CrB2+C powder compacts were reacted by vacuum sintering in a W heater furnace be of quasibinary eutectic character: and analyzed by XPD, LOM. Te = 1877˚C. Starting materials were CrB2 powder <5 μm containing 28.8 mass% B, B-powder containing 0.0036% Fe, 0.0003% Si, 0.01% Cu, 0.0004% Al, 0.0006% Pb, and lampblack C. The results of [1965Lev] were obtained by X-ray diffractometry, metallography and micro-optical pyrometric melting point measurements on metal boride-carbon powder mixtures reacted inside a bore of a graphite tube directly heated by electric current under vacuum. After the melting run, the samples were quenched at a rate of 100 K·s–1 and examined. Melting temperature was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
[1966Mar] Specimens with Cr+B+C were prepared in an electric furnace with a graphite heater. Wet chemical analyses and XPD. Cr and B were determined in HCl. The amount of carbon was calculated from the difference between its contents in the products before and after the HCl treatment. The chemical resistance against acids and alkalies was studied as well as against dry oxygen at temperatures from 600 to 1200˚C.
Interaction of chromium with boron and carbon over wide range of compositions and from 1300 to 1800˚C. A ternary compound Cr7BC4, insoluble in HCl was detected.
[1967Kon] Determination of the crystal structure from X-ray single crystal photographs. R (F0kl) = 0.083.
A single crystal fragment was isolated from a sinter cake reacting Cr7C3 with CrB2 or CrB or B.
[1967Nes] [1971Nes]
Physical properties were measured: electrical resistivity, thermal conductivity, thermal emf, Hall coefficient, magnetic susceptibility.
The ternary chromium boron carbide was prepared by hot-pressing powder containing in mass% Cr86.8B2.5C11.6 in graphite dies coated with BN in argon at 1600˚C and a pressure of 19.6 MPa.
[1969Mar] Specimens with Cr+B+C were prepared in Lower borides with a high metal content an electric furnace with a graphite heater. are unstable with respect to C and react to The reactions of Cr borides with C and carbides or boron carbides. carbides were studied by means of wet chemical analyses and XPD. DOI: 10.1007/978-3-540-88053-0_16 ß Springer 2009
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1973Pap] About 150 samples were prepared by From XPD a phase triangulation was reacting powder compacts in vacuumattempted. A new low temperature form sealed quartz vials at 1000˚C for 1 month. of CrB was identified. Starting materials were 99.8 mass% Cr, 98 mass% B and lampblack C. XPD. [1974Pra] [1978Pra]
About 72 samples in the Cr rich region of the diagram (>50 at.% Cr) were prepared by high frequency or arc melting high purity elements (99.999 mass% Cr, 99.9 mass% B) or Cr-B and Cr-C master alloys containing 10.8 mass% B or 13.2 mass% C, respectively. The fused alloys were annealed at 1450˚C for 30 h and quenched. XPD, LOM, thermal and chemical analysis and microhardness measurements.
Determination of the phase equilibria in the region CrB-Cr-C3C2 on the basis of XPD, LOM, thermal and chemical analysis and microhardness measurements. Determination of the lattice parameters for the solid solutions Cr23(B,C)6, Cr7(B,C)3 at 1450˚C.
[1982Rog] Reaction sintering at 1600˚C (for 8 h under vacuum in a W-metal furnace) of a powder compact (50 g inside a closed C container) with intermediate crushing and recompacting. Starting materials for the nominal composition Cr59B8C33 were powders of 99.95 mass% Cr, reactor grade C, 98.15 mass% enriched B10 isotope (impurities <6000ppm). Neutron powder diffraction.
Determination of the crystal structure and particularly the light atom sites of Cr3(B0.44C0.56)C0.85 (filled Re3B-type) via a neutron powder diffraction study at RT (λn = 0.120 nm).
[1985Lan]
Powder blends were hotpressed in Phase triangulation for the system B-C-Cr graphite dies (1850 to 2150˚C at 30 MPa) (<70 at.% Cr) at 1100˚C derived from X-ray and annealed for 100 h at 1100˚C. and metallographic observations. X-ray and metallographic observations. Vickers hardness, bending strengths and abrasive wear were determined.
Investigation of the quasieutectic system [1989Ord] About 13 samples were prepared from CrB2+B4C with eutectic point at 70 to 75 powders of CrB2 and B4C (which were heated in a vacuum of 13 mPa to 2000˚C mol% CrB2 and Te of 2150˚C. to reduce the free C content below 0.3 mass%). Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon. An internal heater allowed to measure the liquidus points.
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. Table 1 (continued) Reference
Method/Experimental Technique
[1993Vit] [1994Wit]
Partial and integral enthalpies of mixing of liquid components in the Cr-rich part of the system (xCr > 0.6) have been derived from isoperibolic high temperature calorimetry under 5.103 Pa He. Starting materials were 99.83 mass% Cr, 99.999% C and 99.5% crystalline B. The heats of formation were measured by successive introduction of samples at standard temperature into a liquid bath (40 g of pure metal, binary or ternary alloy) equipped with a stirrer.
Temperature/Composition/Phase Range Studied 12, 9 and 6 alloys were measured within the sections (Cr0.91B0.09)1–xCx, (Cr0.82B0.18)1-xCx and (Cr0.72B0.28)1–xCx, respectively.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(Cr) cI2 < 1863 [Mas2] Im 3m < 1878 ± 20 [1987Ere] W (βB) < 2092
(αB)
(C)gr < 3827 (S.P.)
hR333 R 3m βB
hR36 R 3m αB
Lattice Parameters [pm] a = 288.48
at 25˚C [Mas2] dissolves 0.05 at.% B at 1450˚C [1974Pra]
a = 1093.30 c = 2382.52
[1993Wer]
a = 1092.2 c = 2381.1 a = 1096.37 c = 2384.77
at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at 2.7 at.% Cr, “CrB41” [V-C2]
a = 490.8 c = 1256.7
hP4 a = 246.12 P63/mmc c = 670.90 C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78
DOI: 10.1007/978-3-540-88053-0_16 ß Springer 2009
Comments/References
MSIT1
presumably metastable phase, preparation below 1000˚C [1971Amb] pure B, single crystal [1994Cha] at 25˚C [Mas2]
[1967Low] at 2.35 at.% Cmax (2350˚C) linear ∂a/∂x, ∂c/∂x, [1967Low] Landolt‐Bo¨rnstein New Series IV/11E1
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16
. Table 2 (continued)
Phase/ Temperature Range [˚C] ‘B4C’ < 2450
Pearson Symbol/ Space Group/ Prototype hR45 R 3m B13C2
Lattice Parameters [pm]
Comments/References
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 a = 560.1 c = 1213.6 a = 560.0 c = 1213.0
[1989Ord] from alloy with 10 mol% CrB2, quenched from 2420˚C [1989Ord]
B25C
tP68 P 42m or P42/nnm B25C
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog]
Cr2B < 1870
oF48 Fddd Mn4B (Mg2Cu)
a = 425.35 b = 741.33 c = 1470.6
[1992Rog]a)
Cr5B3 < 1900
tI32 I4/mcm Cr5B3
a = 546.40 c = 1011.0
[1992Rog]
CrB(h) 2095 1000
oC8 Cmcm CrB
a = 296.89 b = 786.89 c = 293.33
[1992Rog]
CrB(r) ≲ 1000
tI16 I41/amd αMoB
a = 294.93 c = 1572.8
[1992Rog] transposition type structure
Cr3B4 < 2075
oI14 Immm Ta3B4
a = 295.25 b = 298.56 c = 1302.2
[1992Rog]a)
Cr2B3
oC20 Cmcm V2B3
a = 302.64 ± 0.05 b = 1811.5 ± 0.4 c = 295.42 ± 0.04
[V-C2]
CrB2 < 2200
hP3 P6/mmm AlB2
a = 297.32 c = 307.25
[1992Rog]
a = 297.1 c = 306.6
from alloy with 20 mol% B4C, quenched from 2170˚C [1989Ord]
a = 286.82 b = 474.99 c = 547.88
[1992Rog]a)
CrB4 < 1450
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B–C–Cr
. Table 2 (continued)
Phase/ Temperature Range [˚C] Cr23C6 < 1576 [Mas2] < 1617 ± 4 [1987Ere] Cr23(BxC1–x)6
Cr7C3 < 1766 [Mas2] < 1781 ± 4 [1987Ere]
Pearson Symbol/ Space Group/ Prototype cF116 Fm 3m Cr23C6
oP40 Pnma Mn7C3 or Cr7C3
Lattice Parameters [pm] a = 1065.0
20.7 to 21.7 at.% C Cr rich [V-C2]
a = 1066.2 a = 1069.5
0 ≤ x ≤ 0.3, quenched from 1450˚C, linear ∂a/∂x [1974Pra] x = 0 [1974Pra] x = 0.3 [1974Pra]
a = 452.6 ± 0.5 b = 701.0 ± 0.5 c = 1214.2 ± 0.5
Cr7(BxC1–x)3 a = 452.6 b = 700.8 c = 1209.7 a = 450.2 b = 704.1 c = 1220.2 Cr3C2 < 1811 [Mas2] < 1829 ± 4 [1987Ere]
oP20 Pnma Cr3C2
a = 553.29 ± 0.05 b = 282.90 ± 0.02 c = 1147.19 ± 0.07
Cr3(BxC1–x)2 a = 553.7 b = 283.2 c = 1148 a = 552.3 b = 283.7 c = 1149 *τ1, Cr3(B,C)C1–x < 1710 - 1550
oC20 Cmcm filled Re3B
Comments/References
a = 287.0 b = 926.0 c = 698.2 a = 285.7 b = 923.3 c = 696.7
28.5 to 31 at.% C [V-C2]a)
0 ≤ x ≤ 0.35, quenched from 1450˚C, linear ∂a/∂x, ∂b/∂x, ∂c/∂x [1974Pra] x=0
x = 0.35
at 39.5 to 40 at.% C [V-C2]
0 ≤ x ≤ 0.05, quenched from 1450˚C, [1974Pra] x=0
x = 0.05
for “Cr7BC4” ρexp = 6.30 Mg·m–3 [1967Kon] for Cr3(B0.44C0.56)C0.85 [1982Rog]
a)
note: structure setting standardized according to Typix [1994Par].
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B–C–Cr
. Table 3 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
Cr
B
L Ð CrB2 + ‘B4C’
2150
e2(max)
L
20
L Ð CrB2 + (C)gr
1875
e4(max)
-
C
72.4
7.6
-
-
-
L Ð CrB2 + ‘B4C’ + (C)gr
< 1875
E1
-
-
-
-
L Ð CrB2 + ‘B4C’ + (βB)
< 1830
E2
-
-
-
-
L + CrB Ð Cr5B3 + (C)gr
1800
U1
-
-
-
-
L + Cr5B3 Ð Cr2B + (C)gr
1725
U2
L
70
22
8
L + Cr3C2 ÐCr7C3 + (C)gr
1675
U3
L
68
10
22
Cr7(C,B)3
70
11
19
Cr3(C,B)2
58
2
40
L + (C)gr Ð Cr2B + Cr7C3
1640
U4
L
71.5
18.5
10
Cr2B
67
33
Cr7(C,B)3
70
10
20
<0.5
Cr2B + (C)gr Ð Cr5B3 + Cr7C3
< 1640
U5
-
-
-
-
Cr7C3+(C)gr Ð Cr5B3+Cr3C2
1550
U6
-
-
-
-
L + Cr23C6 Ð (Cr) + Cr7C3
1515
U7
L
85
5
10
LÐ (Cr) + Cr2B + Cr7C3
1475
E3
Cr23(C,B)6
79.5
Cr7(C,B)3
70
L
85
6 10 7.5
14.5 20 7.5
Cr2B
67
33
Cr7(C,B)3
70
10
20
<0.5
‘B4C’ + CrB2 + ($B) Ð CrB4
1450
P1
-
-
-
-
Cr5B3 + (C)gr Ð CrB + Cr3C2
1450
U8
-
-
-
-
(Cr)+Cr7C3 Ð Cr2B+Cr23C6
< 1475
U9
-
-
-
-
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B–C–Cr
. Fig. 1 B-C-Cr. Phase diagram of the C-Cr system
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. Fig. 2 B-C-Cr. Phase diagram of the B-Cr system
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B–C–Cr
. Fig. 3 B-C-Cr. Phase diagram of the B-C system
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. Fig. 4 B-C-Cr. The quasibinary system ’B4C’-CrB2
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B–C–Cr
. Fig. 5a B-C-Cr. Reaction scheme, part 1
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. Fig. 5b B-C-Cr. Reaction scheme, part 2
B–C–Cr
Landolt‐Bo¨rnstein New Series IV/11E1
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B–C–Cr
. Fig. 6 B-C-Cr. Partial liquidus surface projection
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. Fig. 7 B-C-Cr. Isothermal section at 1450˚C
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B–C–Cr
. Fig. 8 B-C-Cr. Partial isothermal section at 1100˚C for the region with <60 at.% Cr
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. Fig. 9a B-C-Cr. Integral enthalpies of mixing in kJ·mol–1; reference state Crliq, Bcryst, Ccryst
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B–C–Cr
. Fig. 9b B-C-Cr. Integral enthalpies of mixing in kJ·mol–1; reference state Crliq, BHSL, CHSL
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. Fig. 10 B-C-Cr. Partial enthalpies of mixing for Crliq, BHSL, CSCL along the sections with ratio xB/xC=0/1, 0.1/0.9, 0.2/0.8 and 0.25/0.75 (curves 1 to 4)
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B–C–Cr
. Fig. 11 B-C-Cr. Glass forming tendency for Cr rich liquids. For explanations see text
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[1965Lev]
[1966Mar]
[1967Kon]
[1967Low] [1967Nes]
[1968Alp]
[1969Mar] [1971Amb] [1971Nes]
[1973Pap] [1974Pra] [1975Kry]
[1975Mar]
[1976Mar] [1978Pra] [1979Pan]
[1980Vla]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Portnoi, K.I., Samsonov, G.V., “Investigation of Refractory Samples” (in Russian), Issled. po Zharoproch. Splavam, Akad. Nauk SSSR, Inst. Met. im. A.A. Baikova, 5, 192–198 (1959) (Experimental, Phase Diagram, Phase Relations, Phys. Prop., 9) Levinskii, Y.V., Salibekov, S.E., Levinskaya, M.K., “Reaction of Chromium, Molybdenum and Tungsten Borides with Carbon”, Sov. Powder Metall. Met. Ceram., 12(36), 1004–1009 (1965), translated from Poroshk. Metall. (Kiev), 12(36), 56–62 (1965) (Crys. Structure, Experimental, 5) Markovskii, L.Y., Vekshina, N.V., Kondrashev, Y.D., “Chromium Boron Carbide”, J. Appl. Chem. USSR, 39(5), 923–926 (1966), translated from Zh. Prikl. Khim., 39(5), 973–977 (1966) (Crys. Structure, Experimental, 8) Kondrashev, Y.D., “Crystal Structure of Chromium Boroncarbide”, Sov. Phys. Crystallography, 11(4), 492–493 (1967), translated from Kristallografiya, 11(4), 559–561 (1967) (Crys. Structure, Experimental, 6) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Neshpor, V.C., “Certain Physical Properties of Chromium Boroncarbide”, Inorg. Mater., 3(12), 1894–1896 (1967), translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 3(12), 2170–2173 (1967) (Experimental, Phys. Prop., 18) Alper, A.M., Doman, R.C., McNally, R.N., “Fusion-Cast Carbide-Boride-Graphite Ceramics” in “Proc. Fourth International Conference on Science of Ceramics”, Maastricht, Netherlands, European Ceramic Association., 23–27 April, 1967, Stewart, G.H. (Ed.), British Ceramic Society, Stoke-on-Trent, 389–420 (1968) (Experimental, Review, 73) Markovskii, L.Y., Vekshina, N.V., Bezruk, E.T., “Reaction of Borides with Carbon and Carbides” (in Russian), Khim. Svoistva Metody Anal. Tugoplavkih Soedin., 143–148 (1969) (Experimental, 8) Amberger, E., Ploog, G., “Formation of Pure Boron Lattice” (in German), J. Less-Common Met., 23, 21–31 (1971) (Crys. Structure, Experimental, 18) Neshpor, V.S., Vekshina, N.V., Nikitin, V.P., Markovskii, L.Y., “Physical Properties of Be, Cr and La-Borocarbides”, Inorg. Mater., 7(12), 1931–1934 (1971), translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 7(12), 2170–2174 (1971) (Experimental, Phys. Prop., 21) Papesch, G., Nowotny, H., Benesovsky, F., “Studies in the Systems Cr-B-C, Mn-B-C and Mn-Ge-C” (in German), Monatsh. Chem., 104, 933–942 (1973) (Crys. Structure, Experimental, 21) Pradelli, G., “Research on Cr Borocarbides” (in Italian), Met. Ital., 66(10), 551–556 (1974) (Experimental, Morphology, Phase Diagram, Phase Relations, 12) Krylov, Y.I., Bronnikov, V.A., Krysina, V.G., Pristavko, V.V., “Possibilities of Producing Thermite Mixtures Based on B4C-metal and SiC-metal Composites”, Powder Metall. Met. Ceram., 14(12), 1000–1003 (1975), translated from Poroshk. Metall. (Kiev), 12(156), 57–60 (1975) (Phase Relations, Experimental, Thermodyn., 11) Marek, E.V., Dudnik, E.M., Makarenko G.N., Remenyuk, E.A., “Physical Properties of Boron Carbide Alloys with Vanadium and Chromium Additives” (in Russian), Poroshk. Metall. (Kiev), (2), 54–6 (1975) (Experimental, Phys. Prop., 4) Marek, E.V., Makarenko, G.N., “Production and Certain Properties of B-C-V and B-C-Cr Alloys”, Karbidy i Splavy na ikh Osnove, Akad. Nauk Ukrain. SSR, 195–198 (1976) (Experimental, Phys. Prop., 3) Pradelli, G., “On the System Cr-B-C” (in Italian), Metall. Ital., 70(5), 223–226 (1978) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, Phys. Prop., #, *, 7) Panasyuk, A.D., Maslennikova, V.R., Makarenko, G.N., Marek, E.V., “Contact Reaction of Boron Carbide-Based Materials with Liquid Nickel”, Inorg. Mater., 15(4), 503–505 (1979), translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 15(4), 643–646, 1979 (Experimental, Morphology, 7) Vlasova, M.V., Kakazey, N.G., Kosolapova, T.Y., Makarenko, G.N., Marek. E.V., “The Structure of Paramagnetic Centers and the Formation of Defects in the B-C, B-C-Ti and B-C-Cr Systems”, J. Mater. Sci., 15(4), 1041–1048 (1980) (Experimental, Crys. Structure, Magn. Prop., 9)
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26
16 [1982Rog]
[1983Sch]
[1984Hol]
[1985Lan]
[1987Ere]
[1989Ord]
[1990Ase]
[1992Lee] [1992Rog]
[1993Rad] [1993Vit] [1993Wer]
[1994Cha]
[1994McH]
[1994Par]
[1994Wit]
[1995Vil]
[1995Rab]
[1996Kas]
B–C–Cr Rogl, P., Kunsch, B., Ettmayer, P., Nowotny, H., Steurer, W., “A Neutron Diffraction Study of Cr3(11B0.44C0.56)C0.85 and Cr3C(C0.52N0.48)”, Z. Kristallogr., Kristallgeom., Kristallphys., Kristallchem., 160, 275–284 (1982) (Crys. Structure, Experimental, 19) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Lange, D., Holleck, H., “Wear Resistant Materials Based On Boron Carbide” (in German), Proc. 11th Plansee Seminar, Reutte, Austria, Contribution HM50, 747–759 (1985) (Phase Diagram, Phase Relations, Experimental, *, 19) Eremenko, V.N., Velikanova, T.Y., Bondar, A.A., “Phase Diagram of the Cr-Mo-C System. 1. Phase Equilibria in the Area of Crystallization of Alloys of the Mo-Mo2C-Cr7C3-C Partial System”, Sov. Powder Metall. Met. Ceram., 26(5), 409–413 (1987), translated from Poroshk. Metall. (Kiev), 5(293), 70–76 (1987) (Phase Diagram, Phase Relations, Experimental, 14) Ordan’yan, S.S., Dmitriev, A.I., “Reaction in the B4C-CrB2 System”, Inorg. Mater., 25(4), 593–595 (1989), translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 25(4), 685–687 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, #, 3) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U. K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Lee, B.J., “On the Stability of Cr-Carbides”, Calphad, 16(2), 121–149 (1992) (Theory, Review, 83) Rogl, P., “The B-N-Cr System” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, OH., 20–25 (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, Review, 11) Radev, D., Zahariev, Z., “Oxidation Stability of B4C-MexBy Composite Materials”, J. Alloys Compd., 197 (1), 87–90 (1993) (Experimental, Morphology, Phys. Prop., 14) Vitusevich, V.T., “Enthalpy of Formation of Cr-B-C Melts”, Russ. Metall. (Engl. Transl.), (3), 31–34 (1993), translated from Izv. RAN, Met., (3), 35–38 (1993) (Experimental, Thermodyn., 9) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-rhombohedral Boron”, Phys. Stat. Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) Chakrabarti, D.J., Laughlin, D.E., “B-Cu (Boron-Copper)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 74–78 (1994) (Review, Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., 24) McHale, A.E., “VI. Boron Plus Carbon Plus Metal”, in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 180 (1994) (Phase Relations, Review, 2) Parthe´, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1–4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Witusiewicz, V.T., “Thermodynamic Properties of Liquid Alloys of 3d Transition Metals with Metalloids (Silicon, Carbon and Boron)”, J. Alloys Compd., 203, 103–116 (1994) (Experimental, Thermodyn., 89) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA, 5323–5327 (1995) (Review, Phase Diagram, Phase Relations, Crys. Structure, 7) Rabitsch, K., Sitte, R., Ebner, R., Jeglitsch, F., “Solidification Behavior in Extremely Fast Solidified FeCr-B-C Alloys”, Prakt. Metallographie, Sonderband, 26, 545–556 (1995) and references therein (Experimental, Theory, 15) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170)
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[1998Rog]
[1999May]
[2002Yam1]
[2002Yam2]
[2002Yam3] [2003Yam1]
[2003Yam2] [2005Tan]
[2007Mat]
[2008Bon]
[Mas2] [V-C2]
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Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Chromium” in “Phase Diagrams of Ternary Metal-BoronCarbon Systems”, Effenberg, G. (Ed.), MSI, ASM International, Materials Park, Ohio, USA, 36–50 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 32) Mayr, W., Lengauer, W., Ettmayer, P., Rafaja, D., Bauer, J., Bohn, M., “Phase Equilibria and Multiphase Reaction Diffusion in the Cr-C and Cr-N Systems”, J. Phase Equilib., 20(1), 35–44 (1999) (Experimental, Crys. Structure, Phase Diagram, Interface Phenomena, Phase Relations, 25) Yamada, S., Hirao, K., Sakaguchi, S., Yamauchi, Y., Kanzaki, S., “Microstructure and Mechanical Properties of B4C-CrB2 Ceramics”, Key Eng. Mater., 206–213(1), 811–814 (2002) (Experimental, Mechan. Prop., 14) Yamada, S., Hirao, K., Yamauchi, Y., Kanzaki, S., “Densification Behaviour and Mechanical Properties of Pressureless-Sintered B4C-CrB2 Ceramics”, J. Mater. Sci., 37(23), 5007–5012 (2002) (Experimental, Mechan. Prop., 26) Yamada, S., Hirao, K., Yamauchi, Y., Kanzaki, S., “Sintering Behavior of B4C-CrB2 Ceramics”, J. Mater. Sci. Lett., 21(18), 1445–1447 (2002) (Experimental, Mechan. Prop., 12) Yamada, S., Hirao, K., Yamauchi, Y., Kanzaki, S., “Mechanical and Electrical Properties of B4C-CrB2 Ceramics Fabricated by Liquid Phase Sintering”, Ceram. Intern., 29(3), 299–304 (2003) (Experimental, Mechan. Prop., Electr. Prop., 27) Yamada, S., Hirao, K., Yamauchi, Y., Kanzaki, S., “B4C-CrB2 Composites with Improved Mechanical Properties”, J. Eur. Ceram. Soc., 23(3), 561–565 (2003) (Experimental, Mechan. Prop., 25) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th Intl. Symp. Boron, Borides and Related Compounds, Hamburg (Germany), Aug. 21–26, 142 (2005) (Experimental, Phase Relations, 3) Matsushita, J., Shimao, K., Machida-Y., Takao, T., Izumi, K., Sawada, Y., Kwang, Bo-Shim., “Sintering and Mechanical Properties of Chromium Boride - Chromium Carbide Composites” Mater. Sci. Forum., 534–536, 1077–1080 (2007) (Experimental, Mechan. Prop., 12) Bondar, A.A., “C-Cr (Carbon-Chromium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2008) (Crys. Structure, Phase Diagram, Phase Relations, 23) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Hafnium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction With melting temperatures near 4000˚C, materials from the B-C-Hf system rank among ultrahigh-temperature ceramics. Particularly the composite HfB2/B4C has attracted attention because of a unique combination of materials properties such as high thermal conductivity, good thermal shock resistance at moderate thermal expansion and oxidation resistance. Besides an early investigation of the phase relations within the isothermal section at 1500˚C [1961Now] and a cursory study of the HfB2-C quasibinary eutectic by [1965Lev], the first detailed and most thorough investigation of the B-C-Hf system is due to [1966Rud], who established phase equilibria in the 1400˚C isothermal section as well as for three isopleths HfB2-C, HfB2-HfC0.9 and HfB2-B4.5C. [1966Rud] furthermore determined the liquidus surface and presented a reaction scheme. More recent independent studies of the isopleths HfB2 HfC1–x by [1977Ord] and HfB2 - ’B4C’ by [1989Ord] are in general agreement with the findings of [1966Rud], despite melting temperatures recorded by [1977Ord, 1989Ord] seem to be persistently lower than those of [1966Rud]. Compilations of the most relevant data on the topology of the B-C-Hf system have been published by [1969Rud, 1974Upa, 1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information in literature for the B-C-Hf system up to 1996 was compiled in a general review of phase relations for metal-boron-carbon systems [1998Bit] including a thermodynamic calculation of the entire B-C-Hf system [1997Bit1, 1997Bit2]. For further details on the thermodynamic assessment, see [1999Rog, 2000Rog]. Experimental details for all investigations in the B-C-Hf system are summarized in Table 1.
Binary Systems The binary boundary systems B-Hf and C-Hf were accepted from [Mas2]. They basically are consistent with the versions presented by [1966Rud] except for small changes, in order to correspond to the accepted (βHf)/(αHf) transition temperature ([1966Rud] at 1800˚C, [Mas2] at 1743˚C) and melting temperature of zirconium-free hafnium ([1966Rud] at 2218 ± 6˚C, [Mas2] at 2231˚C). A thermodynamic calculation of the B-Hf diagram is from [1988Rog] with a refinement of this modeling by [1997Bit2] (Fig. 1). A thermodynamic calculation of the C-Hf diagram was performed by [1997Bit1] (Fig. 2). A partial low temperature phase diagram C-Hf was calculated by [1991Gus] employing the order parameters functional method and essentially concerns the formation of ordered superlattice phases, Hf3C2 and Hf6C5, deriving from NaCl type HfC1–x. These phases were calculated to form below 520˚C. Landolt‐Bo¨rnstein New Series IV/11E1
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The B-C system is adopted from a recent assessment and thermodynamic calculation by [1998Kas, 1996Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The peritectic boron rich reaction, L+‘B4C’Ð(βB), as modelled by [1996Kas], was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan]. Literature data concerning the formation and crystal structure of solid phases pertinent to the B-C-Hf system are listed in Table 2.
Solid Phases No ternary phases were observed throughout the B-C-Hf ternary system. Except for the monocarbide, for which lattice parameters in the ternary system appeared considerably increased with respect to those of the binary system (see Table 2), mutual solid solubilities of the binary boride and carbide phases were found to be very small [1961Now, 1966Rud, 1965Lev, 1977Ord, 1989Ord]. Lattice parameters of alloys annealed at 1500˚C were presented by [1961Now] in a graph for the region around HfC1–x. They are shown in Fig. 3 in comparison to the much smaller solubility limits of B in HfC1–x at 1400˚C reported by [1966Rud]. Precipitation of HfB2 from the carbon saturated monocarbide solution Hf(C, B)1–x was observed to occur considerably faster than from substoichiometric compositions and could not be suppressed at cooling rates lower than 100 K·sec–1. A maximum boron exchange of 12 at.% B was observed at a carbon defect of about 7 at.% at 3140˚C [1966Rud]. At 1400˚C the B solubility amounts to about 2.5 at.% B over the whole homogeneity range of HfC1–x. Crystallographic data for all solid phases in the B-C-Hf system are listed in Table 2.
Quasibinary Systems Three quasibinary sections of eutectic type (HfB2-C, HfB2-HfC0.9 and HfB2-‘B4C’) were established by [1966Rud] (see Figs. 4, 5, 6 and Table 3). It should be noted, that compositional data in all isopleths given by [1966Rud] refer to 1 mole of atoms in contrast to the chemical formulae shown in his figures (i.e. label HfB2 reads as Hf0.33B0.67). The eutectic nature of the HfB2-C system was earlier reported by [1965Lev]. However, the eutectic temperature as measured by optical pyrometry on pre-reacted powders through a borehole in a directly heated graphite tube, 2337 ± 30˚C, is considerably lower than 2515 ± 10˚C given by [1966Rud]. Also the eutectic point at 24 mol% HfB2 compares badly with the value of 35 ± 2 mol% HfB2 (originally given as 62 mol% Hf0.33B0.67) of [1966Rud]. Similarly, the reinvestigation of the section HfB2-HfC1–x by [1977Ord] reveals a lower eutectic temperature (2980 ± 40˚C) than the value derived by [1966Rud] (3140±15˚C). A likely explanation for these discrepancies may be insufficient correction for non-black body conditions in the experiments of [1965Lev] and [1977Ord]. The eutectic composition was reported by [1977Ord] to be 48 mol% HfC1–x, whilst [1966Rud] found 45 ± 2 mol% HfC0.9 (originally given as 34 mol% Hf0.526C0.474). The solid solubility of HfB2 in HfC1–x was claimed to be about 3 mass% HfB2 (2.8 mol% HfB2) by [1977Ord], whereas a maximal solid solubility of 7 mol% HfB2 in HfC0.9 (originally given as 10 mol% Hf0.33B0.67 in Hf0.526C0.474) was observed by [1966Rud] from samples rapidly quenched in tin. According to the carbon composition of the hafnium monocarbide starting powder close to stoichiometry (6.26 mass% C = HfC0.99), the DOI: 10.1007/978-3-540-88053-0_17 ß Springer 2009
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section chosen by [1977Ord] was in fact slightly off the true quasibinary connecting HfB2 and congruently melting HfC0.94. This fact may partly account for the differences in the temperature as well as in the solubility, which was shown by [1966Rud] to drastically decrease when approaching the stoichiometric limit at HfC0.99. Agreement exists on the experimental findings for the HfB2-‘B4C’ quasibinary section among data by [1966Rud] (68 ± 4 mol% B4.5C, originally given as 78 mol% B0.817C0.183 at 2330 ± 25˚C) and the reinvestigation by [1989Ord] (78 mol% ‘B4C’ at 2380 ± 30˚C). It shall be noted, that taking the sections along the tie line of the reaction does not constitute true quasibinary sections as seen for instance from the slight deviation from a congruent melting in HfB2-HfC1–x.
Invariant Equilibria A reaction scheme for the ternary B-C-Hf system including nine observed ternary invariant equilibria was provided by [1966Rud]. Table 3 lists the compositions of the phases at the fourphase isothermal reactions as given by [1966Rud] and compares experimentally derived data with the results of the thermodynamic calculation by [1997Bit1, 1998Bit, 1999Rog, 2000Rog]. Figure 7 shows the Scheil diagram corresponding to the thermodynamic assessment. As far as the very boron rich region is concerned, the thermodynamic calculation favors a transition type reaction at 2091˚C, L + (βB) Ð ‘B4C’ + HfB2, rather than a ternary eutectic at 1950˚C, L Ð (βB) +‘B4C’ + HfB2, as reported by [1966Rud]. This discrepancy essentially results from a new assessment of the B-C system by [1996Kas] revealing a peritectic reaction L + ‘B4C’ Ð (βB) at 2103˚C rather than a eutectic L Ð (βB) + ‘B4C’ at 2080˚C as given by [1966Rud]. The experimentally reported very sharp drop of more than 100˚C and within less than 2 at.% from the binary reaction isotherms into the ternary eutectic (at 1950˚C [1966Rud]) seems unlikely and could not be modeled thermodynamically. It shall be noted that the peritectic binary reaction L + ‘B4C’ Ð (βB) was recently confirmed from floating zone measurements on several carbon-doped boron samples [2005Tan].
Liquidus, Solidus and Solvus Surfaces The calculated liquidus surface projection is shown in Fig. 8 in comparison with data points for the invariant equilibria established by [1966Rud].
Isothermal Sections Based on the experimentally established isothermal sections at 1400˚C [1966Rud], 1500˚C [1961Now] as well as on the liquidus projection and the phase relations derived for the three isopleths HfB2-C, HfB2-HfC0.94 and HfB2-B4.5C, [1966Rud] constructed a series of isothermal sections at 1800, 1940, 2000, 2050, 2300, 2400, 2800, 3100 and at 3200˚C. The isothermal section at 1400˚C [1966Rud] and at 1500˚C [1961Now] are almost identical thus only the 1400˚C section is shown here (Fig. 9). A detailed evaluation of the lattice parameters of the solid solution Hf(C,B)1–x from alloys annealed at 1500˚C was provided by [1961Now], although it has to be mentioned, that the limits of the solid solubility of B in HfC1–x differ significantly for both sections, a fact not being accounted for by temperature Landolt‐Bo¨rnstein New Series IV/11E1
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dependence. For this assessment it was assumed that the later work by [1966Rud] with the smaller solubility limits superseded the larger values earlier presented by [1961Now]. Several isothermal sections at 1400, 1954, 2063, 2091, 2350 and 3100˚C were calculated by [1997Bit1, 1998Bit, 1999Rog], some are shown in Figs. 10 to 14.
Temperature – Composition Sections Isopleths Hf - B0.5C0.5, HfB - C and HfC - B and an isometric view of the ternary B-C-Hf system were constructed by [1966Rud] on the basis of the experimental data derived from the 1400˚C, 1500˚C isothermal sections, the liquidus surface projection and the quasibinary sections HfB2 - C, HfB2 - B4.5C, HfB2 - HfC0.90. In order to portray the phase interactions in the ternary system a series of isopleths was calculated [1997Bit1, 1998Bit, 1999Rog, 2000Rog]: Hf - B0.817C0.183 (Fig. 15) HfB - C (Fig. 16), HfC - B (Fig. 17), Hf51.6C48.4 - B4.5C (Fig. 18) and Hf - BC (Fig. 19).
Thermodynamics A thermodynamic calculation of the ternary B-C-Hf system was based on thermodynamic assessments of the binary systems B-C [1996Kas], B-Hf [1997Bit1] and C-Hf [1997Bit2] as well as relying on the phase diagram data from [1966Rud] for the optimization of the thermodynamic parameters. Technological interest in dense ceramic parts HfB2+‘B4C’ with various degrees of 10B/11B isotope enrichment as nuclear reactor control or moderator substances, demands detailed knowledge on optimal sintering conditions. Therefore thermodynamic calculations provided isopleths HfB2 + ‘B4C’ with variable B/C ratios (Figs. 20a to 20d) [2000Rog].
Notes on Materials Properties and Applications Interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - ‘B4C’ [1993Ord] has been analyzed for transition elements M = Ti, Zr, Hf, V, Nb and Ta and correlations were found between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms. Microhardness measurements on samples in both the isopleths HfB2 - HfC1–x and HfB2 ‘B4C’ near the eutectic composition were found to reveal values well below the linear combination of the binary compounds and to be strongly dependent on the eutectic crystallization conditions. For the HfB2 - HfC1–x eutectic, microhardness ranges from 15.5 GPa for the finely dispersed eutectic to 19.6 GPa for macrocrystalline colonies [1977Ord]. For the HfB2 ‘B4C’ eutectic microhardness was reported to be 32 to 33 GPa [1989Ord].
Miscellaneous The chemical reactivity of Hf powder (0.3 to 20 μm) in combination with ‘B4C’ was investigated at 1500˚C: loose powders were found to react with little weight gain to HfB2+HfC DOI: 10.1007/978-3-540-88053-0_17 ß Springer 2009
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similar to the reaction HfC+2B [2004Blu]. Hot-pressed ‘B4C’+Hf powders with fine particle size distribution constitute a neutron absorber material having a high resistance to mechanical damage and more particularly high resistance to crack propagation [2000Pro]. In an attempt to use ‘B4C’ as inert matrix for actinide burning, ceramic composites ‘B4C’+HfB2 were investigated and showed twice as high thermal conductivity as ‘B4C’ and yield a higher crack propagation resistance; in order to simulate Pu, the Hf was replaced by Ce [2001Gos]. Homogeneous mixing of ‘B4C’ and 25 vol% Hf powders (uniaxial hot-pressing at 1970 to 2270 K (1697 to 1997˚C), 50 MPa) yielded a segregation in Hf clusters and the formation of HfB2 [2001Gos]. Precursor routes to produce group IV metal borides and metalboride/carbide composites were tested: pyrolysis of a Hf/polyhexenyldecaborane polymer dispersion under argon at 1650˚C gave a 92.6% ceramic yield of the expected 98.0% - XPD confirmed the formation of HfC and HfB2 [2004For]. Using a compact diffraction reaction chamber, [2006Won] studied with time resolved X-ray diffraction the chemical dynamics at the combustion front for the directly ignited reactions Hf + C Ð HfC, Hf + 2B = HfB2 and 3Hf + ‘B4C’ Ð 2HfB2 + HfC. Combustion front velocities were 6, 5, and 5 mm·s–1, respectively. The adiabatic combustion temperatures calculated (>3000 K) exceeded the Tm of Hf in all reactions.
. Table 1 Investigations of the B-C-Hf Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
[1961Now] Hot-pressing of 60 binary and ternary powder compacts in C-cartridges at 1600 to 2500˚C. A series of alloys was prepared by reaction sintering in Ar or vacuum respectively. Samples containing HfC and HfB2 were reacted at 1750˚C 13 to 38 h prior to final anneal at 1500˚C for 3 to 16 h. Alloys containing free Hf were annealed for 8–16 h at 1700 and 1500˚C. Samples near HfB were additionally annealed at 1250˚C for 27 h. Starting materials were HfH2 containing 2.2 mass % Zr and 0.18% O), amorphous boron (96 mass% residue O, C), lampblack C. Master alloy HfC (6.36 mass% Ctotal, 0.25 mass% free C). Some alloys were arc melted. XPD and LOM.
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Temperature/Composition/Phase Range Studied Isothermal section at 1500˚C. Study of the homogeneity region of Hf(C,B)1–x at 1500˚C from lattice parameter dependencies.
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. Table 1 (continued) Reference
Method/Experimental Technique
[1965Lev]
HfB2 was prepared via the boron carbide process at 2000 to 2200 K (1727 to 1927˚C) MO2+B4C+CÐMB2+CO. Starting materials were powders of HfO2, B4C (containing B2O3, C), lampblack C (ash <0.05 mass%). HfB2 was prepared from Hf powder (2.8% Zr, 0.05% Fe, 0.01% Mg, 0.07% Si, 0.009 mass% Ti), boron powder (impurities were: 0.0036% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb) by vacuum sintering (Hf+2B) powder compacts for 1 h at 1800˚C and analyzed by XPD, LOM.
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Temperature/Composition/Phase Range Studied Interaction of HfB2 with C yielded a quasieutectic system HfB2+C with eutectic point at 24 mol% HfB2 and TE of 2610 K (2337˚C). Thermal analysis was performed by direct electrical heating of a graphite tube filled with HfB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
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. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1966Rud]
X-ray powder diffraction, metallographic, melting point and differential thermoanalytic (>4/sec) techniques on about 200 vacuum sintered, hot pressed as well as electron beam or argon arcmelted specimens. Starting materials were high purity elemental powders (Hf-sponge with a particle size less than 120 mm containing 4.1 at.% Zr, 680 ppm Nb, less than 300 ppm O and less than 300 ppm N; spectrographic grade graphite powder with less than 100 ppm of metal and sulfur impurities; boron powder of 99.55 mass% minimum purity containing 0.25 mass% Fe and 0.1 mass% carbon) as well as prereacted master alloys of HfB2 (65.0±3 at.% B with 0.02 mass% C, sum of O, N, H <200 ppm). Samples were prepared by short duration hot pressing in graphite dies at temperatures between 1800˚C and 2200˚C and after removing the surface reaction zones they were directly used in as pressed condition for Pirani melting point (under 2.5·105 Pa He) or DTA (graphite container under 105 Pa of He). Samples were analyzed for free and combined carbon, boron as well as oxygen and nitrogen contaminants; the latter never exceeded 250 ppm.
Investigation of the constitution of the B-C-Hf phase diagram. Selected alloys from the metal rich region (>85 at.% Hf) intended for melting point or DTA studies were electron beam or arc melted prior to the runs. Whereas specimens for DTA and melting point analyses were directly equilibrated in the equipment prior to the runs, specimens for the isothermal sections were generally annealed in a tungsten mesh furnace for 92 h at 1400˚C under a vacuum of 5·10–3 Pa or above 1600˚C under helium (12 h at 1750˚C for specimens in the region HfB2 - HfC1–x-B-C or 2 h at 2000˚C for specimens HfC1–x + 5 at.% B under 1.05·10–5 Pa He). Alloys from the region Hf - Hf0.8C0.2 - Hf0.8B0.2 were usually melted and annealed at 1750˚C for 1–3 h in a vacuum of 2·10–3 Pa [1966Rud]. Selected alloys were equilibrated in the melting point furnace and quenched in liquid tin. As far as metallography is concerned, anodic oxidation in an electroetching process using 10% oxalic acid was said to provide excellent phase contrast for alloys containing excess metal (metal phase - light blue, monoboride - brown, diboride unaffected). Specimens from the region Hf0.7C0.3 - Hf0.4B0.3C0.3 - Hf0.4C0.6 were dip etched in 10% aqua regia and HF solutions.
[1977Ord]
Starting materials: HfC (93.3 mass% Hf, Ctot= 6.31 mass% C, 0.05 mass% free C, 0.35 mass% (O2+N2)) and HfB2 (88.3 mass% Hf, Btot = 10.5 mass%, 0.03 mass% free B, 0.06 mass% (O2+N2)). Specimens were obtained from extrusion of plasticized masses in form of cylinders (3 mm diameter · 50 mm length), presintered at 2100˚C for 2 h in vacuum prior to heat treatment at 2300˚C. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the quasieutectic system HfB2+HfC1–x with eutectic point at 52 mol% HfB2 and TE of 2980˚C on 8 samples.
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. Table 1 (continued) Reference
Temperature/Composition/Phase Range Studied
Method/Experimental Technique
[1989Ord]
Samples were prepared from powder Investigation of the quasieutectic system compacts of HfB2 and B4C (which was HfB2+‘B4C’ with eutectic point at 22 mol% vacuum annealed at 2000˚C to reduce HfB2 and TE of 2380˚C on 10 samples. C-content to 0.2 mass% free C). Specimens in form of cylinders (3 mm diameter · 50 mm length) were compacted with aid of 12% aqueous starch solution, presintered at 2300˚C for 2 h in vacuum prior to measurement. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
[2005Tan]
Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)) the liquidus and solidus curves to the L+(βB) field has been derived via chemical analysis.
Confirmation of peritectic type of reaction L+B4+xCÐ(βB) via determination of the liquidus and solidus curves to the L+(βB) field.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C] (βB) < 2092
(C)d
Pearson Symbol/ Space Group/ Prototype hR333 R3m βB
Lattice Parameters [pm] a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1095.57 c = 2402.44
cF8 a = 356.69 Fd3m C (diamond)
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Comments/References [1993Wer] at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x for HfB45 [1981Cre] at 25˚C, 60 GPa [Mas2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Comments/References
a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78 a = 246.8 c = 670.9
at 25˚C [Mas2]
hP2 P63/mmc Mg
a = 319.46 c = 505.10 a = 320.4 c = 507.5 a = 322 c = 513 a = 321 c = 510 a = 321 c = 509
at 25˚C [Mas2]
(βHf) 2231 - 1743
cI2 Im3m W
a = 361.0
[Mas2]
‘B4C’ < 2450
hR45 R3m B13C2
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5
(C)gr < 3827 (S.P.)
(αHf) < 1743
B25C
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tP68 P42m or P42/nnm B25C
a = 560.0 c = 1213.6 a = 560.1 c = 1213.6 a = 875.3 ± 0.4 c = 509.3 ± 1.5
[1967Low] at 2.35 at.% Cmax (2350˚C) linear ∂a/∂x, ∂c/∂x, [1967Low] from alloy Hf12B24C64, quenched from 2923˚C, sample also contains HfB2 [1966Rud]
from alloy Hf92B3C5, quenched from 2312˚C, contains also HfB [1966Rud] from alloy Hf85B5C8, quenched from 2315˚C, contains also HfC1–x [1966Rud] from alloy Hf86B3C11, quenched from 2386˚C, contains also HfC1–x [1966Rud] from alloy Hf86B12C2, quenched from 1945˚C, contains (βHf) and HfB [1966Rud]
at ‘B4C’ [1989Ord] from alloy HfB2 + 88.6 mol% ‘B4C’ [1989Ord] [V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Bit]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] HfB < 2100±20
Pearson Symbol/ Space Group/ Prototype oP8 Pbnm FeB
or cF8 Fm3m NaCl HfB2 < 3380 ± 20
hP3 P6/mmm AlB2
Lattice Parameters [pm] a = 492.38 b = 652.4 c = 322.35 a = 652 b = 322 c = 492 a = 462 ± 2
a = 314.28 c = 347.69 a = 314.31 c = 347.80 a = 314.3 c = 347.9 a = 314.5 c = 347.6 a = 314.2 c = 347.8 a = 314.3 c = 347.8 a = 314.2 c = 347.3
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Comments/References Tm = 2104˚C [1997Bit2] [1992Rog]
from alloy Hf76B20C4, quenched from 1986˚C, contains also (αHf) [1966Rud] [V-C2] impurity stabilized [Mas2]
Tm = 3377˚C [1997Bit2] Hf rich [1992Rog] B rich [1992Rog] from alloy Hf44B48C8, quenched from 2938˚C, contains HfC1–x and (αHf) [1966Rud] from alloy Hf38B49C13, quenched from 3145˚C, contains HfC [1966Rud] from powderized melt HfB2 + C (from kX-units) [1965Lev] from alloy Hf15B55C35, quenched from 2310˚C, contains C and ‘B4C’ [1966Rud] from alloy HfB2 + 8.7 mol% ‘B4C’ [1989Ord]
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B–C–Hf
. Table 2 (continued)
Phase/ Temperature Range [˚C] HfC1–xa) < 3950 ± 40
Pearson Symbol/ Space Group/ Prototype cF8 Fm3m NaCl
Lattice Parameters [pm]
Comments/References Tm = 3942˚C [1997Bit2] 34.1 to 49.5 at.% C [1966Rud] at 34 at.% C [1966Rud] at 49.5 at.% C [1966Rud] from alloy Hf83B5C12 quenched from >2400˚C, contains (Hf) and HfB [1966Rud] from alloy Hf58B20C22, quenched from 2860˚C, contains HfC1–x, HfB2 and (αHf) [1966Rud] from alloy Hf50B29C21, quenched from 3118˚C contains HfB2 [1966Rud] from alloy Hf30B35C35, quenched from 2733˚C contains HfC1–x and C [1966Rud] from alloy Hf57B10C33 tin quenched from 3100˚C, single phase [1966Rud]
a = 460.8 a = 464.0 a = 462.7 a = 464.0 a = 464.3 a = 463.9 a = 464.9
a)
Note: Hf3C2 and Hf6C5, as claimed by [1991Gus], were said to be NaCl type derivative superlattice structures with symmetries Hf6C5 (C2, C2/m or P31) or Hf3C2 (Immm or P3m1), respectively.
. Table 3 Invariant equilibria in the B-C-Hf system; comparison of experimental data from [1966Rud] with thermodynamic calculation Calculated data Composition, at.%
Experimental data Composition, at.% Reaction LÐ HfB2 + HfC1–x
Type
Phase Hf
e2(max) L HfB2
C
40
44
16
3140±15 40.8 42.9
34
>64
<2
33.3 66.7
HfC1–x 55 L Ð HfB2 + (C)gr
LÐHfB2+HfC1–x + (C)gr
Landolt‐Bo¨rnstein New Series IV/11E1
e3(max) L
E1
T [˚C]
B
12 33
Hf
B
C 16.3
37.5
21
42
37
2515±10 20.2 40.7
39.1
HfB2
>33
>64
<3
33.3 66.7
0.0
(C)gr
<1
2
>97
L
24
38
38
HfB2
33
>64
HfC1-x 51 (C)gr
<1
0.9
99.1
20.9 39.3
39.7
<3
33.3 66.7
0.0
5
44
50.1
4.4
45.6
2
>97
0.0
0.9
99.1
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2480
3111
0.00
50.5 12.0
0.0
T [˚C]
2514.2
2513.6
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B–C–Hf
. Table 3 (continued) Calculated data Composition, at.%
Experimental data Composition, at.% Reaction L Ð HfB2 + ‘B4C’
L Ð HfB2 + ‘B4C’ + (C)gr
L + ‘B4C’ Ð HfB2 + (βB)
L Ð HfB2 + ‘B4C’+(βB)
L + HfB2 Ð HfB + HfC1–x
Type
Phase Hf
e4(max) L
E2
U1
E
U2
7
U3
L Ð(αHf)+(βHf)+HfB E3
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T [˚C]
C
78 15 2330±25
Hf
B
C
4.2 80.0
15.8
HfB2
33
>65
33.3 66.7
0.0
‘B4C’
<1
81 >18
0.0 81.9
18.1
L
9
66 25 2260
6.6 64.9
28.5
HfB2
33
>64
33.3 66.7
0.0
‘B4C’
<1
80 >19
0.0 80.4
19.6
(C)gr
<1
2
>97
0.0
2.1
97.9
L
-
-
-
1.1 98.4
0.5
‘B4C’
-
-
-
0.0 90.2
9.8
HfB2
-
-
-
33.3 66.7
0.0
(βB)
-
-
-
0.0 98.6
1.4
L
2
96 2
-
-
-
HfB2
33
>66
-
-
-
‘B4C’
<1
89 >10
-
-
-
(βB)
<1
98 <1
-
-
-
<2
<3
-
1950
<1
L
79
18 3
78.6 18.9
2.5
HfB2
34
>64
<2
33.3 66.7
0.0
HfB
50
>49
<1
50.0 50.0
0.0
9
31
61.2
HfC1-x 60 L+HfC1–xÐ HfB +(αHf)
B
L
84
HfC1-x 62
14 2
2050
1940
7.8
31.0
82.1 15.6
2.3
6
32
62.8
6.5
30.8
HfB
51
>48
<1
50.0 50.0
0.0
(αHf)
90
1
9
92.3
0.4
7.3
L
87.5 11 1.5 1850
85.0 14.9
0.1
(αHf)
96
3
97.7
1.3
2.0
(βHf)
>97.5 1.5 <1
99.0
0.96
0.04
HfB
50
50.0 50.0
1
49
MSIT1
<1
T [˚C] 2404
2318
2091
-
2063
1954
1880
0.0
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. Fig. 1 B-C-Hf. Calculated B-Hf phase diagram
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B–C–Hf
. Fig. 2 B-C-Hf. Calculated C-Hf phase diagram
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. Fig. 3 B-C-Hf. Lattice parameters of Hf(C,B)1–x (in pm) from alloys annealed at 1500˚C [1961Now]; phase boundaries at 1500˚C [1961Now] (solid lines) and at 1400˚C [1966Rud] (dashed lines)
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B–C–Hf
. Fig. 4 B-C-Hf. Calculated isopleth “HfB2-C” with experimental data from [1965Lev] and [1966Rud]
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. Fig. 5 B-C-Hf. Calculated isopleth HfB2 - HfC1–x at the maximum in the liquid trough with experimental data from [1966Rud] and [1977Ord]. Note the top axes indicating experimental sections for sample location
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B–C–Hf
. Fig. 6 B-C-Hf. Calculated isopleth HfB2 - ’B4C’ at the maximum in the liquid trough with experimental data from [1966Rud] and [1989Ord]. Note the top axes indicating experimental sections for sample location
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. Fig. 7 B-C-Hf. Reaction scheme
B–C–Hf
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20
17
B–C–Hf
. Fig. 8 B-C-Hf. Calculated liquidus projection at 100 K intervals and invariant temperatures. Experimental data by [1966Rud] are indicated by symbols
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. Fig. 9 B-C-Hf. Experimental isothermal section at 1400˚C
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B–C–Hf
. Fig. 10 B-C-Hf. Calculated isothermal section at 1954˚C, temperature of the four-phase equilibrium: L + HfC1–x⊊Ð⊊HfB + (αHf)
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. Fig. 11 B-C-Hf. Calculated isothermal section at 2063˚C, temperature of the four-phase equilibrium: L + HfB2 Ð HfB + HfC1–x
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24
17
B–C–Hf
. Fig. 12a B-C-Hf. Calculated isothermal section at 2091˚C, temperature of the four-phase equilibrium: L + ’B4C’ Ð HfB2 + (βB) (see Fig. 12b for an enlarged view of the B rich region)
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. Fig. 12b B-C-Hf. Calculated isothermal section at 2091˚C, temperature of the four-phase equilibrium: L + ’B4C’ Ð HfB2 + (βB), enlarged view of the B rich region
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26
17
B–C–Hf
. Fig. 13 B-C-Hf. Calculated isothermal section at 2350˚C
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. Fig. 14 B-C-Hf. Calculated isothermal section at 3100˚C
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17
B–C–Hf
. Fig. 15 B-C-Hf. Calculated isopleth from Hf to B0.817C0.183
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. Fig. 16 B-C-Hf. Calculated isopleth HfB - C
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B–C–Hf
. Fig. 17 B-C-Hf. Calculated isopleth HfC - B
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. Fig. 18 B-C-Hf. Calculated isopleth HfC1–x - ’B4.5C’
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17
B–C–Hf
. Fig. 19 B-C-Hf. Calculated isopleth Hf -B0.5C0.5
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. Fig. 20a B-C-Hf. Comparison of calculated isopleths for various sections from HfB2 to B0.92C0.08
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B–C–Hf
. Fig. 20b B-C-Hf. Comparison of calculated isopleths for various sections from HfB2 to B0.86C0.14
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. Fig. 20c B-C-Hf. Comparison of calculated isopleths for various sections from HfB2 to B0.82C0.18
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17
B–C–Hf
. Fig. 20d B-C-Hf. Comparison of calculated isopleths for various sections from HfB2 to B0.80C0.20
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References [1961Now]
[1965Lev]
[1966Rud]
[1967Low] [1969Rud]
[1974Upa]
[1977Ord]
[1980Ord]
[1981Cre] [1983Sch]
[1984Hol]
[1988Rog] [1989Ord]
[1990Ase]
[1991Gus]
[1992Rog]
[1993Wer]
Nowotny, H., Rudy, E., Benesovsky, F., “Investigations in the Systems: Hafnium - Boron - Carbon and Zirconium - Boron - Carbon” (in German), Monatsh. Chem., 92, 393–402 (1961) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 16) Levinskii, Y.V., Salibekov, S.E., “Interaction of Titanium, Zirconium and Hafnium Diborides with Carbon”, Russ. J. Inorg. Chem., 10(3), 319–320 (1965), translated from Zh. Neorg. Khim., 10(3), 588–591 (1965) (Experimental, Crys. Structure, Phase Diagram, 6) Rudy, E., Windisch, S., “Part II. Ternary Systems, Vol. XIII. Phase Diagrams of the Systems Ti-B-C, Zr-B-C and Hf-B-C” in “Ternary Phase Equilibria in Transition Metal - Boron - Carbon - Silicon Systems”, Report AFML-TR-65–2, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, Part II. Vol. XIII, 1–212 (1966) (Experimental, Phase Diagram, Phase Relations, 96) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Rudy, E., “Part V. Compendium of Phase Diagram Data, Section III.K.3. Hf-B-C System” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 and 33(615)-67-C-1513, Air Force Materials Laboratory, WrightPatterson Air Force Base, OH, Part V, 635–654 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 3) Upadkhaya, G.S., “Nature of the Phase Diagram of some Transition Metals with Boron” in “Bor: Poluchenie, Struktura, Svoistva”, Mater. Mezhdunar. Simp. Boru, 4th Meeting Date 1972, Metsniereba, Tbilisi, 2, 115–123 (1974) (Review, Phase Diagram, Phase 17) Ordanyan, S.S., Unrod, V.I., Lutsenko, A.E., “Reaction in the System HfC-HfB2”, Inorg. Mater. (Engl. Trans.), 13, 451–453 (1977), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 13(3), 546–547 (1977) (Experimental, Morphology, Phase Diagram, 4) Ordanyan, S.S., “Laws of Interaction in the Systems MIV,VC -MIV,VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411, (1980) (Experimental, Thermodyn., 14) Crespo, A.J., Tergenius, L.E., Lundstro¨m, T., “The Solution of 4d, 5d and Some p Elements in Rhombohedral Boron”, J. Less-Common Met., 77, 147–150 (1981) (Experimental, Crys. Structure, 12) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Rogl, P., Potter, P.E., “A Critical Review and Thermodynamic Calculation of the Binary System Hf-B”, Calphad, 12(3), 207–218 (1988) (Review, Thermodyn., Phase Diagram, 43) Ordanyan, S.S., Dmitriev, A.I., “Interaction in the B4C-HfB2 System”, Powder Metall. Met. Ceram., 28(5), 424–426 (1989), translated from Poroshk. Metall., 5(317), 99–101 (1989) (Crys. Structure, Experimental, Morphology, Phase Diagram, Review, Theory, Thermodyn., 5) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides” Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, , Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State, 32(9), 1595–1599 (1991) (Theory, Phase Diagram, Phase Relations, Thermodyn., 18); see also Gusev, A.I., “Physical Chemistry of Nonstoichiometric Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow (1991) (Review, Thermodyn., Crys. Structure, Phase Diagram, 102) Rogl, P., “The System Hf-B-N” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, Ohio, USA (1992) (Experimental, Crys. Structure, Phase Diagram, Review, Phase Relations, 8) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51)
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38
17 [1993Ord] [1994McH]
[1995Vil]
[1996Kas] [1997Bit1]
[1997Bit2]
[1998Bit]
[1998Kas]
[1999Rog] [2000Rog] [2000Pro]
[2001Gos] [2004For]
[2004Blu]
[2005Tan]
[2006Won]
[Mas2] [V-C2]
B–C–Hf Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, 1, 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, 18) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceram. Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 183–188 (1994) (Phase Diagram, Phase Relations, Review, 4) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA (1995) (Review, Phase Diagram, Phase Relations, Crys. Structure, 5) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Bittermann, H., Rogl, P., “Critical Assessment and Thermodynamic Calculation of the Binary System Hafnium - Carbon (Hf-C)”, J. Phase Equilib., 18(4), 344–356 (1997) (Thermodyn., Phase Diagram, Crys. Structure, Phase Relations, 63) Bittermann, H., Rogl, P., “Critical Assessment and Thermodynamic Calculation of the Ternary System Boron - Hafnium - Carbon (B-Hf-C)”, J. Phase Equilib., 18(1), 24–47 (1997) (Thermodyn., Phase Diagram, Crys. Structure, Phase Relations, 39) Bittermann, H., Rogl, P., “The System Boron - Carbon - Hafnium” in “Phase Diagrams of Ternary MetalBoron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 102–141 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 22) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides”, Int. J. Refract. Met. Hard Mater., 17, 27–32 (1999) (Crys. Structure, Experimental, Phase Relations, 6) Rogl, P., Bittermann, H., “On the Ternary System Hafnium-Boron-Carbon”, J. Solid State Chem., 154, 257–262 (2000) (Assessment, Calculation, Phase Relations, 17) Provot, B., Deschanels, X., Bry, P., “Neutron Absorber Material based on Boron Carbide and Hafnium and Method for Making same” (Patent written in French) PCT Int. Appl. 45 pp. (2000) (Mechan. Prop.) Gosset, D., Provot, B., “Boron Carbide as a Potential Inert Matrix: an Evaluation”, Prog. Nucl. Energy, 38 (3-4), 263–266 (2001) (Experimental, Mechan. Prop., Morphology, 8) Forsthoefel, K., Sneddon, L.G., “Precursor Routes to Group 4 Metal Borides, and Metal Boride/Carbide and Metal Boride/Nitride Composites”, J. Mater. Sci., 39(19), 6043–6049 (2004) (Experimental, Phase Relations, Phys. Prop., 35) Blum, Y.D., Kleebe, H.J., “Chemical Reactivities of Hafnium and Its Derived Boride, Carbide and Nitride Compounds at Relatively Mild Temperature”, J. Mater. Sci., 39(19), 6023–6042 (2004) (Experimental, Morphology, Phase Relations, Thermodyn., 40) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), August 21–26, 142 (2005) (Experimental, Phase Diagram, Phase Relations, 4) Wong, J., Larson, E.M., Waide, P.A., Frahm, R., “Combustion Front Dynamics in the Combustion Synthesis of Refractory Metal Carbides and di-Borides Using Time-Resolved X-ray Diffraction”, J. Synchrotron Radiat., 13(pt.4), 326–335 (2006) (Experimental, Phys. Prop., 30) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Molybdenum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl, Kostyantyn Korniyenko, Tamara Velikanova
Introduction The system boron-carbon-molybdenum is one of the few refractory metal boron carbide systems that exhibit a ternary compound, which forms a stable equilibrium with carbon at temperatures higher than 2000˚C. Based on early studies [1952Gla, 1953Ste, 1955Gea] on the interaction between molybdenum borides, molybdenum carbides, boron carbide and graphite, a first phase triangulation in the B-C-Mo system at 1775˚C was attempted by [1955Bre], proposing stable two-phase equilibria MoB2-‘B4C’, MoB2-C, MoB-C, “Mo5B3”-C, Mo2B-C and Mo2B-Mo2C. The existence of the phases labelled as “Mo5B3” or “Mo3B2” (U3Si2 type) [1952Gla, 1953Ste, 1955Bre] was not confirmed in later investigations by [1963Rud, 1965Lev] and therefore are neither part of the binary B-Mo system nor of the ternary B-C-Mo system in the temperature regions investigated. It should be mentioned, that early recordings of the melting behaviour of the binary molybdenum borides [1953Ste] suffered from the fact, that whenever graphite supports or substrates were used in these measurements, the metal boride-graphite interaction was recorded instead. Later papers give detailed information on the phase relations in the ternary system from the investigations by [1963Rud] and [1965Lev], providing isothermal sections at 1300˚C and 1800˚C [1963Rud], information on the compatibility of the molybdenum borides and graphite in the range from 1825 to 2125˚C [1965Lev] as well as information on the “eutectic nature” of the sections Mo2B5–x-C, MoB-C and Mo2BC-C [1965Lev]. A liquidus surface projection was established by [2002Kor]. Agreement exists on the formation of a ternary boron carbide Mo2BC1–x [1963Rud, 1963Jei, 1965Lev, 1969Smi, 1970Sal, 1977Bov, 1981Lej, 2002Kor]. But Mo2BC1–x was not revealed from a diffusion couple Mo + ‘B4C’ after 10 h at 1700˚C [1980Mor]. Mo2BC1–x was said to melt congruently at 2800 ± 100˚C [1963Rud], 2100 ± 20˚C [1970Sal], 2306 ± 20˚C [2002Kor] and to form a quasibinary eutectic with carbon [1965Lev, 2002Kor]. Accordingly, there is no stable two-phase equilibrium for MoB-MoC and Mo2B-C as claimed earlier by [1960Eng] and [1952Gla, 1955Bre], respectively. Compilations of the most relevant data on the topology of the B-C-Mo system were published by [1984Hol, 1983Sch, 1994McH, 1995Vil]. A previous MSIT assessment for the B-C-Mo system including literature up to 1996 was made in the context of a general review of phase relations for metal-boron-carbon systems [1998Rog]. Experimental details for all investigations in the B-C-Mo system are summarized in brief in Table 1.
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Binary Systems The accepted B-C phase diagram is shown in the evaluation of the B-C-Cr system in the present volume. It corresponds to an assessment and thermodynamic modeling by [1998Kas], based on [1996Kas]; the disputed peritectic boron-rich reaction L + ‘B4C’ Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan]. The B-Mo binary system taken from [1992Rog] is included in the present volume in the evaluation of the Al-B-Mo system. The system C-Mo is basically accepted from [Mas2]. Modifications concerning the Mo2C sub-carbide refer to recent in-situ-neutron diffraction studies by [1988Epi], who revealed a second-order transition between the disordered high temperature Mo2C (hP3) to ordered MoC2 (hP9) at temperatures, which smoothly vary from 1650˚C at the Mo-rich boundary towards 1960˚C at the carbon-rich boundary of the sub-carbide. At variance to the detailed version by [1965Rud, 1967Rud], who claimed the existence of two isotypic sub-carbides αMo2C and βMo2C forming a eutectic at 2469 ± 6˚C (the α-βMo2C separation was interpreted [1965Rud] on the basis of two energetically different interstitial lattice sites), the authors of [1988Vel, 1988Epi] assumed only one sub-carbide region at elevated temperatures. Further agreement between [Mas2] (reporting the version of [1988Vel]) and [1988Epi] exists on the first-order transition of the (hP9) – sub-carbide to the orthorhombic low temperature modification (oP16, the space group Pbcn). The large volume orthorhombic superstructure Mo2C (r2) at the carbon-rich side of the sub-carbide as well as the temperature of the eutectoid decomposition Mo2C (hP9) Ð (Mo) + Mo2C (oP16) [1988Vel], however, are at variance with the diagram shown by [1988Epi]. Except for the phase transitions in Mo2C, a further significant difference between the constitutional diagram proposed by [1988Vel] and the earlier version by [1965Rud, 1967Rud] is the melting behaviour of ηMoC1–x. According to the experimental data of [1965Rud], ηMoC1–x was observed to melt congruently at 39 at.% C and 2550 ± 5˚C, whereas [1988Vel] proposed a maximum transition point δMoC1–x Ð ηMoC1–x at 38 at.% C and 2530 ± 20˚C in an assumed combination with a metatectic reaction δ Ð η + L at 37.5 at.% C and 2520˚C, which certainly needs further verification. For this assessment we follow the version of [1988Epi, 1988Vel]. Literature data concerning the formation and crystal structure of binary and ternary solid phases pertinent to the B-C-Mo system are listed in Table 2.
Solid Phases Mo2BC is the only ternary compound in the B-C-Mo system. [1963Jei, 1963Rud, 1965Lev, 1969Smi, 1970Sal, 1977Bov, 1981Lej, 2002Kor] confirmed the existence of the ternary compound Mo2BC1–x with a unique structure type. The crystal structure was first established from X-ray single crystal photographs by [1963Jei]. X-ray single crystal counter data refinement [1969Smi] (RF = 0.035) established precise data for the light atom arrangement in agreement with earlier assumptions from structural chemical principles [1963Jei]. Chemical analysis of the homogeneous bulk material prepared by argon arc-melting revealed a small carbon defect according to the formula Mo2BC0.97(2) [1969Smi], earlier suspected by [1963Jei]. Planar defects in Mo2BC1–x have been studied by [1977Bov] by means of TEM on single crystalline material revealing various intergrowth (Waldsley) defects among the MoC1–x (NaCl type) and MoB (βCrB type) structural units forming the crystal structure of Mo2BC corresponding to DOI: 10.1007/978-3-540-88053-0_18 ß Springer 2009
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the formal topochemical description Mo2BC Ð MoB + MoC. Attempts to grow large single crystals of Mo2BC gave strong evidence for a C-deficiency in Mo2BC1–x, which resulted in the formation of parasitic Mo2C [2003Lei]. This phase precipitated during the growth and was found isotropically distributed within the individual grains of the crystal. Also a very pronounced trend for facetting during growth was observed. The formation of a spiral like pattern on the cut surface can be understood as an indication of significant composition fluctuations during the crystal growth [2003Lei]. From the interactions between molybdenum borides, molybdenum carbides, boron carbides and carbon, the authors [1952Gla, 1953Ste, 1955Bre, 1955Gea, 1963Rud, 1965Lev, 1970Sal, 2002Kor] essentially agree on small or negligible mutual solid solubilities. As far as the defect type molybdenum monocarbides ηMoC1–x and δMoC1–x are concerned, amounts as small as 0.2 at.% B were said to efficiently stabilize these compounds to temperatures considerably lower than their decomposition temperatures in the C-Mo binary [1979Hol, 1982Kra, 1997Ath]. Maximal solid solubilities of C in (Mo) have been measured by [1981Sch] to be 0.018 at.% C at 1400˚C and 1 at.% C at 2150˚C (logcC.max = 3.92 – 9490/T (K); c in mass ppm). A small homogeneous region was observed for the ternary compound Mo2BC1–x in the temperature range investigated [2002Kor]. Crystallographic data for the binary boundary phases and the compound Mo2BC1–x are summarized in Table 2.
Quasibinary Systems There are numerous reports on the mutual stability for the pairs Mo2B-Mo2C [1955Gea], Mo2BC1–x-C; the latter was claimed to form a quasibinary eutectic at 1825˚C [1965Lev, 1970Sal]. Similarly Mo2B5–x-C and MoB-C were claimed to be “quasibinary” and melting was said to occur at 2175˚C and 2300˚C respectively [1965Lev]. From the accepted B-Mo binary, however, it becomes obvious that Mo2B5–x (and MoB2) are melting incongruently thereby being inconsistent with a true quasibinary Mo2B5–x (or MoB2)-C. Similar arguments probably hold for the section MoB-MoC [1960Eng]. It has to be noted, that some of the early melting point determinations on binary molybdenum borides are suffering from contamination with graphitic substrates or container materials actually recording molybdenum boridegraphite interactions instead. Investigation of the solidification behaviour in the B-C-Mo ternary system by [2002Kor] established seven quasibinary eutectic maxima shown in Fig. 2 and listed in Figs. 1a, 1b as part of a Scheil diagram. All sections are in agreement with the phase triangulation of [1963Rud, 1965Lev]. The pair Mo2BC1–x-C was claimed to form a quasibinary eutectic at 1825˚C [1965Lev, 1970Sal], in contrast to higher temperatures (2128±20˚C and 2193 ± 40˚C) for neighbouring invariant equilibria, determined by [2002Kor]. Attempts to grow directionally solidified eutectic structures for the MoB2-B4C system revealed that a composition Mo19.5B72.2C8.3 (70 mol% MoB2) is close to a eutectic structure with disordered phase distribution with a chinese script like nature [2004Pad].
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Invariant Equilibria Solidification behaviour of the alloys in the entire B-C-Mo system is presented in a Scheil diagram summarizing invariant equilibria involving a liquid phase [2002Kor] (Table 3 and Figs. 1a, 1b). For a discussion of invariant quasibinary reactions see section “Quasibinary Systems”. Some disagreement exists on the congruent melting point of Mo2BC1–x, which was reported at 2800 ± 100˚C by [1963Rud], a value of 2100 ± 20˚C was assigned by [1970Sal], but more recent data by [2002Kor] arrived at 2306 ± 20˚C.
Liquidus, Solidus and Solvus Surfaces The liquidus surface is plotted in Fig. 2 referring to LOM, Pirani, EMPA and X-ray data of [2002Kor]. It comprises eleven ternary invariant reactions with the participation of liquid and thirteen surfaces of primary crystallization.
Isothermal Sections Isothermal sections at 1300˚C and 1800˚C are shown in Figs. 3 and 4, respectively. Data essentially refer to the experimental findings by [1963Rud] with small changes to conform to the accepted binaries. The phase triangulation resulting from the interaction studies of [1965Lev] in the temperature range 1827 to 2127˚C agrees with the data given by [1963Rud]. A diffusion couple Mo + B4C [1980Mor], heated for 10 h at 1700˚C, however, revealed two sublayers consisting of Mo2B and MoB, respectively, with dispersed carbon particles in contrast to the experimentally established two-phase equilibrium regions Mo2B + Mo2C, MoB + Mo2C and MoB + Mo2BC1–x [1963Rud]; see also Figs. 3 and 4.
Thermodynamics No experimental thermodynamic data are presently available for the ternary system. Attempts to estimate upper limits for the heat of formation of molybdenum borides from compatibility studies among molybdenum borides and carbon were not successful [1955Bre].
Notes on Materials Properties and Applications Details on synthesis and properties of Mo2BC1–x were given by [1970Sal] i.e. microhardness (17.7 GPa at a load of 50 g), coefficient of thermal expansion (7.02(1)·10–6 K–1 in the range from 20 to 600˚C, measured in a quartz dilatometer), modulus of elasticity (382 GPa as a maximum value for a specimen with 4 % porosity), electrical and thermal conductivity (σ = 1.10(5) M(Ωm)–1, χ = 16.0(1) W·m–1·K, extrapolated for 20˚C and zero porosity, and thermoelectric force (–4.2 μV·K–1 at 20˚C to –2.2 μV·K–1 at 750˚C). Resistivity at 10 K was 0.60 μΩm for a single crystal [1981Lej]. The rate of oxidation was reported [1970Sal] to increase markedly at 750 to 800˚C, owing to fusion of MoO3 and faster diffusion of oxygen DOI: 10.1007/978-3-540-88053-0_18 ß Springer 2009
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through the liquid phase. Powders of ‘Mo2BC’ with a particle size of < 90 μm started to oxydize in a flow of pure oxygen above 320 K (thermogravimetric measurements) [1981Lej]. Hardness of Mo2BC1–x at elevated temperatures was reported to be quite high (HV = 1780 at 25˚C and HV = 950 at 1000˚C) [1979Hol]. These values compare well with those given by [1981Lej] for polycrystalline material (HV = 1823 ± 175) or 2092 ± 60 for the (101) plane of a Czochralski-grown single crystal of Mo2BC obtained from a triarc furnace at > 2500˚C. Mo2BC1–x is a type II superconductor with Tc = 7.0 to 5.3 K [1967Tot]. [1981Lej] reported Tc = 7.5 K for a single crystal specimen with critical fields HC1 ≈ 100Oe and HC2 ≈ 27 kOe at 4.2 K. The superconducting transition temperature of boron stabilized δMoC1–xBx was said to linearly increase with increasing ratio C/Mo but to decrease with growing boron content: Tc = 10.6 K for MoC0.67B0.01, Tc = 12.1 K for MoB0.01C0.73 but 11.9 K for MoB0.05C0.73 [1982Kra]. A maximum Tc = 12.2 K was measured for MoB0.01C0.82 [1997Ath].
Miscellaneous Forming a boron-enriched surface layer improves the ductility of molybdenum in the first wall of fusion reactors although boron-coating and heating was not as effective as carbon-doping [1981Hir]. Vacuum brazes were developed for molybdenum which can be used in seal joint applications up to 1870 K (1597˚C). The braze materials included (contents in mass%), 50MoB-50MoC and 85Mo-15MoB2 in the form of powder mixtures. Melting temperature for 50MoB-50MoC, according to estimations, lies in the interval 2225–2480 K (1952 to 2207˚C), and for 85Mo-15MoB2 is equal to 2225 K (1952˚C). Brazing temperatures were 2330 K (2057˚C) (during less than 30 s) and 2425 K (2152˚C) (600 s), respectively [1978Lun]. Peak reflectivity measurements of Mo/B4C multilayer mirrors have been performed using line and synchrotron radiation for peak reflectivities from 10 to 25 % between 9 and 20 nm and bandpasses as small as 0.3 nm [1990Zwi]. Improvement of the thermo-mechanical properties of boron carbide based composites (B4C-Mo) containing more than 70 vol% B4C was reported by [1994Gos]. [1991Kus] studied the kinetics of electrochemical precipitation of molybdenum carbide MoC1–x on the surface of electro-conductive abrasive material B4C by electrolysis of ionic melts as a function of temperature, current density and duration of the process as well as correlation of these parameters with properties of metallized material. Deposition of molybdenum carbide on the surface of boron-carbide was successful in Na2WO4-MoO3-LiCO3 or KCl-NaCl-Na2MoO4-Na2CO3 melts at temperatures of 1073–1173 K (800 to 900˚C) and cathode current densities of 10–100 A·m–2 [1996Mal]. First-principles full-potential linear muffin-tin orbital calculations have been used to study the effect on the cohesion and electronic structure of cubic δMoC when 25% of carbon is substituted by boron [1999Hug].
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. Table 1 Investigations of the B-C-Mo System Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1952Gla]
Hot-pressing of powder mixtures (powders of Mo, Mo2C and B, ‘B4C’) in graphite dies at 1450 to 2100˚C. XPD
XPD on ten hot-pressed samples. Study of interation among Mo, B, C
[1955Bre]
Reaction between metal borides and graphite. Sintering of powder compacts in Mo-crucibles under 0.5 bar argon for 50 min at 1777˚C. XPD
XPD on three samples (Mo+B+C; 2Mo+B +0.5C; 2Mo+5B+C). From powder spectrum of Mo+B+C indication for the formation of ternary compounds
[1955Gea] Argon arc melting on a Cu-hearth. XPD
XPD, detection of simple mixtures Mo2C + Mo2B
[1960Eng] Samples were prepared by sintering, arc melting and in a plasma arc furnace and implosive shock technique
Interaction in the system MoB-MoC investigated by melting point analysis, XRD, LOM, oxidation resistance
[1963Jei]
Determination of the crystal structure of Mo2BC on single crystal X-ray photographs
Crystal fragment isolated from a regulus of congruently melting Mo2BC. XRD
[1963Rud] 214 samples including 150 ternary alloys were prepared by short duration hot pressing in graphite dies at 1200 – 2500˚C followed by subsequent homogenization in vacuum of 10–1 Pa argon at 1900, 1700, 1500, 1300˚C for 4 to 35 h. Some of the alloys were also arc-melted. Starting materials were powders of Mo (99.95 mass% pure), carbon (purissimum), boron (94% pure) and Mo2C (containing 5.99 mass% of carbon of which 0.05 mass% was free carbon) prereacted from metalcarbon powder blends in a carbon tube furnace at 1600 to 2000˚C. XPD and LOM.
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Isothermal section at 1800 and 1300˚C. Stability of Mo borides against C. Influence of B on the stability of Mo carbides.
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. Table 1 (continued) Temperature / Composition / Phase Range Studied
Reference
Method / Experimental Technique
[1965Lev]
MoB and Mo2B5 powders were prepared by vacuum sintering powder compacts for 5 h at 1525˚C in a W heater furnace and analysed by XPD, LOM. Starting materials were powders of < 5 μm: MoB2 containing 0.001 mass% Ni, 0.35 % O, 0.001 % SiO2 and B-powder containing 0.0036 % Fe, 0.0003 Si, 0.01 Cu, 0.0004 Al, 0.0006 Pb
Interaction of MoB and Mo2B5 with C was claimed to be of the quasibinary eutectic character in both cases. TE (MoB-C)= 2300˚C; TE (Mo2B5-C)= 2180˚C. Mo2BC formed in samples containing Mo90B5C5 (mass%). Mo2BC + C was said to be stable up to 1830˚C. The results of [1965Lev] were obtained by X-ray diffractometry, metallography and micro-optical pyrometric melting point measurements on metal boride-carbon powder mixtures reacted inside a bore of a graphite tube directly heated by electric current under vacuum. After the melting run, the samples were quenched at a rate of 100 K·s–1 and examined. Melting temperature was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube
[1967Tot]
Superconductivity was measured on specimens from powders hotpressed in graphite dies at 2000–2500˚C
XRD
[1969Smi] Crystal fragment isolated from an arc melted button of Mo2BC [1970Sal]
Mo2BC was prepared from two powder mixtures (2Mo+B+C; 8Mo+B4C+3C) which were cold compacted and vacuum annealed. Reaction started at 1100˚C and was complete after 3 h at 1300˚C. Starting materials were powders of Mo (< 5 μm, 99.6 mass% Mo), B4C (< 30 μm, 78 % B, 21 % C) and lampblack C. Physical property measurements were made on specimens hotpressed in graphite dies at 2000±20 K (1727±20˚C) for 5 min
Determination of the crystal structure of Mo2BC and particularly of the light atom sites from single crystal X-ray counter data Investigation of Mo2BC: melting point, pycnometric density, microhardness at a load of 50 g, coefficient of thermal expansion (measured in a quartz dilatometer), modulus of elasticity for a specimen with 4 % porosity, oxidation in air by gravimetry, electrical conductivity (in the range from 20–750˚C), thermo-emf (at 20, 400, 750˚C), thermal conductivity at RT
[1980Mor] Diffusion couples of Mo+C, heated for 3.6 Investigation of interaction via diffusion 105 s at 1700˚C. LOM, EMPA, SEM-TEM zones [1981Lej]
Czochralski grown single crystal of Mo2BC. Study of magnetization as f (H, T). XPD, LOM, EMPA
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. Table 1 (continued) Reference
Method / Experimental Technique
[2002Kor]
More than 10 ternary samples prepared with various techniques from high purity powders of components: molybdenum (99.9 and 99.85 mass%), boron (Ventron GmbH, D), boron carbide B4C (Johnson Matthey & Co, UK) and carbon (with purity of 99.99 mass%). For melting point measurements by the Pirani technique specimens were prepared in form of cylindrical polycrystalline rods by sintering at ca 2000˚C of ceramic green bodies from isostatically compressed blends of Mo, B4C, B and C powders. Powder blends were compacted in form of the final Pirani shape in steel dies using a small amount of xylol as densification aid. The green bodies (with a lateral hole directly pressed into the sample) were slowly heated (to prevent violent selfheating via exothermic reactions) in high vacuum to 1600˚C for 12 to 15 hours. For LOM and EMPA, several alloys were prepared by argon arc melting of sintered pellets
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Temperature / Composition / Phase Range Studied Melting points were determined on about four Pirani samples of each composition in a precise micro-pyrometer calibrated for the temperature region from 2000 to 2900˚C. The microstructure of the cast alloys was inspected using light optical microscopy (LOM) on flat surfaces prepared by grinding (SiC-paper) and polishing the resin-mounted alloys using diamond pastes down to 1/4 μm grain size. Quantitative composition analyses were performed on a CAMEBAX SX50 wavelength dispersive spectrograph (EMPA) comparing the X-ray emissions of the three elements in the alloys with those from elemental standards of Mo, Si and LaB5.85 for boron after applying a deconvolution and ZAF-correction procedure. Experimental parameters employed were: acceleration voltage of 15 kV, sample current of 20 nA and spectrometer crystals such as PET for Mo-Lα, PC1 for C-Kα and PC3 for the B-Kα radiation. Determination of liquidus and solidus in the whole range of compositions. Determination of reaction isotherms for reactions involving liquid. Derivation of Schulz-Scheil diagram for the whole range of compositions.
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. Table 1 (continued) Reference
Temperature / Composition / Phase Range Studied
Method / Experimental Technique
[2004Pad] Rods of 8–10 mm diameter and 50 to 100 mm length were prepared from powder compacts (99.9 mass% Mo and B, 99.9 % B4.3C which were arc melted and crushed), annealed at 1500˚C in vacuum prior to zone melting in an induction furnace under 0.2 MPa argon. Compositions chosen were: Mo11B76C13 ( 46 mol% MoB2), Mo12.5B7.5C15.5 ( 50 mol% MoB2), Mo14B75C11 ( 56 mol% MoB2), Mo15.8B73.7C10.5 ( 60 mol% MoB2), Mo17.5B7.3C9.5 ( 65 mol% MoB2), Mo19.5B72.2 C8.3 ( 70 mol% MoB2)
Directional crystallisation of B4C-MoB2 eutectic compositions investigated by XPD, SEM and EMPA revealing chinese script eutectic structures
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Mo) < 2623
cI2 Im 3m W
a = 314.70 a = 314.73
[Mas2] from a sample with 27 at.% C quenched from 2000˚C, contains also Mo2C [1965Rud]
(βB) < 2092 [Mas2] < 2075 calculated [1998Rog]
hR333 R 3m βB
a = 1093.30 c = 2382.52
[1993Wer]
a = 1093.02 c = 2381.66 a = 1092.2 c = 2381.1 a = 1093.03 c = 2382.17
pure B [1976Lun]
(C) gr < 3827 (sublimation point)
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at 1.1 at.% C [1993Wer], linear ∂a/∂x, ∂c/∂x at MoB99.5 [V-C2] at 25˚C [Mas2]
[1967Low] at 2.35 at.% Cmax(2350˚C), linear ∂a/∂x, ∂c/∂x [1967Low]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(C)d
cF8 a = 356.69 Fd 3m C (diamond)
‘B4C’ < 2450
hR45 R 3m B13C2
a = 565.1 to 560.7 9 to 20 at.% C c = 1219.6 to 1209.5 [1990Ase]
B25C
tP52 P 42m B25C
a = 872.2 c = 508.0
[V-C2] also B51C1, B49C3, all metastable?
Mo2B < 2280
tI12 I4/mcm CuAl2
a = 554.8 c = 474.06
[1992Rog]
βMoB 2600 - 1800
oC8 Cmcm CrB
a = 314.02 b = 848.9 c = 307.1
[1992Rog]
αMoB < 2180
tI16 I41/amd αMoB
a = 310.68 c = 1696.18
[1992Rog]
MoB2–x 2375 - 1517
hP3 P6/mmm AlB2
a = 303.78 c = 306.03
at 62 to 65 at.% B [1992Rog]
Mo2B5–x < 2140
hR21 R 3m Mo2B5
a = 301.17 c = 2094.9
at 66 to 70 at.% B [1992Rog]
Mo1–xB3 < 1807
hP20 P63/mmc W1–xB3
a = 520.36 c = 635.02
at 80 at.% B [1992Rog]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] Mo2C (h1) 2537 - 1650
Pearson Symbol/ Space Group/ Prototype hP3 P63/mmc defect NiAs
Lattice Parameters [pm] a0 = 299.6 to 301.2 c0= 473.1 to 478.6 a = 299.6 c = 473.8 a = 300.7 c = 475.7 a = 301.2 c = 477.3 a = 299.0 to 301.1 c = 473.3 to 477.1 a = 299.0 c = 472.8 a = 300.6 c = 473.4 a = 300.7 c = 477.8 a = 301.0 c = 478.0 a = 299.0 to 301.0 c = 473.0 to 477.8
Mo2C (h2)
hP9
a = 525
1960 - 1190
P 31m W2C (a)
c = 479
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Comments/References 27 to 36 at.% C [1988Vel] Listed also as L’3 structure, (Fe2N type, βMo2C) [1988Epi] at 255˚C [V-C2]
at 725˚C [V-C2] at 1125˚C [V-C2] at 30 to 34 at.% C, 1700˚C [V-C2] at 30 at.% C, quenched from 2000˚C [1965Rud] at 32.8 at.% C, quenched from 2000˚C [1965Rud] at 33.5 at.% C, quenched from 2000˚C [1965Rud] at 34 at.% C, quenched from 2000˚C [1965Rud] at 30 to 34 at.% C, 2200˚C [V-C2] at 1500˚C, denoted as εMo2C [1988Epi]
at 1700˚C [1988Epi] at 1900˚C [1988Epi] linear ∂a/∂x, ∂c/∂x [1988Epi]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C] Mo2C (r1) ≤ 1380
oP16 Pbcn αMo
Lattice Parameters [pm] pffiffiffi a = 473.0 ≈ co 3 b = 602.7 ≈ 2ap 0 ffiffiffi c = 519.8 ≈ a0 3 a = 473.5 b = 602.5 c = 521.0 a = 473.2 b = 600.6 c = 520.3 a = 473.2 b = 604.8 c = 518.8 a = 476.2 b = 607.2 c = 521.6 a = 480.0 b = 609.0 c = 520.3
Comments/References denoted as β’-Mo2C [1988Vel] at Mo2C, 20˚C [1988Epi], listed as ζFe2N type at Mo2C0.94, 20˚C [1988Epi]
at 227˚C [V-C2]
at 727˚C [V-C2]
at Mo2C, 1350˚C [1988Epi]
Mo2C (r2) < 1220
-
a = 946.6 ≈ 2c0 b = 2415.2 ≈ 8a0 pffiffiffi c = 4167.5 ≈ 8a0 3
at 33.5 at.% C denoted as β’’Mo2C [1988Vel]
ηMoC1–x 2530 - 1647
hP12 P63/mmc ηMoC1–x
a = 301.2 c = 1463.4 a = 300.8 c = 1463 a = 301.2 c = 1465 a = 301.0 c = 1464
at 39 at.% C [1988Vel]
δMoC1–x 2605 - 1956
MoC < 1220
cF8 Fm 3m NaCl
hP2 P 6m2 WC
DOI: 10.1007/978-3-540-88053-0_18 ß Springer 2009
a = 426.6 to 428.1 a = 426.7 a = 428.1 a = 290.6 c = 282.2 a = 289.8 c = 280.9
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at 38 at.% C [1965Rud] at 39 at.% C [1965Rud] at Mo3.05C1.95 [V-C2] 37 to 43 at.% C [Mas2] 39.7 to 43 at.% C [1988Vel] at 41 at.% C [1965Rud] at 43 at.% C [1965Rud] at 50 at.% C [1988Vel] [V-C2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
*τ, Mo2BC1–x oC16 < 2800 ± 100 Cmcm [1963Rud] Mo2BC < 2100 ± 20 [1970Sal] < 2306 ± 20 [2002Kor]
a)
Comments/References [1963Rud] ρexp. = 8.54·103 kg·m–3 [1970Sal] ρexp. = 8.72·103 kg·m–3 [1981Lej] [1963Jei, 1969Smi]
a = 308.6 b = 1735 c = 304.7 a = 308.8 b = 1734.8 c = 304.6
[1977Bov]
partially disordered V2N type, cited by various authors as partially disordered εFe2N type, εMo2C [1988Epi]
. Table 3 Invariant Equilibria Composition (at.%) T [˚C]
Reaction
Type
Phase
B
C
Mo
L Ð δMoC1–x + Mo2BC1–x
-
e2 (max)
L
9
35
56
L Ð βMo2C + Mo2BC1–x
-
e3 (max)
L
12
29
59
L+ηMoC1–xÐδMoC1–x+βMo2C
-
U1
L
1
36
63
L Ð δMoC1-x+βMo2C+Mo2BC1–x
-
E1
L
9
33
58
L Ð βMoB + Mo2BC1–x
-
e6 (max)
L
1
19
50
L Ð βMoB + (C)gr
2300±25
e7 (max)
L
36
36
28
L + βMoB Ð MoB2 + (C)gr
-
U2
L
51
21
28
L Ð MoB2 + ‘B4C’
-
e8 (max)
L
72
8
20
L Ð βMoB + βMo2C
-
e9 (max)
L
25
17
58
L Ð MoB2 + ‘B4C’ + (C)gr
2265±15
E2
L
61
17
22
L Ð βMoB+βMo2C+Mo2BC1–x
2265±5
E3
L
26
18
56
L Ð (C)gr + Mo2BC1–x
-
e10 (max)
L
21
36
57
L Ð (C)gr + βMoB + Mo2BC1–x
2193±22
E4
L
27
30
43
L + βMoB Ð βMo2C + Mo2B
-
U3
L
22
15
63
L Ð (C)gr+Mo2BC1–x + δMoC1–x
2140±25
E5
L
11
41
48
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B–C–Mo
. Table 3 (continued) Composition (at.%) T [˚C]
Reaction
Type
Phase
B
C
Mo
L + MoB2 Ð Mo2B5–x + ‘B4C’
-
U4
L
79
4
17
L Ð (Mo) + Mo2B + βMo2C
2128±26
E6
L
14
10
76
L + ‘B4C’ Ð (βB) + Mo2B5–x
-
U5
L
92
2
6
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. Fig. 1a B-C-Mo. Reaction scheme, part 1
B–C–Mo
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B–C–Mo
. Fig. 1b B-C-Mo. Reaction scheme, part 2
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. Fig. 2 B-C-Mo. Liquidus surface projection
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18
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B–C–Mo
. Fig. 3 B-C-Mo. Isothermal section at 1300˚C
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. Fig. 4 B-C-Mo. Isothermal section at 1800˚C
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B–C–Mo
References [1952Gla] [1953Ste] [1955Bre] [1955Gea] [1960Eng]
[1963Jei] [1963Rud]
[1965Lev]
[1965Rud]
[1967Low] [1967Rud]
[1967Tot]
[1969Smi] [1970Sal]
[1976Lun] [1977Bov] [1978Lun]
[1979Hol]
[1980Mor]
[1981Hir]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron Systems”, Trans. AIME - J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Phase Relations, Experimental, Electr. Prop., 19) Steinitz, R., Binder, I., “New Ternary Boride Compounds”, Powder Met. Bull., 6(4), 123–125 (1953) (Crys. Structure, Experimental, Phase Relations, 6) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–405 (1955) (Phase Relations, Thermodyn., Experimental, Review, 19) Geach, G.A., Jones, F.O., “Interactions in Mixtures of Hard Metals at Very High Temperatures”, Plansee Proc., 1955, 80–91 (1955) (Crys. Structure, Phase Relations, Experimental) (cited from abstract) Engelke, J.L., Halden, F.A., Farley, E.P., “Synthesis of New High-Temperature Materials”, U.S. Dept. Com., Office Tech. Serv., P. B. Rept., 161.720, 1–39 (1960) (Crys. Structure, Experimental, *) (cited from abstract) Jeitschko, W., Nowotny, H., Benesovsky, F., “The Crystal Structure of Mo2BC” (in German), Monatsh. Chem., 94, 565–568 (1963) (Crys. Structure, Experimental, 4) Rudy, E., Benesovsky, F., Toth, L., “Studies of the Ternary Systems of the Group Va and VIa Metals with Boron and Carbon” (in German), Z. Metallkd., 54, 345–353 (1963) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 43) Levinskii, Yu.K., Salibekov, S.E., Levinskaya, M.K., “Reaction of Chromium, Molybdenum and Tungsten Bordies with Carbon”, Sov. Powder Metall. Met. Ceram. 12(36), 1004–1009 (1965) translated from Poroshk. Metall. (Kiev), 12(36), 56–62 (1965) (Crys. Structure, Experimental, *, 5) Rudy, E., Windisch, S., Chang, Y.A., “Ternary Phase Equilibria in Transition Metal-Boron-CarbonSilicon Systems; Part I. Related Binary Systems, Vol. 1. Mo-C Systems”, Air Force Materials Laboratory Report AFML-TR-65–2, 1(l), 1–159 (1965) (Crys. Structure, Phase Diagram, Experimental, 75) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Rudy, E., Windisch, S., Stosick, A.J., Hoffman, J.R., “Constitution of Binary Molybdenum-Carbon Alloys”, Trans. Met. AIME, 239, 1247–1267 (1967) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, 38) Toth, L.E., “High Superconducting Transition Temperatures in the Molybdenum Carbide Family of Compounds”, J. Less-Common Met., 13, 127–129 (1967) (Crys. Structure, Experimental, Electr. Prop., 16) Smith, G.S., Tharp, A.G., Johnson, Q., “Determination of the Light-Atom Positions in Mo2BC”, Acta Crystallogr. B, 25B, 698–701 (1969) (Crys. Structure, Experimental, 12) Salibekov, S.E., Levinskii, Yu.V., Lobankov, V.U., “Synthesis and Properties of Molybdenum Borocarbide”, Inorg. Mater., 6(9), 1430–1432 (1970), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 6(9), 1622–1624 (1970) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Phys. Prop., *, 6) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Bovin, J.-O., O’Keeffe, M., Stenberg, L., “Planar Defects in Mo2BC. An Electron Microscope Study”, J. Solid State Chem., 22, 221–231 (1977) (Crys. Structure, Experimental, 11) Lundberg, L.B., Turner, W.C., Hoffman, C.G., “Brazing Molybdenum and Tungsten for High Temperature Service”, Welding J., 57(10), 311s-318s (1978) (Morphology, Experimental, Phys. Prop., 19) Holleck, H., “Molybdenum as a Possible Substitute for Tungsten in Hard Materials” (in German), Metall, 10, 1064–1069 (1979) (Phase Diagram, Phase Relations, Experimental, Review, Mechan. Prop., 27) Morozumi, S., Kikuchi, M., Sugai, S., Hayashi, M., “Reaction Between Molybdenum and Carbon and Several Carbides” (in Japanese), Nippon Kinzoku Gakkaishi, 44(12), 1402–1413 (1980) (Phase Relations, Experimental, 6) Hiraoka, Yu., Fukutomi, M., Okada, M., “Improvement of Ductility of Molybdenum Sheet by Forming a Surface Layer Containing Boron”, J. Nucl. Mater., 99, 327–330 (1981) (Morphology, Experimental, Phys. Prop., 13)
DOI: 10.1007/978-3-540-88053-0_18 ß Springer 2009
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B–C–Mo [1981Lej]
[1981Sch] [1982Kra]
[1983Sch]
[1984Hol]
[1988Epi]
[1988Vel]
[1990Ase]
[1990Zwi]
[1991Kus]
[1992Rog]
[1993Wer]
[1994Gos]
[1994McH] [1995Vil]
[1996Kas] [1996Mal]
[1997Ath]
18
Lejay, P., Chevalier, B., Etourneau, J., Hagenmuller, P., “The Borocarbides Mo2–xWxBC (0 ≤ x ≤ 1.1), a New Family of Refractory Superconducting Materials”, Synthetic Mater., (4), 139–145 (1981) (Crys. Structure, Experimental, Magn. Prop., 19) Schulze, K., Kim, H.-J., Jehn, H., “Influence of Nb-additions to the Solubility Limit of Carbon in Molybdenum”, Z. Metallkd., 72, 490–494 (1981) (Phase Relations, Experimental, 14) Krauss, W., Politis, C., “Superconductivity of Metastable Single Phase δMoC1–x” in “Superconductivity in d- and f-Band Metals”, Buckel, W., Weber, W., (Eds.), Kernforschungsinstitut Karlsruhe, 4 (1982) (Crys. Structure, Experimental, 4) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Phase Diagram, Phase Relations, Review, Mechan. Prop., 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundliche - Technische Reihe, Gebru¨der Borntra¨ger, Berlin, 6, 260–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Epicier, T., Dubois, J., Esnouf, C., Fantozzi, G., Convert, P., “Neutron Powder Diffraction Studies of Transition Metal Hemicarbides M2C1–x -II. In Situ High Temperature Study on W2C1–x and Mo2C1–x”, Acta Met., 36, 1903–1921 (1988) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, 33) Velikanova, T.Y., Kubliy, V.Z., Khaenko, B.V., “Transformation in Solid State and Phase Equilibria in the Mo-C System” (in Russian), Poroshk. Metall., (11), 61–67 (1988) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, 11) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides”, Freer R., (Ed.), “The Physics and Chemistry of Carbides, Nitrides and Borides”, Kluwer Academic Publishers, Dordrecht, 97–111 (1990) (Crys. Structure, Experimental, Review, 14) Zwicker, A.P., Regan, S.P., Finkenthal, M., Moos, H.W., Saloman, E.B., Watts, R., Roberts, J.R., “Peak Reflectivity Measurements of W/C, Mo/Si, and Mo/B4C Multilayer Mirrors in the 8–190 C Range Using Both Kα Line and Synchrotron Radiation”, Appl. Optics, 29(25), 3694–3698 (1990) (Crys. Structure, Experimental, Phys. Prop., 12) Kushkhov, Kh.B., Malyshev, V.V., Mazur, L.L., Shapoval, V.I., “Deposition of Molybdenum Carbide on Surface of Electroconducting Refractory Carbides by Electrolysis of Ionic Melts” (in Russian), Poroshk. Metall., (5), 49–52 (1991) (Morphology, Experimental, Kinetics, Phys. Prop., 6) Rogl, P., “The System B-N-Mo” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), Ohio, USA: ASM, Materials Park, 64–67 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Experimental, Review, *, 10) Werheit, H., Kuhlmann, U., Laux, M., Lundstroem, T., “Structural and Electronic Properties of Carbon-doped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Gosset, D., Decroix, G.M., Kryger, B., “Improvement of Thermo-Mechanical Properties of Boron-Rich Compounds”, Proc. 11th Int. Symp. Boron, Borides and Related Compounds, Tsukuba, 1993, Jpn. J. Appl. Phys., Series 10, 216–219 (1994) (Crys. Structure, Experimental, Phys. Prop., 6) McHale, A.E., “VI. Boron Plus Carbon Plus Metal”, Phase Equilibria Diagrams, Phase Diagrams for Ceramists, 10, 189–190 (1994) (Phase Diagram, Phase Relations, Review, 2) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, ASM International, Materials Park, Ohio, USA, Vol. 5, 5353–5357 (1995) (Phase Diagram, Crys. Structure, Review, 3) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., *, 170) Malyshev, V.V., Novoselova, I.A., Gab, A.I., Pisanenko, A.D., Shapoval, V.I., “High Temperature Electrochemical Synthesis of Molybdenum Carbide on the Surface of Dielectrics and Semiconductors in Ionic Melts” (in Russian), Zh. Prikl. Khim., 69(8), 1314–1320 (1996) (Morphology, Experimental, Kinetics, Phys. Prop., 16) Athanasiou, N.S., “Structural Instability and Superconductivity of the Defect Cubic Structure δMoC1–x”, Mod. Phys. Lett. B, 11(21-22), 939–947 (1997) (Crys. Structure, Phase Relations, Experimental, Electr. Prop., 16)
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22
18 [1998Kas]
[1998Rog]
[1999Hug]
[2002Kor]
[2003Lei] [2004Pad]
[2005Tan]
[Mas2] [V-C2]
B–C–Mo Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Calculation, 0) Rogl, P., “Mo-B-C (Molybdenum - Boron - Carbon)” in “Phase Diagrams of Ternary Metal-BoronCarbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Ohio, USA, 182–195 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 30) Hugosson, H.W., Nordstroem, L., Jansson, U., Johansson, B., Eriksson, O., “Theoretical Studies of Substitutional Impurities in Molybdenum Carbide”, Phys. Rev. B., 60(22), 15123–15130 (1999) (Morphology, Theory, Phys. Prop., 20) Korniyenko, K., Rogl, P., Leithe-Jasper, A., Bohn, M., Seidl, Yu.E., Tanaka, T., Velikanova, T., “Melting Surfaces in Refractory System Mo-B-C”, Research at University of Vienna, Austria, unpublished work, (2002) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, *, 33) Leithe-Jasper, A., Tanaka, T., Rogl, P., Research at NRIM, Tsukuba, Japan, unpublished work, (2003) (Crys. Structure, Experimental) Paderno, V., Paderno, Yu., Filippov, V., Liashchenko, A., “Directional Crystallization of B4C-NbB2 and B4C-MoB2 Eutectic Compositions”, J. Solid State Chem., 117, 523–528 (2004) (Morphology, Phase Relations, Experimental, *, 14) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, research presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), August 21–26, 2005, (Phase Relations, Experimental) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Nitrogen Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Vasyl Tomashik
Introduction The interest in the B-C-N ternary system is caused by attempts to synthesize new dense phases harder than diamond, cBN and B4C [1992Rie, 1994Rie, 1997Sol1, 1998Pop1]. Besides, several types of semiconductors can be expected for these materials, depending on their composition and structure [1989Liu, 1992Saa]. On the other hand, boron carbonitrides could be used as starting materials for the preparation of heterodiamond [1996Nak]. Cubic and hexagonal BC2N are expected to be more thermally and chemically stable than diamond and harder than cBN [2001Mat1]. This possibility makes them the most interesting class of compounds that can replace the expensive diamond in many mechanical applications. No systematic experimental investigations of phase equilibria in the B-C-N system are reported [2002Sei]. Mainly materials of the sections B4C-BN and hBN-graphite were investigated. The results reported by different authors are rather contradictory and do not allow an unambiguous answer to the questions whether the synthesis product is a ternary compound, a solid solution of C in cBN or a mechanical mixture of diamond and cBN. The solubility of B4C in BN at 1900–2800˚C and at pressures above 0.5 GPa was investigated by [1977Sir] and the solubility between graphite and graphitic BN was determined by [1991Ruy]. [1975And] noted that a ternary compound is formed in the B4C-BN system. [1977But] obtained cubic BN containing up to 10 at.% C using high temperatures and high pressures near the triple point of the p-T diagram for BN. The results of [1977Sir] and [1977But] were included in the review of [1994McH]. Boron carbonitride is the phase with a variable composition and could be described by the formula BC1–xN [1975Pri]. According to [1972Bad, 1991Ruy, 1992Bil, 1992Rie, 1994Rie] B-C-N materials could be considered as solid solutions of carbon and BN. Some isothermal sections of the B-C-N ternary system were calculated by [1993Wen] and some tie lines in this ternary system were determined by [1991Ruy]. B-C-N materials have been synthesized by various techniques. They could be obtained by nitriding a mixture of amorphous B and carbon black [1971Kos]. CVD using BCl3 and NH3 as B and N precursors, respectively, and CCl4 [1981Bad], CH4 [1989Moo, 1990Bes, 1990Yam], C2H2 [1987Kan, 1989Moo, 1990Sau, 1991Sau, 1992Sau] or C3H8 [1989Moo] as C precursors, CVD from BCl3 and acetonitrile [1989Kou, 1992Mor, 1993Sas, 1996Wat, 1998Pop2] or acrylonitrile [1993Kaw] and laser-assisted CVD in a gas atmosphere containing C2H4, B2H6 and NH3 [1999Mor] could be also used for the obtaining such materials. The another possibilities to obtain the B-C-N materials are a synthesis by condensation of polyacrylonitrile-based polymer with BCl3 [1993Kaw], pyrolysis of appropriate precursors [1988May1, 1988May2, 1990May, 1991Rie, 1992Bil, 1995Lev, 1998Kal, 2004Kom] and high temperature shock compression of nanodispersed powders of the ternary graphite-like BCxNy phases Landolt‐Bo¨rnstein New Series IV/11E1
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followed by sharp quenching [1994Kak, 1996Kom, 1998Yam, 1999Kom, 2000Kur, 2001Sol1, 2001Sol2, 2001Sol3, 2002Sol]. A new phase with unknown composition was obtained at the interaction of the diamond power and hBN [1993Pai]. Amorphous B-C-N powders were also prepared by mechanical milling or by static and dynamic compression of hBN and graphite as starting materials [1996Lun, 1998Yao, 1999Yao, 1999Zha2, 2001Hua, 2001Ono, 2002Lan, 2002Zha, 2004Hub]. Such materials could be obtained also using the mixtures of the boric acid with melamine in different ratios with the subsequent heating in N2 atmosphere [2000He, 2001He], from a charge based on H3BO3 and CO(NH2)2 with addition of 60 mass% saccharose nitrided at 1300 and 1400˚C for 3 h [2005Gri] or by the help of a solvothermal reaction of CH3CN, BCl3 and Li3N using benzene as a solvent [2004Hua]. These materials had different compositions, but none of them was in the nitrogen rich part of the B-C-N phase diagram [1998Pop1]. The synthesis of dense phases of the B-C-N ternary system is possible only as a result of nonequilibrium processes [1997Sol1]. [1993Sas] indicated that all synthetic strategies to obtain “BCN” graphite (BxCyNx) are unavoidably complicated by the statistical probability of the resulting material containing graphite and hBN beside the desirable B-C-N materials. An associated difficulty is that of distinguishing a genuine B-C-N material from an intimate mixture of graphite and hBN. Their diffraction patterns are substantially superimposable with each other, especially for the materials of low crystallinity. The dense forms of B-C-N system could be obtained from samples with high carbon content, at pressures beyond 10 GPa, preferably employing nonequilibrium processes and beyond the melting points of carbon and BN [2001Ono]. The compression-annealed BN-C materials consist of separated domains of pyrolitic boronated graphite and pyrolitic BN, while as-deposited samples may possibly be a singlephase mixture of B, C and N [1989Moo]. As-deposited materials containing more than 20% C were found to be more highly oriented than unannealed graphite and the crystallinity of these materials was greatly enhanced by an uniaxial compression annealing. The codeposited materials appear to be slowly changing composition with temperature and time. A set of carbonitride phases has been synthesized by nitridation of B4C powder at 2100˚C under various pressures of N2 [1994And]. B-C-N powders could be obtained by simultaneous nitridation of boric acid and carbonization of saccharose in molten urea followed by annealing in N2 atmosphere at 1500˚C [1995Hub, 1997Sol1, 2005Sol] and graphite-like BC3.88N was prepared by a thermal CVD from a BCl3-CH3CN-Ar mixture using a Ni substrate at 952˚C [1997Sol1]. The fraction of compressed turbostratic graphite-like structure decreases with pressure [2005Sol]. Finally, the total disappearance of the initial phase is observed at 25 GPa. This transformation is fully reversible and the formation of covalent bonds between the layers does not occur. The graphitic BC2N was prepared by a vapor-phase reaction of BCl3 and CH3CN [1994Nak] and a carbodiimide based sol-gel synthesis of B4CN4 was reported in [2003Voe]. Amorphous BC2N powders have been prepared by high energy milling of a mixture of graphite and hBN and then spark plasma sintering method at 1600˚C was used to recrystallize the amorphous precursor at high temperature [2006Luo]. hBC2N could be prepared by the interaction of CCl4, BBr3, Li3N and sodium at 400˚C [2007Sun]. [1996Noz1, 1996Noz2] indicated that a layered BC2N material has a huge number of polymorphic structures, which depend on the atomic arrangements. It has been found that the vibrational modes of BC2N in the higher energy region are sensitively influenced by the intralayer atomic arrangements. According to their calculations the stable structure of BC2N DOI: 10.1007/978-3-540-88053-0_19 ß Springer 2009
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is formed so as to increase the number of both C-C and B-N bonds [1996Noz1, 1996Noz2, 2006Aze1]. The multiple-scattering approach was applied by [1999Wib] to predict electronenergy-loss near-edge structures of BC2N. This material is shown to be energetically more favorable when segregating into thicker diamond and cBN layers [1999Zha1]. The electronic structure of layered materials based on graphite and hBN was calculated via extended-Hueckel crystal orbital band calculations [1990LaF], using ab initio quantum methods [1992Saa], by the help of a non-self-consistent, ab initio based tight-binding moleculardynamics method [1998Wid] and using an ab initio pseudopotential density functional method [2004He, 2006Sun, 2007Wu]. [1994Nak] indicated that not a cubic BC2N ternary compound but a mixture of cBN and diamond exists as thermodynamically stable phases in the B-C-N ternary system at 2000–2400˚C and up to 7.7 GPa. According to the data of [2000Kur] the resulting diamondlike phases are supersaturated BN-C solid solutions having carbon concentrations of about 30 at.%. Laser-heating experiments at different pressures have shown that the formation of cBC2N is observed only above 18 GPa [2004Sol]. At 14.5 GPa and 1730˚C gBC2N decomposes to form a mixture of cBN and diamond. On further decrease in pressure down to 11.0 GPa, thermal decomposition of gBC2N proceeds to form cBN and disordered graphite. [1992Bow] noted that according to the calculation B24C12N24 should be stable as a hybrid analogue of buckminsterfullerene. According to ab initio calculation by [2001Sun] all the cubic BC2N structures are metastable and tend to separate into diamond and cBN. It is not clear if the cBN-diamond system is a true solid solution in the thermodynamic sense [1993Lam]. The BC3N3 ternary compound was obtained by [2000Wil]. According to the data of [1994Nak] it is difficult to distinguish ternary compounds from a microcrystalline mixture of graphite and hBN. [1997Sol1] noted that neither solid solutions of C in cBN nor dense ternary compounds of the B-C-N system were revealed. Details of the experimental studies are reported in Table 1.
Binary Systems The binary systems B-C and C-N are consistent with the critical assessments of [1996Kas]. The B-N system is accepted from [2008Rec]. Only one intermediate phase BN exists in this system. Boron nitride has four crystalline structural modifications: cubic (cBN), wurtzite (wBN), hexagonal (hBN) and rhombohedral (rBN). In addition, there are two other ordered BN phases: E-BN, obtained by explosion (E) of a mixture of hBN and aBN, compressed hBN attributable to a monoclinic lattice distortion of hBN. Additionally two disordered BN phases were reported: turbostratic BN (tBN) and amorphous BN (aBN).
Solid Phases The solid phases with crystallographic data, including the metastable phases, are listed in Table 2. The evidence for the formation of the ternary compounds in the B-C-N system is hardly exhaustive and is open to doubt [1996Lun]. It is evident that substantial difficulties are encountered in the identification and characterization of phases consisting of light elements Landolt‐Bo¨rnstein New Series IV/11E1
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with nearly the same atomic numbers arranged in similar honeycomb networks. Thus it is an arduous task to distinguish between a substantional ternary network and a nanometric-scale mixture of graphite and hBN both turbostratically distorted. Samples prepared in different ways contain phases that differ in structure, which is in agreement with the fact that they are highly defective, non-equilibrium phases. The results that have been reported by different authors are rather contradictory, and currently, there is no unambiguous answer to the question of whether the synthesis products are ternary compounds, solid solutions of carbon in cBN, or mechanical mixture of diamond and cBN [1997Sol2]. Thermodynamically stable ternary phases do not exist in the B-C-N equilibrium phase diagram [2001He, 2006Aze2, 2006Pan] (just [2001Nic] indicated that graphite-like BC2N phase is thermodynamically stable). However, it is anticipated that some metastable phases including solid solutions and compounds should exist under definite extreme conditions. According to [2005Gri] the ternary compounds in the B-C-N system are thermally unstable and at the sintering temperature of 1900˚C are converted into BN and B4C phases. Lattice parameter measurements indicate a slight solubility of B4C in BN, but no solubility of BN in B4C for samples prepared at 2250˚C [1992Ruh] (the solubility of B4C in BN is equal to 7-20 mol% [1977Sir]). No ternary compounds were found in the B4C-BN system [1992Ruh]. No mutual solubility between graphite and cBN was found by [2001Hua]. [2002Gag] indicated that the solubility limit of graphite in the hBN structure is 15 at.%. Under equilibrium conditions the BN solubility in diamond is quite limited [2002Lan]. BC3N3 (B(CN)3) was isolated as a Lewis acid by [2000Wil]. This compound is thermally stable up to about 450˚C. The thermal dissociation via elimination of (CN)2 is likely to yield a stoichiometric material with BCN composition. gB4CN4 is metastable and does not change upon long-term annealing at 1000–1200˚C and annealing at 2000˚C furnishes nearly pure B4C [2003Voe]. According to the data of [1975Pri] the boron carbonitride is a turbostratic strongly deformed structure of BN with the graphite introduced in the lattice. High temperatures and high pressures lead to a decomposition of this structure. The formation of ternary compounds in the B-C-N system was never found in the experiment of [2001Hua].
Invariant Equilibria The complete reaction scheme (p = 100 kPa) calculated using CALPHAD method with some amendments to ensure consistency with the B-N binary system from [2008Rec] is presented in Fig. 1 [2002Sei]. One maximum where boron carbide is formed from gas and liquid was found by calculation. Two transition reactions and two degenerated reaction were found. The calculated temperature of the U2 reaction is not well established experimentally. In situ investigations of the reaction of BN and graphite are not reported and therefore only temperatures indirectly derived from X-ray analysis of quenched materials are documented. [2002Sei] assumed that the temperature is lower than that given by [1995Fil] (>2450˚C) and higher that the one given by [1972Bad] (2200˚C) or [1992Bil] (2200).
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Isothermal Sections Some isothermal sections have been calculated using the Program Thermo-Calc and including the formation of the BCN compound [1993Wen]. These sections are shown in (Figs. 2–6) without the BCN compound which was not taken into account by [2002Sei]. According to the data of [1991Ruy] BN-C join is stable to 1900˚C and above this temperature there is a B4C-N tie line. Geometrically, a B4C-BN tie line seems likely.
Temperature – Composition Sections The section between graphite and BN, including the invariant reaction U1, D1 and U2 (Fig. 1), is shown in Fig. 7 after [1996Kas].
Potential Diagrams A thermodynamic dataset was developed by [1998Sei] for the B-C-N system and potential phase diagrams showing the phase equilibria as a function of nitrogen pressure and temperature were calculated. Figure 8 shows the phase relations for a B:C < 4 (atomic ratio) and Fig. 9 shows the phase relations for B:C > 4. With the B:C ratios > 4 it is possible to stabilize BN-B4C composites.
Notes on Materials Properties and Applications Boron carbonitride ceramics can be widely used in high temperature engineering as such products show high thermal shock resistance and good refractory properties [1986Dub]. The unique physical and technological properties of BCxNy materials recommend their use in very different fields of modern industry as a refractory electroinsulating ceramic in the form of bath-linings, tubes for transporting molten metals, alloys, slags and salts, covers for thermocouples, coating plates and insulators for temperatures up to 2000–2500˚C [1986Dub], lightemitting and thermoelectric materials, transistors working at high temperature, lightweight electrical conductors, materials for molecular sieves and deceleration materials for nuclear reactor, catalysts, high-temperature lubricant [1997Kaw] as well as a negative electrode matrix for rechargeable Li batteries [1992Mor, 1997Kaw]. Boron carbonitride amorphous coatings were synthesized by the pulsed laser deposition [1997Hou]. These coatings were found to be adherent on a variety of substrates with Si interlayer, which have varied effect on adhesion depending on its thickness. The incorporation of nitrogen makes B4C smoother and despite the fact that the friction coefficients are of the same order of magnitude, the wear is more important in BCxNy films than in B4C films [2001Fre]. Thin BCxNy films could be obtained with high hardness and excellent tribological behavior [2001Mar]. Hardness is higher for films with lower nitrogen content but friction coefficient decreases with increasing nitrogen concentration and has very low values when nitrogen content reaches 20 at.%. By adjusting the nitrogen content in the films, coatings with tailored tribological properties in ranges of great technological interest could be designed [2001Mar]. Thermal diffusivity of the amorphous BCxNy films decreases Landolt‐Bo¨rnstein New Series IV/11E1
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with increasing carbon content above 30 at.% and a clear increase in thermal diffusivity with film thickness has been obtained [2002Cha]. Thermal diffusivity varied from 0.55 to 0.7 cm2·s–1 as the thickness of the films increased from 200 to 750 nm. These films adhere strongly to the dielectric films and the critical debond energy GC exceeds 10 J·m–2 [2005Eng]. The films with compositions BC0.37N0.15, BC0.11N0.49 and BC0.92N0.07 are characterized by the dielectric constants of 4.1, 4.2 and 3.8, respectively. [2000Gag] noted that BCxNy films have been grown by ion beam assisted evaporation of graphite and B4C targets. cBC2N may be the first of a new family of ternary B-C-N superhard materials with hardness and elastic modulus higher than those of cBN [2001Sol1, 2001Sol2, 2001Sol3, 2002Sol]. According to [2002Zha] a nominal hardness of BC2N and BC4N is 62 and 68 GPa, respectively. Young’s modulus and modulus of shearing of cBC2N are equal to 980 ± 40 GPa and 447 ± 18 GPa, respectively [2001Sol3]. [2001Mat2] noted that according to the calculations two orthorhombic BC2N crystals are characterized by the bulk moduli 459.41 and 408.95 GPa. BC2N is a graphite-like material and a small bandgap semiconductor (Eg = 0.03eV) [1989Kou]. According to [1998Yao, 1999Yao] this material is a semiconductor with band gap energy of 0.11 eV for temperatures ranging from room temperature to 290˚C and a semimetal for temperatures between 290 and 470˚C. A chalcopyrite BC2N is a wide gap semiconductor with a direct band gap of about 3.3 eV and B0 = 364.7 GPa [2006Sun]. According to first-principles calculations BCxNy monolayers behave as a semiconductor or metal with band gap energy ranging from 0 to 2.45 eV depending on the atomic arrangement [2006Aze1]. It could be intercalated by S2O6F2, Li, Na and K [1989Kou]. BC2N nanotubes and nanostructures have proved to behave as efficient field elements as well as blue and violet light emitting materials [2002Ter]. BCxNy nanoscales materials (nanotubes and nanofibers) may find important applications in the fabrication of nanotransistors, robust nanocomposites, conducting polymers, storage components and field emission sources. These layered materials could be intercalated with alkali metals, which might be useful in the fabrication of batteries or other nanoelectronic devices. The nanocomposite (BN)0.5C0.5 possessed the highest hardness equalling twice that of nanocrystalline graphite and three times of nanocrystalline BN [1999Zha2]. Zero-pressure bulk modulus of the metastable phase gBC4N is equal to 18.1 ± 0.2 GPa and its pressure derivative 6.6 ± 0.1 [1997Sol1, 1997Sol2]. The first-principal pseudopotential method was employed to predict the mechanical properties of the BC2N phases [2001Mat1]. It was shown that B0 = 459.41 and 420.13 GPa and B0’ = 2.11 and 3.40 for orthorhombic and trigonal BC2N, respectively [2001Mat1] (B0 = 83.3 GPa [1999Zha1]). The calculation zero pressure bulk modulus of the rhombohedral BC2N is found to be 438 ± 14 GPa [1997Tat]. This material has an indirect gap of 3.97 eV. According to the data of [2001Sol1] bulk modulus of cBC2N is 282 GPa (B0 = 412 GPa [2002Sol]). The bulk modulus of tetragonal B2CN is equal to 332.1 GPa [2004He]. The calculated bulk moduli of BN rich solid solutions (BN)x(C2)1–x (x > 0.5) are lower than those of diamond, the value of ideal mixing and even cBN [1999Zhe] (B0 = 355 ± 19 GPa for the (BN)0.67C0.33 composition [1995Kni]). The bulk modulus of BC2.5N, which is stable up to 700˚C, is 401 GPa [1999Kom]. The conductivity of (BC3N)n graphite-like material increased in the range 25–700˚C, implying that it behaves as a semiconductor and thermoelectric measurements indicate that it is p type [1993Kaw]. The activation energy was 6.29·10–3 eV. Electronic-structure calculations allow to predict that hypothetical ordered mixed crystals of cubic BN and diamond show a very pronounced band-gap bowing [1993Lam]. DOI: 10.1007/978-3-540-88053-0_19 ß Springer 2009
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Boron carbonitride vigorously reacts with Nb, Ta, Mo and W forming boride phases [1975Bor]. The reactions begin at that temperature at which incongruent evaporation of this material becomes appreciable. BC3N3 is a promising precursor for high-pressure and high-temperature synthesis of novel phases with structures related to Si3N4, diamond and graphite [2000Wil]. Geometry optimization and studies of the electronic properties were carried out using the pseudo-potential method in the framework of the local density functional theory for the two- and threedimensional structures [2004Bet]. They respectively lead to propose a precursor (2D), a β structure and new ultra hard materials with B0 358 GPa for BC3N3 high-pressure phase. The energy of formation of superlattices C2n(BN)n for n = 1, 3, 5 were calculated to be 0.89, 0.95 and 0.94 ± 0.01 eV/(interface unit-cell area) [1989Lam]. The positive values indicate the thermodynamic instability of these superlattices towards disproportionation. The controllability of tensile stress in BN films could be improved by introducing a small amount of carbon into the BN matrix [1990Yam]. The radiation resistance of deposited films at 400˚C was improved five times better than that of BN deposited by low-pressure CVD at similar temperature.
Miscellaneous The ab initio calculations of the ground-state properties of cubic BN-C solid solutions indicate that such solutions could be non-ideal with the equilibrium lattice constants larger than the predicted values of ideal mixing and the positive energies of formation [1999Zhe]. BC3N3 (B(CN)3) have been prepared from the thermal elimination of NCSiMe3 from B (CN)3NCSiMe3 [2000Wil]. BCxNy films could be deposited by dual cathode sputtering of BN and graphite targets [2002Cha] or in He using a low power density, and in N2 using a high power density [1997Heg]. The plasma-activated CVD parameters sensitively influence the structures and chemical compositions of such films derived from single-source precursors. The film compositions strongly depend on the development of inner stresses, which, in turn, depends on the momenta of incident ions. A wide range of ion momenta can be obtained by varying the carrier gas and the applied power density [1997Heg]. According to the data of [2004Bys] and [2002Gag] such films were deposited by N+ ion-beam-assisted pulsed laser deposition technique on glass substrates at different temperatures up to 600˚C using B4C targets. These films could be also prepared by CVD at 360˚C and 132 Pa using dimethylamine borane with no coreactant, NH3 or C2H4, producing different composition films [2005Eng]. The XPS study suggested that B atoms in the deposited films display in a wide variety of chemical bonds (B-C, B-N and B-C-N) [2006Udd]. The highly oriented graphite-like B-C-N hybrids are stable at low B content and the BCxNy films synthesized at room temperature have random orientation. The thermal stability of such films is characterized by the absence of any mass loss after 1 h at 1700˚C while graphitization and formation of small amounts of B4C are observed [1991Sau]. The C-rich domain exhibits a graphite turbostratic structure with a well-developed twodimensional crystalline order greater than graphitizable pyrocarbons heat-treated at 2000˚C and a coherence length of 20 nm. BCxNy hybrid thin films were grown from ion beam plasma of borazine (B3N3H6) on graphite substrate at room temperature, 600 and 900˚C [2005Udd]. It was found that the Landolt‐Bo¨rnstein New Series IV/11E1
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hybrid formation is enhanced at high temperature. It is possible to control the composition of these hybrids by changing the ion fluence and the temperature during ion implantation. BCxNy nanotubes were also obtained by the plasma rotating electrode process [2005Kim]. It was determined that with rotation of the anode and/or increase in metal concentration, the number and length of these nanotubes increase. Using a first-principles pseudopotential approach with a localized basis for the wave functions three possible atomic arrangements have been investigated for the monolayer of the gBN hybrid, BC2N [1989Liu]. The structure with the lowest total energy has a nearestneighbor environment which seems to optimize the chemical bond energy by maximizing the number of C-C and B-N bonds. Of the three structures considered, the one with the highest cohesive energy lacks inversion symmetry and has a local-density monolayer gap of about 1.6 eV and the one metallic structure has a slightly lower cohesive energy. The crystallite size for synthetic BC2N sample, obtained from hBN and graphite, was in the range of 4–8 nm [2002Zha].
. Table 1 Investigations of the B-C-N Phase Relations, Structures and Thermodynamics Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1971Kos]
XRD, chemical analysis
1800–2000˚C / B4C-BN
[1972Bad]
XRD, chemical analysis
Up to 2000˚C / BN-C
[1975And]
XRD, metallography, resistivity, coefficient of Up to 2000˚C / B4C-BN thermal expansion and frequency dependence of the penetrability and dielectric loss measurements
[1975Bor]
XRD, metallography, microhardness measurements
1500–1900˚C / 53 mass%B + 9 mass% C + 36 mass% N
[1975Pri]
XRD, chemical analysis, IR spectrometry
1400–2000˚C / BC1–xN
[1977But]
XRD, metallography, microhardness measurements, X-ray microanalysis
BN-C
[1977Sir]
XRD, metallography, microhardness measurements
1900–2800˚C and more than 0.5 GPa / B4C-BN
[1981Bad]
XRD, TEM, EPMA
Up to 3330˚C / BN-C
[1986Dub]
XRD
Up to 2020˚C / BN-C
[1987Kan]
XRD, TEM, electron diffraction, XPS, IR spectroscopy, chemical analysis, conductivity measurements
400–700˚C / B-C-N
[1988May1] XRD, metallography, NMR, IR and Raman spectroscopy
Up to 400˚C / B-C-N
[1988May2] XRD, DTA, TGA, chemical analysis, IR spectroscopy Up to 800˚C / B-C-N [1989Kou]
XRD, EELS, AES, electron diffraction, electron microscopy
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Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1989Moo]
XRD, TEM, electron diffraction, X-ray photoelectron and Raman spectroscopy, optical reflectivity, thermal conductivity, thermopower and electrical resistivity measurements, NMR
1500–1900 and 2500–2700˚C / BN-C
[1990Bes]
XRD, metallography, EPMA, AES
1380˚C / B4C-BN at 10–50 mol% BN
[1990May]
XRD, XPS, AES, ESCA, IR and Raman spectroscopy, Up to 800˚C / B-C-N electrical conductivity measurements
[1990Sau]
XRD, ESCA
800–1000˚C / BN-C
[1990Yam]
IR spectroscopy, chemical analysis, ellipsometry, stress and Young’s modulus measurements,
Up to 800˚C / B-C-N
[1991Rie]
XRD, TEM, EELS, NMR
1050˚C / B-C-N
[1991Ruy]
XRD, metallography, SEM, DTA, TGA
Up to 1500˚C / B-C-N
[1991Sau]
XRD, XPS, TEM, EPMA, Rutherford backscattering
Up to 2500˚C / B-C-N
[1992Bil]
XRD, TGA, mass spectroscopy, ESCA
Up to 2200˚C / B-C-N
[1992Mor]
XRD, chemical analysis, electrochemical properties Up to 1000˚C / B-C-N measurement
[1992Ruh]
XRD, TEM/STEM, metallography, flexural strength 1900–2250˚C / B4C-BN determination
[1992Sau]
XRD, X-ray microprobe analysis, XPS, Rutherford Up to 2500˚C / B-C-N backscattering, HREM, TEM, density and electrical conductivity measurements
[1993Kaw]
XRD, electron diffraction
Up to 1000˚C / BC3N
[1993Pai]
XRD, SEM, EDS
Up to 3230˚C and 100 GPa / BN-C
[1993Sas]
XRD, SEM, EELS spectroscopy
Up to 1600˚C / B-C-N
[1994And]
XRD, Raman spectroscopy
1900–2300˚C / B4C-N2
[1994Kak]
XRD, TEM
1000–1500˚C / BC2N
[1994Nak]
XRD, metallography, HRSEM, AES
2000–2400˚C and up to 7.7 GPa / B-C-N
[1995Hub]
XRD, SEM, IR spectroscopy, chemical analysis
Up to 2150˚C / BN-C
[1995Kni]
XRD, Raman and IR spectroscopy
1730–2230˚C and 30–50 GPa / BN-C
[1996Kom]
XRD, NMR, HRTEM, EELS spectroscopy, chemical analysis, electron diffractometry
Room temperature and 35 GPa / BN-C
[1995Lev]
XRD, IR spectroscopy, index of refraction, hardness Up to 800˚C / B-C-N and Young’s modulus measurements
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1996Lun]
XRD, NMR, Raman spectroscopy
2100–2300˚C / B-C-N
[1996Nak]
XRD, HRSEM, AES
Up to 2400˚C and 7.7 GPa / BN-C
[1996Wat]
XRD, XPS, SIMS
850˚C / BC2N
[1997Heg]
XRD, TEM, AES, ESCA, densitometry, microindentation, FTIR, AFM
Room temperature / B-C-N
[1997Hou]
XRD, EDXA, profilometry, microhardness measurements
Room temperature / B-C-N
[1997Sol1, 1997Sol2]
High-pressure high-temperature XRD, chemical analysis
Up to 1830˚C and up to 7 GPa / BC4N and BC3.88N compositions
[1998Kal]
XRD, IR and Raman spectroscopy, SEM, AES, energy dispersive X-ray analysis
2000–2200˚C / B-C-N
[1998Pop2] XRD, SEM, XPS, FTIR, SIMS, Raman spectroscopy
800–1050˚C / B-C-N
[1998Yam]
TEM, EELS, electron diffraction
B-C-N
[1998Yao]
XRD, TEM, IR spectroscopy, conductivity measurements
Up to 1200˚C / B-C-N
[1999Kom]
XRD, EPMA, NMR, TEM, TGA, IR spectroscopy, chemical analysis
Up to 3230˚C / B-C-N
[1999Mor]
XRD, EPMA, XPS, SEM, metallography
B-C-N
[1999Yao]
XRD, electron diffraction, IR spectroscopy, chemical analysis
Up to 900˚C / B-C-N
[1999Zha2] XRD, IR and Raman spectroscopy, microhardness measurements
Up to 1200˚C / B-C-N
[2000Gag]
XANES spectroscopy
Up to 150˚C / B-C-N
[2000He]
XRD, HRTEM with EDX
Up to 1600˚C / B-C-N
[2000Kur]
XRD, EPMA, TEM, chemical analysis
Up to 3230˚C and 20 or 30 GPa / B-C-N
[2000Wil]
XRD, NMR, IR and EELS spectroscopy
Up to 450˚C / BC3N3
[2001Fre]
TEM, ion beam analysis, IR spectroscopy, AFM, Raman scattering measurements
Room temperature / B4C + N2
[2001He]
XRD, XPS, IR spectroscopy
Up to 1600˚C and 5.5 GPa / B0.54C0.28N0.30
[2001Hua]
XRD, HRTEM, EELS, electron diffraction, inductive coupled plasma emission spectroscopy
Up to 2300˚C and 7.7 GPa / B-C-N
[2001Mar]
SEM, AFM, SIMS, microhardness and Young’s modulus measurements
Room temperature / B-C-N
[2001Nic]
XRD, TEM, SEM, EELS, EDX analysis
1300–1500˚C / BC2N
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[2001Ono]
XRD, chemical analysis, X-ray emission spectroscopy, hardness measurement
1200–1900˚C and 6–16 GPa / B-C-N
[2001Sol1]
XRD, EPMA, hardness measurement
Up to 1830˚C / BC2N and BC4N
[2001Sol2, 2001Sol3]
XRD, TEM, metallography, micro- and nanoindentation
Up to 1830˚C / BC2N
[2002Cha]
XPS,SEM, ellipsometry, thermal diffusivity measurement
Room temperature / B-C-N
[2002Gag]
HRTEM, X-ray absorption near-edge spectroscopy, Room temperature / B-C-N Raman and IR spectroscopy
[2002Lan]
XRD, HRTEM, EELS
Up to 3230˚C and 30 GPa / B-C-N
[2002Sol]
XRD, EPMA, HRTEM, micro- and nanoindentation
Up to 3230˚C and 30 GPa / (BN)xC1–x at 0.33 ≤ x ≤ 0.61)
[2002Zha]
Synchrotron XRD, HRTEM, SEM, Raman and EELS spectroscopy, hardness measurements
Up to 1830˚C and 20 GPa / BC2N and BC4N
[2003Voe]
IR and Raman spectroscopy, NMR, XRD, SEM, TEM, Up to 2000˚C / B4CN4 chemical analysis, thermal gravimetry with mass spectroscopy, EDX
[2004Bys]
TEM, HREM, EELS
Up to 600˚C / B-C-N
[2004Hua]
XRD, XPS, IR spectroscopy, transmission electron diffraction
300˚C and up to 7 MPa / B-C-N
[2004Hub]
Raman spectroscopy
1830˚C and 25 GPa / B-C-N
[2004Kom]
XRD, XPS, chemical analysis, IR spectroscopy
1500˚C / B-C-N
[2004Sol]
XRD with synchrotron radiation
2730˚C and 30 GPa / BN-C
[2005Eng]
XPS, dielectric constant evaluation
360˚C / B-C-N
[2005Gri]
XRD, chemical analysis, IR spectroscopy
Up to 1900˚C / B-C-N
[2005Kim]
SEM, XRD, HRTEM
6130–6930˚C / B-C-N
[2005Sol]
XRD
Room temperature and up to 30 GPa / BN-C
[2005Udd]
XPS, quadruple mass spectroscopy
Up to 850˚C / B-C-N
[2006Luo]
XRD, TEM, EELS
Up to 1900˚C / BC2N
[2006Udd]
XPS, NEXAFS
Up to 800˚C / B-C-N
[2007Sun]
XRD, FTIR, TEM, EELS, energy dispersive spectrometry
400˚C / BC2N
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. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βB) < 2092
hR333 R 3m βB
a = 1093.02 c = 2381.66
pure B (99.9999%) [1976Lun]
(γN) < –253
tP4 P42/mnm γN
a = 395.7 c = 510.9
at 3.3 GPa [Mas2]
(βN) hP4 –210.0042 - (–237.54) P63/mmc βN
a = 405.0 c = 660.4
[Mas2] triple point
(αN) < –237.54
cP8 Pa3 αN
a = 566.1
[Mas2]
(C)d
cF8 Fd 3m C (diamond)
a = 356.69
at 25˚C, 60 GPa [Mas2]
(C)gr < 3827
hP4 P63/mmc C (graphite)
a = 246.12 c = 670.90
at 25˚C [Mas2] sublimation point
’B4C’
hR45 R 3m B13C2
a = 560.12 ± 0.04 c = 1207.34 ± 0.13
[V-C2, 1996Kas, 1998Kas]
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5
hBN < 2397
hP4 P63mc BN
a = 250.4 c = 666.1
[2008Rec]
cBN
cF8 F 43m ZnS
a = 361.53 ± 0.04
[2008Rec]
wBN
hP4 P63/mmc ZnS
a = 255.0 ± 0.5 c = 423 ± 1
[2008Rec]
rBN
hR6
a = 250.4 c = 999.1
[2008Rec]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C]
Lattice Parameters [pm]
Comments/References
Compressed hBN
mC4 C2/c or Cc
a = 433 b = 250 c = 310 to 330 β = 92 to 95˚
[2008Rec]
B3CN3 or B(CN)3 < 450
t**
a = 821.4 c = 1383.2
[2000Wil]
(BN)xC1–x unstable phase
cF* Fd 3m
a = 357.2 ± 1.0 a = 357.6 ± 0.2 a = 358.2 ± 1.0 a = 358.2 ± 0.2 a = 364.2 ± 0.2
x = 0.11, at 30 GPa [2002Lan] x = 0.2, at 8–10 GPa [2001Ono] x = 0.25, at 30 GPa [2002Lan] x = 0.26 [1981Bad] x = 0.33, [2001Sol1, 2001Sol2, 2001Sol3, 2002Sol] x = 0.33, [2002Zha] x = 0.33, [1994Kak, 1999Kom] x = 0.33, [2002Lan] x = 0.33, [2002Lan] x = 0.33, at 15 ± 2.1 GPa [1995Kni] x = 0.33, at 30.5 ± 3.5 GPa [1995Kni] x = 0.33, at 65.6 ± 1.4 GPa [1995Kni] x = 0.33, at 95.5 ± 3.2 GPa [1995Kni] x = 0.45, at 30 GPa [2002Sol] x = 0.5, at 35 ± 4 GPa [1995Kni] x = 0.53, at 30 GPa [2002Sol] x = 0.6, at 30–40 GPa [1995Kni]
a = 359.5 ± 0.07 a = 360.5 ± 0.1 a = 360.0 ± 1.0 a = 359.3 ± 0.4 a = 355.1 ± 0.3 a = 349.5 ± 0.6 a = 345.7 ± 0.6 a = 338.5 ± 0.6 a = 359.8 ± 0.4 a = 360.2 ± 0.3 a = 360.4 ± 0.4 a = 359.6 ± 0.3 (BN)xC1–x unstable phase
h**
gBC4N unstable phase gBC3.88N unstable phase
h**
BC4N unstable phase
c** ? ZnS
a = 358.6 ± 0.9
[2002Zha]
BCN unstable phase
-
a = 360.6 ± 0.5 a = 359.6 ± 0.5
[2000Kur] [2000Kur]
Landolt‐Bo¨rnstein New Series IV/11E1
a = 246.2 ± 0.2 c = 679 ± 1 a = 247.5 ± 0.1 c = 676 ± 2
x = 0.26 [1981Bad]
a = 247.2 ± 0.2 c = 1090 ± 20 a = 248 ± 3 c = 1061 ± 10
[1997Sol1, 2001Sol1, 1997Sol2]
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x = 0.33 [1998Yam]
[1997Sol1]
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B–C–N
. Table 2 (continued) Phase/ Temperature Range [˚C] BC2N unstable phase
BC2N unstable phase
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
oP32 a = 355.36 P2221 b = 359.86 Ni3{AsO4}2{H2O} c = 355.28 a = 357.9 b = 357.9 c = 361.2 α = 90˚ β = 90˚ γ = 89.32˚ oP16 Pmm2
Comments/References calculation [2001Mat1, 2001Mat2]
[1999Zha1]
a = 252.80 b = 250.24 c = 358.71 a = 253.6 b = 251.0 c = 360.5
calculation [2001Mat2]
calculation [2006Sun]
BC2N unstable phase
oI16 Imm2
a = 254.0 b = 624.2 c = 451.1
calculation [2007Wu]
BC2N unstable phase
tP* P 3m1
a = 249.55 c = 419.23
[2001Mat1]
BC2N unstable phase
hI* I 42d
a = 361.3 c = 714.7
calculation [2006Sun]
BC2N unstable phase
hP* P 4m2
a = 253.7 c = 363.2
calculation [2006Sun]
BC2N unstable phase
h**
a = 243 c = 664
[2007Sun]
gBC2N unstable phase
oC* Cmcm
a = 247.2 ± 0.2 c = 727 a = 248.20 ± 0.07 c = 662. 0 ± 0.3
[2001Sol1]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C] B2CN unstable phase
o*
tP* P 42m B0.82C0.18N unstable phase
Landolt‐Bo¨rnstein New Series IV/11E1
-
Lattice Parameters [pm]
Comments/References
a = 477.6 b = 458.5 c = 362.9 a = 354.4 b = 354.4 c = 389.3
[2001He]
a = 251 c = 666
[2000He]
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[2004He]
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B–C–N
. Fig. 1 B-C-N. Reaction scheme
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. Fig. 2 B-C-N. Isothermal section at 4030˚C
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B–C–N
. Fig. 3 B-C-N. Isothermal section at 3630˚C
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. Fig. 4 B-C-N. Isothermal section at 3130˚C
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B–C–N
. Fig. 5 B-C-N. Isothermal section at 2130˚C
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. Fig. 6 B-C-N. Isothermal section at 1530˚C
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B–C–N
. Fig. 7 B-C-N. Temperature - composition section BN-C
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. Fig. 8 B-C-N. ln(pN2)-T diagrams in the B-C-N system.B: C<4
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B–C–N
. Fig. 9 B-C-N. ln(pN2)-T diagrams in the B-C-N system.B: C>4
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References [1971Kos]
[1972Bad]
[1975And] [1975Bor]
[1975Pri]
[1976Lun] [1977But]
[1977Sir]
[1981Bad] [1986Dub] [1987Kan] [1988May1] [1988May2] [1989Kou]
[1989Lam]
[1989Liu] [1989Moo]
[1990Ase]
[1990Bes] [1990LaF]
Kosolapova, T.Ya., Makarenko, G.N., Serebryakova, T.I., Prilutskiy, E.V., Khorpyakov, O.T., Chernysheva, O.I., “About Nature of Boron Carbonitride. I. Investigation of Conditions for Obtaining of Boron Carbonitride” (in Russian), Poroshk. Metal. 1(97), 27–33 (1971) (Experimental, Phase Relations, Crys. Structure, 4) Badyan, A., Nemyski, T., Appenkheymer, S., Ol’kusnik, E., “Crystal Strycture in The System BoronCarbon-Nitrogen” (in Russian), in “Khim. Svyaz’ v Poluprovodn. I Polumetallakh.”, Nauka I Tekhn. Publish., Minsk, 362–365 (1972) (Experimental, Phase Relations, 4) Andreeva, T.V., Dubovik, T.V., “Nature of Boron Carbonitride” in “Konfigurats. Lokaliz. Elektronov v Tverd. Tele”, Naukova Dumka, Kiev, 222–227 (1975) (Experimental, Phase Relations, 10) Borisova, A.L., Martsenyuk, I.S., “Reaction of Boron and Aluminum Nitrides, and Materials Based on them, with Refractory Metals”, Powder Metall. Met. Ceram., 14(10), 822–826 (1975), translated from Poroshk. Metal., 10(154), 51–56 (1975) (Experimental, Morphology, 10) Prilutskii, E.V., Makarenko, G.N., Serebryakova, T.I., “Structure and Properties of the Boron Carbonitride” (in Russian), Vysokotemperatur. Karbidy, Naukova Dumka, Kyev, 84–89 (1975) (Experimental, Crys. Structure, Phase Relations, 2) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Experimental, Crys. Structure, 10) Butylenko, A.K., Samsonov, G.V., Timofeeva, I.I., Makarenko, G.N., “Doping of Cubic Boron Nitride with Carbon” (in Russian), Pisma Zhur. Tekh. Fiz., 3(4), 186–188 (1977) (Experimental, Phase Relations, 10) Sirota, N.N., Zhuk, M.M., Mazurenko, A.M., Olekhnovich, A.I., “About Solubility of Boron Carbide in Cubic Boron Nitride” (in Russian), Vesti Akad. Navuk Belarus. SSR, Ser. Fiz. Mat. Navuk, (2), 111–112 (1977) (Experimental, Phase Relations, 3) Badzian, A.R., “Cubic Boron Nitride–Diamond Mixed Crystals”, Mater. Res. Bull., 16(11), 1385–1393 (1981) (Experimental, Crys. Structure, Morphology, 12) Dubovick, T.V., Andreeva, T.V., “High Temperature Boron Carbonitride Ceramics”, J. Less-Common Met., 117(1-2), 265–269 (1986) (Experimental, Crys. Structure, Phys. Prop., 9) Kaner, R.B., Kouvetakis, J., Warble, C.E., Sattler, M.L., Barlett, N., “Boron-Carbon-Nitrogen Materials of Graphite-Like Structure”, Mater. Res. Bull., 22(3), 399–404 (1987) (Experimental, Phase Relations, 6) Maya, L., “Semiconducting Amorphous Films Containing Carbon Nitrogen and Boron”, J. Electrochem. Soc., 135(5), 1278–1281 (1988) (Experimental, Phase Relations, 18) Maya, L., “Aminoborane Polymers as Precursors of C-N-B Ceramic Materials”, J. Am. Ceram. Soc., 71 (12), 1104–1107 (1988) (Experimental, Phase Relations, 12) Kouvetikas, J., Sasaki, T., Shen, C., Hagiwara, R., Lerner, M., Krishnan, K.M., Bartlett, N., “Novel Aspects of Graphite Intercalation by Fluorine and Fluorides and New B/C, C/N and B/C/N Materials Based on the Graphite Network”, Synth. Metals, 34(1-3), 1–7 (1989) (Experimental, Kinetics, Phase Relations, 12) Lambrecht, W.R.L., Segall, B., “Electronic Structure of (Diamond C)/(Sphalerite BN) (110) Interfaces and Supperlattices”, Phys. Rev. B, 40(14), 9909–9919 (1989) (Calculation, Phase Relations, Electronic Structure, 34) Liu, A.Y., Wentcovitch, R.M., Cohen, M.L., “Atomic Arrangement and Electronic Structure of BC2N”, Phys. Rev. B, 39(3), 1760–1765 (1989) (Calculation, Crys. Structure, Electronic Structure, 17) Moore, A.W., Strong, S.L., Doll, G.L., Dresselhaus, M.S., Spain, I.L., Bowers, C.W., Issi, J.P., “Properties and Characterization of Codeposited Boron Nitride and Carbon Materials”, J. Appl. Phys., 65(12), 5109–5118 (1989) (Experimental, Phase Relations, 28) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Besmann, T.M., “Chemical Vapor Deposition in the Boron-Carbon-Nitrogen System”, J. Am. Ceram. Soc., 73(8), 2498–2501 (1990) (Calculation, Experimental, Crys. Structure, Morphology, 7) La Femina, J.P., “Electronic Band Structure of Graphite-Boron Nitride Alloys”, J. Phys. Chem., 94, 4346–4351 (1990) (Calculation, Electronic Structure, 8)
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26
19 [1990May] [1990Sau]
[1990Yam]
[1991Rie] [1991Ruy] [1991Sau]
[1992Bil] [1992Bow] [1992Mor]
[1992Rie] [1992Ruh]
[1992Saa]
[1992Sau]
[1993Kaw]
[1993Lam]
[1993Pai]
[1993Sas]
[1993Wen] [1994And] [1994Kak]
[1994McH]
B–C–N Maya, L., Harris, L.A., “Pyrolitic Deposition of Carbon Films Containing Nitrogen and/or Boron”, J. Am. Ceram. Soc., 73(7), 1912–1916 (1990) (Experimental, Phase Relations, 15) Saugnac, F., Marchand, A., “On the Mechanism of Chemical Vapor Deposition of Carbon-BoronNitrogen Compounds Between 800 and 1000˚C” (in French), Compt. Rend. Acad. Sci. Paris, Ser. 2, 310 (Ser. II), 187–192 (1990) (Experimental, Phase Relations, 5) Yamada, M., Nakaishi, M., Sugishima, K., “Improvements of Stress Controllability and Radiation Resistance by Adding Carbon to Boron Nitride”, J. Electrochem. Soc., 137(7), 2242–2246 (1990) (Experimental, Mechan. Prop., 13) Riedel, R., Bill, J., Passing, G., “A Novel Carbon Material Derived from Pyridine-Borane”, Adv. Mater., 3(11), 551–552 (1991) (Experimental, Phase Relations, 14) Ruys, A.J., Sorrell, C.C., “Ternary Phase Equilibria In the System Zr-N-B-C at 1500˚C”, Key Eng. Mater., 53-55, 92–100 (1991) (Experimental, Calculation, Phase Relations, 93) Saugnac, F., Teyssandier, F., Marchand, A., “New Compounds Obtained by LPCVD in the B-C-N Chemical System”, J. Phys. IV, France, 1(C2), C2_673-C2_680 (1991) (Experimental, Crys. Structure, Phase Relations, 13) Bill, J., Riedel, R., Passing, G., “Amin-Boran als Precursoren fuer Borcarbidnitrid” (in German), Z. Anorg. Allg. Chem., 610(4), 83–90 (1992) (Experimental, Phase Relations, 14) Bowser, J.R., Jelski, D.A., George, T.F., “Stability and Structure of C12B24N24: A Hybrid Analogue of Buckminsterfullerene”, Inorg. Chem., 31(2), 154–156 (1992) (Calculation, Phase Relations, 17) Morita, M., Hanada, T., Tsutsumi, H., Matsuda, Y., Kawaguchi, M., “Layered-Structure BC2N as a Negative Electrode Matrix for Rechargeable Lithium Batteries”, J. Electrochem. Soc., 139(5), 1227–1230 (1992) (Experimental, Phase Relations, Electrochemistry, 14) Riedel, R., “Materials Harder than Diamond?”, Adv. Mater., 4(11), 759–761 (1992) (Experimental, Phase Relations, 22) Ruh, R., Kearns, M., Zangvil, A., Xu, Y., “Phase and Property Studies of Boron Carbide–Boron Nitride Composites”, J. Am. Ceram. Soc., 75(4), 864–872 (1992) (Experimental, Phase Relations, Mechan. Prop.,16) Saalfrank, P., Ruemler, W., Hummel, H.-U., Ladik, J., “Energetics and Gap Engineering in Alternating Layer and Intralayer Substituted Boron-Nitrogen-Carbon Compounds”, Synth. Met., 52(1), 1–19 (1992) (Calculation, Crys. Structure, Phase Relations, 29) Saugnac, F., Teyssandier, F., Marchand, A., “Characterisation of C-B-N Solid Solutions Deposited from a Gaseous Phase Between 900˚ and 1050˚C”, J. Am. Ceram. Soc., 75(1), 161–169 (1992) (Experimental, Crys. Structure, Phase Relations, 23) Kawaguchi, M., Kawashima, T., “Synthesis of a New Graphite-Like Layered Material of Composition BC3N”, J. Chem. Soc., Chem. Commun., 1133–1134 (1993) (Experimental, Phase Relations, Crys. Structure, 9) Lambrecht, W.R.L., Segall, B., “Anomalous Band-Gap Behavior and Phase Stability of c-BN-Diamond Alloys”, Phys. Rev. B, 47(15), 9289–9296 (1993) (Calculation, Crys. Structure, Phase Diagram, Phys. Prop., Thermodyn., 40) Paisley, M.J., Horie, Y., Davis, R.F., Dan, K., Tamura, H., Sawaoka, A.B., “Novel Phases in ShockCompacted Mixtures of Diamond and Boron Nitride”, J. Mater. Sci. Lett., 12, 1768–1770 (1993) (Experimental, Phase Relations, 4) Sasaki, T., Akaishi, M., Yamaoka, S., Fujiki, Y., Oikawa, T., “Simultaneous Crystallization of Diamond and Cubic Boron Nitride from the Graphite Relative BC2N under High Pressure/High Temperature Conditions”, Chem. Mater., 5(5), 695–699 (1993) (Experimental, Phase Relations, 26) Wen, H.M.Sc., “Thermodynamic Calculations and Constitution of the Al-B-C-N-Si-Ti System” (in German), Thesis, University Stuttgart, 1–183 (1993) (Calculation, Phase Diagram, Thermodyn., 223) Andreev, Yu.G., Lundstroem, T., “High-Temperature Synthesis and Investigation of Hexagonal Boron Carbonitride”, J. Alloys Compd., 210(1-2), 311–317 (1994) (Experimental, Phase Relations, 26) Kakudate, Y., Yoshida, M., Usuba, S., Yokoi, H., Fujiwara, S., Kawaguchi, M., Sako, K., Sawai, T., “Shock Synthesis of a Hybride of Diamond and Cubic Boron Nitride”, Trans. Mater. Res. Soc. Jap., 14B, 1447–1450 (1994) (Experimental, Crys. Structure, 15) McHale, A.E., “VI. Boron Plus Carbon Plus Nitrogen, B-C-N”, Phase Equilibria Diagrams, Phase Diagrams for Ceramists, 10, 209 (1994) (Review, Phase Relations, 1)
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B–C–N [1994Nak]
[1994Rie] [1995Fil] [1995Hub] [1995Kni]
[1995Lev]
[1996Kas] [1996Kom]
[1996Lun]
[1996Nak]
[1996Noz1] [1996Noz2] [1996Wat] [1997Heg]
[1997Hou] [1997Kaw] [1997Sol1]
[1997Sol2] [1997Tat] [1998Kal]
[1998Kas]
[1998Pop1]
19
Nakano, S., Akaishi, M., Sasaki, T., Yamaoka, S., “Segregative Crystallization of Several Diamond-Like Phases from Graphitic BC2N without an Additive at 7.7 GPa”, Chem. Mater., 6, 2246–2251 (1994) (Experimental, Crys. Structure, 31) Riedel, R., “Novel Ultrahard Materials”, Adv. Mater., 6(7-8), 549–560 (1994) (Review, Phase Relations, 95) Filipozzi, L., Derre, A., Conrad, J., Piraux, L., Marchand, A., Carbon, 33, 1747 (1995) as quoted by [2002Sei] Hubacek, M., Sato, T., “Preparation and Properties of a Compound in the B-C-N System”, J. Solid State Chem., 114(1), 258–264 (1995) (Experimental, Phase Relations, Crys. Structure, 21) Knittle, E., Kaner, R.B., Jeanloz, R., Cohen, M.L., “High-Pressure Synthesis Characterization and Equation of State of Cubic C-B-N Solid Solutions”, Phys. Rev. B, 51, 12149–12156 (1995) (Experimental, Phase Relations, Crys. Structure, 43) Levy, R.A., Mastromatteo, E., Grow, J.M., Paturi, V., Kuo, W.P., Boeglin, H.J., Shalvoy, R., “Low Pressure Chemical Vapor Deposition of B-N-C-H Films from Triethylamine Borane Complex”, J. Mater. Res., 10(2), 320–327 (1995) (Experimental, Phase Relations, 31) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institut, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Komatsu, T., Nomura, N., Kakudate, Y., Fujiwara, S., “Synthesis and Characterization of a ShockSynthesized Cubic B-N-C Solid Solution of Composition BC2.5N”, J. Mater. Chem., 6(11), 1799–1803 (1996) (Experimental, Phase Relations, Crys. Structure, 12) Lundstro¨m, T., Andreev, Y.G., “Superhard Boron-Rich Borides and Studies of the B-C-N System”, Mater. Sci. Eng. A - Struct. Mater. Prop. Microstruct. Process., 209(1-2), 16–22 (1996) (Experimental, Phase Relations, 49) Nakano, S., Akaishi, M., Sasaki, T., Yamaoka, S., “Characterization of Several Cubic Phases Directly Transformed from the Graphitic BC2N”, Mater. Sci. Eng. A., 209(1-2), 26–29 (1996) (Experimental, Phase Relations, Crys. Structure, 6) Nozaki, H., Itoh, S., “Structural Stability of BC2N”, J. Phys. Chem. Solids, 57(1), 41–49 (1996) (Calculation, Crys. Structure, 9) Nozaki, H., Itoh, S., “Lattice Dynamics of BC2N”, Phys. Rev. B, 53(21), 14161–14170 (1996) (Calculation, Crys. Structure, 15) Watanabe, M.O., Itoh, S., Mizushima, K., Sasaki, T., “Bonding Characterization of BC2N Thin Films”, Appl. Phys. Lett., 68(21), 2962–2964 (1996) (Experimental, Phase Relations, 16) Hegemann, D., Riedel, R., Dressler, W., Oehr, C., Schindler, B., Brunner, H., “Boron Carbonitride Thin Films by PACVD of Single-Source Precursors”, Chem. Vap. Deposition, 3(5), 257–262 (1997) (Experimental, Phase Relations, 33) Hou, Q., Gao, J., “Micro-Hardness and Adhesion of Boron Carbon Nitride Coatings”, Mod. Phys. Lett. B, 11(16-17), 749–756 (1997) (Experimental, Phase Relations, Mechan. Prop., 20) Kawaguchi, M., “B/C/N Materials Based on the Graphite Network”, Adv. Mater., 9(8), 615–625 (1997) (Review, Phase Relations, Crys. Structure, Phys. Prop., 59) Solozhenko, V.L., “Phase Formation in the B-C-N System at High-Pressures and Temperatures: in situ Studies”, Eur. J. Solid State Inorg. Chem., 34(7-8), 797–807 (1997) (Experimental, Crys. Structure, Thermodyn., 24) Solozhenko, V.L., Turkevich, V.Z., Sato, T., “Phase Stability of Graphitelike BC4N up to 2100 K and 7 GPa”, J. Am. Ceram. Soc., 80(12), 3229–3232 (1997) (Experimental, Phase Relations, Thermodyn., 20) Tateyama, Y., Ogitsu, T., Kusakabe, K., Tsuneyuki, S., Itoh, S., “Proposed Synthesis Path for Heterodimaond BC2N”, Phys. Rev. B, 55(16), R10161-R10164 (1997) (Calculation, Crys. Structure, 23) Kalss, W., Haubner, R., Lux, B., “Deposition of B-C-N Coatings from Trisdimethylaminoboran by Hot-Filament and Microwave Plasma Activation”, Int. J. Refract. Hard Met., 16, 233–241 (1998) (Experimental, Phase Relations, 27) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Oficial Publications of the European Communities, Belgium, 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Popov, C., Saito, K., Yamamoto, K., Ouchi, A., Nakamura, T., Ohana, Y., Koga, Y., “Synthesis of Nitrogen-Rich B-C-N Materials from Melamine and Boron Trichloride”, J. Mater. Sci., 33, 1281–1286 (1998) (Review, Phase Relations, 27)
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19 [1998Pop2]
[1998Sei]
[1998Wid]
[1998Yam] [1998Yao] [1999Kom]
[1999Mor]
[1999Wib]
[1999Yao]
[1999Zha1] [1999Zha2]
[1999Zhe]
[2000Gag] [2000He]
[2000Kur]
[2000Wil]
[2001Fre]
[2001He]
[2001Hua]
B–C–N Popov, C., Saito, K., Ivanov, B., Koga, Y., Fujiwara, S., Shanov, V., “Chemical Vapor Deposition of BC2N Films and their Laser-Induced Etching with SF6”, Thin Solid Films, 312(1-2), 99–105 (1998) (Experimental, Phase Relations, Morphology, 21) Seifert, H.J., Lukas, H.L., Aldinger, F., “Development of Si-B-C-N Ceramics Supported by Phase Diagrams and Thermochemistry”, Ber. Bunsen-Ges. Phys. Chem., 102(9), 1309–1313 (1998) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 16) Widany, J., Verwoerd, W.S., Frauenheim, Th., “Density-Functional Based Tight-Binding Calculations on Zinc-Blende Type BC2N Crystals”, Diamond Relat. Mater., 7(11-12), 1633–1638 (1998) (Calculation, Crys. Structure, Electronic Structure, 20) Yamada, K., “Shock Synthesis of a Graphitic Boron-Carbon-Nitrogen System”, J. Am. Ceram. Soc., 81(7), 1941–1944 (1998) (Experimental, Crys. Structure, 8) Yao, B., Chen, W.J., Liu, L., Ding, B.Z., Su, W.H., “Amorphous B-C-N Semiconductor”, J. Appl. Phys., 84(3), 1412–1415 (1998) (Experimental, Phase Relations, Semicond., 11) Komatsu, T., Samedima, M., Awano, T., Kakudate, Y., Fujiwara, S., “Creation of Superhard B-C-N Heterodiamond Using an Advanced Shock Wave Compression Technology”, J. Mater. Proc. Techn., 85 (1-3), 69–73 (1999) (Experimental, Crys. Structure, 13) Morjan, I., Conde, O., Oliveira, M., Vasiliu, F., “Structural Characterization of CxByNz (x = 0.1 to x = 0.2) Layers Obtained by Laser-Driven Synthesis”, Thin Solid Films, 340(1-2), 95–105 (1999) (Experimental, Phase Relations, Crys. Structure, 21) Wibbelt, M., Kohl, H., Kohler-Redlich, Ph., “Multiple-Scattering Calculations of Electron-Energy-Loss Near-Edge Structures of Existing and Predicting Phases in the Ternary System B-C-N”, Phys. Rev. B, 59 (18), 11739–11745 (1999) (Calculation, Crys. Structure, 58) Yao, B., Liu, L., Su, W.H., “Formation, Characterization and Properties of a New Boron-CarbonNitrogen Crystal”, J. Appl. Phys., 86(5), 2464–2467 (1999) (Experimental, Phase Relations, Crys. Structure, 14) Zhang, R.Q., Chan, K.S., Cheung, H.F., Lee, S.T., “Energetics of Segregation in β-C2BN”, Appl. Phys. Lett., 75(15), 2259–2261 (1999) (Calculation, Crys. Structure, Phase Relations, Thermodyn., 22) Zhang, Y.F., Tang, Y.H., Lee, C.S., Bello, I., Lee, S.T., “Nanocrystalline C-BN Synthesized by Mechanical Alloying”, Diamond Relat. Mater., 8(2-5), 610–613 (1999) (Experimental, Mechan. Prop., Thermodyn., 12) Zheng, J.-C., Huan, C.H.A., Wee, A.T.S., Wang, R.-Z., Zheng, Y.-M., “Ground-State Properties of Cubic C-BN Solid Solutions”, J. Phys.: Condens. Matter, 11(3), 927–935 (1999) (Calculation, Phase Relations, Crys. Structure, 41) Gago, R., Jimenez, I., Albella, J.M., “Boron-Carbon-Nitrogen Compounds Grown by Ion Beam Assisted Evaporation”, Thin Solid Films, 373(1-2), 277–281 (2000) (Experimental, Phase Relations, 18) He, J.L., Tian, Y.J., Yu, D.L., Wang, T.S., Xiao, F.R., Li, A.D., Li, D.C., Peng, Y.G., Li, L., “Chemical Synthesis of Crystalline Hexagonal B-C-N Compound”, J. Mater. Sci. Lett., 19(22), 2061–2063 (2000) (Experimental, Crys. Structure, Phase Relations, 16) Kurdyumov, A.V., Solozhenko, V.L., Gubachek, M., Borimchuk, N.I., Zelyavskii, V.B., Ostrovskaya, N. F., Yarosh, V.V., “Shock Synthesis of Ternary Diamond-Like Phases in the B-C-N System”, Powder Metall. Met. Ceram., 39(9-10), 467–473 (2000), translated from Poroshk. Metal., (9–10), 53–61 (2000) (Experimental, Crys. Structure, 22) Williams, D., Pleune, B., Kouvetakis, J., Williams, M.D., Andersen, R.A., “Synthesis of LiBC4N4, BC3N3, and Related C-N Compounds of Boron: New Precursors to Light Element Ceramics”, J. Am. Chem. Soc., 122, 7735–7741 (2000) (Experimental, Crys. Structure, 17) Freire, F.L., Reigada, D.C., Prioli, R., “Boron Carbide and Boron-Carbon Nitride Films Deposited by DC-Magnetron Sputtering: Structural Characterization and Nanotribological Properties”, Phys. Stat. Sol., A, 187(1), 1–12 (2001) (Experimental, Phase Relations, Mechan. Prop., 31) He, J.L., Tian, Y.J., Yu, D.L., Wang, T.S., Liu, S.M., Guo, L.C., Li, D.C., Jia, X.P., Chen, L.X., Zou, G.T., Yanagisawa, O., “Orthorhombic B2CN Crystal Synthesized by High Pressure and Temperature”, Chem. Phys. Lett., 340(5-6), 431–436 (2001) (Experimental, Crys. Structure, 21) Huang, J., Zhu, Y.T., Mori, H., “Structure and Phase Characteristics of Amorphous Boron-CarbonNitrogen under High Pressure and High Temperature”, J. Mater. Res., 16(4), 1178–1184 (2001) (Experimental, Phase Relations, 20)
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Martinez, E., Lousa, A., Esteve, J., “Micromechanical and Microtribological Properties of BCN Thin Films Near the B4C Composition Deposited by R.F. Magnetron Sputtering”, Diamond Relat. Mater., 10 (9-10), 1892–1896 (2001) (Experimental, Mechan. Prop., 12) Mattesini, M., Matar, S., “First-Principles Characterisation of New Ternary Heterodiamond BC2N Phases”, Comput. Mater. Sci., 20(1), 107–119 (2001) (Calculation, Mechan. Prop., Crys. Structure, 55) Mattesini, M., Matar, S.F., “Search for Ultra-Hard Materials: Theoretical Characterization of Novel Orthorhombic BC2N Crystals”, Int. J. Inorg. Mater., 3(7), 943–957 (2001) (Calculation, Crys. Structure, Electronic Structure, Mechan. Prop., 55) Nicolich, J.P., Hofer, F., Brey, G., Riedel, R., “Synthesis and Structure of Three-Dimensionally Ordered Graphitlike BC2N Ternary Crystals”, J. Am. Ceram. Soc., 84(2), 279–282 (2001) (Experimental, Crys. Structure, 15) Onodera, A., Matsumoto, K., Hirai, T., Goto, T., Motoyama, M., Yamada, K., Kohzuki, H., “Synthesis of Dense Forms of B-C-N System Using Chemical - Vapor - Deposition/ High - Pressure Process”, J. Mater. Sci., 36(3), 679–684 (2001) (Experimental, Phase Relations, 36) Solozhenko, V.L., Andrault, D., Fiquet, G., Mezouar, M., Rubie, D.C., “Synthesis of Superhard Cubic BC2N”, Appl. Phys. Lett., 78(10), 1385–1387 (2001) (Experimental, Crys. Structure, Mechan. Prop., 23) Solozhenko, V.L., Dub, S.N., Novikov, N.V., “Mechanical Poperties of Cubic BC2N, a New Superhard Phase”, Diamond Relat. Mater., 10(12), 2228–2231 (2001) (Experimental, Mechan. Prop., 25) Solozhenko, V.L., Novikov, N.V., “Cubic Boron Carbonitride - a New Superhard Phase” (in Russian), Dop. Nat. Akad. Nauk Ukr., (11), 81–86 (2001) (Experimental, Phase Relations, Crys. Structure, 12) Sun, H., Jhi, S.-H., Roundy, D., Cohen, M.L., Louie, S.G., “Structural Forms of Cubic BC2N”, Phys. Rev. B, 64(9), 094108_1–094108_6 (2001) (Calculation, Crys. Structure, Phys. Prop., 19) Chattopadhyay, S., Chen, L.C., Chien, S.C., Lin, S.T., Wu, C.T., Chen, K.H., “Phase and Thickness Dependence of Thermal Diffusivity in a-SiCxNy and a-BCxNy”, Thin Solid Films, 420-421, 205–211 (2002) (Experimental, Interface Phenomena, 26) Gago, R., Jimenez, I., Agullo-Rueda, F., Albella, J.M., Czigany, Zs., Hultman, L., “Transition from Amorphous Boron Carbide to Hexagonal Boron Carbon Nitride Thin Films Induced by Nitrogen Ion Assistance”, J. Appl., Phys., 92(9), 5177–5182 (2002) (Experimental, Phase Relations, 30) Langenhorst, F, Solozhenko, V.L., “ATEM-EELS Study of New Diamond-Like Phases in the B-C-N System”, Phys. Chem. Chem. Phys., 4(20), 5183–5188 (2002) (Experimental, Experimental, Crys. Structure, 27) Seifert, H.J., Aldinger, F., “Phase Equilibria in the Si-B-C-N System”, Struct. Bonding, 101, 1–58 (2002) (Review, Phase Diagram, Phase Relations, Thermodyn., 275) Solozhenko, V.L., “Synthesis of Novel Superhard Phases in the B-C-N System”, High Pressure Res., 22 (3-4), 519–524, (2002) (Experimental, Phase Relations, Crys. Structure, Mechan. Prop., 17) Terrones, M., Grobert, N., Terrones, H., “Synthetic Routes to Nanoscales BxCyNz Architectures”, Carbon, 40(10), 1665–1684 (2002) (Review, Crys. Structure, Mechan. Prop., Morphology, Phys. Prop., Thermodyn., 79) Zhao, Y., He, D. W., Daemen, L. L., Shen, T.D., Schwarz, R. B., Zhu, Y., Bish, D. L., Huang, J., Zhang, J., Shen, G., Qian, J., Zerda, T. W., “Superhard B-C-N Materials Synthesized in Nanostructured Bulks”, J. Mater. Res., 17(12), 3139–3145 (2002) (Experimental, Crys. Structure, Mechan. Prop., 33) Voelger, K.W., Kroke, E., Gervais, Ch., Saito, T., Babonneau, F., Riedel, R., Iwamoto, Y., Hirayama, T., “B/C/N Materials and B4C Synthesized by a Non-Oxide Sol-Gel Process”, Chem. Mater., 15(3), 755–764 (2003) (Experimental, Kinetics, Phase Relations, 51) Betranhandy, E., Matar, S.F., Weihrich, R., Demazeau, G., “Potential New Candidates for Hard Material within the Ternary XC3N3 (X = B, Al, Ga) Stoichiometry”, Compt. Rend. Chimie, 7(5), 529–535 (2004) (Calculation, Crys. Structure, Phys. Prop., 34) Bysakh, S., Chattopadhyay, K., Ling, H., Wu, J.D., Dong, C., Wang, Y.Q., Duan, X.F., Kuo, K.H., “A Study of Nanostructures of Thin Films in B-C-N System Produced by Pulsed Laser Deposition and Nitrogen Ion-Beam-Assisted Pulsed Laser Deposition”, J. Mater. Res., 19(3), 759–767 (2004) (Experimental, Kinetics, Phase Relations, 13) He, J.L., Guo, L.C., Wu, E., Luo, X.G., Tian, Y.J., “First-Principles Study of B2CN Crystals Deduced from the Diamond Structure”, J. Phys.: Condes. Matter., 16(46), 8131–8138 (2004) (Calculation, Crys. Structure, Electronic Structure, Phys. Prop., 21)
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[2006Pan] [2006Sun]
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[Mas2] [V-C2]
B–C–N Huang, F.L., Cao, C.B., Xiang, X., Lv, R.T., Zhu, H.S., “Synthesis of Hexagonal Boron Carbonitride Phase by Solvothermal Method”, Diamond Relat. Mater., 13(10), 1757–1760 (2004) (Experimental, Phase Relations, 27) Hubble, H.W., Kudryashov, I., Solozhenko, V.L., Zinin, P.V., Sharma, S.K., Ming, L.C., “Raman Studies of Cubic BC2N, a New Superhard Phase”, J. Raman Spectrosc., 35(10), 822–825 (2004) (Experimental, Phase Relations, 32) Komatsu, T., “Bulk Synthesis and Characterization of Graphite-Like B-C-N and B-C-N Heterodiamond Compounds”, J. Mater. Chem., 14(2), 221–227 (2004) (Experimental, Phase Relations, 31) Solozhenko, V.L., “In situ Studies of High-Pressure Phase Transformations in the B-C-N System”, High Pressure Res., 24(4), 499–509 (2004) (Experimental, Phase Relations, Crys. Structure, 26) Engbrecht, E.R., Fitzpatrick, P.R., Junker, K.H., Sun, Y.-M., White, J.M., Ekerdt, J.G., “Adhesion of Chemical Vapor Deposited Boron Carbo-Nitride to Dielectric and Copper Films”, J. Mater. Res., 20(8), 2218–2224 (2005) (Experimental, Electr. Prop., Morphology, 22) Grigor’ev, O.N., Lyashenko, V.I., Timofeeva, I.I., Rogozinskaya, A.A., Tomila, T.V., Dubovik, T.V., Panashenko, V.M., “Study of the Synthesis of the Ternary Compound B-N-C”, Powder Metall. Met. Ceram., 44(9-10), 415–419 (2005) (Experimental, Phase Relations, 13) Kim, Ch.S., Choi, S.H., Kim, W.J., Jo, S.J., Lee, S.J., Song, K.M., Baik, H.K., “Enhancement of Formation in Boron-Carbon-Nitrogen Nanotubes by Plasma Rotating Electrode Process”, Jpn. J. Appl. Phys., 44(32), L1034-L1037 (2005) (Experimental, Morphology, 15) Solozhenko, V.L., Kurakevych, O.O., “Reversible Pressure-Induced Structure Changes in Turbostratic BN-C Solid Solutions”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., B61(5), 489–503 (2005) (Calculation, Experimental, Crys. Structure, Phase Relations, 26) Uddin, M.N., Shimoyama, I., Baba, Y., Sekiguchi, T., Nagano, M., “X-Ray Photoelectron Spectroscopic Observation of B-C-N Hybrides Synthesized by Ion Beam Deposition of Borazine”, J. Vac. Sci. Technol. A, 23(3), 497–502 (2005) (Experimental, Phase Relations, 41) Azevedo, S., de Paiva, R., “Structural Stability and Electronic Properties of Carbon-Boron Nitride Compounds”, Europhys. Lett., 75(1), 126–132 (2006) (Calculation, Phase Relations, Electronic Structure, 25) Azevedo, S., “Energetic Stability of B-C-N Monolayer”, Phys. Lett. A, 351(1-2), 109–112 (2006) (Calculation, Phase Relations, 21) Luo, X., Zhang, J., Guo, X., Zhang, G., He, J., Yu, D., Liu, Z., Tian, Y., “Synthesis of B-N-C Nanocrystalline Particle by Mechanical Alloying and Spark Plasma Sintering”, J. Mater. Sci., 41(24), 8352–8355 (2006) (Experimental, Crys. Structure, Morphology, 18) Pan, Z., Sun, H., Chen, C., “Ab initio Structural Identification of High Density Cubic BC2N”, Phys. Rev. B, 73(21), 214111_1–214111_4 (2006) (Calculation, Phase Relations, Crys. Structure, 25) Sun, J., Zhou, X.-F., Qian, G.-R., Chen, J., Fan, Y.-X., Wang, H.-T., Guo, X., He, J., Liu, Z., Tian, Y., “Chalcopyrite Polymorph for Superhard BC2N”, Appl. Phys. Lett., 89(15), 151911_1–151911_3 (2006) (Calculation, Crys. Structure, Electronic Structure, 28) Uddin, M.N., Shimoyama, I., Baba, Y., Sekiguchi, T., Nath, K.G., Nagano, M., “Synthesis and Characterization of Oriented Graphitelike B-C-N Hybrid”, J. Appl. Phys., 99(8), 084902_1–084902_5 (2006) (Experimental, Electronic Structure, Morphology, 27) Sun, G., Liu, Z.-Y., He, J.-L., Yu, D.-L., Tian, Y.-J., “Chemical Synthesis of C3N and BC2N Compounds”, Chin. Phys. Lett., 24(4), 1092–1094 (2007) (Experimental, Crys. Structure, 29) Wu, Q.-H., Hu, Q.-K., Luo, X.G., Yu, D.-L., Li, D.-C., He, J.-L., “First-Principles Study of Structural and Electronic Properties of Layered B2CN Crystals”, Chin. Phys. Lett., 24(1), 180–183 (2007) (Calculation, Crys. Structure, Electronic Structure, 28) Record, M.-Ch., Tedenac, J.-C., “B-N (Boron-Nitrogen)”, MSIT Binary Evaluation Program, Effenberg, G. (Ed.), to be published in MSIT Workplace, MSI, Stuttgart (2008) (Review, Phase Diagram, 50) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Niobium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl, Kostyantyn Korniyenko, Tamara Velikanova
Introduction Alloys of the boron-carbon-niobium system are among the most refractory materials known. Applications not only for cutting tools and abrasive purposes but also for heat shields in supersonic outer space carriers are envisaged. Based on early studies [1952Gla, 1955Bre] on the interaction between niobium carbides, niobium borides, B4C and carbon on samples hot-pressed at T > 1500˚C, phase relations in the B-C-Nb system were investigated by a series of authors [1963Rud, 1965Lev, 1977Ord, 1984Ord, 1985Zak, 1987Ord, 1993Ord, 2004Pad] and were presented in terms of an isothermal section at 1750˚C [1963Rud], three eutectic quasibinaries: NbB2-C [1965Lev], NbB2NbC1–x [1977Ord, 1978Ord] and NbB2 - ‘B4C’ [1987Ord, 2004Pad], as well as by two partial isopleths in the Nb-rich corner for Nb + 0.4 mass% C, and Nb + 1 mass% B [1985Zak]. These results were obtained employing X-ray powder diffraction (XPD) [1952Gla, 1955Bre, 1963Rud, 1965Lev, 1977Ord, 1985Zak, 1987Ord, 2004Pad], light optical micrographic analysis (LOM) [1963Rud, 1965Lev, 1985Zak, 1987Ord], scanning electron microscopy (SEM) [2004Pad], pyrometric melting point measurements [1965Lev, 1977Ord, 1987Ord], chemical analyses [1963Rud, 1977Ord], electron microprobe analyses (EMPA) [2004Pad] as well as by DTA and Auger spectroscopy [1985Zak]. Parts of this research were summarized in [1983Sch], [1984Hol] and [1994McH]. A full status of all information in literature for the B-C-Nb system up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog]. Although earlier investigations did not encounter any ternary B-C-Nb compounds, four novel ternary compounds, τ1 to τ4, were reported [2001Hil]. These compounds were synthesized on the basis of Al/Cu-flux experiments and their structure types were determined from X-ray single crystal four-circle intensity data. The most recent systematic reinvestigation of phase relations in the entire B-C-Nb system is due to [2004Kor] providing also a Scheil diagram on the solidification. Experimental details for all investigations in the B-C-Nb system are summarized in brief in Table 1.
Binary Systems The binary system C-Nb, as presented in Fig. 1, combines the earlier version of [Mas2] with results of a reinvestigation by means of diffusion couples, XPD, EMPA [1996Len, 1997Len, 1998Wie] (note the congruent melting of NbC1–x at 45 at.% C). For the low temperature region [1991Gus, 2000Gus] calculated the phase relations considering two superlattice phases (carbon/vacancy ordering), Nb3C2 and Nb6C5, forming in the defect region of NbC1–x. The binary B-Nb system was reassessed from data by [1966Rud, 1985Zak, 1992Rog, 2003Bor, Landolt‐Bo¨rnstein New Series IV/11E1
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2005Nun] (see Fig. 2). Although [1985Zak] claimed that the formation of Nb3B2 is associated with oxygen contamination during prolonged alloy annealing, a reinvestigation by [2003Bor] confirmed the existence of Nb3B2 in a sluggish peritectoid reaction: (Nb) + NbB Ð Nb3B2 as earlier suggested by [1966Rud] at 2080 ± 40˚C. Agreement exists on the eutectic reaction L Ð (Nb) + NbB, which was reassured by [2003Bor], was given by [1966Rud] at 2165 ± 10˚C, 19 ± 2 at.% B and was reported by [1985Zak] at 2170 ± 20˚C, 12 at.% B. However, EMPA data located the (Nb) + NbB eutectic at 16 at.% B [2003Bor]. An unsolved puzzle is the existence of Nb2B3, reported by [1991Oka] from Cu-flux experiments, but not yet confirmed in bulk samples. A thermodynamic modeling of the B-Nb system is due to [2007Pec]. The B-C system corresponds to an assessment and thermodynamic modelling by [1998Kas, 1996Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron rich reaction L+ ‘B4C’ Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan]. Literature data concerning the formation and crystal structure of binary and ternary solid phases pertinent to the B-C-Nb system are listed in Table 2.
Solid Phases In contrast to earlier investigations, which did not encounter any ternary B-C-Nb compounds, four novel ternary compounds were claimed to exist from synthesis via Cu/Al flux [2001Hil]. These compounds adopt unique structure types (for details see Table 2) forming a structural series (NbB)2(NbB2)n(NbC)m combining structural sub-units of NbC, NbB and NbB2. However, a recent investigation in the temperature region from 1750˚C to the melting (LOM, X-ray and EMPA) by [2004Kor] only revealed the compound Nb3B3C (labeled τ1), which melts congruently at 2970˚C. At present it is unclear what served as a stabilizer for the other three (metastable?) compounds reported by [2001Hil]. As-cast alloys corresponding to the compositions of the phases, Nb7B6C3 (τ2), Nb4B3C2 (τ3), as well as Nb7B4C4 (τ4) contained primary NbC1–x , followed by a ternary eutectic τ1 + NbC1–x + Nb3B4 or with last portions of melt solidifying in a ternary eutectic NbC1–x + Nb5B6 + NbB (see also Figs. 3, 4). Mutual solid solubilities among binary compounds in the B-C-Nb system were reported to be very low [1963Rud, 1987Ord, 2004Kor]. The maximal solubility of boron in the niobium carbides, according to the data of [1963Rud] at 1750˚C, is not higher than 1 at.%. Despite solubility of NbB2 in NbC1–x was observed to be rather small at moderate temperatures, 3 mol% NbB2 were reported to dissolve in NbC1–x at 2300˚C rising to 7 mol% NbB2 at 2600˚C with linear increase of the NbC lattice parameter [1977Ord, 1978Ord, 2004Kor]. The joint solubility of boron and carbon in niobium at 2080˚C was reported to be 1.67 at.% B and 1.5 at.% C [1985Zak].
Quasibinary Systems A total of eight quasibinary sections have been encountered so far. Figures 3, 4 present the eutectic type quasibinary sections NbB2-NbC1–x [1977Ord, 1978Ord] and B4.5C-NbB2 [1987Ord, 2004Pad] with some corrections in order to comply with the most reliable, accepted binary phase diagrams. In particular, all alloys prepared by the authors of [1977Ord, 1978Ord] from the NbB2-NbC1–x system with contents up to 7 mol% NbC1–x DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009
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and for more than 25 mol% NbC1–x reveal melting temperatures that are considerably higher than 2600˚C, however, the samples, annealed at subsolidus temperature, were metallographically two-phase NbB2 + NbC1–x. The carbon content of 11.3 mass% C of NbC1–x used by [1977Ord], in fact corresponded to Nb0.504C0.496 which does not comply with congruent melting of Nb0.55C0.45 (Fig. 3); nevertheless the extrapolated melting point was correctly shown to be at 3600˚C (NbC0.82) [1977Ord]. The significantly lower melting point recorded by [1977Ord] for NbB2 ( 2950˚C), is in contrast to the value of 3036 ± 30˚C given in the accepted B-Nb binary in Fig. 2. Thus the true quasibinary eutectic temperature might in fact be slightly higher. Indeed melting point data by [2004Kor] from the alloy Nb39B43C18 (at.%) near the tie line NbB2-NbC1–x recorded a significantly higher temperature (2888 ± 19˚C, Table 3). The existence of a quasibinary eutectic e7, L Ð NbB + NbC1–x, at 2800˚C was established by [2004Kor] from the Nb50B40C10 (at.%) Pirani specimen. The E2e7U2 curve falls from this point both to the eutectic point E2 (eutectic L Ð NbB + NbC1–x+ Nb5B6) at 2728 ± 12˚C, and to the peritectic point U2 at 2375 ± 15˚C (Figs. 5a, 5b and 6). For the quasibinary eutectic system, B4.5C-NbB2, the authors of [1987Ord, 2004Pad] showed a fiber like directionally solidified eutectic structure with NbB2 needle-like single crystals (whiskers of 1–2 μm diameter). The content of NbB2 in the eutectic point is 32.7 mol% (12.3 at.% C), slightly less than that reported by [1987Ord] (12 at.% C corresponding to 35–37 mol% NbB2; TE = 2250 ± 30˚C). Pirani data [2004Kor] measured on Nb11B78C11 (at.%), yield a slightly higher temperature of 2290 ± 25˚C and a slightly higher C content of 13.5 at.% C (e12). The U3e12E6 curve, falls from the e12 point to the point U3 which corresponds to the liquid taking part in the invariant ternary transition reaction LU3 + ‘B4C’ Ð NbB2 + (βB) (Fig. 6). With respect to the peritectic binary reaction L + ‘B4C’ Ð (βB) [2005Tan], the probability of a transition reaction U4 near the B-apex is higher than the formation of a ternary eutectic. This was already proven in several Calphad calculations for the systems B-C-T (T = Ti, Zr, Hf, V, W) [1998Rog]. For the quasibinary eutectic L Ð NbB2 + (C)gr, e9, the reinvestigation by [2004Kor] gave a slightly higher melting point TE = 2707 ± 9˚C in contrast to older data from [1965Lev] (TE = 2650˚C); furthermore the position of the eutectic shifted from 32 mol% NbB2 ([1965Lev]) to Nb23.5B47.5C29 ([2004Kor]; Table 3, Fig. 6).
Invariant Equilibria Solidification behavior of the niobium-boron-carbon alloys is presented in a Scheil diagram summarizing the invariant equilibria involving a liquid phase [2004Kor] (Table 3 and Fig. 5a and 5b). With respect to early melting point data by [1965Lev] on the quasibinary eutectic e9 L Ð NbB2 + (C)gr (TE = 2650˚C), the melting point recently determined by [2004Kor] employing the Pirani technique on the specimen Nb22B45C33 (at.%) appeared somewhat higher, namely TE = 2707 ± 9˚C (Table 3). Primary carbon forms on solidification of alloy Nb25B40C35 (at.%), followed by the quasibinary eutectic L Ð NbB2 + (C)gr and finally by the ternary eutectic L Ð NbB2 + (C)gr + NbC1–x for which EMPA yielded Nb27B38C35 (in at.%). Composition of the ternary eutectic LE6 Ð NbB2 + (C)gr + ‘B4C’ at 2243 ± 11˚C was established from EMPA on the alloy Nb15B65C20 (at.%) the microstructure of which shows primary dendrites of NbB2, quasibinary eutectic L Ð ‘B4C’ + NbB2 and the ternary eutectic L Ð ‘B4C’ + NbB2 + (C)gr, which solidifies as last portion of liquid in the alloy Nb5B65C30 Landolt‐Bo¨rnstein New Series IV/11E1
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(at.%). The τ1 phase participates in three invariant four-phase eutectic processes, and in three eutectic invariant reactions forming quasibinary sections τ1-NbC1–x , τ1-NbB2 and τ1-Nb3B4 (see Fig. 6). Two modifications of Nb2C are not distinguished in the diagrams presented in the present evaluation.
Liquidus, Solidus and Solvus Surfaces The liquidus surface, plotted in Fig. 6, refers to LOM, Pirani, EMPA as well as X-ray data of [2004Kor]; it comprises ten invariant reactions with participation of liquid and eleven surfaces of primary crystallization. The experimental results confirmed essentially the general form of the B-C-Nb liquidus projection presented earlier [1998Rog]. Among the four ternary phases τ1 to τ4 reported by [2001Hil] (Nb3B3C, Nb4B3C2, Nb7B6C3 and Nb7B4C4) only τ1 (Nb3B3C) was detected in as-cast alloys. The liquidus in the niobium rich corner essentially corresponds to the data of [1985Zak] (L Ð (Nb) + Nb2C + NbB, TE = 2080 ± 20˚C, at 12.3 at.% B, 5.8 at.% C). More recent data by [2004Kor] locate this eutectic in all alloys that contain 55 at.% and more niobium at 2060˚C with a slight shift to Nb75B18C7, Fig. 6). The highest temperatures on the B-C-Nb liquidus surface occur in the surface of primary crystallization of carbon. This surface is not a complete liquidus surface because alloys near carbon show direct sublimation from solid to gas. The E3e4E4 monovariant curve was placed between the alloys Nb39B43C18 and Nb35B40C25 (in at.%) as they contained small amounts of primary phases (NbB2 and NbC1–x , respectively). The process with participation of liquid, NbB2 and Nb3B4 phases, starting in the B-Nb system as incongruent L + NbB2 Ð Nb3B4, changes its character to congruent L Ð NbB2 + Nb3B4. Fields of primary crystallization of Nb3B4 and Nb5B6 phases become wider with increasing carbon content, especially of the Nb3B4 phase. Positions of these fields were determined on the basis of combined analysis of LOM, EMPA and X-ray data, in particular, for the alloy Nb47B48C5 (at.%), in which the Nb5B6 phase is primary [2004Kor].
Isothermal Sections The isothermal section at 1750˚C, shown in Fig. 7, is mainly based on data of [1963Rud], however, was amended by two niobium borides established later (Nb5B6 and Nb2B3 [1991Oka]). From Fig. 7 one can see that carbon forms equilibria with ‘B4C’, NbC1–x and NbB2. This fact confirms the observations of [1955Bre, 1963Rud] that all niobium borides with boron contents lower than NbB2 are unstable when heated in combination with carbon.
Temperature – Composition Sections Two partial isopleths, namely Nb + 0.4 mass% C and Nb + 1 mass% B, were reported by the authors of [1985Zak] assuming the non-existence of “impurity stabilized” Nb3B2. With respect to the reaffirmed Nb3B2 [2003Bor] these isopleths are not shown. DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009
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Thermodynamics No experimental thermodynamic data are presently available for the ternary system. Upper limits for the heat of formation of NbB2 and the lower niobium borides were estimated from compatibility studies among niobium borides and carbon [1955Bre].
Notes on Materials Properties and Applications Properties of niobium carbide/niobium boride for designing structural materials to be used under extreme conditions were reviewed by [2001Bur]. Niobium carbide and niobium boride may be used in dispersion-strengthened nanocomposites prepared by mechanical alloying of Cu with Nb, B, C, which will produce a dispersion of stable carbide and boride reinforcement particles within a nanostructured copper matrix [2006Liv]. Microhardness (under a load of 100 g) was recorded for samples from the quasibinary systems NbB2-NbC1–x [1977Ord, 1978Ord] and NbB2-‘B4C’ [1987Ord]. In both cases microhardness values are significantly below the linear combination of the binary constituents (13.7 to 19.6 GPa for NbB2/NbC1–x and 28.5 to 31 GPa for NbB2/‘B4C’). The variation is due to the degree of dispersion of the phases decreasing with increasing dispersion (“superplasticity effect”). Solubility of NbB2 in NbC1–x significantly rises the microhardness from 21 GPa (1 mass% NbB2) to 30 GPa for 7 mass% NbB2 in NbC1–x [1977Ord]. A powder mixture 2Nb+2B+C, that was ball milled for 105 min spontaneously self ignited in air and produced via self propagating high temperature synthesis (SHS) a fine grained mixture of NbC+NbB2; Spark Plasma Sintering (SPS) of this mixture at 1800˚C yielded a homogeneous microstructure of few μm-sized grains but a Vickers harness of 19.8 GPa which compared favorably with a sinter compact from commercial powder mixtures without SHS-SPS (14.3 to 16.6 GPa) [2006Tsu, 2007Tsu]. Interaction in the quasibinary eutectic systems MC-MB2 [1980Ord] and MB2-B4C [1993Ord] has been analyzed for transition elements M = Ti, Zr, Hf, V, Nb and Ta and correlations were found between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms. The bending strength limit of sintered composites NbB2/NbC1–x was shown to exhibit a pronounced maximum of 400 MPa at about 50 mass% NbB2 at 300 K; the maximum reduces to below 300 MPa and shifts to about 30 mass% NbB2 at higher temperatures between 1200 and 2000 K (927 to 1727˚C) [1984Ord].
Miscellaneous Detailed solubility limits of B in Nb were given by [1985Zak] as 0.23 mass% B at 2170˚C, 1.16 mass% at 1950˚C, 0.09 mass% at 1600˚C and 0.03 mass% B at 1200˚C. Preparation of oxygenfree exothermic mixtures of B4C-Nb cermets was reported by [1975Kry].
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. Table 1 Investigations of the B-C-Nb Phase Relations, Structures and Thermodynamics Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied XPD on 6 hot-pressed samples. Compatibility of NbB2 with C and with “B4C”
[1952Gla]
Hot-pressing of powder mixtures (powders of NbH2, NbC and NbB2) in graphite dies at 1500 to 3000˚C. XPD
[1955Bre]
Reaction between metal borides and XPD on two samples (Nb + B + 0.5C; Nb + graphite. Sintering of powder compacts in 2B + C) Mo-cruciles under 0.5 bar argon for 50 min at 1777˚C. XPD
Isothermal section at 1750˚C [1963Rud] Hot-pressing of 60 binary and ternary powder compacts in C-cartridges at 1200 to 2500˚C followed by subsequent anneal in a W-tube vacuum furnace (2.5 Pa) for 9 h at 1750˚C. Starting materials: amorphous boron (94 mass% residue O, C, Fe), lampblack C, Nb (98 mass% Nb, 0.8 mass% Ta, 0.3% Ti, 0.3% Fe, traces of oxide), NbC (11.5 mass% Ctotal, 0.40 mass% free C). XPD, LOM. [1965Lev]
NbB2 prepared by vacuum sintering (Nb +2B) powder compacts for 5 h at 1400 to 1700˚C and analyzed by XPD, LOM. Starting materials were Nb powder of 5 to 10 μm containing 0.06 mass% impurities, 0.08% Fe, 0.15% Pb, 0.05% Si, 0.16% O. Boron impurities were: 0.0036% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb. Thermal analysis by direct electrical heating of a graphite tube filled with NbB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
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Interaction of NbB2 with C yielded a quasieutectic system NbB2+C with eutectic point at 32 mol% C and TE of ca. 2650˚C.
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1977Ord] Samples prepared from NbB2 (synthesized Investigation of the quasieutectic system at 1700˚C in vacuum from >99.85 mass% NbB2+NbC1–x with eutectic point at 55% Nb and amorphous >97mass% B) and NbB2 and TE of 2600˚C on 14 samples. NbC (synthesized at 1900˚C in vacuum from Nb and lampblack C; Ctot = 11.40 mass% C, 0.10 mass% free C 0.15 mass% (O2+N2)). Specimens in form of cylinders (3 mm diameter x 50 mm were compacted with aid of 12% aqueous starch solution, presintered at 2100˚C for 2 h in vacuum prior to heat treatment at 2300˚C. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon. [1978Ord] See [1977Ord], same tasks
See [1977Ord]
[1985Zak] Alloys prepared by electron beam melting of Nb (< 0.01 mass% C, 0.007 mass% N, 0.01 mass% O, 0.02 mass% Ta), lampblack B (> 99.0 mass% B), NbC and NbB2. 50 g ingots were arc melted under He. Alloys containing < 1 mass% C and B were compressed in vacuum with 50–60% deformation. Specimens were annealed: 1950˚C 25 h, 1600˚C 100 h, 1200˚C 150 h, with intermediate quenching in liquid Ga. As-cast and annealed alloys were investigated by metallography, XPD, Auger spectroscopy and DTA (under He, 1.5 g sample in HfO2-crucibles).
Niobium corner covering the range Nb-NbB-Nb2C up to 2.4 mass% B and 1.8 mass% C. Polythermal sections though the Nb-NbB-Nb2C system at 0.4 mass% C and at 1.0 mass% B.
[1987Ord] Samples were prepared from NbB2 (synthesized at 1700˚C in vacuum from > 99.85 mass% Nb and amorphous > 97 mass% B) and B4C which was annealed in vacuum at 2000˚C to reduce C content to 0.2 mass% free C. Specimens in form of cylinders were compacted with aid of 12% aqueous starch solution, presintered at 2100˚C for 2 h in vacuum prior. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the quasieutectic system NbB2+‘B4C’ with eutectic point at 35–37 mol% NbB2 and TE of 2250˚C on 12 samples.
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[2001Hil]
Compounds forming a structural series Determination of the crystal structures (NbB)2(NbB2)n(NbC)m were synthesised (Nb3B3C, Nb4B3C2, Nb7B6C3 and Nb7B4C4) from Cu/Al flux in Al2O3 crucibles under from X-ray counter data. argon starting from Cu:Al:B:C:Nb ratios 20:5:2:1:4 (Nb3B3C); 20:5:0.2:0.1:0.2 (Nb4B3C2); 20:5:2:1:2 (Nb7B6C3); 20:5:0.1:0.1:0.2 (Nb7B4C4). The mixtures were heated to 1600˚C at 400˚/h, kept for 24 h and cooled to 1500˚C at 100˚/h, to 1100˚C at 1˚/h and to RT at 150˚/h. Single crystals were obtained after dissolution of the flux in halfconcentrated HNO3. From EDX only Nb was found as a metal. In all cases NbC was found as a side product.
[2004Kor]
More than 30 ternary samples prepared with various techniques (A, B, C) from high purity powders of components: niobium (99.9 and 99.85 mass%), boron (Ventron GmbH, D), boron carbide B4C (Johnson Matthey & Co, UK) and carbon (with purity of 99.99 mass%). For melting point measurements by the Pirani technique [1923Pir] specimens were prepared in form of cylindrical polycrystalline rods by sintering at 2000˚C of ceramic green bodies from isostatically compressed blends of Nb, B4C, B and C powders. Powder blends were compacted in form of the final Pirani shape in steel dies using a small amount of xylol as densification aid. The green bodies (with a lateral hole directly pressed into the sample) were slowly heated (to prevent violent self-heating via exothermic reactions) in high vacuum to 1600˚C for 12 to 15 h. For LOM and EMPA, several alloys were prepared by argon arc melting of sintered pellets
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Melting points were determined on bout 4 Pirani samples of each composition in a precise micro-pyrometer calibrated for the temperature region from 2000 to 2900˚C. The microstructure of the cast alloys was inspected using light optical microscopy (LOM) on flat surfaces prepared by grinding (SiC-paper) and polishing the resin-mounted alloys using diamond pastes down to 1/4 μm grain size. Quantitative composition analyses were performed on a CAMEBAX SX50 wavelength dispersive spectrograph (EMPA) comparing the X-ray emissions of the three elements in the alloys with those from elemental standards of Mo, Si and LaB5.85 for boron after applying a deconvolution and ZAF-correction procedure. Determination of liquidus and solidus in the entire range of compositions. Determination of isothermal section at 1750˚C. Determination of reaction isotherms for reactions involving liquid. Derivation of Schulz-Scheil diagram for the entire diagram. Revision of isopleths NbB2-NbC1–x, ‘B4C’(B4.5C)-NbB2
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[2004Pad] Rods of 8–10mm diameter and 50 to 100 mm length were prepared from powder compacts (99.5 mass% NbB2 and 99.9% B4.3C) annealed at 1500˚C in vacuum prior to zone melting in an induction furnace under 0.2 MPa argon. Compositions chosen were: Nb11B76C13 (46 mol% NbB2) and Nb7.5B77C15.5 (33 mol% NbB2). XPD, SEM, EMPA.
Directional crystallization of B4C-NbB2 eutectic compositions investigated by XPD, SEM and EMPA revealing NbB2 needles (whiskers of 1–2 mm diameter)
[2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)) the liquidus and solidus curves to the L + (βB) field has been derived via chemical analysis
Confirmation of peritectic type of reaction L + B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L + (βB) field
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Nb) < 2469
cI2 Im 3m W
a = 330.04
at 25˚C [Mas2]
(βB) < 2092
hR333 R3m βB
a = 1093.30 c = 2382.52
[1993Wer]
a = 1093.02 c = 2381.66 a = 1092.2 c = 2381.4 a = 1091.91 c = 2382.24
pure B [1976Lun]
(αB)
hR36 R3m αB
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at 1.1 at.% C [1993Wer], linear da/dx, dc/dx at NbB99.5 [1992Rog] presumably metastable phase, preparation below 1000˚C [1971Amb] pure B, single crystal [1994Cha]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
(C)gr hP4 a = 246.12 < 3827 P63/mmc c = 670.90 (sublimation point) C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78
Comments/References at 25˚C [Mas2]
[1967Low] at 2.35 at.% Cmax (2350˚C), linear da/ dx, dc/dx [1967Low]
(C)d
cF8 a = 356.69 Fd 3m C (diamond)
at 25˚C, 60 GPa [Mas2]
‘B4C’ < 2450
hR45 R 3m B13C2
a = 565.1 to 560.7 c = 1219.6 to 1209.5
9 to 20 at.% C [1990Ase]
a = 556.0 c = 1212.0 a = 556.1 c = 1212.0
sample quenched from 2400˚C [1987Ord] sample ‘B4C’ + 14.5 mol% NbB2 quenched from 2420˚C [1987Ord]
B25C
tP68 P 42m or P42/nnm B25C
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog]
Nb3B2 < 2080
tP10 P4/mbm U3Si2
a = 619.79 c = 329.26
[1992Rog] oxygen stabilized? [1985Zak]
NbB < 2917
oC8 Cmcm CrB
a = 329.74 b = 872.38 c = 316.69 a = 330.2 b = 875.5 c = 316.8 a = 330.0 b = 873.3 c = 316.6 a = 329.61 b = 872.24 c = 316.53
[1992Rog]
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as-cast, at 32.9 at.% B [1985Zak] at 32.9 at.% B, 25 h at 1950˚C [1985Zak] [1991Oka]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] Nb5B6 < 2870
Nb3B4 < 2935
Pearson Symbol/ Space Group/ Prototype oC22 Cmmm V5B6
oI14 Immm Ta3B4
Nb2B3
oC20 Cmcm V2B3
NbB2 < 3036
hP3 P6/mmm AlB2
Lattice Parameters [pm] a = 315.30 b = 2227.44 c = 330.49 a = 315.67 b = 2276.7 c = 330.34 a = 314.1 b = 2275.6 c = 330.6 a = 314.51 b = 1410.62 c = 330.19 a = 314.28 b = 1407.6 c = 330.33 a = 330.58 b = 1948.1 c = 312.93 a = 311.26 c = 326.27 a = 308.61 c = 330.69 a = 311.15 c = 326.57 a = 310.37 c = 332.37 a = 311.2 c = 327.0 a = 312.3 c = 329.0 a = 314.0 c = 331.5 a = 309.5 c = 330.0 a = 308.9 c = 330.2 a = 310.0 c = 330.0
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Comments/References [1992Rog]
[1991Oka]
as-cast, together with Nb3B4, τ1 and NbC1–x [2004Kor] [1992Rog]
[1991Oka]
[1991Oka]
65 to 70 at.% B Nb rich [1992Rog] B rich [1992Rog] Nb rich [1991Oka] B rich [1991Oka] at 227˚C [V-C2] at 727˚C [V-C2] at 1227˚C [V-C2] for NbB2, quenched from 3000˚C [1987Ord] for samples NbB2 and NbB2 + 1 mol % NbC1–x quenched from 2600˚C [1977Ord] for sample NbB2 + 7.7 mol% ‘B4C’, quenched from 2620˚C [1987Ord]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] Nb2C(h2) 3080 - 2450
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
hP4 P63/mmc defect NiAs
28 to 35.5 at.% C; labeled as γNb2C [1996Len, 1997Len, 1998Wie] defect structure hP3 [1998Rog] for NbC0.36 quenched from 1750˚C [1963Rud] for Nb2C quenched from 1750˚C [1963Rud] at 11.11 at.% C, as-cast [1985Zak]
a = 311.6 c = 495.8 a = 312.4 c = 496.8 a = 311.9 c = 495.6 a = 311.5 c = 495.4 Nb2C(h1) 2530 - 1195
hP9 P31m W2C1)
Nb2C(r) < 1195
oP12 Pnma Nb2C42) (MnO2)
Nb4C3–x < 1575 ± 25
hR24 R 3m V4C3
Nb6C5 < 1050
mC22 C2/m Nb6C53)
hP36 or P31 Nb6C5
DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009
Comments/References
at 11.11 at.% C, 150 h annealing at 1200˚C [1985Zak]
a = 541.69 c = 479.19
27 to 36.6 at.% C; labeled as βNb2C [1996Len, 1997Len, 1998Wie] [V-C2]
a = 1091.0 b = 309.54 c = 497.46
at 34.6 to 36.6 at.% C; labeled as αNb2C [1967Rud] (impurity stabilized?)
a = 314 ± 1 c = 3010 ± 10
40.1 to 40.7 at.% C [1996Len, 1997Len, 1998Wie] [V-C2] defect structure tP52 [1998Rog]
a = 544.7 b = 943.5 c = 544.7 β = 109.47˚ a = 546.4 c = 1542.2
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at 45 at.% C [Mas2] [V-C2]
[V-C2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] NbC1–x < 3600
Pearson Symbol/ Space Group/ Prototype cF8 Fm 3m NaCl
Lattice Parameters [pm] a = 443.0 a = 447.0 a = 446.81 a = 448.44 a = 443.13 a = 446.90 to 447.36
* τ1, Nb3B3C < 2970
oC28 Cmcm Nb3B3C
a = 326.47 b = 2871.0 c = 312.85 a = 325.9 b = 2875.5 c = 312.7
Comments/References 40 to 49.6 at.% C [Mas2] at 41 at.% C, quenched from 1750˚C [1963Rud] at 48 at.% C, quenched from 1750˚C [1963Rud] 48 at.% C, T = 20˚C [V-C2] 48 at.% C, T = 616˚C [V-C2] 41 at.% C, T = 20˚C [V-C2] from samples NbC1–x + NbB2, quenched from 2600˚C; linear increase from 0 to 7 mol% NbB2 [1977Ord] [2001Hil]
as-cast [2004Kor]
* τ2, Nb4B3C2
oC36 Cmcm Nb4B3C2
a = 322.88 b = 3754.4 c = 313.32
[2001Hil] metastable?
* τ3, Nb7B6C3
oC32 Cmmm Nb7B6C3
a = 313.41 b = 3316.1 c = 324.28
[2001Hil] metastable?
* τ4, Nb7B4C4
oI30 Immm Nb7B4C4
a = 315.442 b = 321.66 c = 3226.0
[2001Hil] metastable?
partially ordered V2N type, cited by various authors as partially ordered εFe2N type [1963Rud]. [1967Rud] indexed Nb2C (r) on the basis of the ζFe2N type with cell parameters a = 1089.5, b = 496.8, c = 1235 pm. This phase was assumed to be impurity (oxygen?) stabilized and was thus removed from the diagram in [Mas2]. 3) Nb6C5 as claimed by [1991Gus], was said to be a NaCl-type derivative superlattice structure with symmetry C2, C2/m or P31. 1) 2)
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. Table 3 Invariant Equilibria Composition (at.%) Reaction
T [˚C]
Type
Phase
B
C
Nb
L Ð Nb3B4 + τ1
2940
e2 (max)
L
44.5
12.5
43
L Ð NbB2 + τ1
2910
e3 (max)
L
44.4
13.2
42.4
L Ð NbB2 + τ1 + Nb3B4
2900 ± 15
E1
L
44.8
12.5
42.7
L Ð NbB2 + NbC1–x
2888 ± 19
e4 (max)
L
43
16
41
L Ð Nb3B4 + NbC1–x
2830 ± 30
e6 (max)
L
44.
10
46
L Ð NbB + NbC1–x
2800
e7 (max)
L
40
8.5
51.5
NbB2
66
1
33
NbC1-x
4
40
56
14.5
43
L Ð τ1 + NbC1–x
2790 ± 10
e8 (max)
L
42.5
L+Nb3B4 Ð Nb5B6 + NbC1–x
2772 ± 13
U1
L
44
9.5
46.5
L Ð Nb5B6 + NbC1–x + NbB
2728 ± 12
E2
L
42.5
8.0
49.5
L Ð NbB2 + (C)gr
2707 ± 9
e9 (max)
L
47
29
24
L Ð NbB2 + τ1 + NbC1–x
2700 ± 20
E3
L
43
15.5
41.5
L Ð NbB2 + NbC1–x + (C)gr
2570 ± 10
E4
L
40
31.5
28.5
L + NbC1–x Ð NbB + Nb2C
2375 ± 15
U2
L
28.5
14.0
57.5
L Ð NbB2 + ‘B4C’
2290 ± 25
e12 (max)
L
78
13
9
L Ð Nb3B4 + τ1 + NbC1–x
2340 ± 20
E5
L
43
12.5
44.5
L Ð NbB2 + ‘B4C’ + (C)gr
2243 ± 11
E6
L
62.5
22.5
15
L + ‘B4C’ Ð NbB2 + (βB)
< 2103
U3
L
98
1
1
L Ð (Nb) + NbB + Nb2C
2060
E7
L
18
7
75
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. Fig. 1 B-C-Nb. The C-Nb phase diagram
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. Fig. 2 B-C-Nb. The B-Nb phase diagram. Phase boundaries from [1985Zak] are shown as dash lines, from [1963Rud] as dotted lines
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. Fig. 3 B-C-Nb. The NbB2-NbC1–x phase diagram. Dashed lines from [1977Ord]; dotted lines from [2004Kor]
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. Fig. 4 B-C-Nb. The ‘B4C’ (B4.5C)-NbB2 phase diagram
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. Fig. 5a B-C-Nb. Reaction scheme, part 1
B–C–Nb
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. Fig. 5b B-C-Nb. Reaction scheme, part 2
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. Fig. 6 B-C-Nb. Liquidus surface projection
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. Fig. 7 B-C-Nb. Isothermal section at 1750˚C
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References [1923Pir] [1952Gla] [1955Bre] [1963Rud]
[1965Lev]
[1966Rud]
[1967Low] [1967Rud] [1971Amb] [1975Kry]
[1976Lun] [1977Ord]
[1978Ord]
[1980Ord]
[1983Sch]
[1984Hol]
[1984Ord]
[1985Zak]
[1987Ord]
Pirani, M., Alterthum, H., “Method for the Determination of the Melting Point of Metals which Fuse at High Temperatures” (in German), Z. Elektrochem., 29, 5–8 (1923) (Phase Relations, Experimental, 5) Glaser, F.W., “Contribution to the Metal-Carbon-Boron Systems”, Trans. AIME - J. Metals, (4), 391–396 (1952) (Crys. Structure, Experimental, Electr. Prop., 19) Brewer, L, Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–405 (1955) (Phase Relations, Thermodyn., Experimental, Review, 19) Rudy, E., Benesovsky, F., Toth, L., “Studies of the Ternary Systems of the Group Va and VIa Metals with Boron and Carbon” (in German), Z. Metallkd., 54, 345–353 (1963) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 43) Levinskii, Yu.V., Salibekov, S.E., Levinskaya, M.K., “Interaction of Diborides of V, Nb, Ta with Carbon” (in Russian), Poroshk. Metall., (Kiev), (5), 66–69 (1965) (Phase Diagram, Phase Relations, Experimental, *, 5) Rudy, E., Windisch, S., “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Tech. Rep. AFML-TR-65–2, Air Force Materials Laboratory, Wright-Patterson Air Force Base OH, Part I, Vol. X, 1–103 (1966) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, *, 64) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Rudy, E., Brukl, C.E., “Lower-Temperature Modifications of Nb2C and V2C”, J. Am. Ceram. Soc., 50, 265–268 (1967) (Crys. Structure, Experimental, 14) Amberger, E., Ploog, G., “Formation of Pure Boron Lattice” (in German), J. Less-Common Met., 23, 21–31 (1971) (Crys. Structure, Experimental, 18) Krylov, Y.I., Bronnikov, V.A., Krysina, V.G., Pristavko, V.V., “On the Possibility of Creating Thermite Mixtures Based on Compositions of Boron Carbide (B4C) Materials and Silicon Carbide Materials” (in Russian), Poroshk. Metall., (Kiev), (12), 57–60 (1975) (Morphology, Experimental, Phys. Prop., 5) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Ordan’yan, S.S., Stepanenko, E.K., Unrod, V.I., “Reactions in the System NbC-NbB2”, Inorg. Mater., 13(2), 312–314 (1977), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 13(2), 373–375 (1977) (Crys Structure, Phase Diagram, Experimental, Mechan. Prop., Phase Relations, 4) Ordan’yan, S.S., Stepanenko, E.K., Unrod, V.I., “Reactions in the NbC-NbB2 System” (in Russian), in “Vysokotemp. Boridy Silitsidy”, Kosolapova, T.Ya. (Ed.), Naukova Dumka, Kiev, USSR, 64–67 (1978) (Crys Structure, Phase Diagram, Experimental, Mechan. Prop., Phase Relations, 4) Ordanyan, S.S., “Laws of Interaction in the Systems MIV,VC - MIV,VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Thermodyn., Experimental, Theory, Electronic Structure, 14) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Phase Diagram, Phase Relations, Review, Mechan. Prop., 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Ordan’yan, S.S., Stepanenko, E.K., Sokolov, N.V., “Strength of Sintered Niobium Carbide - Niobium Boride (NbC-NbB2) Composite Materials” (in Russian), Izv. Vyss. Uchebn. Zaved., Khim., Khim. Tekhnol. USSR, 27(10), 1201–1203 (1984) (Morphology, Experimental, Mechan. Prop., 4) Zakharov, A.M, Pshokin, V.P, Ivanova, E.I., “Niobium Corner of the System Nb-B-C”, Russ. Metall., (5), 192–195 (1985), translated from Izv. Akad. Nauk SSSR, Met., (5), 193–196 (1985) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, *, 10) Ordan’yan, S.S, Dmitriev, A.I., Bizhev, K.T., Stepanenko, E.K., “Methods of Examination and Properties of Powder Material Interaction in B4C-Me(V)B2 Systems”, Sov. Powder Metall. Met. Ceram., 298(10), 834–836 (1987), translated from Poroshk. Metall. (Kiev), 298(10), 66–69 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *, 5)
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20 [1990Ase]
[1991Gus]
[1991Oka]
[1992Rog]
[1993Ord]
[1993Wer]
[1994Cha]
[1994McH]
[1996Kas] [1996Len]
[1997Len]
[1998Kas]
[1998Rog]
[1998Wie]
[2000Gus] [2001Bur] [2001Hil] [2003Bor]
[2004Kor]
B–C–Nb Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides” Freer, R. (Ed.), Dordrecht: Kluwer Academic Publishers, 97–111 (1990) (Crys. Structure, Experimental, Review, 14) Gusev, A.I. “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State., 32(9), 1595–1599 (1991) (Phase Diagram, Phase Relations, Thermodyn., Theory, *, 18) Okada, S., Hamano, K., Lundstroem, T., Higashi, I., “Crystal Growth of the New Compound Nb2B3, and the Borides NbB, Nb5B6, Nb3B4 and NbB2 Using the Copper-flux Method” in “AIP Conference Proceedings 231 on Boron-rich Solids”, Albequerque, USA, 1990, New York, AIP, 590–593 (1991) (Crys. Structure, Experimental, 12) Rogl, P., “The System B-N-Nb” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), ASM, Materials Park, Ohio, USA, 68–72 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, *, 6) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, (1), 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, Electronic Structure, 18) Werheit, H., Kuhlmann, U., Laux, M., Lundstroem, T., “Structural and Electronic Properties of Carbon-doped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Chakrabarti, D.J., Laughlin, D.E., “B-Cu (Boron-Copper)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 74–78 (1994) (Review, Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., 24) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 189–190 (1994) (Phase Diagram, Phase Relations, Review, 2) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Lengauer, W., Wiesenberger, H., Joguet, M., Rafaja, D., Ettmayer, P., “Chemical Diffusion in Transition Metal-Nitrogen Systems” in “The Chemistry of Transition Metal Carbides and Nitrides”, Oyama, S.T. (Ed.), Oxford, Blacky Academic, 91–106 (1996) (Phase Diagram, Phase Relations, Experimental, *, 29) Lengauer, W., Wiesenberger, H., Mayr, W., Bidaud, E., Berger, R., Ettmayer, P., “Phase Stabilities of Transition Metal Carbides and Nitrides Investigated by Reaction Diffusion”, J. Chim. Phys., 94, 1020–1025 (1997) (Morphology, Experimental, Interface Phenomena, *, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Niobium” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 197–213 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, #, 25) Wiesenberger, H., Lengauer, W., Ettmayer, P., “Reaction Diffusion and Phase Equilibria in the V-C, Nb-C, Ta-C and Ta-N systems”, Acta Mater., 46(2), 651–666 (1998) (Phase Diagram, Experimental, Interface Phenomena, *, 30) Gusev, A.I., “Order-Disorder Phase Transformations in Strongly Non-Stoichiometric Compounds” (in Russian), Physics Usbekhi, 3(1), 1–4, 34–37 (2000) (Phase Diagram, Thermodyn., Theory, *, 12) Burkhanov, G.S., “Structural Materials on the Basis of Rare Materials” (in Russian), Metally, (5), 57–61 (2001) (Morphology, Review, Phys. Prop., 46) Hillebrecht, H., Gebhardt, K., “Crystal Structures from a Building Set: the First Boride Carbides of Niobium” (in German), Angew. Chem., 113(8), 1492–1495 (2001) (Crys. Structure, Experimental, 17) Borges, L.A., Jr., Coelho, G.C., Nunes, C.A., Suzuki, P.A., “New Data on Phase Equilibria in the Nb-rich Region of the Nb-B System”, J. Phase Equilib., 24(2), 140–146 (2003) (Crys. Structure, Experimental, *, 14) Korniyenko, K., Rogl, P., Velikanova, T., Leithe-Jasper, A., Bohn, M., Tanaka, T., “The System BoronCarbon-Niobium”, Research at the University of Vienna, 2000 to 2004, OEAD-Report, Nov. 2004, unpublished work, 1–39 (2004) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, #, 27)
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[2006Tsu]
[2007Tsu]
[2007Pec]
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Paderno, V., Paderno, Yu., Filippov, V., Liashchenko, A., “Directional Crystallization of B4C-NbB2 and B4C-MoB2 Eutectic Compositions”, J. Solid State Chem., 117, 523–528 (2004) (Morphology, Phase Relations, Experimental, *, 14) Nunes, C.A., Kaczorowski, D., Rogl, P., Baldissera, M.R., Suzuki, P.A., Coelho, G.C., Grytsiv, A., Andre, G., Bouree, F., Okada, S., “The NbB2-Phase Revisited: Homogeneity Range, Defect Structure, Superconductivity”, Acta Mater., 53, 3679–3687 (2005) (Crys. Structure, Experimental, Electr. Prop., *, 33) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg, Germany, August 21–26, 142 (2005) (Phase Relations, Experimental, *, 4) Livramento, V., Marques, M.T., Correia, J.B., Almeida, A., Vilar, R., “Dispersion-Strengthened Nanocomposites Prepared with Niobium Carbide and Niobium Boride by Mechanical Alloying”, Mater. Sci. Forum, 514-516(Pt. 1, Advanced Materials Forum III), 707–711 (2006) (Morphology, Experimental, Mechan. Prop., 10) Tsuchida, T., Kakuta, K., “Fabrication of SPS Compacts from NbC-NbB2 Powder Mixtures Synthesized by the MA-SHS in Air Process” J. Alloys Compd., 415(1-2), 156–161 (2006) and ibid 398(1-2), 67–73 (2005) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 14) Tsuchida, T., Kakuta, K., “MA-SHS of NbC and NbB2 in Air from the Nb/B/C Powder Mixtures”, J. Eur. Ceram. Soc., 27(2-3), 527–530 (2007) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 7) Pecanha, R.M., Ferreira, F., Coelho, G.C., Nunes, C.A., Sundman, B., “Thermodynamic Modeling of the Nb-B System”, Intermetallics, 15, 999–1005 (2007) (Thermodyn., Phase Diagram, Phase Relations, 31) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Kostyantyn Korniyenko
Introduction Phase relations in the B-C-Si system are of great interest above all because boron, carbon and silicon are the basic elements for the development of technically important refractory ceramics and hard materials [2002Sei]. The attention of investigators of the constitution of the ternary B-C-Si system has been concentrated mainly on understanding the sintering mechanisms of SiC with boron in combination with carbon and the sintering of boron carbide with silicon [1990Tel]. Attempts of the construction of the phase diagram has been carried out by [1972Gug, 1972Kie] (partial liquidus surface projection, schematic isothermal sections as well as a series of temperature-composition sections) and by [1964Sec, 1969Sha] (the B4CSiC section). Publications concerning experimental studies of phase relations, crystal structures and thermodynamics as well as the techniques applied are listed in Table 1. However, all the available experimental data reported in literature are quite contradictory and insufficient to produce a clear interpretation of the phase relationships in the system. With a view to solving this problem, thermodynamic calculations involving the published experimental data were undertaken by [1982Doe, 1986Lil, 1994Gou1, 1994Gou2, 1995Gou, 1996Kas, 2002Sei] but the results obtained are still in need of further experimental verification. The thermodynamic properties, namely related to the vaporization behavior of B-C-Si alloys, as well as a thermodynamic analysis of the boron dissolution in silicon carbide were obtained experimentally by [1964Ver] and by [1965Mee1, 1977Saf], respectively. Reviews of the literature data relating to phase equilibria in the ternary are presented in [1972Kie, 1983Sch, 1996Kas, 2002Sei]. A review of the thermodynamic properties has been provided by [1996Sin].
Binary Systems The accepted B-C boundary binary system comes from a thermodynamic assessment and modelling carried out by [1996Kas]. A thermodynamic description for the system also appears in [1998Kas], the main differences between the two being in the modelling of the ‘B4C’ and (βB) phases, but for the purposes of this assessment, the former is accepted as the fit to the experimentally determined eutectic temperature is superior. The binary boundary B-Si system shown in Fig. 1 is accepted from [1998Fri] (based on an original publication in 1991 [1991Fri] and reproduced in [1995Lim]). The C-Si boundary binary system (Fig. 2) is accepted according to [1996Gro, 1998Gro].
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Solid Phases Crystallographic data of the solid B-C-Si phases, and their concentration and temperature ranges of stability are presented in Table 2. The joint solubility of two components in boron, carbon or silicon has not been reported. Ternary compounds of compositions SiB5C2 and Si2C2B3 were suggested by [1960Por] but their existence has not been confirmed in subsequent studies. [1996Lij] and [1996Lih] hot pressed mixtures of B4C, SiC and C at temperatures of 1800, 2000 and 2200˚C and a pressure of 25 MPa for 30 min under N2. From the X-ray diffraction studies, they found new ternary phases with crystal structures different from those of the phases in the binary systems. These ternary phases were denoted as P and M (labeled in Table 2 as τ1 and τ2, respectively) but as their appearance has not been confirmed by any other method of investigation it would seem that these two phases are metastable. [1996Kas] modeled the silicon solubility in ‘B4C’ by taking into account Si2-units occupying C-sites in the linear C-B-C chain as described by [1994Wer]. Structural characteristics and chemical bonds of Si-doped boron carbides were studied in [1999Xin] through calculations of different structural unit models by using a self-consistent-field discrete variation Xα method. The calculations show that the most energetically stable configuration for Si doped boron carbide occurs when the Si atom substitutes for B or C atoms at the end of boron carbide chain, and this allows occupation of interstitial sites. However, it is energetically difficult for Si to substitute B or C atom in the center of chain or in the icosahedral structural unit.
Invariant Equilibria Temperatures, reaction types and phase compositions relating to the invariant equilibria of the system are listed in Table 3; the corresponding reaction scheme is shown in Fig. 3. The data are based mainly on the review [2002Sei], which describes the Calphad assessment of [1996Kas]. However, some differences in the detail exist between the two publications. Also, two other thermodynamic calculations were published previously; [1982Doe, 1995Lim], the latter is also a Calphad assessment. But for the reaction scheme here, the data of [2002Sei] were preferred as the boundary binary systems used in the assessment agree in the most part with those accepted in the present report. At the same time, the compositions of the phases taking part in invariant equilibria as calculated by [1996Kas] are listed in Table 3, except for the data for the solid phases taking part in the reaction involving the liquid, SiBn, SiB6 and ‘B4C’ phases. This reaction is reported in [1996Kas] as transition, whereas [2002Sei] lists it as degenerate. All of the listed phase compositions in which the concentration of one of the components is zero, are marked as approximate. Existence of the quasibinary eutectic reaction L Ð βSiC + ‘B4C’ was reported at 2300 ± 20˚C by [1964Sec] following micrographic and X-ray studies. [1972Gug, 1972Kie], determined the eutectic temperature to be at 2240˚C using a similar array of techniques. However, [1969Sha] suggested that this reaction is in fact an invariant four-phase eutectic equilibrium L Ð βSiC + ‘B4C’ + (C)gr, which takes place at 2245 ± 5˚C. The discrepancies in these data result from an incorrect interpretation of the experimental data, in particular, the large scatter of the melting point values for the ‘eutectic-containing compositions’. Instead of the above ternary eutectic reaction, a transition reaction L + (C)gr Ð ‘B4C’ + βSiC was calculated by DOI: 10.1007/978-3-540-88053-0_21 ß Springer 2009
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[2002Sei] at 2295˚C. In total, three transition and three degenerate four-phase equilibria were proposed, but further experimental verification of their character is necessary.
Liquidus, Solidus and Solvus Surfaces Figure 4 presents the liquidus surface projection involving the stable phase equilibria, consistent with the reaction scheme (Fig. 3). The primary crystallization fields and isotherms are taken from the thermodynamic assessment of [1996Kas]. It can be seen that the boron and silicon solid solution regions as well as binary B-Si phases are pressed against the corresponding boundary side. The experimental data of [1972Gug] (the similar work is [1972Kie]) were used in the assessment of [1996Kas] and it was noted that calculated liquidus temperatures agree within the limits of experimental error. A liquidus surface projection was also reported in [1982Doe] but obsolete versions of the boundary binary systems were used in the calculation. A solidus surface projection is shown in Fig. 5 consistent with the reaction scheme and the compositions of the phases taking part in the invariant equilibria as calculated by [1996Kas] (Table 3).
Isothermal Sections A partial isothermal section for 2477˚C in the range of compositions adjacent to the βSiC phase (30 to 70 at.% Si, up to 1 at.% B) was proposed by [1996Kas] on the basis of thermodynamic calculations. Quite good agreement with the data of [1969Sha] in relation to the solubility of boron in βSiC at this temperature (about 0.1 mass% or 0.15 at.%) is observed. The isothermal section for 2327˚C was calculated by [1982Doe]. It is presented in Fig. 6 with slight changes relating to the constitution of the accepted binary boundary systems, in particular, with respect to the homogeneity range of the ‘B4C’ phase. The calculated isothermal section for 2227˚C taken from [2002Sei] is shown in Fig. 7. On comparing with the section for 2327˚C, it can be seen that the liquid phase field extends only from the B-Si side. The isothermal section for 2127˚C proposed by [1982Doe] is identical to the that for 2227˚C given by [2002Sei] with respect to constitution, but certain variations in the phase ranges exist. The isothermal section for 1927˚C was calculated by [1982Doe] across the whole range of composition. For boron contents up to 70 at.% (Fig. 8) it is similar to those for 2227 and 2127˚C, while phase fields containing (βB) and SiBn appear in the boron-rich corner. Because the B-Si binary system at the boron end was revised subsequent to the publication of [1982Doe], the corresponding part of the section is omitted in Fig. 8. [1982Doe] predicted similar phase equilibria taking place at 1727˚C. Schematic isothermal sections at 1900˚C and 1700˚C proposed by [1972Kie] are in good agreement with the calculated results of [1982Doe] for similar temperatures. [1996Lij] investigated phase equilibria of composites of approximate composition Si6.2B38.1C55.7 (in at.%) from the C-B4C-SiC region of the system. The samples were prepared by hot pressing at temperatures of 1800, 2000 and 2200˚C. According to X-ray diffraction data, the phases (C)gr, ‘B4C’ and βSiC were found in the specimens prepared at all sintering temperatures but at 2200 and 2000˚C the τ1 or τ2 phases, respectively (Table 2), were also identified. These phases seem to be metastable because they were not detected by TEM. Landolt‐Bo¨rnstein New Series IV/11E1
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Isothermal sections at 1227˚C and 1127˚C were calculated by [1996Kas] and [2002Sei], respectively. The character of phase relationships at both these temperatures is similar. In particular in the B rich corner, phase equilibria involving SiB3 appear, in contrast to higher temperatures. The section at 1127˚C according to [2002Sei] is shown in Fig. 9. [1994Gou1, 1994Gou2] presented a thermodynamic evaluation of the isothermal section at 1127˚C, taking into account their own results of phase codeposition using the chemical vapor deposition (CVD) technique. However, the character of the proposed equilibria differ from those reported in [2002Sei], and moreover show the section only schematically. As [2002Sei] presents the phase diagram data at various temperatures and these data are in a good mutual agreement, they should be considered to be more reliable.
Temperature – Composition Sections As was noted in the section “Invariant Equilibria”, the quasibinary reaction L Ð βSiC + ‘B4C’ along the SiC-B4C section was reported on the basis of experimental data by [1964Sec] (commented in [1964Sch, 1965Sec]) and by [1972Gug, 1972Kie] at 2300 ± 20˚C or at 2240˚C, respectively. At the same time, [1969Sha] observed an invariant four-phase equilibrium L Ð βSiC + ‘B4C’ + (C)gr at 2245 ± 5˚C. Thermodynamic calculation of the SiC-B4C temperature-composition section was carried out by [1982Doe], and later by [1996Kas] for the SiC-B4.6C section. This change is related with correction of the ‘B4C’ phase composition corresponding to the congruent melting point as appearing in the later version of the B-C boundary binary system. The calculated SiC-B4.6C temperature-composition section is shown in Fig. 10 according to [1996Kas, 2002Sei], accepted in this report as preferable. Properties of the SiC-B4C composites are presented in [2002Gun, 2003Aka, 2004Ueh, 2005Lee, 2005Tka, 2007Lat], the SiC/B4C multilayer coatings were studied by [1996Her]. The calculated temperature-composition sections SiC-B and the isopleth at 80 at.% B (both according to [1996Kas]) as well as the Si-B4.18C [2002Sei] are presented in Figs 11, 12 and 13, respectively. All the section are brought into conformity with accepted boundary binary systems and the reaction scheme.
Thermodynamics The vaporization behavior was studied by [1964Ver] using a mass spectrometer. The gaseous molecule SiBC was identified on the basis of measurements of mass, isotopic distribution, intensity profile in the molecular beam and appearance potential. The determined reaction enthalpies are listed in Table 4. The values of the free energy functions derived from the usual statistical thermodynamic formulas for the SiBC molecule at the temperatures of 1727, 1827, 1927 and 2027˚C as well as it’s atomization energy ΔH˚0 (at.) calculated using the third-law method are presented in Table 5. The atomization energy value demonstrates that this molecule is very strongly bonded. The partial pressures are listed in Table 6. Thermodynamic behavior of the gaseous phase during the alloying of silicon carbide with boron were investigated by [1977Saf]. The SiC single-crystals alloyed with boron were obtained by a diffusional annealing technique in the temperature range 1600 to 2550˚C. The partial heat of dissolution of boron in silicon carbide was reported as 415.9 kJ·mol–1. Thermodynamic aspects of the dissolution process of the acceptor impurity boron in silicon DOI: 10.1007/978-3-540-88053-0_21 ß Springer 2009
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carbide are also reported in [1986Lil]; the temperature dependences of the partial pressures of the interacting components were calculated for the temperature interval 1527 to 2727˚C. The calculation of the Fermi level in p-SiC (B) also was carried out. Thermodynamic calculations of phase equilibria were carried out by in [1982Doe, 1996Kas, 2002Sei]; the work of [1996Kas] is a Calphad assessment of the system, which was reported in the review by [2002Sei]. The thermodynamic description of [1996Kas] is based on data for the pure elements as stored in the SGTE database [SGTE]. Datasets for the boundary binary systems were taken from [1991Fri, 1996Gro, 1996Kas]. Because of a lack of data and the probably small energetic differences out of measurable quantities, a single analytical Gibbsenergy description was used to describe the α- and β- modifications of SiC [1996Gro]. The review by [2002Sei] reports the ternary assessment by [1996Kas], although some of the details of the calculation are slightly different to the original work. This may be owing to different unary data being used in the calculations given by [2002Sei], but no indication of this is actually given. The calculated elements of phase diagram are described in the relevant parts of the present report. The results of an experimental study and calculation of the thermodynamic parameters for chemical vapor codeposition (CVD) in a hot-wall reactor of alloys of the B-C-Si system are reported in [1994Gou1, 1994Gou2] (for 1127˚C) and in [1995Gou] (for the temperature range 927 to 1127˚C). The initial gaseous mixture consisted of methyltrichlorosilane, boron trichloride and hydrogen. The codeposits were first obtained on graphite, but the final purpose was to extend the process to composite materials. A total pressure was 0.395 bar was applied, with a total flow rate was between 0.1 and 1.0 g·min–1. It was concluded in [1995Gou] that low temperatures and high mass flow rates favor the deposition of uniform coatings but different results can be obtained as a function of the inlet gaseous composition since different kinetic limitations can arise which favor a boron or a silicon excess in the coating. A theoretical thermodynamic analysis of the SiC-B4C section [1996Sin] indicated that the adiabatic temperatures in combustion synthesis can be reduced significantly by the addition of filler (SiC or ‘B4C’ phases). The ‘B4C’ phase was shown to be much more effective in reducing the adiabatic temperature than SiC. In their opinion, these parameters can be very helpful in the selection of optimum processing conditions for the synthesis and densification of SiC-B4C composites.
Notes on Materials Properties and Applications B-C-Si alloys are of great practical interest due to their excellent mechanical properties (hardness, strength, etc.), chemical stability, high-temperature stability as well as high semiconducting properties. Among the phases taking part in equilibria in this system, silicon carbide and boron carbide are the most promising materials for use in different fields of technology. The addition of boron to silicon carbide increases its hardness, heat resistance and polishing ability at the same time as retaining the high oxidation resistance. On the other hand, the addition of silicon to boron carbide improves its sintering behavior and mechanical properties [1978Ekb]. The ceramic eutectic SiC-B4C is of interest [1979Hon] because of the improved mechanical properties and high temperature stability. These constituents also serve as the basis for engineered multilayer coatings (SiC/B4C) offering significant potential for improved tribological properties [1996Her] as well as for heterojunction diodes owing to their electronic properties [2001Ade]. Landolt‐Bo¨rnstein New Series IV/11E1
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The experimental techniques applied and properties studied are listed in Table 7. In a review of the mechanical properties of nanostructured superhard materials [2001And] a value of Vickers microhardness of 63 GPa was cited for a Si0.35B12C2.9 monolayer film [1998Bad].
Miscellaneous Investigations of the scaling resistance of B-C-Si alloys with a boron content of 80 at.% at 1000˚C carried out by [1965Mee2] show that the kinetics of oxidation of these alloys exhibits a parabolic diffusional character with a gradual decrease in the oxidation rate and the formation of a thick and strongly adhering glasslike film of the borosilicide type. According to the findings of [1977Gor], investigations of the absorption spectra associated with the excitation of excitons localized on neutral boron in the silicon carbides αSiC and α’’SiC make it possible to discover a series of bands that arise owing to the unequivalent positions of the substitutional impurities in the crystal lattice. A model of the energy spectrum of these excitations is proposed that takes into account valley-orbit interaction and explains the polarization properties of the spectra. It was reported by [1979Hon], that directionally solidified the SiC-B4C eutectic forms lamellar microstructures with no colonies if the solidification rate is below 0.02 m·h–1. [1979Pan] have studied the kinetics of wetting of ‘B4C’ by silicon. Formation of the SiC-based phase was reported. The morphology of the in-situ growth of SiC in the carbonbased composite bodies consisting of βSiC and ‘B4C’ particles was studied in [1994Oga]. A great part of SiC grains was found to be disks showing irregular utlines and rugged surface. Rod shaped grains of SiC were also observed in the fracture surface of the composites, of which cross sections revealed various morphologies such as triangular, rectangular and irregular. A semiempirical analysis of the sizes and signs of the small boron hyperfine interaction constants as determined previously by [1995Adr] using electron nuclear double resonance in boron-doped αSiC shows that the B–Si-C+ model of this centre can account for the highly unusual feature of small isotropic and anisotropic boron hyperfine constants that are positive and negative, respectively. A high-frequency pulsed EPR/ENDOR study at 95 GHz and on the 13 C-enriched single crystals of boron-doped αSiC was reported by [1997Sch]. This study enabled the formulation of a consistent model of the electronic structure of the shallow and deep boron acceptor at low temperatures. With regard to the shallow boron acceptor, it was concluded that about 40% of the spin density was located in the pz-orbital of the carbon that is nearest to boron. Later, a high-frequency (95 GHz) and conventional-frequency (9.3 GHz) pulsed EPR/ENDOR study of the deep boron acceptor in αSiC was presented in [1998Dui]. The results suggest a model in which the deep boron acceptor consists of a boron atom in a silicon position with an adjacent carbon vacancy. EPR spectra of deep boron in βSiC and αSiC crystals have been measured and studied by [1998Bar]. An ENDOR investigation has been performed on the shallow boron acceptor in the βSiC by [1998Hof]. Hyperfine and quadrupole parameters were determined with high precision for both isotopes (10B and 11B). The oxidation behavior of B4C-SiC/C composites of various compositions at temperatures up to 1500˚C was studied in [1999Guo, 2003Fan, 2003Nar], and the results indicated that the composites exhibited variable oxidation resistance at high temperatures depending on composition and oxidation temperature. The variance of self-healing properties was attributed to
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the difference in the compositions, and the properties of the decarbonized layers including wetting ability, viscosity, volatility and oxygen permeability. The sintering of nano crystalline αSiC by doping with boron carbide was carried out in [2002Dat]. The maximum density of the sintered αSiC was obtained at a concentration of 0.5 mass% of boron carbide with 1 mass % of carbon, sintered at 2050˚C for 15 min under vacuum (3 mbar). It was reported that carbon reduced the silica layer and enhanced the bulk self diffusion coefficient of silicon carbide by many orders of magnitude. It also inhibits grain growth of SiC crystals. The optimum carbon content which gives rise to high density was found to be 1 mass%. According to [2002Gun], the mass gain of directionally solidified B4C-SiC composites at 750˚C owing to oxidation is about 1/3 to 1/5 less than that of monolithic B4C. The surface perpendicular to the growth direction showed slightly better oxidation resistance than that parallel to the growth direction. The diffusion of B impurities in βSiC was analyzed by [2002Rur] using firstprinciples electronic structure calculations. Through molecular dynamics, it was found that substitutional B at a Si lattice site is readily displaced by a nearby Si interstitial by the process known as a kick-out mechanism, in agreement with recent experimental results. This is in contrast to the situation in Si, where B has recently been shown to diffuse via a mechanism involving interstitial lattice sites. The kinetics of βSiC sintering in the presence of boron at 2150˚C was studied in [2003Sto]. It was found that for a carbon concentration of 3 mass%, the optimum addition of boron was 0.2 to 0.5 mass%. The flux growth of SiC crystals from a SiCB4C eutectic melt was carried out by [2004Epe] at temperatures 2300–2350˚C. Both selfnucleated SiC platelets of up to 5 mm in diameter and epitaxial layers grown by physical vapor transport (PVT) on αSiC substrates have been produced. αSiC crystals with low boron contents have been obtained by the same technique by [2006Fan]. The effects of growth conditions, diffusion barrier coatings and hot zone materials on B incorporation into the αSiC crystals were evaluated. Following electron-microscope studies, the processes and mechanisms of structural and phase transformations in the diamond particle Si-B4C contact zone involved during sintering of a composite having an initial composition of ‘B4C’ + diamond + (αSi) are discussed in [2006Shu]. It has been shown that as the sintering temperature increases in the range from 1300 to 2000˚C, the formation of the zone microstructure is determined by the following sequence of the processes: the formation of the nano-dispersed secondary boron carbide on the diamond particle surface, thickening of this layer through the growth of boron carbide grains caused by the boron diffusion from the matrix component to the diamond particle surface, transformation of the resultant layer of anisometric grains to form at the first stage a fine-grained sub-layer in contact with diamond, which later undergoes a complete reconstruction to form a boron carbide-silicon carbide fine-grained composition. Molecular dynamics simulation of the amorphous B-C-Si system was carried out by [2006Ye] in order to investigate the diffusion behavior and analyze the influence of the addition of B on the thermal stability and creep resistance at high temperature of the amorphous system. The results show that the self-diffusion of boron tends to ascend apparently up to 1800˚C, which accounts for phase separation of the amorphous state would take place at about 1800˚C. Below this temperature, the B-C-Si system will retain the thermal stability and good creep resistance. An investigation of the chemical vapor deposition (CVD) process for the B-C-Si system is presented in [2007Ber]. By adding a porous substrate with a high internal surface in the hot zone of the reactor, the consumption of specific species is enhanced, revealing the effective precursors of the solid. In order to better understand the mechanisms of the solid formation,
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correlations are indicated between the gas phase analysis, the deposition kinetics and the deposit physicochemical characteristics. Two chemical processes (at 800–900 and at 900–1000˚C) compete against each other.
. Table 1 Investigations of the B-C-Si Phase Relations, Structures and Thermodynamics Temperature / Composition / Phase Range Studied
Reference
Method / Experimental Technique
[1960Por]
Sintering by hot pressing at 1700–2300˚C; polishing; chemical etching; optical microscopy; X-ray diffraction; melting points measurements (Samsonov and Petrash’s method [1955Sam])
[1964Sec]
Cold and hydrostatic pressing; firing and cooling; 2250˚C, 2300˚C, 2350˚C; the SiCmelting points measurements (optical B4C section pyrometry); X-ray diffraction; optical microscopy; chemical analysis
[1964Ver]
Knudsen cell mass spectrometry
The Si-SiC-B4C-B partial system, the Si-B4C section
1507–2227˚C
[1965Mee1] Precipitation from the gaseous phase; X-ray diffraction; optical microscopy; chemical analysis
50 at.% C
[1965Mee2] Sintering by hot pressing at 1700–2000˚C; X-ray diffraction; optical microscopy; chemical analysis
80 at.% B
[1967Dok]
Hot pressing; arc melting; step annealing at The Si-SiC-B4C-B partial system 1660–1100˚C; chemical etching; X-ray diffraction; optical microscopy; chemical analysis
[1969Nie]
Reduction by hydrogen; precipitation from the gaseous phase; X-ray diffraction; optical microscopy; chemical analysis
SiB5C2, Si2C2B3
[1969Sha]
Crystal growth; firing and quenching; X-ray diffraction; optical microscopy
2200–2550˚C; the Si-B4C-C partial system
[1970Sha]
Recrystallization; crystal growth; X-ray diffraction; emission spectroscopy
αSiC with boron additions
[1971Kal]
Heat pressing; sintering; chemical etching; X-ray diffraction; optical microscopy
SiC-B section adjoining range of compositions
[1972Gug]
Hot pressing and sintering; X-ray diffraction; chemical analysis; optical microscopy; fusion; direct observation of melting in a furnace
1700–2300˚C, whole range of compositions
[1972Kie]
Hot pressing and sintering; X-ray diffraction; chemical analysis; optical microscopy; fusion; direct observation of melting in a furnace
1700–2300˚C, whole range of compositions
[1975Bin]
Hot pressing and sintering; SEM; X-ray diffraction 1750–1950˚C; (SiC) with 1 mass% of (‘B4C’) additions
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1975Kal]
Hot pressing; sintering; X-ray diffraction; optical microscopy
SiC-B section adjoining range of compositions
[1977Saf]
Diffusional annealing; neutron-activation analysis 1600–2550˚C
[1983Wal]
Polymer pyrolysis; DTA; thermogravimetric analysis (TGA); X-ray diffraction
SiC and SiC/B4C ceramics
[1986Mor]
Annealing; SEM; TEM; Auger electron spectroscopy (AES); electron energy loss spectrometry (EELS)
1400, 1500, 1600, 1800, 2000˚C; αSiC with boron additions
[1990Tel]
X-ray diffraction; thermal analysis; EELS
B-rich corner
[1994Gie]
Self-propagating high-temperature synthesis (SHS), X-ray diffraction
SiC + (0.25–20) mass% B; (Si)+SiC; ‘B4C’ + SiC; (Si) + ‘B4C’ mixtures
[1994Gou1] Chemical vapor deposition (CVD) from an initial gaseous mixture (methyltrichlorosilane + boron trichloride + hydrogen)
1127˚C
[1994Gou2] CVD from an initial gaseous mixture (methyltrichlorosilane + boron trichloride + hydrogen)
1127˚C
[1994Wer]
Optical absorption; infrared phonon spectroscopy; Raman spectroscopy
‘B4C’ with Si additions
[1995Gou]
Chemical vapor deposition (CVD); X-ray diffraction 927–1127˚C; codeposition of B, C and Si
[1996Bar]
EPR; luminescence spectroscopy; optically αSiC with boron additions detected magnetic resonance (ODMR); deep level transient spectroscopy (DLTS)
[1996Her]
Sputter deposition; grazing incidence X-ray B4C/SiC multilayer coatings scattering (GIXS); atomic force microscopy (AFM); X-ray diffraction; SEM
[1996Lih]
Hot pressing; sintering; TEM; electron diffraction
2000˚C; Si6.2B38.1C55.7 (at.%) composite
[1996Lij]
Hot pressing; sintering; TEM; energy-dispersive spectrometry (EDS)
1800˚C, 2000˚C, 2200˚C; Si6.2B38.1C55.7 (at.%) composite
[1997Gor]
Self propagating high-temperature synthesis (SHS); SEM
SiC-B4C section (0–30 mass% B)
[1997Sch]
EPR; electron nuclear double resonance (ENDOR); αSiC with boron additions
[1998Ang]
Hot pressing; carbonization; graphitization; X-ray ≤1500˚C; C/B4C/SiC composites diffraction; TGA
[1998Bar]
EPR spectroscopy
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[1998Dui]
Sublimation sandwich method; EPR; ENDOR; X-ray αSiC with boron additions diffraction
[1998Hof]
EPR; ENDOR; X-ray diffraction
βSiC with boron additions
[1998Li]
Arc melting; SEM; X-ray diffraction
B4C-SiBn (n ≥ 14)-Si composites
[1999Guo]
Grinding; sintering; thermal gravimetric/ differential thermal analysis (TG/DTA); SEM
≤1500˚C; B4C-SiC/C composites
[1999Li]
Pressing; arc melting; X-ray diffraction; SEM; TEM SiB4-B4C, SiB6-B4C, SiB14-B4C sections
[2001Ade]
Plasma-enhanced chemical-vapor deposition (PECVD); X-ray diffraction
SiC-B4C heterojunction diode;
[2001Mag]
Sintering at 1950–2200˚C; SEM; optical microscopy; EDS; X-ray diffraction
αSiC-B4C composite
[2002Gun]
Isostatic pressing; floating zone method; X-ray diffraction; SEM; TEM
SiC-B4C eutectic composites
[2003Aka]
Isostatically pressing; floating zone method; X-ray SiC-B4C eutectic composites diffraction; SEM; TEM; EMPA
[2003Nar]
Mixing, pressing, arc melting; oxidation; in situ Raman spectroscopy; X-ray diffraction; SEM
800–1500˚C; 25–60 mol% SiC - B4C
[2003Shu]
Toroid-type high pressure apparatus synthesis; 7.7 MPa, 1400–1600˚C; X-ray diffraction
Si-B4C-based cermets
[2003Sto]
Ball milling; pressing; sintering; SEM; TEM
αSiC ceramics doped by boron
[2004Ole]
Infiltration of a porous compact; EMPA
Si-B4C composites
[2004Pai]
Sintering at 2000–2100˚C; X-ray diffraction; SEM
αSiC ceramics doped by boron
[2004Rog]
Arc melting; annealing with quenching; X-ray diffraction; SEM; EMPA; wavelength-dispersive spectrometry (WDS); single crystal extraction
SiBn doped by carbon
[2004Ueh]
Hot pressing; X-ray diffraction; SEM; TEM
SiC-B4C composites
[2005Lee]
Vibration milling; hot pressing; X-ray diffraction
SiC-B4C composites
[2005Tka]
Compaction of a composite powder; sintering; SiC-B4C composites hot pressing up to 2150˚C; TEM; SEM; EMPA; X-ray diffraction
[2006Hay]
Uniaxial compaction at the pressure of 25 MPa; Reaction bonded ‘B4C’ with silicon additions sintering at 1900–2100˚C; infiltration; optical microscopy; SEM; energy-dispersive spectrometry (EDS); length-dispersive spectrometry (LDS); TEM; X-ray diffraction
[2006Shu]
Sintering at 1300–2000˚C; EMPA
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. Table 1 (continued) Reference
Method / Experimental Technique
Temperature / Composition / Phase Range Studied
[2007Ber]
CVD; Fourier transform infrared Spectroscopy (FTIR) in-situ analysis
B-C-Si ceramics
[2007Hwa]
Ball-miling; sintering at 1400˚C; polishing; X-ray powder diffraction; SEM; EDX
1400˚C; B-Si + 5 mass% B4C
[2007Lat]
Unbalanced radio-frequency (rf) magnetron sputtering; X-ray diffraction; SEM; TEM; Raman spectroscopy; Auger- electron spectroscopy
SiC and B4C monolayers; SiC/B4C multilayer coatings
[2007Mic]
Chemical vapor infiltration (CVI); Raman spectroscopy; X-ray diffraction; TEM; polarized light optical microscopy
1200–22˚C, B-C-Si microcomposites
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] (βB) < 2092
(C) gr < 3827 (sublimation point)
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
hR333 R 3m βB
a = 1093.30 c = 2382.52
hP4 P63/mmc C (graphite)
a = 246.12 c = 670.90
Comments/References [1993Wer] Dissolves up to 1.5 at.% C at 2098˚C [1996Kas] Dissolves up to 2.1 at.% Si at 2037˚C [1998Fri] [Mas2]
Dissolves up to 2.3 at.% B at 2382˚C [1996Kas]
(αSi) (I) cF8 < 1414 at 1.013 bar Fd 3m C (diamond)
a = 543.06
at 25˚C [Mas2, V-C2] Dissolves up to 1.1 at.% B at 1384.5˚ C [1998Fri] Dissolves up to 0.7 at.% C at 1270˚C [1998Fri]
(βSi) (II) > 9.624 bar
tI4 I41/amd βSn
a = 468.6 c = 258.5
at 25˚C [Mas2, V-C2]
(γSi) (III) > 16.208 bar
cI16 Im 3m γSi
a = 663.6
at 25˚C [Mas2, V-C2]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(δSi) hP4 > 16.208 - 1.013 bar P63/mmc αLa
Lattice Parameters [pm] a = 380 c = 628
Comments/References at 25˚C [Mas2, V-C2]
‘B4C’ < 2458
hR45 R 3m B13C2
from 8.9 at.% B at 2098˚C to 18.9 at.% B at 2382˚C [1996Kas] a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 labeled as B4+δC [2002Sei]; dissolves 4.3 at.% Si at 2000˚C [1972Kie] a = 560.7 in a Si10B80C10 (at.%) alloy, hotc = 1220.6 pressed [1965Mee2] a = 558 to 561 in B4C-SiC alloys hot-pressed and c = 1199 to 1211 sintered at 1800–2100˚C [1972Kie] a = 555.92; 556.37 reaction bonded ‘B4C’ [2006Hay] c = 1205.2; 1233.8 a = 563.8 in a Si0.5B13C2.5 specimen obtained c = 1231.5 by SHS [1994Gie] a = 560.03 in a Si6.2B38.1C55.7 (at.%) composite c = 1208.0 sintered at 2000˚C [1996Lih]
SiB3 < 1270
hR42 R 3m B6P
a = 631.9 c = 1271.3
[V-C2] 73 to 74 at.% B [1998Fri]
SiB6 < 1850
oP340 Pnnm SiB6
a = 1439.7 b = 1831.8 c = 991.1
[V-C2] 85.4 to 86.2 at.% B [1998Fri]
SiBn < 2037
hR36 R 3m βB hR339 R 3m FeB49
SiBnCx βSiC < 2824
[Mas2], n ≈ 23; 94.1 to 98.5 at.% B [1998Fri] a = 1101 c = 2390
[V-C2]
a = 1101.52 c = 2386.25
n = 30.4, x = 0.35, single crystal data [2004Rog]
cF8 a = 435.81 F 43m ZnS (sphalerite)
a = 435.3
a = 436
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50 at.% C [V-C2] labeled elsewhere as SiC of the 3C type dissolves 16 at.% B at 1800˚C [1972Kie] dissolves 3 at.% B (SHS obtained specimens) [1994Gie] in a Si30B20C50 (at.%) alloy, T = 1600˚C [1965Mee1]; Si50C50 (at.%) specimen sintered at T = 1950˚C in the SiC-B4C composite sintered at 1500˚C [2002Gun] Landolt‐Bo¨rnstein New Series IV/11E1
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B–C–Si
. Table 2 (continued) Phase/ Temperature Range [˚C] αSiC
Pearson Symbol/ Space Group/ Prototype hP12 P63mc αSiC
Lattice Parameters [pm] a = 308.07 c = 1511.74
Comments/References 50 at.% C [V-C2] labeled elsewhere as SiC of the 6H type, also known as moissanite; metastable; dissolves 0.2 mass% B at 2450–2500˚C [1970Sha] in the Si6.2B38.1C55.7 (at.%) composite sintered at 2000˚C [1996Lih]
a = 307.3 c = 1508 α’SiC
hP42 P 3m1 α’SiC
a = 307 c = 528.7
50 at.% C [V-C2] metastable; tentative structure
α’’SiC
hR48 R 3m α’’SiC
a = 308.2 c = 604.9
50 at.% C [V-C2] labeled elsewhere as SiC of the 15R type; metastable
α’’’SiC
hR186 R 3m α’’’SiC
a = 307 c = 2341.7
50 at.% C [V-C2] labeled elsewhere as SiC of the 15R type; tentative structure; metastable
τ1
o**
a = 850 b = 900 c = 490
in the Si6.2B38.1C55.7 (at.%) composite sintered at 2200˚C, together with the (C)gr, ‘B4C’ and βSiC phases [1996Lij]; metastable?
τ2
m**
a = 900 b = 590 c = 540 β = 119.3˚
in the Si6.2B38.1C55.7 (at.%) composite sintered at 2000˚C, together with the (C)gr, ‘B4C’ and βSiC phases; labelled as M [1996Lih]; metastable?
. Table 3 Invariant Equilibria Composition (at.%) Reaction L + (C)gr Ð ‘B4C’ + βSiC
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Type
Phase
U1
L
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B
C
Si
61.7
25.0
13.3
(C)gr
2.0
98.0
0.0
‘B4C’
81.0
19.0
0.0
βSiC
0.1
49.9
50.0
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B–C–Si
. Table 3 (continued) Composition (at.%) T [˚C]
Reaction L + (βB) Ð ‘B4C’ + SiBn
2005
Type
Phase
B
C
Si
U2
L
93.0
0.2
6.8
(βB)
98.4
1.4
0.2
‘B4C’
91.0
8.9
0.1
SiBn
96.8
0.0
3.2
L + SiBn Ð SiB6, ‘B4C’
1850
D1
L
62.1
0.0
37.8
L + βSiC Ð ‘B4C’ + (αSi)
1396
U3
L
L Ð (αSi) + SiB6, ‘B4C’
(αSi) + SiB6 Ð SiB3, ‘B4C’
1384
1198
D2
D3
4.7
0.0
95.3
βSiC
0.7
49.3
50.0
‘B4C’
85.8
14.0
0.2
(αSi)
0.6
0.0
99.4
L
8.1
0.0
91.9
(αSi)
0.9
0.0
99.1
SiB6
85.4
0.0
14.6
‘B4C’
86.3
12.6
1.1
(αSi)
85.5
0.0
14.5
SiB6
0.5
0.0
99.5
SiB3
73.8
0.0
26.2
‘B4C’
86.8
13.1
0.1
. Table 4 Thermodynamic Data of Reaction or Transformation Quantity, per mole of atoms [kJ, mol, K]
Reaction or Transformation
Temperature [˚C]
1/3{SiBC + SiC2 Ð BC2 + Si2C}
1507–2227˚C 36.8 ± 5.9
1/3{SiBC + Si Ð Si2C + B}
1507–2227˚C –37.6 ± 7.1
1/3{SiBC + Si Ð Si2 + B + (C)}
1507–2227˚C 12.5 ± 6.7
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Comments [1964Ver] Knudsen effusion. The state of the species is gas.
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. Table 5 Thermodynamic Properties of Single Phases Phase
Temperature Range [˚C]
Property, per mole of atoms [J, mol, K]
SiBC
1727
–{(GT0 – H˚0)/T} = 298.9
1827 1927 2027
–{(G
T
– H˚0)/T} = 301.4
–{(G
T
– H˚0)/T} = 303.9
–{(G
T
– H˚0)/T} = 306.4
0 0 0
Comments [1964Ver] Knudsen effusion. The state of the species is gas.
ΔH˚0 (at.) = 1032460 ± 25100
. Table 6 Vapor Pressure Measurements Phase(s) SiBC
Temperature [˚C]
Pressure [bar]
Comments
1976
log10 (pSiBC) = –6.62
[1964Ver] Knudsen effusion.
2032
log10 (pSiBC) = –6.34
The state of the species is gas.
2071
log10 (pSiBC) = –6.11
2028
log10 (pSiBC) = –6.77
1877
log10 (pSiBC) = –6.77
1893
log10 (pSiBC) = –6.46
1810
log10 (pSiBC) = –6.66
. Table 7 Investigations of the B-C-Si Materials Properties Reference [1960Por]
Method / Experimental Technique
Type of Property
Hydrostatic measurements; pycnometry; Specific gravity; thermo-emf; mechanical and electrical properties tests microhardness; electrical resistivity
[1965Mee1] Microhardness measurements
Microhardness
[1965Mee2] Microhardness and electrical resistivity measurements
Microhardness; specific electrical resistivity
[1967Dok]
PMT-3 microhardness tests
Microhardness
[1969Nie]
Mechanical properties tests
Microhardness
[1971Kal] [1975Kal]
PMT-3 microhardness tests; electrical resistivity and density measurements
Microhardness; specific electrical resistivity; density
[1972Gug] [1972Kie]
Vickers microhardness tests
Microhardness
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B–C–Si
. Table 7 (continued) Reference [1975Bin]
Method / Experimental Technique
Type of Property
Sintering under pressure
Densification behavior
[1978Ekb]
Strength, Vickers microhardness tests
Strength, microhardness
[1979Hon]
Microhardness tests (Leitz microhardness Microhardness tester, Knoop diamond indenter)
[1979Pan]
Vickers microhardness tests
Microhardness
[1983Wal]
Strength, Vickers microhardness tests
Flexural strength, microhardness
[1987Bou]
Microhardness tests (Leitz microhardness Microhardness; Young’s modulus; tester, Knoop diamond indenter); ductility dynamic method (Grindo Sonic); point method
[1994Wer]
Optical absorption; infrared phonon spectroscopy; Raman spectroscopy
Absorption coefficient; reflectivity; absorption index; Raman intensity
[1995Adr]
Electron nuclear double resonance
Boron acceptor behavior
[1995Mas]
PMT-3 mechanical tests
Microhardness; microbrittleness; microstrength; crack resistance; density
[1996Bar]
EPR spectroscopy (X-band and Q-band spectrometer); photoluminescence (PL) spectroscopy
EPR and PL spectra
[1996Her]
Double crystal diffraction topography (DCDT); grazing incidence X-ray scattering (GIXS); atomic force microscopy (AFM); SEM
Residual stress; structure of films
[1997Gor]
Electrical resistivity and strength measurements
Electrical resistivity; bending strength
[1997Shi]
Strength, electrical resistance, thermal emf measurements
Compressive strength, specific electrical resistance, the coefficient of thermal emf, volt-ampere characteristic; density
[1998Ang]
Thermogravimetric analysis (TGA) (Seiko SSC 5200 apparatus); thermomechanical analysis (TMA) (Perkin-Emler thermal analyzer); density and porosity tests; Knoop hardness measurements; 4 point bend test
Density; hardness; flexural strength; flexural modulus; porosity
[1998Bad]
Vickers microhardness tests
Microhardness
[1998Bar]
EPR spectroscopy
EPR spectra
[1998Dui]
EPR spectroscopy; electron-nuclear EPR, ENDOR, ESEEM spectra double resonance (ENDOR); electron spin echo envelope modulation (ESEEM)
[1998Li]
Four-probe method; ac method; laser-flash technique
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Electrical conductivity; the Seebeck coefficient; thermal conductivity Landolt‐Bo¨rnstein New Series IV/11E1
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21
. Table 7 (continued) Reference
Method / Experimental Technique
Type of Property
[1999Li]
Four-probe method; ac method; laserflash technique; Archimede’s method
[2000Gri]
Indentation and SENVB methods; three- Strength; Young’s modulus; hardness; point bending test; measuring the rate of crack resistance; thermal stress propagation of an acoustic signal; Vickers hardness tests
[2001Ade]
Current and voltage measurements (Keithley 2400-C SourceMeter and TestPoint software)
[2001Mag]
Archimede’s method; four-point bending Density; flexural strength; Weibull test; hardness tests (Leitz hardness tester) modulus; hardness; fracture toughness
[2002Dar]
Tensile creep tests (PSB 100 Hydropuls Schenck machine); SEM; TEM
Creep curves
[2002Gun]
DC four-probe method; laser-flash technique; Vickers microhardness tests
Electrical conductivity; thermal conductivity; microhardness
[2002Shi]
Microhardness and electrical conductivity Microhardness; electrical conductivity measurements
[2003Shu]
Vickers indentation method (PMT-3 tester)
Microhardness; hardness; fracture toughness
[2004Pai]
DC four-probe method; laser-flash technique
Electrical conductivity; thermoelectric power; Seebeck coefficient; thermal conductivity; thermal diffusivity; density
[2004Ueh]
Vickers indentation method; voltage measurements
Hardness; fracture toughness; Seebeck coefficient
[2005Lee]
Vickers indentation method; wear tests (reciprocating ball-on-disk tester TE77)
Hardness; wear resistance
[2005Tka]
Mechanical properties tests
Microhardness; cracking resistance; bending strength; density
[2006Shu]
Mechanical tests
Strength
[2006Ye]
Molecular dynamics simulation; mechanical tests
Thermal stability; creep resistance
[2007Hwa]
Vickers hardness measurements
Hardness
[2007Lat]
Vickers indentation method; scratch test Hardness; residual stress; critical load of (CSEM-Revetest) Failure; Young’s modulus
[2007Mic]
Tensile tests; strain measurements (compliance calibration technique);
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Electrical conductivity; the Seebeck coefficient; thermal conductivity; density
Current-voltage curves
Young’s modulus; thermal expansion; microcomposite’s failure stress; fracture behavior
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B–C–Si
. Fig. 1 B-C-Si. The B-Si phase diagram
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. Fig. 2 B-C-Si. The C-Si phase diagram
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B–C–Si
. Fig. 3 B-C-Si. Reaction scheme
20
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. Fig. 4 B-C-Si. Calculated liquidus surface projection
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B–C–Si
. Fig. 5 B-C-Si. Calculated solidus surface projection
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. Fig. 6 B-C-Si. Calculated isothermal section at 2327˚C
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B–C–Si
. Fig. 7 B-C-Si. Calculated isothermal section at 2227˚C
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. Fig. 8 B-C-Si. Calculated partial isothermal section at 1927˚C
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B–C–Si
. Fig. 9 B-C-Si. Calculated isothermal section at 1127˚C
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. Fig. 10 B-C-Si. Calculated temperature - composition section SiC-B4.6C
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B–C–Si
. Fig. 11 B-C-Si. Temperature - composition section SiC-B
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. Fig. 12 B-C-Si. Calculated temperature - composition section at 80 at.% B
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B–C–Si
. Fig. 13 B-C-Si. Calculated temperature - composition section Si-B4.18C
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21
References [1955Sam]
[1960Por]
[1964Sch]
[1964Sec]
[1964Ver] [1965Mee1]
[1965Mee2]
[1965Sec] [1967Dok]
[1969Nie]
[1969Sha] [1970Sha] [1971Kal]
[1972Gug]
[1972Kie]
[1975Bin] [1975Kal]
Samsonov, G.V., Petrash, Ye.V., “Some Physical-Chemical Properties of Titanium Boride and Nitride Alloys” (in Russian), Metalloved. Term. Obrab. Met., (4), 19–24 (1955) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 10) Portnoi, K.I., Samsonov, G.V., Solonnikova, L.A., “Alloys of the Boron-Silicon-Carbon System”, Russ. J. Inorg. Chem. (Engl. Transl.), 5(9), 988–993 (1960), translated from Zh. Neorg. Khim. SSSR, 5 (9), 2032–2041 (1960) (Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., Phys. Prop., 25) Schaffer, P.T.B., Hannam, A.L., “Comments on Phase Equilibria in the System Boron Carbide-Silicon Carbide by D.R. Secrist”, J. Am. Ceram. Soc., 47(11), 594–595 (1964) (Crys. Structure, Phase Relations, Assessment, 9) Secrist, D.R., “Phase Equilibria in the System Boron Carbide-Silicon Carbide”, J. Am. Ceram. Soc., 47(3), 127–130 (1964) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 3) Verhaegen, G., Stafford, F.E., Drowart, J., “Mass Spectrometric Study of the Systems Boron-Carbon and Boron-Carbon-Silicon”, J. Chem. Phys., 40(6), 1622–1628 (1964) (Thermodyn., Experimental, 31) Meerson, G.A., Kiparisov, S.S., Gurevich, M.A., Fen-Sian, D., “Obtaining and Investigation of the Properties of Solid Solutions on a Pseudobinary Section SiC-BC by the Method of Joint from a Gaseous Phase” (in Russian), Poroshk. Metall., (2), 15–21 (1965) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 6) Meerson, G.A., Kiparisov, S.S., Gurevich, M.A., Fen-Sian, D., “Investigation of Conditions for Obtaining of Solid Alloys of the Pseudobinary System B4C-B4Si” (in Russian), Poroshk. Metall., (3), 62–68 (1965) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Kinetics, Mechan. Prop., 9) Secrist, D.R., “Reply to Comments on “Phase Equilibria in the System Boron Carbide-Silicon Carbide”, J. Am. Ceram. Soc., 48(4), 215–215 (1965) (Crys. Structure, Phase Relations, Assessment, 5) Dokukina, I.V., Kalinina, A.A., Sokhor, M.I., Shamrai, F.I., “On the Problem of Chemical Compounds in the System Silicon-Boron-Carbon” (in Russian), Izv. Akad. Nauk SSSR, Neorgan. Mater., 3(4), 630–637 (1967) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 8) Niemyski, T., Appenheimer, S., Panczyk, J., Badzian, A., “Vapor Phase Crystallization of B-Si-C Phase”, J. Cryst. Growth, 5(5), 401–404 (1969) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., 6) Shaffer, P.T.B., “The SiC Phase in the System SiC-B4C-C”, Mater. Res. Bull., 4(3), 213–220 (1969) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 12) Shaffer, P.T.B., “Solubility of Boron in Alpha Silicon Carbide”, Mater. Res. Bull., 5, 519–521 (1970) (Crys. Structure, Morphology, Phase Relations, Experimental, 4) Kalinina, A.A., Sokhor, M.I., Shamrai, F.I., “Investigation of the Alloys of the System Si-B-C” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 7(5), 778–785 (1971) (Crys. Structure, Morphology, Phase Relations, Experimental, Phys. Prop., 15) Gugel, E., Kieffer, R., Leimer, G., Ettmayer, P., “Investigation in the Ternary System Boron-CarbonSilicon”, Nat. Bur. Stand. Spec. Pub., Sol. State Chem., Proc. 5th Mater. Res. Symp., 364, 505–513 (1972) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *) as quoted by [1996Kas] Kieffer, R., Gugel, E., Leimer, G., Ettmayer, P., “Investigations in the System Boron-Carbon-Silicon“ (in German), Berich. Deut. Keram. Gesel., 49(2), 41–46 (1972) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, Mechan. Prop., *, 39) Bind, J.M., Biggers, J.V., “Hot-Pressing of Silicon Carbide with 1% Boron Carbide Addition”, J. Am. Ceram. Soc., 58(7-8), 304–306 (1975) (Crys. Structure, Morphology, Experimental, Phys. Prop., 19) Kalinina, A.A., Sokhor, M.I., “Phase Composition and Some Properties of Silicon-Boron-Carbon System Alloys Adjacent to the Section Slicon Carbide-Boron” (in Russian), Vysokotemperatur. Karbidy, Naukova Dumka, Kiev, 96–99 (1975) (Crys. Structure, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., Phys. Prop., 6)
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21 [1977Gor]
[1977Saf]
[1978Ekb]
[1979Hon]
[1979Pan]
[1982Doe]
[1983Sch]
[1983Wal]
[1986Lil]
[1986Mor]
[1987Bou]
[1990Ase]
[1990Tel]
[1991Fri]
[1993Wer]
[1994Gie]
[1994Gou1]
B–C–Si Gorban’, I.S., Krokhmal, A.P., “Structure of the Spectrum of Excitons Bound to Neutral Boron Atoms in Silicon Carbide”, Sov. Phys. - Solid State (Engl. Transl.), 19(5), 733–755 (1977) (Morphology, Experimental, Theory, Electronic Structure, 8) Safaraliev, G.K., Tairov, Yu.M., Tsvetkov, V.F., “Thermodynamic Analysis of Solubility and Transformation Coefficient of Boron in Silicon Carbide” in “Svoistva Legirovannykh Poluprovodnikov” (in Russian), Nauka, Moscow, 53–58 (1977) (Thermodyn., Calculation, Experimental, 6) Ekbom, L.B., “Effect of Increased Boron Content on the Sintering Behaviour and Mechanical Properties of Boron Carbide”, unknoun source, a copy is available, 183–189 (1978) (Morphology, Experimental, Mechan. Prop., 6) Hong, J.-D., Spear, K.E., Stubican, V.S., “Directional Solidification of SiC-B4C Eutectic: Growth and Some Properties”, Mater. Res. Bull., 14(6), 775–783 (1979) (Morphology, Experimental, Kinetics, Mechan. Prop., 13) Panasyuk, A.D., Oreshkin, V.D., Maslennikova, V.R., “Kinetics of the Reactions of Boron Carbide with Liquid Aluminium, Silicon, Nickel and Iron”, Sov. Powder Metall. Met. Ceram., 199(7), 487–490 (1979), translated from Poroshk. Metall., 199(7), 79–83 (1979) (Morphology, Experimental, Kinetics, Mechan. Prop., 9) Doerner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, University Stuttgart, Institut fuer Metallkunde, 1–194 (1982) (Phase Diagram, Phase Relations, Thermodyn., Calculation, *, 126) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Phase Diagram, Phase Relations, Review, 154) Walker, B.E., Rice, R.W., Becher, P.F., Bender, B.A., Coblenz, W.S., “Preparation and Properties of Monolithic and Composite Ceramics Produced by Polymer Pyrolysis”, Amer. Ceram. Soc. Bul., 62(8), 916–923 (1983) (Crys. Structure, Morphology, Experimental, Phys. Prop., 9) Lilov, S.K., “Thermodynamic Study of the Solubility Process of Boron in Silicon Carbide, Grown from the Vapour Phase”, J. Phys. Chem. Solids, 47(3), 245–250 (1986) (Thermodyn., Calculation, Review, Electronic Structure, 16) More, K.L., Carter, C.H., Bentley J. Jr., Wadlin, W.H., LaVanier L., Davis, R.F., “Occurrence and Distribution of Boron-Containing Phases in Sintered α-Silicon Carbide”, J. Am. Ceram. Soc., 69(9), 695–698 (1986) (Crys. Structure, Morphology, Phase Relations, Experimental, 17) Bougoin, M., The´venot, F., Dubois, J., Fantozzi, G., “Synthesis and Thermomechanical Properties of the Ceramic Dense Composites Boron Carbide-Silicium Carbide” (in French), J. Less-Common Met., 132(2), 209–228 (1987) (Morphology, Experimental, Phys. Prop., 55) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides”, Freer R. (Ed.), “The Physics and Chemistry of Carbides, Nitrides and Borides”, Kluwer Academic Publishers, Dordrecht 97–111 (1990) (Crys. Structure, Experimental, Review, 14) Telle, R., “Structure and Properties of Si-Doped Boron Carbide” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Kluwer Academic Publishers, Dordrecht (Germany), 249 (1990) (Crys. Structure, Phase Diagram, Phase Relations, Experimental) as quoted by [2002Sei] Fries, S., Lim, S-K., Lukas, H.L., “Optimisation of the Al-C, Al-Sn, B-Si, Mg-Zn, Ci-C and Zn-Sn Binary Systems as well as of the Al-Sn-Zn Ternary System” in “Leuven Proceedings, COST 507, New Light Alloys, Part C Thermodynamic Evaluation and Calculation”, Effenberg, G. (Ed.)/, /Commission of the European Communities, Part C, D9–1–5 (1991) (Phase Diagram, Thermodyn., Assessment, Calculation, 0) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of Carbondoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Gierlotka, S., Oleksyn, O., Palosz, B., “Si-C-B System Synthesized by SHS Method: Phase Diagram and Crystal Structure Evolution”, Mater. Sci. Forum, 166-169, 529–538 (1994) (Crys. Structure, Phase Relations, Experimental, 11) Goujard, S., Vandenbulcke, L., Bernard, C., “Thermodynamic Study of the Chemical Vapour Deposition in the Si-B-C-H-Cl System”, Calphad, 18(4), 369–385 (1994) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, *, 35)
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[1994Oga] [1994Wer] [1995Adr]
[1995Gou]
[1995Lim]
[1995Mas]
[1996Bar]
[1996Gro] [1996Her]
[1996Kas]
[1996Lih]
[1996Lij]
[1996Sin] [1997Gor]
[1997Sch]
[1997Shi]
[1998Ang]
21
Goujard, S., Vandenbulcke, L., Bernard, C., Blondiaux, G., Debrun, J.L., “Thermodynamic and Experimental Study of the Chemical Vapor Codeposition in the Silicon-Boron-Carbon System at 1400 K”, J. Electrochem. Soc., 141(2), 452–460 (1994) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, *, 48) Ogawa, I., “Morphologies of SiC Growing in C/SiC/B4C Composites” (in Japanese), J. Ceram. Soc. Jpn., 102(8), 802–803 (1994) (Morphology, Experimental, 4) Werheit, H., Kuhlmann, U., Laux, M., “Solid Solutions of Silicon in Boron-Carbide-Type Crystals”, J. Alloys Compd., 209, 181–187 (1994) (Crys. Structure, Phase Relations, Experimental, Phys. Prop., 14) Adrian, F.J., Greulich-Weber, S., Spaeth, J.-M., “Further Evidence for the B–Si-C+ Model of the Boron Acceptor in 6H Silicon Carbide from a Theoretical Analysis of the Hyperfine Interactions”, Solid State Commun., 94(1), 41–44 (1995) (Morphology, Experimental, Electronic Structure, Phys. Prop., 9) Goujard, S., Vandenbulcke, L., Bernard, C., “On the Chemical Vapor Deposition of Si/B/C-Based Coatings in Various Conditions of Supersaturation”, J. Eur. Ceram. Soc., 15(6), 551–561 (1995) (Thermodyn., Calculation, Experimental, 19) Lim, S.K, Lukas, H.L., “Thermodynamic Optimisation of the System B-C-Si and its Boundary Systems” in “Hochleistungskeramik, Herstellung, Aufbau und Eigenschaften”, Deutsche Forschungsgemeinschaft, Petzow, G., Tobolski, J., Telle, R. (Eds.), VCH, Weinheim, 605–616, (Phase Diagram, Thermodyn., Assessment, Calculation, 57) Maslennikova, V.R., Belkina, A.A., Panasyuk, A.D., Struk, L.I., Smirnov, V.P., “Effect of High Pressure on the Structure and Properties of Materials Based on Boron Carbide”, Powder Metall. Met. Ceram., 34(9,10), 491–495 (1995), translated from Poroshk. Metall., (9/10), 3–7 (1995) (Morphology, Experimental, Mechan. Prop., Phys. Prop., 6) Baranov, P.G., Mokhov, E.N., “Electron Paramagnetic Resonance of Deep Boron in Silicon Carbide”, Semicond. Sci. Technol., 11(4), 489–494 (1996) (Crys. Structure, Experimental, Theory, Electronic Structure, Magn. Prop., Optical Prop., 25) Groebner, J., Lukas, H.L., Aldinger, F., “Thermodynamic Calculation of the Ternary System Al-Si-C”, Calphad, 20(2), 247–254 (1996) (Phase Diagram, Thermodyn., Calculation, 37) Hershberger, J., Ying, T., Kustas, F., Fehrenbacher, L., Yalisove, S.M., Bilello, J.C., “Residual Stress, Atomic Structure, and Growth Morphology in B4C/SiC Multilayer Coatings”, Surf. Coat. Technol., 86-87(1–3), 237–242 (1996) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 23) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institut, Dissertation, Stuttgart, 1–225 (1996) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calculation, Experimental, Review, #, 170) Lihua, Z., Lijun, W., Qizhong, H., Qiaoqin, Y., Shaoli, L., “Structure of C-B4C-SiC Composites with Silicon Additive”, J. Mater. Sci. Lett., 15(4), 353 (1996) (Crys. Structure, Morphology, Phase Relations, Experimental, *, 10) Lijun, W., Qizhong, H., Qiaoqin, Y., Lihu, Z., Zhongyu, X., “Effect of Sintering Temperature on Structure of C-B4C-SiC Composites with Silicon Additive”, Scr. Mater., 35(1), 123–127 (1996) (Crys. Structure, Morphology, Experimental, *, 12) Singh, M., “Thermodynamic Analysis for the Combustion Synthesis of SiC-B4C Composites”, Scr. Mater., 34(6), 923–927 (1996) (Thermodyn., Review, Theory, 12) Go´rny, G., Raczka, M., Stobierski, L., Wojnar, L., Pampuch, R., “Microstructure-Property Relationship in B4C-βSiC Materials”, Solid State Ionics, 101-103 (Pt.2), 953–958 (1997) (Crys. Structure, Morphology, Experimental, Electr. Prop., Mechan. Prop., 4) Schmidt, J., Matsumoto, T., Poluektov, O.G., Arnold, A., Ikoma, T., Baranov, P.G., Mokhov, E.N., “High-Frequency EPR Studies of Shallow and Deep Boron Acceptors in 6H-SiC”, Mater. Sci. Forum, Defects in Semiconductors-19, 258-263, 703–708 (1997) (Crys. Structure, Experimental, Theory, Electronic Structure, 33) Shipilova, L.A., Petrovskii, V.Ya., Chugunova, S.I., “Structure Formation and Electrophysical Properties of a Silicon Carbide-Boron Carbide Sintered Composite”, Powder Metall. Met. Ceram., 36(11-12), 652–656 (1997), translated from Poroshk. Metall., (11–12), 97–101 (1997) (Morphology, Experimental, Electr. Prop., Mechan. Prop., 6) Angelovici, M.M., Bryant, R.G., Northam, G.B., Roberts, A.S., Jr., “Carbon/Ceramic Microcomposites, Preparation and Properties”, Mater. Lett., 36(5-6), 254–265 (1998) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 12)
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[1998Bar]
[1998Dui]
[1998Fri]
[1998Gro]
[1998Hof]
[1998Kas]
[1998Li]
[1999Guo]
[1999Li]
[1999Xin]
[2000Gri]
[2001Ade]
[2001And]
[2001Mag]
[2002Dar] [2002Dat]
[2002Gun]
B–C–Si Badzian, A., Badzian, T., Drawl, W.D., Roy, R., “Synthesis and Properties of the B-C-Si and Si-N-C Hard Materials”, Diamond Related Compd., (7), 1519–1524 (1998) (Morphology, Experimental, Mechan. Prop.) as quoted by [2001And] Baranov, P.G., Il‘in, I.V., Mokhov, E.N., “Electron Paramagnetic Resonance of Deep Boron Acceptors in 4H-SiC and 3C-SiC Crystals”, Phys. Solid State, 40(1), 31–34 (1998) (Crys. Structure, Experimental, Electronic Structure, Magn. Prop., 12) von Duijn-Arnold, A., Ikoma, T., Poluektov, O.G., Baranov, P.G., Mokhov, E.N., Schmidt, J., “Electronic Structure of the Deep Boron Acceptor in Boron-Doped 6H-SiC”, Phys. Rev. B, 57(3), 1607–1619 (1998) (Crys. Structure, Experimental, Electronic Structure, Magn. Prop., 42) Fries, S., Lukas, H.L., “System B-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 126–128 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, 1) Groebner, J., Lukas, H.L., Aldinger, F., “System C-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 132–133 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, 1) Hofstaetter, A., Meyer, B.K., Scharmann, A., Baranov, P.G., Ilyin, I.V., Mokhov, E.N., “X-Band ENDOR of Boron and Beryllium Acceptors in Silicon Carbide”, Mater. Sci. Forum, Silicon Carbide, III-Nitrides and Related Materials, 264-268, 595–598 (1998) (Crys. Structure, Experimental, Electronic Structure, 6) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Li, J., Goto, T., Hirai, T., “Microstructure and Thermoelectric Properties of B4C-SiBn-Si Composites Prepared by Arc Melting”, J. Ceram. Soc. Jpn., 106(2), 194–197 (1998) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 16) Guo, Q., Song, J., Liu, L., Zhang, B., “Relationship Between Oxidation Resistance and Structure of B4C-SiC/C Composites with Self-Healing Properties”, Carbon, 37(1), 33–40 (1999) (Morphology, Phase Relations, Experimental, Kinetics, 22) Li, J., Goto, T., Hirai, T., “Thermoelectric Properties of B4C-SiBn (n = 4, 6, 14) In-situ Composites”, Mater. Trans., JIM, 40(4), 314–319 (1999) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 25) Xinmin, M., Cewen, N., Kefeng, C., “Structural Characteristics and Quantum Chemistry Calculation of Si-Doped Boron Carbides” in “Multiscale Modelling of Materials”, Proc. Symp. Mater. Res. Soc., Warrendale, PA, USA, 579–584 (1999) (Crys. Structure, Calculation, Electronic Structure, 12) Grigor’ev, O.N., Gogotsi, G.A., Gogotsi, Yu.G., Subbotin, V.I., Brodnikovskii, N.P., “Synthesis and Properties of Ceramics in the SiC-B4C-MeB2 System”, Powder Metall. Met. Ceram., 39(5-6), 239–250 (2000) (Morphology, Experimental, Mechan. Prop., 15) Adenwalla, S., Welsch, P., Harken, A., Brand, J.I., Sezer, A., Robertson, B.W., “Boron Carbide/n-Silicon Carbide Heterojunction Diodes”, Appl. Phys. Lett., 79(26), 4357–4359 (2001) (Crys. Structure, Experimental, Electr. Prop., 7) Andrievski, R.A., “Superhard Materials Based on Nanostructured High-Melting Point Compounds: Achievements and Perspectives”, Iner. J. Ref. Met. Hard Mater., 19(4-6), 447–452 (2001) (Morphology, Review, Mechan. Prop., 59) Magnani, G., Beltrami, G., Minoccari, G.L., Pilotti, L., “Pressureless Sintering and Properties of αSiCB4C Composite”, J. Eur. Ceram. Soc., 21(5), 633–638 (2001) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 19) Darzens, S., Farizy, G., Vicens, J., Chermant, J.-L., “Microstructure and Creep of SiCf-SiBC”, Key Eng. Mater., 206-213(2), 989–992 (2002) (Morphology, Experimental, Mechan. Prop., 18) Datta, M.S., Bandyopadhyay, A.K., Chaudhuri, B., “Sintering of Nano Crystalline α Silicon Carbide by Doping with Boron Carbide”, Bull. Mater. Sci. (India), 25(3), 181–189 (2002) (Morphology, Experimental, Kinetics, Interface Phenomena, 22) Gunjishima, I., Akashi, T., Goto, T., “Characterization of Directionally Solidified B4C-SiC Composites Prepared by a Floating Zone Method”, Mater. Trans., 43(9), 2309–2315 (2002) (Crys. Structure, Morphology, Experimental, Electr. Prop., Kinetics, Mechan. Prop., Phys. Prop., 24)
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Rurali, R., Godignon, P., Rebollo, J., Ordejon, P., Hernandez, E., “Theoretical Evidence for the KickOut Mechanism for B Diffusion in SiC”, Appl. Phys. Lett., 81(16), 2989–2991 (2002) (Morphology, Review, Theory, Electronic Structure, Interface Phenomena, Kinetics, 20) Seifert, H.J., Aldinger, F., “Phase Equilibria in the Si-B-C-N System”, Struct. Bonding, 101, 1–58 (2002) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Calculation, Review, #, 275) Shipilova, L.A., Petrovskii, V.Ya., “Structure Formation, Electrophysical and Mechanical Properties of an Electrically Conducting Ceramic Composite Based on Silicon and Boron Carbides”, Powder Metall. Met. Ceram., 41(3-4), 147–149 (2002), translated from Poroshk. Metall., 3-4(424), 41–44 (2002) (Morphology, Experimental, Electr. Prop., Mechan. Prop., 3) Akashi, T., Gunjishima, I., Goto, T., “Characterization of Directionally Solidified B4C-TiB2 and B4CSiC Eutectic Composites Prepared by Floating-Zone Metod”, Key Eng. Mater., 247, 209–212 (2003) (Crys. Structure, Morphology, Experimental, 3) Fan, Z., Wei, T., Shi, J., Zai, G., Song, J., Liu, L., Li, J., Chen, J., “New Route for Preparation of SiCB4C/C Composite with Excellent Oxidation Resistance up to 1400˚C”, J. Mater. Sci. Lett., 22(3), 213–215 (2003) (Morphology, Experimental, Kinetics, 10) Narushima, T., Goto, T., Maruyama, M., Arashi, H., Iguchi, Y., “Oxidation of Boron Carbide-Silicon Carbide Composite at 1073 to 1773 K”, Mater. Trans., JIM, 44(3), 401–406 (2003) (Morphology, Phase Relations, Experimental, Kinetics, 33) Shulzhenko, A.A., Stratiychuk, D.A., Gargin, V.G., Belyavina, N.N., “Preparation and PhysicoMechanical Properties of B-C-Si-Based Cermet”, J. Superhard Mater., 25(5), 74–76 (2003), translated from Sverkhtverd. Mater. (Ukraine), (5), 82–84 (2003) (Morphology, Phase Relations, Experimental, Mechan. Prop., 9) cited from abstract Stobierski, L., Gubernat, A., “Sintering of Silicon Carbide. II. Effect of Boron”, Ceram. Intern., 29(4), 355–361 (2003) (Phase Relations, Morphology, Experimental, Kinetics, Transport Phenomena, 13) Epelbaum, B.M., Gurzhiyants, P.A., Herro, Z., Bickermann, M., Winnacker, A., “Flux Growth of SiC Crystals from Eutectic Melt SiC-B4C”, Mater. Sci. Forum, 457-460, 119–122 (2004) (Morphology, Experimental, 10) Oleynik, G.S., Shulzhenko, A.A., Stratiychuk, D.A., Gargin, V.G., Vereshchaka, V.M., “Special Features of the Microstructure of a Composite with an Increased Fracture Toughness Produced in the B-C-Si System at High Pressure”, J. Superhard Mater., 26(4), 13–24 (2004), translated from Sverkhtverd. Mater. (Ukraine), 26(4), 16–28 (2004) (Morphology, Experimental, Mechan. Prop., 45) cited from abstract Pai, Ch.-H., “Thermoelectric Properties of Boron Compound-Doped α-SiC Ceramics”, J. Ceram. Soc. Jpn., 112(2), 88–94 (2004) (Crys. Structure, Morphology, Experimental, Electr. Prop., 12) Roger, J., Babizhetskyy, V., Halet, J.-F., Gue´rin, R., “Boron-Silicon Solid Solution: Synthesis and Crystal Structure of a Carbon-Doped Boron-rich SiBn (n 30) Compound”, J. Solid State Chem., 177(11), 4167–4174 (2004) (Crys. Structure, Experimental, 27) Uehara, M., Shiraishi, R., Nogami, A., Enomoto, N., Hojo, J., “SiC-B4C Composites for Synergistic Enhancement of Thermoelectric Property”, J. Eur. Ceram. Soc., 24, 409–412 (2004) (Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., 5) Lee, K.S., Han, I.S., Chung, Y.H., Woo, S.K., Lee, S.W., “Hardness and Wear Resistance of Reaction Bonded SiC-B4C Composite”, Mater. Sci. Forum, 486-487, 245–246 (2005) (Phase Relations, Experimental, Mechan. Prop., 7) Tkachenko, Yu, G., Britun, V.F., Prilutskii, E.V., Yurchenko, D.Z., Bovkun, G.A., “Structure and Properties of B4C-SiC Composites”, Powder Metall. Met. Ceram., 44(3-4), 196–201 (2005) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 4) Fanton, M.A., Cavalero, R.L., Ray, R.G., Weiland, B.E., Everson, W., Snyder, D., Gamble, R., Oslosky, E., “Growth of SiC Boules with Low Boron Concentration”, Mater. Sci. Forum, 527-529, 47–50 (2006) (Morphology, Experimental, Kinetics, 21) Hayun, S., Frage, N., Dariel, M.P., “The Morphology of Ceramic Phases in BxC-SiC-Si Infiltrated Composites”, J. Solid State Chem., 179, 2875–2879 (2006) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, 11) Shulzhenko, A.A., Stratiychuk, D.A., Oleinik, G.S., Vereshchaka, V.M., “Structural Transformations in the Formation of a Superhard Composite in the B-C-Si System”, J. Superhard Mater., 28(4), 11–26 (2006) (Morphology, Phase Relations, Experimental, Interface Phenomena, Mechan. Prop., 26) cited from abstract
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B–C–Si Ye, Y.-J., Zhang, L.-T., Cheng, L.-F., Xu, Y.-D., “Diffusion Behavior in Amorphous Si-B-C System by Molecular Dynamics Simulation”, J. Inorg. Mater., 21(4), 843–847 (2006) (Morphology, Experimental, Interface Phenomena, Phys. Prop., 13) cited from abstract Berjonneau, J., Langlais, F., Chollon, G., “Understanding the CVD Process of (Si)-B-C Ceramics Through FTIR Spectroscopy Gas Phase Analysis”, Surf. Coat. Technol., 201, 7273–7285 (2007) (Phase Relations, Experimental, Kinetics, 39) Hwang, G.-C., Mastushita, J., Lee, J.-J., “Preparation of Si-B-C System Powder Using Silicon, Boron, and Boron Carbide”, Mater. Sci. Forum, 544-545, 933–936 (2007) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 6) Lattemann, M., Ulrich, S., “Investigation of Structure and Mechanical Properties of Magnetron Sputtered Monolayer and Multilayer Coatings in the Ternary System Si-B-C”, Surf. Coat. Technol., 201(9-11), 5564–5569 (2007) (Morphology, Phase Relations, Experimental, Mechan. Prop., 60) Michaux, A., Sauder, C., Camus, G., Pailler, R., “Young’s Modulus, Thermal Expansion Coefficient and Fracture Behavior of Selected Si-B-C Based Carbides in the 20–1200˚C Temperature Range as Derived from the Behavior of Carbon Fiber Reinforced Microcomposites”, J. Eur. Ceram. Soc., 27(12), 3551–3560 (2007) (Crys. Structure, Calculation, Experimental, Mechan. Prop., Phys. Prop., 21) “Scientific Group Thermodata Europe”, Grenoble Campus, 1001 Avenue Centrale, BP66, F-38402 Saint Martin D’Heres France (Thermodyn., Review) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Tantalum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction Although borides and carbides of tantalum are among compounds with the highest melting temperatures (Tm, TaC = 3985˚C, Tm, TaB = 3090˚C), high temperature phase relations in the B-C-Ta system have not yet been explored in detail. Studies on the interaction between tantalum carbides, tantalum borides, B4C and carbon on samples hot-pressed at 1500 < T < 2900˚C [1952Gla, 1955Bre] were used to define phase relations in the B-C-Ta system particularly for the regions TaC-TaB2-B-C and Ta2C-Ta2B [1955Bre]. Further studies of phase equilibria were presented by [1963Rud, 1965Lev, 1976Ord, 1987Ord] and refer to an isothermal section at 1750˚C [1963Rud] and three eutectic quasibinaries: TaB2-C [1965Lev], TaB2-TaC1–x [1976Ord] and TaB2-‘B4C’ [1987Ord]. These results were experimentally accomplished employing X-ray powder diffractometry [1952Gla, 1955Bre, 1963Rud, 1965Lev, 1976Ord, 1987Ord], micrographic analysis [1963Rud, 1965Lev, 1976Ord, 1987Ord], pyrometric melting point measurements [1965Lev, 1976Ord, 1987Ord] and chemical analysis [1963Rud, 1976Ord]. The most relevant data on the topology of the B-C-Ta system were compiled by [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog]. Experimental details for all investigations in the B-C-Ta system are summarized in Table 1.
Binary Systems The B-Ta binary system is essentially taken from [1992Rog], however, phase relations in the Ta rich part have been revised recently [2006Cha]: the composition of the eutectic, l Ð (Ta) + Ta2B, has been found at 18 at.% B and the composition of the liquid for the peritectic, l + TaB Ð Ta2B, has been located at 22.5 at.% B. Furthermore the decomposition of Ta2B was found to occur between 1900 and 1950˚C. Due to reinvestigation of C-Ta phase relations from 1700 to 2300˚C by means of diffusion couples, XPD, LOM and light atom EMPA [1996Len, 1997Len, 1998Wie] the C-Ta binary phase diagram of [Mas2] has to be modified. The revised version, as presented in [1998Rog], is shown in Fig. 1. According to calculations employing the order parameter functional method, an ordered NaCl-type derivative phase Ta6C5 is supposed to form below 1150˚C [1991Gus]. A thermodynamic estimation of the systems B-Ta, C-Ta is from [1991Kau]. The B-C system is adopted from an assessment and thermodynamic modelling by [1996Kas, 1998Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron rich reaction L+‘B4C’ Ð (βB)
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was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan].
Solid Phases No ternary compounds were found to exist [1952Gla, 1955Bre, 1963Rud, 1965Lev, 1976Ord, 1987Ord] and mutual solid solubilities at temperatures below 2000˚C were said to be insignificantly low. According to [1963Rud] the maximal solubility of B in the tantalum carbides at 1750˚C was given as less than 1 at.% B. Whereas practically no reaction was observed for TaB2TaC1–x composites up to 2100˚C, a maximal solid solubility of < 2.8 mol% TaB2 in TaC1–x(= 3 mass% TaB2) was reported at 2400˚C raising to 6.7 mol% TaB2 at 2730˚C (= 7 mass% TaB2) [1976Ord] (see also “Quasibinary Sections”). Crystal data for the binary boundary phases are listed in Table 2.
Quasibinary Systems Three quasibinary sections of the eutectic type were reported of which TaB2-TaC0.89 is presented in Fig. 2 (after [1976Ord]) and TaB2-B4.5C is shown in Fig. 3 (after [1987Ord]). Small changes were made to comply with the accepted binary systems. Although the TaC1–x starting powder used by [1976Ord] contained 6.11 mass% C (equivalent to a composition in at.% of Ta50.5C49.5) which is at the carbon rich boundary of the monocarbide rather than at its congruent melting point at 47 at.% C (TaC0.89, Fig. 1), the melting point of TaC1–x was given as 3985˚C [1976Ord] corresponding to the maximum melting point of TaC0.89. From a cursory investigation of the TaB2-C system, [1965Lev] suggested a quasibinary eutectic with the absence of significant mutual solid solubilities. The eutectic point was mentioned at about 32 mol% TaB2 and at 2650˚C.
Invariant Equilibria A tentative partial reaction scheme (see Fig. 4, Table 3) for the B, C rich part was assigned [1998Rog] on the basis of the three experimentally established quasibinary isopleths assuming the formation of three ternary eutectic invariants E1, E2, E3, in correspondence with the phase triangulation in the isothermal section at 1750˚C [1963Rud]; (see also Fig. 5). Temperatures given in Fig. 4 are, however, tentative and strongly depend on the accuracy of the experimentally determined eutectic maxima e2, e3, e5. Depending on the relative temperatures of the reactions p1, e6 and E3, the reaction E3 may alternatively be a transition reaction U1: L + ‘B4C’ Ð TaB2 + (βB) at 2070˚C.
Liquidus, Solidus and Solvus Surfaces No experimentally determined liquidus surface is available. Attempts to schematically represent solidification behavior in the B, C rich part of the system on the basis of the experimentally defined quasibinary reactions e2, e3, e5 (see Fig. 4), cast severe doubts on the correct DOI: 10.1007/978-3-540-88053-0_22 ß Springer 2009
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position of the TaB2-C quasibinary eutectic e3 at 2650˚C [1965Lev], which appears too close to graphite and will need further experimental verification.
Isothermal Sections Figure 5 presents the only isothermal section established for the B-C-Ta ternary system at 1750˚C. The phase triangulation, as obtained by [1963Rud], was amended with additional tie lines to include the phases Ta5B6, Ta3B2 and to comply with the accepted binaries. As the high temperature phase Ta2B was shown by many researchers (see i.e. [2006Cha]) to decompose below 1900˚C it is not included in Fig. 5. As a ternary sample with 5 at.% B near Ta4C3–x did not reveal Ta4C3–x [1963Rud], the likely two-phase region TaB + Ta4C3–x is shown with dashed lines. Alternatively, a very narrow three-phase equilibrium Ta2C + Ta4C3–x + TaC1–xBx may prevent Ta4C3–x to appear at higher boron concentrations. The phase triangulation confirms the observation [1952Gla, 1955Bre] that all tantalum borides with boron contents lower than TaB2 are unstable when heated in combination with carbon.
Thermodynamics No experimental thermodynamic data are presently available for the ternary system. Upper limits for the heat of formation of TaB2 and the lower tantalum borides were estimated from compatibility studies among tantalum borides and carbon [1955Bre].
Notes on Materials Properties and Applications Microhardness of samples near the quasibinary eutectic compositions TaB2 - TaC1–x and TaB2 -‘B4C’ was observed to lie markedly below the linear combination of the two constituents (21.6 GPa for TaB2/TaC1–x and 27 GPa for TaB2/‘B4C’) depending strongly on the dispersion of phases, decreasing with increasing dispersion [1976Ord, 1987Ord]. Engeldinger et al. [1977Eng] demonstrated an increase up to 1300˚C of the hot-hardness of TaC1–x hot-pressed at 2100˚C by adding a few percent of boron. Boron was said to exert a twofold effect: (i) increase of hardness within the solubility range TaC1–xBx (given as 3at.% B maximally) and (ii) hardness increase in the two-phase field (TaC1–x + TaB2) as a combined effect of nonstoichiometry and particle hardening (precipitation of needles and platelets of TaB2 at a coherent interface (100)TaB2//(111)TaC).
Miscellaneous The authors of [1980Ord, 1993Ord] analyzed the interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - B4C [1993Ord] for transition elements M = Ti, Zr, Hf, V, Nb and Ta and presented correlations between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms involved.
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Using a compact diffraction reaction chamber, [2006Won] studied with time resolved X-ray diffraction the chemical dynamics at the combustion front for the directly ignited reaction 3Ta + B4C = 2TaB2 + TaC. Combustion front velocity was 2 mm·s–1 and the adiabatic combustion temperature calculated (2476 K) was well below the Tm of all reactants and reaction products. The combustion (involving no liquid phase) completed in 0.6 s. The authors of [2006Sot] studied the conditions to produce thin B-C-Ta films with variable compositions using a hot Ta-filament (>2000˚C) in a CH4+B2C6+H2 environment: from Auger spectroscopy and XPD the films were said to consist of TaC and Ta2B. Via precipitation strengthening Ta(B,C)-additions increased high-temperature strength and low-temperature ductility in chromium [1975Klo]. Additions of C and/or B4C proved to be an efficient densification aid to obtain up to 98% dense sinter-bodies of tantalum monocarbide (0.36 mass% B4C or 0.43 mass% B4C+0.13 mass % C, hot-pressed at 2200˚C) [2007Zha].
. Table 1 Investigations of the B-C-Ta Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1952Gla]
Hot-pressing of powder mixtures (powders of TaH2, TaC, TaB and TaB2,) in graphite dies at 1500 to 2900˚C. XPD on nine hot-pressed samples.
Compatibility of TaB2 with C and with ‘B4C’ below 2900˚C. TaB and C or TaC with B or ‘B4C’ always resulted in TaB2 + C.
[1955Bre]
Reaction between metal borides and graphite. Sintering of powder compacts in Mo crucibles under 0.5 bar argon for 50 min at 1777˚C
No experimental data listed but isothermal section shown.
[1963Rud] 45 alloys were prepared by short duration Isothermal section at 1750˚C. hot-pressing in graphite dies at 1200˚C to 2450˚C starting from well blended powders of boron (94 mass% residue O, C, Fe), lampblack C, tantalum (purity 99.7 mass% Ta, containing 0.21 mass% Nb, 0.014% Fe, 0.003% W, 0.036 mass% C) and pre-reacted tantalum carbide (containing 6.34 mass% C of which 0.14% were free C). The powder compacts were heat treated in a W tube vacuum furnace (2.5 Pa) for 9 h at 1750˚C. XPD and LOM.
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. Table 1 (continued) Reference
Method/Experimental Technique
[1965Lev]
TaB2 was prepared by vacuum sintering (Ta+2B) powder compacts for 5 h at 1400 to 1700˚C and analyzed by XPD, LOM. Starting materials were Ta powder of 5 to 10 μm containing 1 mass% Nb, 0.93% C, 0.128% Fe, 0.15% Pb, 0.07% Si, 0.07% Ti. Boron impurities were: 0.0036 mass% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb.
Temperature/Composition/Phase Range Studied Interaction of TaB2 with C yielded a quasieutectic system TaB2+C with eutectic point at 32 mol% C and TE of ca 2650˚C. Thermal analysis was performed by direct electrical heating of a graphite tube filled with TaB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
[1976Ord] Samples were prepared from TaB2 (>89.25 mass% Ta and 10.61% total B and 0.05% free C, 0.02% O2+N2) and TaC (synthesized from 93.73 mass% Ta, 6.11% total C, 0.08% free C, 0.013% O2+N2). Specimens in form of cylinders (3 mm diameter · 50 mm length) were compacted with aid of 15% aqueous polyvinyl alcohol, pre-sintered at 2000˚C in a stream of pure argon prior to heat treatment at above 2400˚C. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the quasieutectic system TaB2+TaC1–x with eutectic point at 66 mass% TaB2 and TE of 2730±40˚C on 13 samples.
[1987Ord] Samples were prepared from TaB2 and B4C which was vacuum annealed at 2000˚C to reduce C-content to 0.2 mass% free C. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the quasieutectic system TaB2+‘B4C’ with eutectic point at 31–33 mol% TaB2 and TE of 2370±30˚C on 12 samples.
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. Table 1 (continued) Reference
Method/Experimental Technique
[2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)); the liquidus and solidus curves to the L+(βB) field have been derived via chemical analysis.
Temperature/Composition/Phase Range Studied Confirmation of peritectic type of reaction L+B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L+(βB) field.
. Table 2 Crystallographic Data of Solid Phases Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype
Lattice Parameters [pm]
Comments/References
(Ta) < 3020
cI2 Im 3m W
a = 330.3
at 25˚C [Mas2]
(βB) < 2092
hR333 R 3m βB
a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1092.05 c = 2386.73
[1993Wer]
hP4 P63/mmc C (graphite)
a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78
(C)gr < 3827 (S.P.)
at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at TaB99.5 [1992Rog] at 25˚C [Mas2] [1967Low] at 2.35 at.% Cmax (2350˚C) linear ∂a/∂x, ∂c/∂x, [1967Low]
(C)d
cF8 Fd 3m C (diamond)
a = 356.69
at 25˚C, 60 GPa [Mas2]
‘B4C’ < 2450
hR45 R 3m B13C2
a = 565.1 c = 1219.6 a = 560.7 c = 1209.5 a = 556.0 c = 1212.0 a = 556.1 c = 1212.0
9 to 20 at.% C [1990Ase]
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quenched from 2400˚C [1987Ord] sample B4C+87 mass% TaB2 quenched from 2430˚C [1987Ord]
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. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype B25C
tP68 P 42m B25C or tP68 P42/nnm B25C
Lattice Parameters [pm]
Comments/References
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog]
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] defect structure tP52 [1998Rog]
Ta2B 2417 - 1900 [2006Cha]
tI12 I4/mmm CuAl2
a = 577.93 c = 486.38
[1992Rog]
Ta3B2 < 2180
tP10 P4/mbm U3Si2
a = 619.27 c = 330.27
[1992Rog]
TaB < 3090
oC8 Cmcm CrB
a = 327.49 b = 868.16 c = 315.84
[1992Rog]
Ta5B6
oC22 Cmmm V5B6
a = 313.85 b = 2260.2 c = 328.95
[1992Rog]a)
Ta3B4 < 3030
oI14 Immm Ta3B4
a = 313.20 b = 1399.68 c = 328.84
[1992Rog]a)
TaB2 < 3037
hP3 P6/mmm AlB2
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a = 309.73 c = 322.57 a = 305.86 c = 328.92 a = 309.8 c = 322.5 a = 310.0 c = 323.0 a = 311.0 c = 325.0 a = 312.5 c = 326.0 a = 307.8 c = 326.5 a = 307.9 c = 326.4 a = 308.0 c = 326.4
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26 to 37 at.% B [1992Rog] Ta rich [1992Rog] B rich [1992Rog] at 25˚C (298 K) [V-C2] at 227˚C (500 K) [V-C2] at 727˚C (1000 K) [V-C2] at 1227˚C (1500 K) [V-C2] sample quenched from 3100˚C [1976Ord] sample TaB2+5 mass% TaC1–x quenched from 2900˚C [1976Ord] sample TaB2+7.6 mass% B4C quenched from 3000˚C [1987Ord]
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. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype Ta2C(h) 3330 - 2020
Ta2C(r) ≤ 2020
Ta4C3–x ≤ 2170
hP4 P63/mmc defect NiAs
hP3 P 3m1 CdI2
hR24 R 3m V4C3–x
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Lattice Parameters [pm]
a = 310.5 c = 494.0 a = 310.5 c = 494.5
Comments/References 26 to 35.6 at.% C [1996Len, 1997Len, 1998Wie] defect structure hP3 [1998Rog] at Ta2C0.92 [V-C2] quenched from 1750˚C [1963Rud]
a = 310.37 ± 0.04 at 25˚C (298 K) [V-C2] c = 493.94 ± 0.11 a = 310.5 c = 494.1 a = 311.4 c = 495.3 a = 312.6 c = 496.8
a = 311.6 ± 0.5 c = 3000 ± 5
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at 227˚C (500 K) [V-C2] at 627˚C (900 K) [V-C2] at 1127˚C (1400 K) [V-C2] 38.2 to 39.0 at.% C [1996Len, 1997Len, 1998Wie] defect structure hR20 [1998Rog] [V-C2]
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B–C–Ta
. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype TaC1–xb) < 3985
cF8 Fm 3m NaCl
Lattice Parameters [pm]
Comments/References
36.5 to 49.8 at.% C [1996Len, 1997Len,1998Wie] TaC0.74 to TaC1.0, a = 441.3 to 445.4 quenched from 1750˚C [1963Rud] a = 442.43 at TaC0.789, 298 K [V-C2] a = 445.62 ± 0.02 at TaC0.997, 298 K [V-C2] a = 446.07 at TaC0.997, 474 K [V-C2] a = 446.67 at TaC0.997, 696 K [V-C2] a = 447.26 at TaC0.997, 891 K [V-C2] a = 447.95 at TaC0.997, 1087 K [V-C2] a = 445.1 quenched from 3980˚C [1976Ord] a = 445.32 sample TaC1–x+2.85 mol% TaB2 quenched from 3650˚C [1976Ord] a = 445.40 sample TaC1–x+4.75 mol% TaB2 quenched from 3620˚C [1976Ord] a = 445.57 sample TaC1–x+6.68 mol% TaB2 quenched from 3590˚C [1976Ord] a = 445.57 sample TaC1–x + 9.54 mol% TaB2 quenched from 3400˚C [1976Ord]
a)
Note: Crystal setting standardized with program Typix [1994Par]. Note: The symmetry assigned to the (hypothetical) structure Ta6C5 was C2, C2/m or P31 deriving from the parent NaCl type phase TaC1–x as a superlattice structure [1991Gus].
b)
. Table 3 Invariant Equilibria Composition (at.%) Reaction L Ð TaB2 + TaC1–x
T [˚C]
Type
2730±40
Phase
B
C
Ta
e2(max)
L
49.8
11.9
38.3
L Ð TaB2 + (C)gr
2650
e3(max)
L
40
40
20
L Ð TaB2 + TaC1–x + (C)gr
2550
E1
-
-
-
-
L Ð TaB2 +’B4C’
2370±30
e5(max)
L
77.5
14.5
8
L Ð TaB2 +’B4C’ + (C)gr
2150
E2
-
-
-
-
L Ð TaB2 +’B4C’ + (βB)
2000
E3
-
-
-
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. Fig. 1 B-C-Ta. The C-Ta phase diagram
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. Fig. 2 B-C-Ta. Vertical section TaB2-TaC1–x
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. Fig. 3 B-C-Ta. Vertical section TaB2 -‘B4C’
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. Fig. 4 B-C-Ta. Partial reaction scheme (proposed)
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. Fig. 5 B-C-Ta. Isothermal section at 1750˚C
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References [1952Gla] [1955Bre] [1963Rud]
[1965Lev]
[1967Low] [1975Klo] [1976Ord]
[1977Eng]
[1980Ord]
[1983Sch]
[1984Hol]
[1987Ord]
[1990Ase]
[1991Gus]
[1991Kau] [1992Rog]
[1993Ord] [1993Wer]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–406 (1955) (Experimental, Thermodyn., 19) Rudy, E., Benesovsky, F., Toth, L.E., “Investigation of Ternary System Between Va and VIa-Metals with Boron and Carbon” (in German), Z. Metallkd., 54(6), 345–353 (1963) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, 43) Levinskii, Y.V., Salibekov, S.E., Levinskaya, M.K., “Interaction of Diborides of V, Nb, Ta with Carbon” (in Russian), Poroshk. Metall. (Kiev), 5(11), 66–69 (1965) (Experimental, Phase Diagram, Phase Relations, 6) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Klopp, W.D., “A Review of Chromium, Molybdenum, and Tungsten Alloys”, J. Less-Common Met., 42(3) 261–278 (1975) (Experimental, Mechan. Prop., 51) Ordanyan, S.S., Unrod, V.I., Polishchuk, V.S., Storonkina, N.M., “Reactions in the System TaC-TaB2”, Powder Metall. Met. Ceram., 15(9), 692–695 (1976), translated from Poroshk. Metall., 9(165), 40–43 (1976) (Crys. Structure, Kinetics, Morphology, Phase Relations, Phase Diagram, #, 6) Engeldinger, M., Ritzhaupt-Kleissl, H.J., Thuemmler, F., “Hard Materials in the Ta-C-B System”, Sci. Sintering, 9(1), 121–140 (1977) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, Phys. Prop., 29) Ordanyan, S.S., “Laws of Interaction in the Systems MIV, VC-MIV, VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Experimental, Thermodyn., 14) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron Systems” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 264–274 (Review, Crys. Structure, Phase Diagram, Phase Relations, 87) Ordanyan, S.S., Dmitriev A.I., Bizhev, K.T., Stepanenko, E.K., “The Interaction in B4C-MeVB2 Systems”, Powder Metall. Met. Ceram., 26(10), 834–836 (1987), translated from Poroshk. Metall., 10(298), 66–69 (1987) (Morphology, Phase Diagram, Phase Relations, #, 5) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State, 32(9), 1595–1599 (1991) (Theory, Phase Diagram, Phase Relations, Thermodyn., 18) see also Gusev, A.I., “Physical Chemistry of Nonstoichiometric Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow, (1991) (Review, Thermodyn., Crys. Structure, Phase Diagram, Phase Relations, 102) Kaufman, L., “Coupled Thermochemical and Phase Diagram Data for Tantalum Based Binary Alloys”, Calphad, 15(3), 243–259 (1991) (Phase Diagram, Phase Relations, Thermodyn., 27) Rogl, P., “The System B-N-Ta” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM International, Materials Park, Ohio, USA, 97–100 (1992) (Crys. Structure, Thermodyn., Phase Diagram, Phase Relations, Experimental, Review, 7) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, 1, 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, 18) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51)
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22 [1994McH]
[1994Par]
[1995Vil] [1996Len]
[1996Kas] [1997Len]
[1998Kas]
[1998Rog]
[1998Wie]
[2005Tan]
[2006Cha]
[2006Sot]
[2006Won]
[2007Zha]
[Mas2] [V-C2]
B–C–Ta McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 190–191 (1994) (Phase Diagram, Phase Relations, Review, 2) Parthe, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1–4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA, 5366–5370 (1995) (Review, Phase Diagram, Crys. Structure, 7) Lengauer, W., Wiesenberger, H., Joguet, M., Rafaja, D., Ettmayer, P., “Chemical Diffusion in Transition Metal - Nitrogen Systems” in “The Chemistry of Transition Metal Carbides and Nitrides”, Oyama, S.T. (Ed.), Blacky Academic, Oxford, 91–106 (1996) (Phase Diagram, Phase Relations, Experimental, #, 29) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Lengauer, W., Wiesenberger, H., Mayr, W., Bidaud, E., Berger, R., Ettmayer, P., “Phase Stabilities of Transition Metal Carbides and Nitrides Investigated by Reaction Diffusion”, J. Chim. Phys., 94, 1020–1025 (1997) (Experimental, Phase Relations, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Tantalum” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM International, Materials Park, Ohio, USA, 257–268 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 22) Wiesenberger, H., Lengauer, W., Ettmayer, P., “Reaction Diffusion and Phase Equilibria in the V-C, Nb-C, Ta-C and Ta-N Systems”, Acta Mater., 46(2), 451–666 (1998) (Experimental, Phase Diagram, Phase Relations, 30) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC⊊Ð⊊βB”, research presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg, Germany, August 21–26, 142 (2005) (Experimental, Phase Diagram, Phase Relations, 4) Chad, V.M., Ramos, E.C.T., Coelho, G.C., Nunes, C.A., Suzuki, P.A., Flavio, F., Rogl, P., “Evaluation of the Invariant Reactions in the Ta-rich Region of the Ta-B System”, J. Phase Equilib. Diffus., 27(5), 452–455 (2006) (Experimental, Phase Diagram, Phase Relations, Crys. Structure, 10) Soto, G., Silva, G., Contreras, O., “A Study on the Flexibility of the Hot-Filament Configuration and its Implementation for Diamond, Boron Carbide and Ternary Alloys Deposition”, Surf. Coat. Technol., 201 (6), 2733–2740 (2006) (Experimental, Morphology, Phase Relations, 22) Wong, J., Larson, E.M., Waide, P.A., Frahm, R., “Combustion Front Dynamics in the Combustion Synthesis of Refractory Metal Carbides and di-Borides Using Time-Resolved X-ray Diffraction”, J. Synchrotron Radiat., 13(pt.4), 326–335 (2006) (Experimental, Phys. Prop., 30) Zhang, X.H., Hilmas, G.E., Fahrenholtz, W.G., Deason, D.M., “Hot Pressing of Tantalum Carbide with and without Sintering Additives”, J. Am. Ceram. Soc., 90(2), 393–401 (2007) (Experimental, Mechan. Prop., 24) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Vanadium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction Interaction of vanadium carbides and borides is not only of significance for the design of borided vanadium rich steels but also for superhard boride-carbide composites. Although several research groups have studied phase relations in the B-C-V system [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord], hitherto neither high temperature data are available nor a liquidus projection for the complete B-C-V ternary system. Based on a preliminary evaluation [1952Gla] of the chemical interactions among vanadium borides (VB, VB2), carbon and boron carbide in samples hot-pressed at T >1500˚C, an isothermal section at 1450˚C [1963Rud] and three quasibinary eutectic sections: VB2-C [1965Lev], VB2-VC1–x [1982Ord] and VB2-‘B4C’ [1987Ord], were established experimentally. Techniques of analysis employed were: X-ray powder diffractometry [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord], micrographic inspection [1963Rud, 1965Lev, 1982Ord, 1987Ord], melting point analyses [1965Lev, 1982Ord, 1987Ord] and chemical analyses [1963Rud, 1982Ord]. The most relevant data on the topology of the B-C-V system were compiled by [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog] including a thermodynamic extrapolation of the B-C-V ternary system on the basis of binary data known at that time. Some of the calculated sections were reproduced in [1999Rog]. Experimental details for all investigations in the B-C-V system are summarized in Table 1.
Binary Systems The C-V binary equilibria still need reinvestigation particularly with respect to the melting point of VC1–x as well as with respect to the behavior in the region V2C-VC1–x at temperatures below about 1200˚C. Detailed investigation of the C-V phase relations from 1200 to 1900˚C by means of diffusion couples, XPD, LOM and light atom EMPA [1996Len, 1997Len, 1998Wie] invokes a revision of the C-V binary given by [Mas2, 1985Car]. Furthermore, the C-V diagram presented in [Mas2, 1985Car] is essentially based on the experimental data by [1969Rud], although the melting point of vanadium monocarbide was raised from 2648 ± 12˚C at 43.0 ± 0.5 at.% C [1969Rud] to 2800˚C at 45.5 at.% C [Mas2, 1985Car]. Accordingly the eutectic l Ð VC1–x + (C)gr was raised from 2625 ± 12˚C at 49.5 ± 0.5 at.% C [1969Rud] to 2670˚C at 53.5 at.% C [Mas2, 1985Car]. A thermodynamic assessment and calculation of the C-V system by [1991Hua] arrived at 2656˚C and 44.62 at.% C for the melting point of VC1–x and at 2605˚C, 50.6 at.% C for the VC1–x + (C)gr eutectic (see Fig. 1a). The authors of [1991Gus] calculated the formation of ordered NaCl type derivative phases Landolt‐Bo¨rnstein New Series IV/11E1
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V3C2, V6C5 and V8C7 below ca. 1200˚C employing the order parameter functional method. Besides the diffusion data of [1998Wie], phase equilibria studies by XPD on hot-pressed and annealed alloys [1999Lip] were said to yield a peritectoid formation V2C+VC1–xÐ V6C5 at 1212˚C, a peritectoid formation V2C+V6C5ÐV3C2 at 882˚C and a V8C7 superstructure of VC1–x at the C rich boundary below 1107˚C. However, no clear decision was taken, how the V4C3–x phase enters the phase relations; a decomposition of V4C3–x below 1200˚C was suggested [1999Lip], but carbon contents of the ordering phases seem to be too high with respect to the low-carbon phase boundary of the monocarbide as established from diffusion couples [1998Wie]. A tentative revised version of the C-V system, which tries to interrelate all data hitherto reported is shown in Fig. 1b. A thermodynamic calculation of the B-V system is from [1981Spe] with a recent refinement of this modeling by [1998Rog, 2001Fab]. All these calculations reproduced the peritectoid formation V3B4+VB Ð V5B6 at about 1727˚C, as given by [1969Spe]. A high temperature study, however, revealed V5B6 forming from the liquid in a peritectic reaction l + V3B4 Ð V5B6 at about 2600˚C [2004Nun]. Consequently the formation of V5B6 from the liquid implies a change in the peritectic formation of VB: l + V5B6 Ð VB (see Fig. 2). The B-C system is adopted from an assessment and thermodynamic modeling by [1996Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron rich reaction L+‘B4C’⊊Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan]. The crystallographic data of all phases pertinent to the phase equilibria are listed in Table 2.
Solid Phases Besides the very boron rich V0.65B24C [1980Amb] no ternary B-C-V compound was reported [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord]. Furthermore, mutual solid solubilities among vanadium borides and carbides were observed to be very low. The only exceptions are VC1–x and V3B2. For V3B2 a carbon solubility of 3 at.% C at 1450˚C was estimated by [1963Rud] from the significantly lower lattice parameters in the ternary system and accordingly a B/C substitutional solution, V3(B1–xCx)2, was assumed. The solubility of B in VC1–x at 1450˚C was observed not to exceed about 1 at.% B [1963Rud]. Samples from the quasibinary system VB2+VC0.8, however, revealed an extended solid solubility of about 10 mass% VB2 in congruently melting VC1–x at the temperature of the quasibinary eutectic 2120 ±20˚C [1982Ord]. The solubility was said to quickly drop to 5 mass% VB2 at 2000˚C, and at 1800˚C unit cell dimensions of VC0.88 remain unchanged [1982Ord]. Although no variation of lattice parameters was encountered on B4C in B4C-VB2 samples [1987Ord], a slight increase of the unit cell dimensions of B4C alloyed with V was said to be indicative of a small solubility of V in B4C [1975Mar].
Quasibinary Systems Three quasibinary sections of the eutectic type were experimentally established of which two are presented in Fig. 3 (VB2-VC0.862, after [1982Ord]) and Fig. 4 (VB2-‘B4C’, after [1987Ord]), with small changes to comply with the accepted binary systems. With respect to the rather low DOI: 10.1007/978-3-540-88053-0_23 ß Springer 2009
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melting points of the boundary compounds recorded by [1982Ord] (VB2, 2700˚C; VC1–x, 2650˚C) in comparison with the accepted values (VB2, 2750˚C; VC0.85, 2800˚C) (Figs. 1a, 1b and 2), the quasibinary eutectic temperature TE = 2120±20˚C [1982Ord] seems low. Furthermore, VC0.88 was claimed to be the composition exhibiting congruent melting behavior [1982Ord], whilst [Mas2] actually assessed VC0.85. Preliminary data on the VB2-C system state the quasibinary eutectic nature of the isopleth (see Fig. 5) with insignificant mutual solubilities [1965Lev]. The eutectic point was reported at about 30 mol% VB2 and 2450˚C [1965Lev]. A reinvestigation is recommended.
Invariant Equilibria No experimentally established reaction scheme exists for the entire B-C-V ternary system, nevertheless a tentative partial reaction scheme for the vanadium-poor part can be assigned on the basis of the three experimentally observed ‘eutectic’ quasibinary sections assuming the formation of three invariants (E1, E2, U3, see Figs. 6a and 6b) in correspondence with the phase triangulation at 1450˚C [1963Rud]. Figures 6a and 6b present the reaction scheme as a results of a thermodynamic extrapolation of the entire system [1998Rog]. This reaction scheme has been amended taking care of the higher stability of the V5B6 phase (l + V3B4 Ð⊊V5B6 at about 2600˚C [2004Nun]) and assuming that the field of primary crystallization will run parallel to the primary field of V3B4 until V5B6 finally will be consumed in a ternary eutectic reaction L Ð⊊V5B6 + VB + VC1–x. As in an alternative case the L+V5B6 field may peter out before joining the VC1–x field, for instance in a transition reaction L + V5B6 Ð⊊VB + V3B4, a detailed experimental determination of the ternary reactions and of the liquidus surface is recommended.
Liquidus, Solidus and Solvus Surfaces No liquidus surface was derived experimentally. Experimental observations only concern the three eutectic quasibinary maxima, e3, e4, e5 (see Table 3). From a thermodynamic extrapolation of the binaries the liquidus surface was calculated [1998Rog], however, employing the older B-V binary system with V5B6 only stable below 1727˚C: the liquidus surface projection, as calculated but with a narrow primary field of V5B6 inserted (see discussion in section “Invariant Equilibria”), is shown in Fig. 7. Comparison with the experiments revealed doubts on the correct position and temperature of the quasibinary eutectic reaction e4(max) (L Ð⊊VB2C + (C)gr) [1965Lev]; further experimental verification is recommended.
Isothermal Sections Figure 8 is a presentation of the phase relations calculated at 1450˚C [1998Rog], which compare well with the results of [1963Rud] with some additional tie lines that incorporate the novel binary vanadium borides V5B6, V2B3 in consistency with the accepted binary boundary systems. The phase triangulation confirms the observations of [1952Gla] that all vanadium borides with boron contents lower than VB2 are unstable, when heated in combination with carbon. A series of isothermal sections was calculated thermodynamically Landolt‐Bo¨rnstein New Series IV/11E1
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[1998Rog] of which those at 1600 and 2000˚C are shown in Fig. 9 and Fig. 10, respectively. Tie lines to V5B6 were inserted in the latter to comply with the accepted B-V binary.
Temperature – Composition Sections Several isopleths were calculated [1998Rog, 1999Rog], however, employing the older B-V binary system with V5B6 only stable below 1727˚C. The isopleths VC1–x-B4C and VB2-VC1–x (x = 0.082, carbon rich boundary of VC at low temperature) are not affected by V5B6 and are shown in Figs. 11 and 12.
Thermodynamics A thermodynamic calculation of the ternary B-C-V system [1998Rog, 1999Rog] was based on thermodynamic assessments of the binary systems B-C [1996Kas], B-V [1998Rog] and C-V [1991Hua] as well as relying on the phase diagram data from [1963Rud] for the optimization of the thermodynamic parameters. The solid solubilities of both B and C in (V) could be satisfactorily reproduced by ternary ideal mixing terms in its Gibbs energy description. More accurate experimental investigations are needed to clarify the established shortcomings and to refine the current description of the B-C-V system.
Notes on Materials Properties and Applications [1968Alp] discussed phase assemblages, microstructures, and properties of fused carbides and/or borides from refractory systems (also for B-C-V) containing free graphite in terms of compatibility and phase diagrams. All these bodies have excellent thermal-shock resistance. Other properties (such as electrical, thermal, mechanical, chemical) can be modified by choosing different phase assemblages. Some of these materials have been cast into large shapes more than 45 cm long, and they can be machined into articles. The authors of [1998Rad, 2002Rad] fabricated a new class of superhard boron carbidebased materials by pressureless sintering of B4C-VB2 and/or B4C-VC compacts showing a considerable increase in microhardness (76 GPa) and abrasive wear resistance values of the sintered materials (1.8 times as compared to “pure” hot-pressed B4C). Free carbon and vanadium boride in hot-pressed B4C - VB2 - C composites were reported to activate the sintering process and to obtain dense, highly dispersed ceramics with higher hardness and bending strength than the monophase boron carbide ceramic. The new composite material was said to be promising for fabricating wear-resistant and shock-resistant components of various structures and machines [2006Gri].
Miscellaneous The authors of [1980Ord, 1993Ord] analyzed the interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - B4C [1993Ord] for transition elements M = Ti, DOI: 10.1007/978-3-540-88053-0_23 ß Springer 2009
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Zr, Hf, V, Nb, Ta and presented correlations between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms involved. Microhardness values of samples near the quasibinary eutectic compositions VB2 - VC1–x and VB2 -‘B4C’ were observed to lie markedly below the linear combination of the two constituents (11 to 18 GPa for VB2/VC0.88 and 21.4 to 25.2 GPa for VB2/’B4C’) depending strongly on the dispersion of phases, decreasing with increasing dispersion [1982Ord, 1987Ord]. Physical properties of ‘B4C’ - V, VB2 + ‘B4C’ + B and VB + ‘B4C’ + (C)gr cermets were studied: microhardness, microfracture [1975Mar, 1976Mar], thermal expansion [1975Mar, 1976Mar] and electrical resistance [1975Mar].
. Table 1 Investigations of the B-C-Ta Phase Relations, Structures and Thermodynamics Reference [1952Gla]
Method/Experimental Technique Hot-pressing of powder mixtures (powders of B, C, V, VC, VB and VB2,) in graphite dies at 1500 to 3150˚C. XPD on ten hot-pressed samples.
Temperature/Composition/Phase Range Studied Compatibility of VB2 with C and with ‘B4C’ below 2100˚C. VB and C or VC with B or B4C resulted in VB2 + C.
[1963Rud] 30 alloys were prepared by short duration Isothermal section at 1450˚C. hot-pressing in graphite dies at 1200˚C to 2450˚C starting from well blended powders of boron (94 mass% residue O, C, Fe), lampblack C, and vanadium (purity 99.72 mass% V, containing 0.18 mass% Fe, 0.026% N, 0.13% O, 0.055% C and 0.0053 mass% H). The powder compacts were heat treated in a W-tube vacuum furnace (2.5 Pa) for 9 h at 1450˚C. Alloys near VB2 were annealed for 150 min at 1900˚C. XPD and LOM.
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. Table 1 (continued) Reference [1965Lev]
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
VB2 was prepared by vacuum sintering Interaction of VB2 with C yielded a of powder compacts for 5 h at 1400 to quasieutectic system VB2+C with eutectic 1700˚C and analyzed by XPD, LOM. point at 30 mol% C and TE of ca. 2450˚C. Starting materials were VH0.7 powder of 5 to 10 μm containing 0.6 mass% Fe, 0.08% Ti, 0.15% Pb, 0.05% Si, 0.16% O. Boron impurities were: 0.0036% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb. Thermal analysis was performed by direct electrical heating of a graphite tube filled with VB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
[1975Mar] Specimens of V+‘B4C’ with 2–5% porosity V+‘B4C’ were prepared by hot pressing at 1900˚C under Ar for 5–7 min and were then stress relieved at 1600˚C for 12 h in vacuum. Phase analyses by XPD [1982Ord] Samples were prepared by intimate Investigation of the quasieutectic system mixing under ethanol of powders of VC0.88 VB2+VC0.88 with eutectic point at 46 mol% VB2 and TE of 2120 ±⊊20˚C on 13 samples. (synthesized from the elements) and commercial VB2 powders. Specimens (3 mm · 3 mm · 40 mm length) were obtained by spark erosion from blanks pre-sintered at 1800˚C for 2 h in vacuum (0.013 Pa). 1 mass% of Ni was added as a densification aid and was said to have evaporated almost completely during high vacuum sintering. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
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. Table 1 (continued) Reference
Temperature/Composition/Phase Range Studied
Method/Experimental Technique
Investigation of the quasieutectic system VB2+‘B4C’ with eutectic point at 45–48 mol% VB2 and TE of 2170±30˚C on 13 samples.
[1987Ord] Samples were prepared from commercial powders of VB2 and B4C which were vacuum annealed at 10 mPa and 2000˚C to reduce the C-content to 0.2 mass% free C. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Stability of V5B6 forming from the liquid [2004Nun] Investigation of three alloys V48B52, V45.45B54.55, V44B56, prepared by Argon arc l + V3B4 Ð⊊V5B6 melting from the elements (99.75 mass% V, 99.5 mass% B). A part of the samples was vacuum annealed at 2000˚C for 2 h. SEM, XPD. [2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)); the liquidus and solidus curves to the L+(βB) field have been derived via chemical analysis.
Confirmation of peritectic type of reaction L + B4+xC⊊Ð⊊(βB) via determination of the liquidus and solidus curves to the L+(βB) field.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C] (V) < 1910
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Pearson Symbol/ Space Group/ Prototype cI2 Im3m W
Lattice Parameters [pm] a = 302.40
Comments/References at 25˚C [Mas2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] (βB) < 2092
(C)gr < 3827 (S.P.)
Pearson Symbol/ Space Group/ Prototype hR333 R3m βB
Lattice Parameters [pm] a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1094.76 ± 0.07 c = 2384.22 ± 0.23 a = 1097.2 ± 0.3 c = 2390.8 ± 0.9 a = 1094.9 ± 0.3 c = 2384.0 ± 1.0
hP4 a = 246.12 c = 670.90 P63/mmc C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78
(C)d
cF8 a = 356.69 Fd3m C (diamond)
‘B4C’ < 2450
hR45 R3m B13C2
Comments/References [1993Wer] at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at VB20 [V-C2] at VB65 [V-C2] at VB99 [V-C2] at 25˚C [Mas2] [1967Low] at 2.35 at.% Cmax (2350˚C), linear ∂a/∂x, ∂c/∂x, [1967Low] at 25˚C, 60 GPa [Mas2]
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 a = 556.0 c = 1212.0 a = 556.1 c = 1212.0
sample quenched from 2400˚C [1987Ord] for 15.5 mol% VB2, 84.5% B4C, quenched from 2420˚C [1987Ord]
B25C
tP68 P42m or P42/nnm B25C
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog]
V3B2 < 1913
tP10 P4/mbm U3Si2
a = 574.14 c = 302.95
[1998Rog]
a = 572.8 to 573.9 c = 302.6 to 303.0 a = 570.3 c = 302.5
[1963Rud]
V3(B1-xCx)2
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
VB < 2549
oC8 Cmcm CrB
a = 306.03 b = 805.4 c = 297.2
[1998Rog]
V5B6 ≲ 2560
oC22 Cmmm V5B6
a = 297.64 b = 2130.8 c = 305.8
[1998Rog, 2004Nun]
V3B4 < 2600
oI14 Immm Ta3B4
a = 298.1 b = 1322.0 c = 306.06
[1998Rog]
V2B3 < 2610
oC20 Cmcm V2B3
a = 305.88 b = 1842.2 c = 298.46
[1998Rog]
VB2 < 2750 (< 2700±50 [1982Ord])
hP3 P6/mmm AlB2
a = 299.89 c = 305.8 a = 299.8 c = 306.0 a = 300.2 c = 306.2 a = 299.7 c = 306.1 a = 300.2 c = 306.5
[1998Rog]
V1-xB25
tP52 P42/nnm TiB25
a = 882.4 ± 0.9 c = 507.2 ± 1.2
[V-C2] metastable (?)
βV2C(h1) < 2187
hP4 P63/mmc NiAs
a = 288.78 c = 457.43 a = 288.0 to 289.4 c = 445.59 to 459.0
27 to 35.4 at.% C [1996Len] defect NiAs structure hP3 [1998Rog] [V-C2] labelled as εFe2N type sample quenched from 1450˚C [1963Rud]
β’V2C(h2) 1600 - 800
hP9 P31ma) W2C
αV2C < 850
oP12 Pbcn Fe2N
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sample quenched from 2600˚C [1987Ord] sample quenched from 2520˚C [1982Ord] sample VB2-6.7 mol% B4C quenched from 2400˚C [1987Ord] sample VB2-5 mol% VC0.88, quenched from 2320˚C [1982Ord]
a = 500.5 c = 455.1
[1998Rog]
<850˚C, 31–33 at.% C [Mas2] [V-C2]
a = 503.6 b = 456.3 c = 575.0
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C]
V4C3-x hR24 < 1875 - 1220 R3m V4C3 VC1-x < 2800
Lattice Parameters [pm]
at 37.9 at.% C [1998Wie] defect structure hR20 [1998Rog] [1999Lip, V-C2]
a = 291.7 c = 2783
cF8 Fm3m NaCl
37 to 48 at.% C [1996Len] at 293 K, [V-C2] at VC0.75, quenched from 1450˚C [1963Rud] at VC0.96, quenched from 1450˚C [1963Rud] at 17 mass% C, quenched from 2750˚C, [1982Ord] sample VB2+10 mol% VC0.88 quenched from 2300˚C [1982Ord]
a = 415.9 a = 412.8 a = 416.5 a = 416.7 a = 417.1
V3C2 < 882
-
V6C5 < 1212
hP33 P3112orP31 V6C5
V8C7 < 1107
cP60 P4132 V8C7
-
[1991Gus, 1999Lip] 1200˚C, 43 to 46 at.% C, also described as monoclinic C2 or C2/m [1999Lip] [V-C2]
a = 509 c = 1440
a = 833.34 ± 0.06 * τ1, V1-xB24C
tP52 P42/nnm TiB25
Comments/References
a = 885.7 ± 0.9 c = 507.0 ± 0.2
1100˚C, 47 to 48 at.% C superstructure of VC1-x (2nd order transition) [1999Lip] [V-C2] [V-C2] metastable (?), eventual solid solution from binary VB25 or B25C
Note: partially disordered V2N type, cited by various authors as partially disordered εFe2N-type.
a)
. Table 3 Invariant Equilibria (Experimental Data) Composition (at.%) Reaction L Ð VB2 + ‘B4C’
T [˚C]
Type
2170±20
e3 (max)
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. Table 3 (continued) Composition (at.%) Reaction L Ð VB2 + (C)gr
L Ð VB2 + VC1-x
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2120±20
Type e4 (max)
e5 (max)
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VB2
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33
(C)gr
0
100
0
L
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19.9
41.8
VB2
67
0
33
VC1-x
8.7
40.7
50.6
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. Fig. 1a B-C-V. The C-V phase diagram, calculated after [1991Hua]
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. Fig. 1b B-C-V. The C-V phase diagram
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. Fig. 2 B-C-V. The B-V phase diagram
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. Fig. 3 B-C-V. Vertical section VB2-‘B4C’
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. Fig. 4 B-C-V. Isopleth VB2-VC1–x
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. Fig. 5 B-C-V. Isopleth VB2-C
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. Fig. 6a B-C-V. Partial reaction scheme (proposed), part 1
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. Fig. 6b B-C-V. Partial reaction scheme (proposed), part 2
B–C–V
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. Fig. 7 B-C-V. Proposed liquidus surface (calculated by [1998Rog], with field of primary crystallization of V5B6 inserted)
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. Fig. 8 B-C-V. Isothermal section calculated at 1450˚C
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. Fig. 9 B-C-V. Isothermal section calculated at 1600˚C
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. Fig. 10 B-C-V. Isothermal section calculated at 2000˚C
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. Fig. 11 B-C-V. Calculated section VC1–x-‘B4C’
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. Fig. 12 B-C-V. Calculated section VC1–x-VB2
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References [1952Gla] [1963Rud]
[1965Lev]
[1968Alp]
[1967Low] [1969Rud]
[1969Spe] [1975Mar]
[1976Mar]
[1980Amb]
[1980Ord]
[1981Spe] [1982Ord]
[1983Sch]
[1984Hol]
[1985Car] [1987Ord]
[1990Ase]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Rudy, E., Benesovsky, F., Toth, L.E., “Investigation of Ternary System Between Va and VIa-Metals with Boron and Carbon” (in German), Z. Metallkd., 54(6), 345–353 (1963) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, #, *, 43) Levinskii, Y.V., Salibekov, S.E., Levinskaya, M.Kh., “Interaction of Diborides of V, Nb, Ta with Carbon” (in Russian), Poroshk. Metall. (Kiev), 5(11), 66–69 (1965) (Phase Diagram, Phase Relations, Experimental, 6) Alper, A.M., Doman, R.C., McNally, R.N., “Fusion-Cast Carbide-Boride-Graphite Ceramics” in “Proc. Fourth International Conference on Science of Ceramics”, Maastricht, Netherlands, European Ceramic Association., 23–27 April, 1967, Stewart, G.H. (Ed.), British Ceramic Society, Stoke-on-Trent, 389–420 (1968) (Experimental, Review, 73) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Rudy, E., “Part V. Compendium of Phase Diagram Data, System V-C” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65-2, Contact No. USAF 33 (615)-1249 and 33(615)-67-C-1513, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, Part V, 168–170 (1969) (Experimental, Phase Diagram, Phase Relations, 5) Spear, K.E., Gilles, P.W., “Phase and Structure Relationship in the Vanadium Boron System”, High Temp. Sci., 1, 86–97 (1969) (Experimental, Crys. Structure, 17) Marek, E.V., Dudnik, E.M., Makarenko, G.N., Remenyuk, E.A., “Physical Properties of Boron Carbide Alloys with Vanadium and Chromium Additives” (in Russian), Poroshk. Metall. (Kiev), (2), 54–56 (1975) (Experimental, Phys. Prop., 6) Marek, E.V., Makarenko, G.N., “Production and Certain Properties of B-C-V and B-C-Cr Alloys” in “Karbidy i Splavy na ikh Osnove”, Akad. Nauk Ukrain. SSR, 195–198 (1976) (Experimental, Mechan. Prop., Crys. Structure, Phase Relations, 3) Amberger, E., Gerster, H.P., “Structure of the Ternary I-Tetragonal Boride: (B12)4C2Ti1.86 and (B12)4C2V1.29” (in German), Acta Crystallogr., 36(B), 672–675 (1980) (Crys. Structure, Experimental, 13) Ordanyan, S.S., “Laws of Interaction in the Systems MIV, VC-MIV, VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Experimental, Thermodyn., 14) Spear, K.E., Blanks, J.H., Wang, M.S., “Thermodynamic Modeling of the V-B System”, J. Less-Common Met., 82, 237–243 (1981) (Review, Thermodyn., Phase Diagram, Phase Relations, 17) Ordanyan, S.S., Topchii, L.A., Khoroshilova, I.K., Chupov, V.D., “Reactions in the VC0.88-VB2 System”, Powder Metall. Met. Ceram., 21(2), 122–124 (1982), translated from Poroshk. Metall., 2(230), 49–51 (1982) (Crys. Structure, Phase Diagram, Phase Relations, #, 12) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Carlson, O.N., Ghaneya, A.H., Smith, J.F., “The Carbon-Vanadium System”, Bull. Alloy Phase Diagrams, 6(2), 115–124 (1985) (Review, Crys. Structure, Thermodyn., Phase Diagram, Phase Relations, 77) Ordanyan, S.S., Dmitriev, A.I., Bizhev, K.T., Stepanenko, E.K., “The Interaction in B4C-MeVB2 Systems”, Powder Metall. Met. Ceram., 26(10), 834–836 (1987), translated from Poroshk. Metall., 10 (298), 66–69 (1987) (Morphology, Experimental, Phase Diagram, Phase Relations, #, 5) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides”, in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14)
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[1993Ord] [1993Wer]
[1994McH]
[1995Vil] [1996Len]
[1996Kas] [1997Len]
[1998Kas]
[1998Rad] [1998Wie]
[1998Rog]
[1999Lip]
[1999Rog] [2001Fab]
[2002Rad] [2004Nun]
[2005Tan]
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Huang, W., “An Assessment of the V-C System”, Z. Metallkd., 82(3), 174–181 (1991) (Review, Phase Diagram, Phase Relations, 40) Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State, 32(9), 1595–1599 (1991) (Theory, Phase Diagram, Thermodyn., 18), see also Gusev, A.I., “Physical Chemistry of Nonstoichiometric Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow, (1991) (Review, Thermodyn., Crys. Structure, Phase Diagram, Phase Relations, 102) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, 1, 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, 18) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 197–198 (1994) (Phase Diagram, Phase Relations, Review, 2) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA, 5357 (1995) (Review, Phase Diagram, Crys. Structure, 4) Lengauer, W., Wiesenberger, H., Joguet, M., Rafaja, D., Ettmayer, P., “Chemical Diffusion in Transition Metal - Nitrogen Systems” in “The Chemistry of Transition Metal Carbides and Nitrides”, Oyama, S.T. (Ed.), Blacky Academic, Oxford, 91–106 (1996) (Phase Diagram, Experimental, Phase Relations, #, 29) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Lengauer, W., Wiesenberger, H., Mayr, W., Bidaud, E., Berger, R., Ettmayer, P., “Phase Stabilities of Transition Metal Carbides and Nitrides Investigated by Reaction Diffusion”, J. Chim. Phys., 94, 1020–1025 (1997) (Experimental, Phase Relations, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Radev, D.D., Zakhariev, Z., “Structural and Mechanical Properties of Activated Sintered Boron Carbide-based Materials”, J. Solid State Chem., 137(1), 1–5 (1998) (Experimental, Mechan. Prop., 24) Wiesenberger, H., Lengauer, W., Ettmayer, P., “Reaction Diffusion and Phase Equilibria in the V-C, Nb-C, Ta-C and Ta-N Systems”, Acta Mater., 46(2), 651–666 (1998) (Experimental, Phase Diagram, Phase Relations, 30) Rogl, P., “The System Boron - Carbon - Tantalum” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 257–268 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 22) Lipatnikov, V.N., Gusev, A.I., Ettmayer, P., Lengauer, W., “Phase Transitions in Non-Stoichiometric Vanadium-Carbide”, J. Phys.: Con. Matter, 11, 163–184 (1999) (Experimental, Crys. Structure, Thermodyn., Phase Diagram, Phase Relations, 55) Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides”, Int. J. Refract. Met. Hard Mater., 17, 27–32 (1999) (Crys. Structure, Experimental, Phase Relations, 6) Fabrichnaya, O., “B-V (Boron-Vanadium)” in Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group IV: Physical Chemistry. Ed. W. Martiensen, “Thermodynamic Properties of Inorganic Materials compiled by SGTE. Binary systems. Part 2: Elements and Binary Systems from B – C to Cr – Zr”, Vol. 19B2, Franke, P., Neuschu¨tz, D. (Eds.), Springer-Verlag, Berlin, Heidelberg, IV/19B2, 45–48 (Thermodyn., Phase Diagram, Phase Relations, 5) Radev, D.D., Mihailova, B., Konstantinov, L., “Raman Spectroscopy Study of Metal-Containing Boron Carbide-Based Ceramics”, Solid State Sci., 4(1), 37–41 (2002) (Experimental, Mechan. Prop., 24) Nunes, C.A., de Lima, B.B., Coelho, G.C., Rogl, P., Suzuki, P.A., “On the Stability of the V5B6Phase”, J. Alloys Compd., 370, 164–168 (2004) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 8) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), August 21–26, 142 (2005) (Experimental, Phase Diagram, Phase Relations, 4)
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B–C–V Grigor’ev, O.N., Koval’chuk, V.V., Zaporozhets, O.I., Bega, N.D., Galanov, B.A., Prilutskii, E.V., Kotenko, V.A., Kutran’, T.N., Dordienko, N.A., “Synthesis and Physicomechanical Properties of B4CVB2 Composites”, Pow. Metall. Met. Ceram., 45(1-2), 47–58 (2006) (Experimental, Mechan. Prop., 34) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Tungsten Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction As a unique combination of high hardness and ultra-high temperature compounds, the system B-C-W has attracted wide attention starting with early studies of the interaction between tungsten borides, tungsten carbides and carbon [1952Gla, 1955Bre, 1959Sam] summarized in a phase triangulation of the system [1955Bre]. Phase relations in an isothermal section at 1700˚C were examined by [1963Rud]. [1965Lev] presented information on the quasibinaries WB + C and W2B5–x + C, whereas [1966Kip, 1968Kip1, 1968Kip2] studied the isopleths W2B WC, WB - WC, W2B - W2C and derived an isothermal section B-C-W at 1800˚C [1968Kip1]. The most complete and thorough experimental investigation, however, stems from [1969Rud1, 1969Rud2, 1970Rud] who provided a series of isothermal sections at 1500, 2000, 2150, 2320, 2500, 2700 and 2800˚C, several isopleths, i.e. W2B5–x - ‘B4C’, W2B5–x - C, WB - C, WB - W2C, W2B - W2C, the liquidus surface, a reaction diagram and a threedimensional view of the constitution system. Compilations on the most relevant data on the topology of the B-C-W system were provided by [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information in literature for the B-C-W system up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog1]. Thermodynamic modeling of the B-C-W system is due to [1998Rog1, 1998Rog2, 1999Rog]. Experimental details for all investigations in the B-C-W system are summarized in Table 1.
Binary Systems The binary system B-W, as assessed and calculated by [1995Dus] (see Fig. 1) is preferred with respect to the version presented in [Mas2]. The negative curvature of the liquidus line l/l+(W) shown by [1970Rud] is an unlikely feature and was not revealed in the thermodynamic calculation [1995Dus]. In contrast to [1995Dus], who extracted the homogeneity region for W2B5–x from the assessment to be within 67 to 71 at.% B, [1995Oht] claimed, that the compound only exists within the region WBx for 1.80 < x < 1.97 (64 to 66 at.% B). The C-W binary system used is essentially based on the version presented by [1965Rud]. Data on the C solubility in W for the temperature range from 1400˚C to 2000˚C are due to [1973Kuh] and extrapolate to 0.7 to 0.8 at.% C at 2710˚C, the temperature of the (W)+W2C eutectic according to [1991Nag]. The transitions among the various W2C modifications are subject to change according to in situ neutron diffraction data by [1988Epi]. At variance to the C-W phase diagram in [Mas2], where W2C is proposed to decompose below 1250˚C into (W)+WC, phase relations, elucidated by [1988Epi] and pertinent to the variously ordered carbon vacancy sublattices, suggest thermal stability of the W2C-subcarbide at temperatures below Landolt‐Bo¨rnstein New Series IV/11E1
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800˚C. The thermodynamic calculations of the C-W diagram by [1986Gus] and its modification by [1993Jon] both refer to the phase diagram version in [Mas2]. Until more sophisticated experiments will solve the aforementioned controversies, the C-W calculated by [1993Jon] (Fig. 2) is used. The B-C system is adopted from an assessment and thermodynamic modelling by [1996Kas, 1998Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron-rich reaction L + ‘B4C’ Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan].
Solid Phases No ternary compounds were observed throughout the B-C-W ternary system. Mutual solid solubilities of the binary boundary phases were found to be very small (less than 3 at.%) or negligible [1955Bre, 1963Rud, 1965Lev, 1966Kip, 1968Kip1, 1968Kip2, 1970Rud]. Crystallographic data of all solid phases pertinent to the B-C-W ternary system are listed in Table 2.
Quasibinary Systems Five quasibinary sections of the eutectic type, W2B - W2C, WB - C, WB - W2C, W2B5–x - C, W2B5–x - B4.5C, were established by [1970Rud] confirming the eutectic nature of the isopleths WB - C and W2B5–x - C earlier reported by [1965Lev]. However, the eutectic temperature as measured by [1965Lev] with optical pyrometry on prereacted powders through a bore-hole in a directly heated graphite tube, 2187˚C for W2B5–x - C and 2267˚C for WB - C, are considerably lower than those given by [1970Rud]. This fact might be caused by insufficient correction for non-blackbody conditions. The section W2B - WC, studied by [1966Kip, 1968Kip2], was claimed to be “eutectic”, with TE = 2787 ± 20˚C at 50 mol% WC. Observations by [1970Rud] indeed reveal a ternary eutectic L Ð W2C + W2B + WB at 2305˚C and at 37 mol% WC close to the section W2B - WC. The transition reaction WB + W2C Ð W2B + WC experimentally reported at 2150˚C [1970Rud], however, disproves the section to be truly quasibinary (see also Table 3 and sections “Invariant Equilibria” and “Thermodynamics”). The WB-WC eutectic at 28±2 mass% WB was seen in a WB0.49C0.76 alloy slowly crystallized [2006Pad]. As seen from calculations (see below) the entire section WB-WC cannot be a quasibinary diagram.
Invariant Equilibria A reaction scheme based on fifteen observed ternary invariant equilibria was provided by [1970Rud]. Owing to the small mutual solid solubilities of the boundary phases only the compositions of the liquid phase, which are listed in Table 3, were reported. The reaction scheme has been amended in order to correspond to the concentration dependent transition βWB Ð αWB in the ternary system, since [1970Rud] quoted only a reaction WB + W2C Ð W2B + WC at 2150˚C, without specifying the crystallographic modification of the monoboride. Therefore two transition reactions βWB + C Ð αWB + WC, calculated at 2149˚C (U4) and βWB + WC Ð W2B + αWB, calculated at 2131˚C (U6), have to be added (see Table 3). DOI: 10.1007/978-3-540-88053-0_24 ß Springer 2009
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Awaiting further in situ neutron diffraction studies on the W2C subcarbide in order to resolve the transitions recorded by [1970Rud] and [1988Epi] in the temperature range from 2300˚C to 2500˚C, no attempts were made to incorporate these C concentration dependent W2Ctransitions in the reaction scheme presented in Figs. 3a to 3c, as a result of a thermodynamic modeling by [1999Rog].
Liquidus, Solidus and Solvus Surfaces The liquidus projection, as shown in Fig. 4, is based on the results of [1970Rud]. The unlikely positive curvature of the liquidus l/l+(W) in the binary B-W system, as shown by [1970Rud], also affects the liquidus surface in the W rich corner of the B-C-W ternary. The thermodynamic calculation of the B-W system [1995Dus] in contrast revealed a pronounced negative curvature.
Isothermal Sections Isothermal sections were experimentally studied at 1500˚C by [1970Rud], at 1700˚C by [1963Rud] and at 1800˚C by [1968Kip1]. Both the sections of [1970Rud] and [1968Kip1] are in good agreement with each other. On the basis of the phase equilibria and the experimental findings at 1500 and 1700˚C, [1970Rud] presented isotherms for a series of temperatures, 1500, 2000, 2150, 2320, 2350, 2500, 2700 and 2800˚C. Figure 5 presents the X-ray evaluation of the alloy series equilibrated at 1500˚C by [1970Rud]. Figure 6 shows the subsolidus phase relations as derived from as-cast alloys [1970Rud]. The phase triangulation presented by [1955Bre] (no temperature given) applies for the temperature range below the transition reaction U5 at 2143˚C.
Temperature – Composition Sections Besides the experimental sections claimed and listed under quasibinary systems, no further sections were experimentally derived. For thermodynamic calculation of isopleths see section below.
Thermodynamics Estimates of thermodynamic stability of binary borides and carbides (heat of formation data) were extracted from phase reactions in the B-C-W ternary [1955Bre]. Thermodynamic modeling of the B-C-W system in a first attempt [1998Rog1, 1998Rog2] employed the B-C binary with the l Ð (βB) + ‘B4C’ eutectic at 2073˚C [1994Kas] resulting in a ternary boron rich eutectic reaction L Ð W1–xB3 + B + ‘B4C’ at 1899˚C. A later assessment and thermodynamic modeling of the B-C binary [1996Kas] revealed a peritectic reaction l + ‘B4C’ Ð (βB) at 2103˚C, which was experimentally confirmed from a floating zone experiment [2005Tan]. Thermodynamic modeling of the B-C-W system on the basis of the new B-C binary data, however, changed the ternary eutectic into a ternary transition reaction Landolt‐Bo¨rnstein New Series IV/11E1
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in the boron rich corner: L + ‘B4C’ Ð W1–xB3 + (βB) at 1912˚C [1999Rog]. Although the type of reaction changed, the calculated reaction temperature is closer to the experimentally observed value T = 1950˚C. The allowance of a small carbon solubility in WB at higher temperatures provides full consistency with the experimentally observed liquidus trough and particularly saves the quasibinary eutectic nature of the section βWB - W2C. However, compromises had to be taken with respect to the experimental melting point data, which at these temperatures are extremely difficult to access. The deviations of the calculation from experimental values seems to be acceptable when compared to the experimental ambiguity for the B-C system. In most cases the section running through the congruent melting points of the binary phases does not result in a true quasibinary system. Emphasizing on the eutectic maximum point as part of the section and releasing the constraints for the congruent melting of the binary phases usually restored a partially quasibinary character of the section, i.e. the characteristic tie lines lie exactly in the plane of the section, for instance the quasibinary condition in the section W2B5–x - C depends only on the solution range of carbon. Calculated data are presented for some selected isothermal sections at 2150˚C (Fig. 7) and at 2350˚C (Fig. 8) as well as for the isopleths, W2B5–x - B4C (Fig. 9), W2B5–x -C (Fig. 10), WB- B4C (Fig. 11). The complex reactions between W and B4C can be read of from Fig. 12.
Notes on Materials Properties and Applications Combining refractory metal-like borides with high temperature ceramics constitutes attractive high temperature materials. Thus considerable interest is devoted to the manufacture, high temperature behavior and properties of tungsten boride/carbide composites and many articles deal with various techniques on hot-pressing-reaction sintering of dense material as well as with the characterization of mechanical properties and friction/wear behavior. Particularly high density composites containing 30 to 70 vol% W2B5 were fabricated from in situ reaction of fine powders of B4C, WC (and C) yielding high flexural strengths of about 850 MPa, high fracture toughness of up to 9 MPa·m1/2, high hardness up to 12 GPa and low electrical resistivity of 0.6 μΩm [2005Yi, 2006Lv1, 2006Lv2, 2006Lv3, 2006Wen, 2007Lv]. B-C-W composites consisting of B4C, WC, W2B5 and C phases, hot pressed at 1900˚C from a powder mixture of B4C-40WC, gave a high strength of 453 MPa and a rather high fracture toughness (KIC) of 8.7 MPa·m1/2 [2000Wen]. Dense WC based composites (WC+WB+W2C), prepared by reaction sintering in a hot press from B4C-W-WC powders without metal binders, showed high hardness up to 24.5 GPa, high fracture toughness up to 6.1 MPa·m1/2 and high Young’s modulus up to 700 GPa [2002Sug, 2004Sug1, 2004Sug2]. The superior mechanical properties of boron carbide make it an attractive candidate in tribological applications as a wear resistant coating on cemented tungsten carbide substrates. Boron carbide (B4C) was deposited on a tungsten substrate by chemical vapor deposition from a dichloroborane-CH4 gas mixture. The activation energy of the boron carbide formation reaction was ascertained to be 56.1 ± 4.0 kJ·mol–1 [2006Kar]. The solid particle erosion behavior of 12–18 μm thick CVD boron carbide coatings was studied in erosion tests on a high-energy air solid particle erosion rig. Erosion rates and mechanisms were discussed in terms of coating thickness, particle velocity, particle shape and size. Erosion of CVD boron carbide occurs predominantly through a single-stage mechanism by the formation of DOI: 10.1007/978-3-540-88053-0_24 ß Springer 2009
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lateral-radial crack systems that propagate outwards towards the free CVD surface and extend into the coating substrate interface [2005Bos].
Miscellaneous A significant increase of hardness was observed for eutectic WB + WC alloys (27–30 GPa), which were slowly solidified with respect to WC grains (18–24 GPa), the absence of radial crack formation on indentation suggesting plasticity of the composite material [2006Pad]. By use of the quasibinary eutectic B4C - W2B5–x at 2220 ± 20˚C [1991Zak] obtained rather dense and fine crystalline composites by pressureless sintering at 2250˚C starting from B4C + 10 mass% WC. The resistance of alloys against air-oxidation varies almost linearly along the section W2B5–x - WC [1959Sam] (no further details given). Alloys containing less than 10 mol% WC show the highest resistance against oxidation. Porosity, density and electrical resistance were examined on samples with the composition ‘W2B2C’, sintered at various temperatures from 2200 to 2700˚C for various durations (5 to 20 min) [1959Sam]. Samples were furthermore sintered along the section W2B5–x - WC at temperatures from 2350 to 2400˚C for 3 to 5 min [1959Sam]. The short duration of sintering was said to be responsible for nonreaction of the starting material.
. Table 1 Investigations of the B-C-W Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1952Gla]
Hot-pressing of powder mixtures (powders of W, WC, C and B) in graphite dies at 1500 to 2600˚C. XPD on seven hotpressed samples.
Reaction is sluggish at 1400˚C but is complete at 2100˚C. Compatibility of W2B5 with ‘B4C’ up to 2000˚C. W2B5 melts undercomposed at 2980˚C.
[1955Bre]
Reaction between metal borides and graphite. Sintering of powder compacts in Mo crucibles under 0.5 bar argon for 50 min at presumably 1777˚C. XPD
XPD on a sample W2B + C = WB + WC. Compatibility of W2C+WC+W2B. Phase triangulation (at 1777˚C ?).
[1959Sam]
Investigation of the section W2B5-WC and Samples were prepared from WO3 via borothermic reaction with B4C or B2O3 W2B2C. Immiscibility of W2B5-WC. and C or from mixtures WC + W + C. LOM, XPD and microhardness. Oxidation in air, as well as friction and wear were investigated.
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. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1963Rud]
Hot-pressing of 45 binary and ternary powder compacts in C-cartridges at 1200 to 2500˚C followed by subsequent anneal in a W-tube vacuum furnace (2.5 Pa) for 9 h at 1300 and 1700˚C, respectively. Some alloys were arc melted under argon. Starting materials were amorphous boron (94 mass% residue O, C, Fe), lampblack C, W (0.19 mass% O, 0.01% C), WC (6.20 mass% Ctotal, 0.11mass % free C). XPD and LOM.
Isothermal section at 1700˚C. New boride WB12 (or WB4) forms single phase up to 1800˚C but decomposes at higher temperatures.
[1965Lev]
WB and W2B5 powders were prepared by vacuum sintering powder compacts for 5 h at 1525˚C in a W heater furnace and analyzed by XPD, LOM. Starting materials were powders of <5 μm: WB2 containing 0.029 mass% Mo, 0.001% Ni, 0.1% O, 0.01% SiO2 and B-powder containing 0.0036% Fe, 0.0003% Si, 0.01% Cu, 0.0004% Al, 0.0006% Pb. Thermal analysis was performed by direct electrical heating of a graphite tube filled with powder mixtures of WB, W2B5 and C or W+B+C or W+B4C or WC+B (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
Interaction of WB and W2B5 with C yielded a quasieutectic systems W2B5+C with eutectic TE of 2187±30˚C and WB+C with TE of 2267±30˚C (no eutectic concentrations given).
[1966Kip] [1968Kip1] [1968Kip2]
Samples within 10–90% W2B and WC hotpressed in graphite dies. Starting materials were powders of 99.9 mass% W, WC (containing 6.01 mass% Ctotal, 0.03 mass% free C), lampblack 99.79 mass% C and stoichiometric B4C. XPD, LOM, micro-hardness and electric resistance data.
Quasibinary eutectic section W2B - WC with TE of 2290±20˚C at 50 mol% WC. No mutual solid solubility [1966Kip, 1968Kip2]. B-C-W system studied along the quasibinary eutectic sections WB-WC, W2B-WC, W2B-W2C and compositions richer than 50 mol% W. No significant mutual solid solubility. Isothermal section at 1800˚C [1968Kip1]
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24
. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1969Rud1] Samples were prepared mainly by short [1969Rud2] duration hot pressing in graphite dies [1970Rud] at temperatures between 1800˚C and 2200˚C. The samples were directly used in as-pressed condition for Pirani melting point (under 2.5·105 Pa He) or differential thermal analysis (graphite container under 105 Pa of He). Starting materials were (i) high purity elemental powders (i.e. W with particle size smaller than 6.8 μm containing 720 ppm O, <50 ppm Mo; 40 ppm Fe, 20 ppm Ni, 430 ppm N, other metals <60 ppm; spectrographic grade graphite powder with 0.3 ppm Al, 0.1 Cu, 0.2 Fe, 0.1 Mg, 0.2 Si; and boron powder of 99.55 mass% B containing 0.25 mass% Fe and 0.1 mass% C) as well as (ii) prereacted master alloys of W2B5 (from elemental compacts reacted in Ta cans for 2h at 1750˚C in vacuum. After crushing, acid leaching (HCl+H2SO3), washing and drying the powders contained 70.7 at.% B, 0.12 mass% C and impurities: Fe 500 ppm, Si, Mg, Mn, Ca, Cu, Ni, Cr, Mn each 100 ppm, and Ti 600 ppm), and WC (from elemental compacts reacted in a C-furnace at 2100˚C for 4 h under H2. After ball-milling to 60 μm, acid leaching, washing and drying the powders contained 50.7 at.% C, and 150 ppm O). Whereas specimens for DTA and melting point analyses were directly equilibrated in the equipment prior to the runs, specimens for the isothermal sections were generally annealed in a tungsten mesh furnace for 140 h at 1500˚C under a vacuum of 5·10–3 Pa or for 6 h at 2000˚C under 1.05 to 2·105 Pa of helium and rapidly quenched.
Investigation of the constitution of the BC-W phase diagram employing X-ray powder diffraction, metallographic, melting point and differential thermoanalytic (>4/sec, 1 bar He) techniques on hot pressed and sintered as well as argon arc-melted specimens. Selected alloys were equilibrated in the melting point furnace and quenched in a preheated tin bath (300˚C). Samples were chemically analyzed for free and combined carbon, boron, oxygen and nitrogen contaminants. For polishing and etching usually a slurry of alumina in 5% chromic acid was used; electroetching in 2% aqueous solution with NaOH was successful.
[2005Tan]
Confirmation of peritectic type of reaction L + B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L + (βB) field.
Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)). The liquidus and solidus curves to the L + (βB) field have been derived via chemical analysis.
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. Table 1 (continued) Reference
Method/Experimental Technique
[2006Pad]
Samples WB0.12C0.74, WB0.25C0.75, WB0.34C0.32, WB0.49C0.76, WB0.59C0.76, WB0.89C0.75 and (WC)0.9B0.10 were prepared by crystallization of pre-sintered rods made from powder compacts by arcmelting or directional crystallization in vertical crucibles via zone melting in an induction furnace. XPD, SEM, TEM, density, hardness, Young’s modulus and fracture toughness.
Temperature/Composition/Phase Range Studied Directional crystallization of B4C-NbB2 eutectic compositions investigated by XPD, SEM, TEM, density, hardness, Young’s modulus and fracture toughness.
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(W) < 3422
cI2 Im 3m W
a = 316.52 a = 316.65
at 25˚C [Mas2] [1965Rud]
(βB) < 2092
hR333 R 3m βB
a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1093.82 c = 2383.56
[1993Wer]
hP4 P63/mmc C (graphite)
a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78
at 25˚C [Mas2]
(C)gr < 3827 (S.P.)
at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at WB99 [V-C2]
[1967Low] at 2.35 at.% Cmax (2350˚C) linear ∂a/∂x, ∂c/∂x, [1967Low]
‘B4C’ < 2450
hR45 R 3m B13C2
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5
B25C
tP68 P42m or P42/nnm B25C
a = 875.3 ± 0.4 c = 509.3 ± 1.5
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[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog1]
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24
. Table 2 (continued) Phase/ Temperature Range [˚C] γW2C > 1800
βW2C 1800 - 1050
Pearson Symbol/ Space Group/ Prototype hP3 P63/mmc defect NiAs or hP4 P 3m1 anti CdI2 hP9 P 31m W2C(a)
Lattice Parameters [pm]
Comments/References
a = 300.0 c = 472.0
listed as disordered L’3type ("Fe2N")a), γW2C [1988Epi]
a = 298.5 c = 471.6
at 29.5 at.% C [1970Rud]
a = 300.1 c = 472.8
at 33.3 at.% C [1970Rud] quenched from 2000˚C
a = 518.2 c = 472.0 a = 518.09 c = 472.16 a = 518.33 c = 472.40 a = 518.52 c = 472.32 a = 519.4 c = 472.1 a = 519.0 c = 472.4
listed as “ε-W2C” at W2C0.81 [1988Epi] at W2C0.84 [1986Loe] at W2C0.86 [1978Har] at W2C0.89 [1978Har] at W2C0.92 [1988Epi] at W2C1.00 [1988Epi]
αW2C < 1050
oP12 Pbcn αMo2C
a = 472.1 b = 603.0 c = 518.0 a = 473.8 b = 600.9 c = 519.3
[V-C2, 1988Epi, 1994Par] cited as “ξ-Fe2N” by [1970Rud], “ξ-W2C” [1988Epi] at 32.6 at. % C [1970Rud]
WC1–x 2747 - 2530
cF8 Fm 3m NaCl
a = 422.0
at 38 at. % C [1970Rud]
WC < 2776
hP2 P 6m2 WC
a = 290.62 c = 283.68
[1970Rud]
W2B < 2670
tI12 I4/mcm CuAl2
a = 556.7 ± 0.2 c = 474.4 ± 0.2
[V-C2]
a = 557.0 c = 474.4 a = 557.2 c = 474.6
W rich [1965Rud]
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B rich [1965Rud]
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. Table 2 (continued) Phase/ Temperature Range [˚C] αWB < 2170
βWB < 2665
W2B5–x(h) 2365 6 900
Pearson Symbol/ Space Group/ Prototype tI16 I41/amd αMoB
oC8 Cmcm CrB
hP12 P63/mmc2 W2B5–x
W2B5–x(r) < 900
hR21 R 3m Mo2B5–x
W1–xB3 < 2020
hP20 P63/mmc W1–xB3
Lattice Parameters [pm]
Comments/References
a = 309.73 ± 0.02 c = 1695.67 ± 0.25
W rich [V-C2]
a = 312.18 c = 1691.88 a = 310.1 c = 1695.5 a = 312.8 c = 1690.3
B rich [V-C2]
a = 312.4 b = 844.5 c = 306.0 a = 314.2 b = 850.6 c = 306.5
[1963Rud]
a = 298.2 c = 1387.3
at WB1.80 [1995Oht]
a = 298.35 c = 1388.12 a = 298.76 c = 1389.70
W rich [V-C2] for WB1.97 [1995Oht] B rich [V-C2] b)
a = 301.1 ± 0.3 c = 2093 ± 1
[V-C2] metastable?
a = 520.04 c = 633.48 a = 520.05 c = 633.56
W rich [1995Oka] B rich [1995Oka]
W rich [1965Rud]
WB4 [V-C2] W rich [V-C2] B rich [V-C2]
a)
Partially disordered V2N type, cited by various authors as partially disordered “ε-Fe2N” type [1970Rud, 1988Epi]. For high temperature thermal expansion data see [1997Bul].
b)
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. Table 3 Invariant Equilibria; comparison of experimental data from [1970Rud] with calculated data [1999Rog]
Reaction
Reaction type
L + WC1–x Ð W2C + U1 WC
L Ð βWB + (C)gr
L Ð W2B + W2C
e7
e9
L Ð W + W2C + W2B E1
L + (C)gr Ð WC + βWB
L Ð W2C + βWB
L Ð W2B + βWB + W2C
U2
e11
E2
L Ð W2C+WC+βWB E3
Landolt‐Bo¨rnstein New Series IV/11E1
Experimental data
Calculated data
Composition in at.%
Composition in at.%
C
T [˚C]
Phase
W
B
W
B
C
L
58.8
4.1
37.1 2570
60.0 11.3 28.7
WC1–x
62
0.5
37.5
61.7
0
W2C
66
1
33
67.6
1.0 31.4
WC
51
0.5
48.5
50
0
13.0 2360
44.7 41.5 13.9
50
45.2
41.8
βWB
50.4
46.6
(C)gr
0
0
L
67.6
19.9
W2B
66.7
32.4
0.9
W2C
68.8
2.4
28.8
70
L
71.0
17.0
12.0 2355
70.7 18.1 11.2
W
-
-
-
99.5
0.1
W2C
-
-
-
71.1
2.3 26.6
W2B
-
-
-
66.7 33.3
52.7
29.1
(C)gr
-
-
-
0
WC
-
-
-
50
βWB
-
-
-
51.1 45.6
L
100
0
12.5 2370
0.2 99.8
68.1 19.8 12.1 66.7 33.3
18.2 2350
2.3 27.7
0
50.1 32.5 17.3
0
50 3.3
1.6
βWB
53.2
44.5
L
62.0
24.0
W2C
-
-
-
67.7
W2B
-
-
-
66.7 33.3
0
βWB
-
-
-
52.2 45.3
2.5
14.9 2330
58.8 26.3 14.9
30
67.6
2325
1.6 30.8
52.1 45.2
2.7
59.7 26.2 14.1
2324
1.7 30.6
55.0
26.0
W2C
-
-
-
67.4
1.4 31.1
WC
-
-
-
50
0
βWB
-
-
-
51.9 45.0
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2362
0.1 99.9
68.4
19.0 2300
2384
0.4
60.1
14.0 2305
2389
0
W2C
2.3
2400
2.1
L
L
25
50.8 47.1
2587
38.3
L
3
T [˚C]
57.7 25.5 15.8
2323
50 3.1
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. Table 3 (continued) Experimental data
Calculated data
Composition in at.%
Composition in at.%
Reaction
Reaction type
Phase
W
L Ð W2B5–x + (C)gr
e12
L
30.0
63.0
7.0 2275
W2B5–x
31.4
66.6
2.0
0
0
32.4
60.2
βWB
-
W2B5–x
-
(C)gr L Ð βWB+W2B5–x + E4 (C)gr
L
(C)gr L Ð W2B5–x + ‘B4C’
e13
E5
βWB + W2B5–x Ð (C) U3 gr + αWB
βWB + (C)gr Ð αWB U4 + WC
βWB + WC Ð αWB + U6 W2B
DOI: 10.1007/978-3-540-88053-0_24 ß Springer 2009
T [˚C]
100
C
T [˚C]
30.0 62.4
7.6
2285
31.7 65.8
2.5
W
0
B
0.8 99.2
7.4 2240
35.3 56.8
7.9
-
-
50.3 49.0
0.7
-
-
33.0 63.5
3.5
-
-
21.4
72.6
6.0 2220
19.8 73.0
7.2
W2B5–x
30.5
68.7
0.8
30.0 68.0
2.0
0.4
81.8
17.8 10.7 2180
L
18.3
71
W2B5–x
29
69
(C)gr
0.5
‘B4C’
1
βWB W2B5–x
3.5
-
0
0
2
19.8 72.0
8.2
30.0 68.9
1.1
79
20
0.0 80.4 19.64
-
-
-
-
-
-
(C)gr
-
-
-
αWB
-
-
-
βWB
-
-
-
(C)gr
-
-
-
αWB
-
-
-
-
-
-
βWB
51
47
2
W2 C
67
1
32
W2 B
65
34
1
WC
51
0.5
48.5
-
50.3 49.1
0.6 3.2
0.6
51.0 46.5
2.5
50
0
2.4
0
0 50
βWB
-
-
-
WC
-
-
-
αWB
-
-
-
52.1 45.4
W2 B
-
-
-
66.7 33.3 50.0
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-
52.1 45.6 50.0
2143
0.8 31.9
66.7 33.3 50
2.6 50
52.1 45.5 67.3
2149
0.1 99.9
50.9 46.5
2150
2169
0.5 99.5
50.2 49.2
0
<2181
1.7 98.3
32.9 63.9 0
>2181
80.4 19.6
0.0
-
2259
0.5 99.5
96
WC βWB + W2C Ð WC + U5 W2B
C
L
‘B4C’ L Ð W2B5–x + (C)gr + ‘B4C’
B
2.32 2131
0.0 50.0 2.5
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. Table 3 (continued)
Reaction L + W2B5–x Ð W1–xB3
Experimental data
Calculated data
Composition in at.%
Composition in at.%
Reaction type
Phase
U7
L
6.1
92.5
W2B5–x
-
W1–xB3
W
B
C
T [˚C]
W
B
C
T [˚C] 2029
1.4 2000
16.1 83.5
0.4
-
-
29.3 70.7
0
-
-
-
18.2 81.8
0
‘B4C’
-
-
-
0
L Ð W1–xB3 + (βB) + ‘B4C’
L
3.2
96.0
0.8 1950
-
L + ‘B4C’ Ð W1–xB3 + U8 (βB)
L
-
-
-
‘B4C’
-
-
-
0.0 90.3
W1–xB3
-
-
-
18.2 81.8
(βB)
-
-
-
W2C
-
-
-
W2B
-
-
-
66.7 33.3
WC
-
-
-
50
0
50
(W)
-
-
-
100
>0
0
W2C Ð W2B + WC + E6 (W)
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-
-
-
9.2 90.8 >0
0 -
84.7 15.3
66.8
98.7
1912
9.7 0 1.3
0.1 33.2
1244
0
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B–C–W
. Fig. 1 B-C-W. Calculated phase diagram B-W
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. Fig. 2 B-C-W. Calculated phase diagram C-W
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B–C–W
. Fig. 3a B-C-W. Reaction scheme, part 1
16
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. Fig. 3b B-C-W. Reaction scheme, part 2
B–C–W
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B–C–W
. Fig. 3c B-C-W. Reaction scheme, part 3
18
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. Fig. 4 B-C-W. Liquidus surface projection
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B–C–W
. Fig. 5 B-C-W. Isothermal section at 1500˚C with alloy data from [1970Rud]
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. Fig. 6 B-C-W. Phase relations at subsolidus temperatures with alloy data from [1970Rud]
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B–C–W
. Fig. 7 B-C-W. Calculated isothermal section at 2150˚C
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. Fig. 8 B-C-W. Calculated isothermal section at 2350˚C
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B–C–W
. Fig. 9 B-C-W. Calculated isopleth W2B5–x -‘B4C’ with alloy data from [1970Rud]
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. Fig. 10 B-C-W. Calculated isopleth W2B5–x -C with alloy data from [1970Rud]
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B–C–W
. Fig. 11 B-C-W. Calculated isopleth WB-‘B4C’
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. Fig. 12 B-C-W. Calculated isopleth W-‘B4C’
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B–C–W
References [1952Gla] [1955Bre] [1959Sam]
[1963Rud]
[1965Lev]
[1965Rud]
[1966Kip] [1967Low] [1968Kip1]
[1968Kip2]
[1969Rud1]
[1969Rud2] [1970Rud]
[1973Kuh]
[1978Har] [1983Sch]
[1984Hol]
[1986Gus] [1986Loe]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–406 (1955) (Experimental, Thermodyn., 19) Samsonov, G.V., “The Interaction of Ti, Zr and W Borides with their Carbides”, Vopr. Poroshk. Metall. i Prochnosti Materialov, Akad. Nauk Ukr. SSR, (7), 72–98 (1959) (Experimental, Phase Diagram, Phase Relations, 15) Rudy, E., Benesovsky, F., Toth, L.E., “Investigation of Ternary System Between Va and VIa-Metals with Boron and Carbon” (in German), Z. Metallkd., 54(6), 345–353 (1963) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, #, 43) Levinskii, Yu.V., Salibekov, S.E., Levinskaya, M.K., “Reaction of Chromium, Molybdenum and Tungsten Borides with Carbon”, Sov. Powder Metall. Met. Ceram., 12(36), 1004–1009 (1965), translated from Poroshk. Metall. (Kiev), 12(36), 56–62 (1965) (Experimental, Phase Relations, 5) Rudy, E., Windisch, S., “Part I. Related Binary Systems, Vol. III: Systems Mo-B and W-B”, in “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems” Technical Report AFML-TR-65–2, Part I, Vol. III, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, 72 (1965) (Experimental, Phase Diagram, Phase Relations, 20) Kiparisov, S.S., Nikiforov, O.A., Borisova, N.V., “Some Properties of Alloys in the System W2B-WC” (in Russian), Poroshk. Metall. (Minsk), 353–358 (1966) (Experimental), as cited in [C.A.] Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–4 (1967) (Crys. Structure, Experimental, 5) Kiparisov, S.S., Nikiforov, O.A., Yakushin, Yu.S., “Tungsten-Rich Zone of the Tungsten - Boron Carbon System” (in Russian), Sb. Mosk. Inst. Stali Splavov., 45, 121–127 (1968) (Phase Diagram, Experimental), as cited in [C.A.] Kiparisov, S.S., Nikiforov, O.A., Borisova, N.V., “Pseudobinary Section of the Tungsten - Boron Carbon Ternary System” (in Russian), Sb. Mosk. Inst. Stali Splavov, 45, 128–131 (1968) (Experimental), as cited in [C.A.] Rudy, E., “Part V: Compendium of Phase Diagram Data” in “Ternary Phase Equilibria in Transition Metal - Boron - Carbon - Silicon Systems”, Technical Report AFML-TR-65–2, Part V, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, 192–197 (1969) (Experimental, Phase Diagram, Phase Relations, 6) Rudy, E., “Section III.K.4 W-B-C System”, in “Ternary Phase Equilibria in Transition Metal-BoronCarbon-Silicon Systems”, 655–670 (1969) (Crys. Structure, Experimental, Phase Diagram, 3) Rudy, E., “Part V, The Phase Diagram W-B-C” in “Experimental Phase Equilibria of Selected Binary, Ternary and Higher Order Systems”, Report AFML-TR-69–117, Part V, Air Force Materials Laboratory Wright Patterson Air Force Base, Ohio, 1–51 (1970) (Crys. Structure, Phase Diagram, Experimental, #, *, 22) Kuhlmann, H.S., “Determination of the Solubility of Carbon in Tungsten in the Temperature Range from 1400 to 2000˚C” (in German) Tech. Wiss. Abhandl. Osram. Ges., 11, 328–332 (1973) (Phase Diagram, Phase Relations, Experimental), as cited in [C.A.] Harsta, A., Rundquist, S., Thomas, J.O., “A Neutron Powder Diffraction Study of W2C”, Acta Chem. Scand., 32A, 891–892 (1978) (Crys. Structure, Experimental, 10) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Gustafson, P., “Thermodynamic Evaluation of C-W System”, Mater. Sci. Tech., 2, 253–658 (1986) (Phase Diagram, Thermodyn., Review, Phase Relations, 28) Lo¨nnberg, B., Lundstro¨m, T., Tellgren, R., “A Neutron Powder Diffraction Study of Ta2C and W2C”, J. Less-Common Met., 120, 239–245 (1986) (Crys. Structure, Experimental, 17)
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Epicier, T., Dubois, J., Esnouf, C., Fantozzi, G., Convert, P., “Neutron Powder Diffraction Studies of Transition Metal Hemicarbides M2C1–x - II. In Situ High Temperature Study on W2C1–x and Mo2C1–x”, Acta Metall., 36, 1903–1921 (1988) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 33) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Nagender Naidu, S.V., Rama Rao, P., “C-W (Carbon - Tungsten)” in “Phase Diagrams of Binary Tungsten Alloys”, Nagender Naidu, S.V., Rama Rao, P. (Eds.), Indian Inst. Metals, Calcutta, 37–50 (1991) (Review, Phase Diagram, Phase Relations, 100) Zakhariev, Z., Radev, D., “Dense Material Obtained on the Basis of Boron Carbide Sintered without Pressing” in “AIP Conference Proceedings 231 on Boron-Rich Solids”, Emin, D., Aselage, T.L., Switedick, A.C., Morrosin, B., Beckel, B.C. (Eds.), Albuquerque, USA, (1990), published by AIP, New York, 464–467 (1991) (Experimental, Phys. Prop., 8) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) Jonsson, S., “An Assessment of the Ti-W-C System”, Trita-Mac 519, The Royal Inst. Technology, Div. Physical Metallurgy, S-10044 Stockholm, (1993) (Phase Diagram, Phase Relations, Thermodyn.), as cited in [C.A.] Kasper, B., Max - Planck - Institute - PML, unpublished work, Stuttgart (1994) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 198–203 (1994) (Phase Diagram, Phase Relations, Review, 4) Parthe, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1-4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Duschanek, H., Rogl, P., “A Critical Assessment and a Thermodynamic Calculation of the Binary System Boron - Tungsten (B-W)”, J. Phase Equilib., 16(2), 150–161 (1995) (Phase Diagram, Phase Relations, Thermodyn., Assessment, #, 50) Ohtani, S., Ohashi, H., Ishizawa, Y., “Lattice Constants and Nonstoichiometry of WB2–x”, J. Alloys Compd., 221, L8-L10 (1995) (Crys. Structure, Experimental, 6) Okada, S., Kudom, K., Lundstro¨m, T., “Preparation and Some Properties of W2B, δ-WB and WB2 Crystals from High-Temperature Metal Solutions”, Jpn. J. Appl. Phys. 34, 226–231 (1995) (Experimental, Crys. Structure, 23) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA 5357, (1995) (Review, Phase Diagram, Phase Relations, Crys. Structure, 4) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Bulfon, C., Leithe-Jasper, A., Sassik, H., Rogl, P., “Thermal Expansion and Hardness of Czochralski-grown Single Crystal of WB2–x” in “Symposium 7, Materialwissenschaftliche Grundlagen”, Aldinger, F., Mughrabi, H. (Eds.), DGM-Informationsgesellschaft, 191–196 (1997) (Experimental, Crys. Structure, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Tungsten” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, (1998) 372–427 (Experimental, Crys. Structure, Review, Phase Diagram, 32) Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides”, Int. J. Refract. Met. Hard Mater., 17, 27–32 (1999) (Crys. Structure, Thermodyn. Calculation, Phase Relations, Phase Diagram, 6)
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[2006Lv2] [2006Lv3] [2006Wen] [2007Lv] [C.A.] [Mas2] [V-C2]
B–C–W Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides - Constitution, Thermodynamics, Compound Formation and Structural Chemistry” in “Materials Science of Carbides, Nitrides and Borides”, Gogotsi, Y., Andreev, R.A. (Eds.), Kluwer Academic Publishers, 29–46 (1999) (Thermodyn., Calculation, Phase Relations, Phase Diagram, 16) Wen, G., Li, S.B., Zhang, B.S., Guo, Z.X., “Processing of in situ Toughened B-W-C Composites by Reaction Hot Pressing of B4C and WC”, Scr. Mater., 43(9), 853–870 (2000) (Experimental, Mechan. Prop., 21) Sugiyama, S., Taimatsu, H., “Preparation of WC-WB-W2B Composites from B4C-W-WC Powders and their Mechanical Properties”, Mater. Trans., 43(5), 1197–1201 (2002) (Experimental, Mechan. Prop., 21) Sugiyama, S., Taimatsu, H., “Preparation of W-C-B Composites by Reactive Resistance-Heated Hot Pressing”, Mater. Sci. Forum, 449-452, pt.1, 309–312 (2004) (Experimental, Mechan. Prop., 16) Sugiyama, S., Taimatsu, H., “Mechanical Properties of WC-WB-W2B Composites Prepared by Reaction Sintering of B4C-W-WC Powders”, J. Eur. Ceram. Soc., 24(5), 871–876 (2004) (Experimental, Mechan. Prop., 24) Bose, K., Wood, R.J.K., “High Velocity Solid Particle Erosion Behaviour of CVD Boron Carbide on Tungsten Carbide”, Wear, 258(1-4), 366–376 (2005) (Experimental, Mechan. Prop., 31) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), Aug. 21–26, 142 (2005) (Experimental, Phase Relations, 3) Yi, L., Wen, G.W., Song, L.N., Lei, T.Q., Zhou, Y., “Microstructure and Properties of the C/W2B5 Composites Fabricated by Reaction Hot-Pressing”, Rare Metal. Mater. Eng., 34, 381–384 (2005) (Experimental, Morphology, Mechan. Prop., 5) Paderno, Y., Paderno, V., Liashchenko, A., Filipov, V., Evdokimova, A., Martynenko, A., “The Directional Crystallization of W-B-C d-transition Metal Alloys”, J. Solid State Chem., 179, 2939–2943 (2006) (Experimental, Crystal Structure, Phys. Prop., 4) Karaman, M., Sezgi, N.A., Dogu, T., Ozbelge, H.O., “Kinetic Investigation of Chemical Vapor Deposition of B4C on Tungsten Substrate”, Aiche J., 52(12), 4161–4166 (2006) (Experimental, Interface Phenomena, 0) Lv, Y., Wen, G., Zhang, B., Lei, T.Q., “Mechanical Properties and Electrical Conductivity of W-B-C Composites Fabricated by in-situ Reaction”, Mater. Chem. Phys., 97(2-3), 277–282 (2006) (Experimental, Mechan. Prop., Electr. Prop., 19) Lv, Y., Wen, G., Lei, T.Q., “Friction and Wear Behavior of C-based Composites in-situ Reinforced with W2B5”, J. Eur. Ceram. Soc., 26(15), 3477–3486 (2006) (Experimental, Mechan. Prop., 24) Lv, Y., Wen, G., Lei, TQ, “Tribological behavior of W2B5 Particulate Reinforced Carbon Matrix Composites”, Mater. Lett., 60(4), 541–545 (2006) (Experimental, Mechan. Prop., 13) Wen, G., Lv, Y., Lei, T.Q., “Reaction-formed W2B5/C Composites with High Performance”, Carbon, 44(5), 1005–1012 (2006) (Experimental, Mechan. Prop., 23) Lv, Y., Wen, G., Song, L., Lei, T.Q., “Wear Performance of C-W2B5 Composite Sliding Against Bearing Steel”, Wear, 262(5-6), 592–599 (2007) (Experimental, Mechan. Prop., 17) Chemical Abstracts - pathways to published research in the world’s journal and patent literature - http:// www.cas.org/ Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Carbon – Zirconium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction Ultra-high-temperature ceramics, composites and cermets, ZrB2/ZrC1–x and ZrB2/B4C have attracted significant attention because of a unique combination of materials properties such as high thermal conductivity, good thermal shock resistance at moderate thermal expansion and oxidation resistance. Based on early studies of the interaction between zirconium borides, zirconium carbide and carbon [1952Gla, 1955Bre, 1959Sam], phase relations in the B-C-Zr ternary system have been established by [1960Now, 1961Now], who investigated the isothermal section at 1400˚C, and by [1965Lev], who provided preliminary information on the quasibinary eutectic ZrB2-(C)gr. The most complete experimental information, however, is from [1966Rud], comprising a reinvestigation of the isothermal section at 1400˚C, a partial isothermal section for the B rich region at 1900˚C, the determination of the liquidus surface and the investigation of three isopleths, i.e. ZrB2-ZrC0.88, ZrB2-C and ZrB2-B4.5C. On the basis of these experimental results, [1966Rud] derived a reaction scheme for the ternary system and constructed a series of further isotherms at 1800, 2160, 2300, 2400, 2600, 2800 and 3000˚C, as well as a threedimensional isometric view of the ternary system. More recent and independent studies of the isopleths ZrB2-ZrC1–x and ZrB2-‘B4C’ confirmed the quasibinary and eutectic nature of these sections [1975Ord, 1988Ord, 2000Kov]. Whilst the eutectic temperature TE = 2280±30˚C of [1988Ord] for ZrB2-‘B4C’ is in acceptable agreement with the data given by [1966Rud] (TE = 2220±20˚C), the concentration of the eutectic point at 75 mol% ‘B4C’ [1988Ord], or 71.5 mol % ‘B4C’ [2000Kov] is considerably richer in ‘B4C’ than recorded by [1966Rud] (65±5 mol% B4.5C). Besides the redetermination of the ZrB2-ZrC1–x quasibinary eutectic section (TE = 2660±40˚C at 43 mol% ZrC1–x) by [1975Ord], the studies of [1979Tka, 1979Fri, 1982Unr] are mainly concerned with the physical properties in the system ZrB2-ZrC1–x and particularly with the rather low microhardness for the eutectic composition. In this case the eutectic composition is close to the value of 42 mol% Zr0.54C0.46 given by [1966Rud] but the eutectic temperature seems to be rather low when compared to [1966Rud] (2830±15˚C). Compilations of the most prominent features of the B-C-Zr phase diagram were presented by [1969Rud, 1974Upa, 1983Sch, 1984Hol, 1991Ruy, 1994McH, 1995Vil]. A full status of all information in literature for the B-C-Zr system up to 1996 was compiled in a general review of phase relations for metal-boron-carbon systems [1998Dus] including a thermodynamic calculation of the entire B-C-Zr system. For further details on the thermodynamic assessment, see [1999Rog]. Experimental details for all investigations in the B-C-Zr system are summarized in Table 1.
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Binary Systems The C-Zr system is taken from a critical assessment and thermodynamic calculation by [1995Gui] (Fig. 1). A partial low temperature phase diagram C-Zr was calculated by [1991Gus] integrating the formation of ordered superlattice phases ‘Zr2C’, ‘Zr3C2’, ‘Zr6C5’, all deriving from the NaCl type parent phase ZrC1–x. These phases were calculated to form below about 900˚C. A critical assessment and thermodynamic calculation of the B-Zr binary system is from [1988Rog] and in a refined version from [1998Dus] (Fig. 2) including a discussion on the (αZr)/(βZr) transition as well as on the temperature region of existence of ZrB12. The B-C system is adopted from a recent assessment and thermodynamic calculation by [1996Kas, 1998Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The boron rich reaction, L+‘B4C’ Ð (βB), as modelled by [1996Kas], was experimentally confirmed from floating zone experiments on several carbon-doped boron samples [2005Tan]. Crystallographic data for the binary boundary phases are summarized in Table 2.
Solid Phases No ternary B-C-Zr compounds have been reported. Mutual solid solubilities among the binary borides and carbide phases generally were found to be very small [1960Now, 1961Now, 1965Lev, 1966Rud, 1975Ord, 1982Unr, 1988Ord, 2000Kov], except for the zirconium monocarbide, for which lattice parameters in the ternary are considerably increased with respect to those of the binary [1966Rud] (Table 2). The ternary solid solubility of the nonmetal elements in Zr at the temperature of the ternary eutectic E2 (L Ð (βZr) + ZrB2 + ZrC1-x) was said to be smaller than 1 at.% altogether and a peritectoid decomposition of the (αZr) phase on heating was suggested [1966Rud] (see also section “Invariant Equilibria”). ZrB12 was absent in the ternary alloy series annealed at 1400˚C and 1600˚C; moreover ternary alloys annealed at 1900˚C revealed ZrB12 with a lattice parameter practically identical with the binary value of a = 740.8 pm from which a small or negligible homogeneity region was concluded [1966Rud].
Quasibinary Systems Three quasibinary systems of the eutectic type were established by [1966Rud]. The eutectic nature of the ZrB2-C quasibinary was earlier claimed by [1965Lev]. The eutectic temperature, 2227±30˚C as measured by optical pyrometry on pre-reacted powders through a bore hole in a directly heated graphite tube, is, however, remarkably low compared to 2390±15˚C reported by [1966Rud] and probably explains from insufficient correction for non-blackbody conditions in the experiment of [1965Lev]. Furthermore, the eutectic composition at 81 mol% C [1965Lev] is in distinct disagreement with the eutectic at 33±2 mol% C [1966Rud]. Recrystallization of the ZrB2-ZrC1-x quasibinary eutectic at temperatures close to the eutectic line was said to be extremely fast and partially or fully annealed structures resulted, if cooling rates lower than about 30 K·s–1 were employed [1966Rud]. Alloys crossing the homogeneous range of the monocarbide at 5 at.% B were found to be single phase when rapidly cooled from 2900 to 3400˚C; diboride precipitation from the zirconium monocarbide, however, was observed to DOI: 10.1007/978-3-540-88053-0_25 ß Springer 2009
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be much slower than in the corresponding Ti-containing alloys [1966Rud]. The absence of interaction between ZrC1–x and ZrB2 at T ≤ 2100˚C was confirmed from powder X-ray and microhardness investigations on samples compacted from Zr + C + ZrB2 powders presintered at 2100˚C and heated at temperatures above 2300˚C [1975Ord, 1979Fri, 1979Tka, 1982Unr]. Solid solubility of ZrB2 in ZrC1–x in samples quenched from 2400˚C was said to be about 3 mol% ZrB2 and was claimed [1975Ord] to remain virtually identical after reheating to 2100˚C, as seen from the practically unchanged lattice parameter of Zr(C,B)1–x (a = 470.0 pm). Whereas the eutectic composition at 43 mol% ZrC1–x, reported by [1975Ord] is in close agreement with that of [1966Rud] (42 ±⊊2 mol% Zr0.54C0.46), the eutectic temperature recorded at 2660±40˚C [1975Ord] is relatively low, when compared to 2830±15˚C [1966Rud] and may be explained by insufficient correction for nonblack body conditions in pyrometric recording. From the carbon content of the ZrC1–x powder used by [1975Ord], the actual concentration of the section investigated is ZrB2 - ZrC0.95. The lattice parameter of ZrC1–x (a = 470.0 pm), however, will point towards the carbon poor side of the Zr(C,B)1–x solution, where solubility of ZrB2 according to [1966Rud] was observed to be much larger, i.e. 5 mol% Zr0.33B0.67. It shall be noted, that taking the aforementioned sections through the congruent melting points of the binary compounds does not constitute true quasibinary sections as seen for instance from the slight deviation in the sections ZrB2-(C)gr (see Fig. 3) or ZrB2-ZrC1–x (see Fig. 4). Deviations are, however, small and experimentally inaccessible. Acceptable agreement exists on the eutectic nature, the eutectic temperature and to a lesser extent on the eutectic composition in the section ZrB2-B4.5C (Fig. 5) investigated by [1966Rud] (TE = 2220±20˚C, 65 mol% B0.817C0.183) and by [1988Ord] (TE = 2280"30˚C, 75 mol% B4.5C). From a directionally solidified eutectic alloy (zone melted) [2000Kov] determined the eutectic composition from chemical analysis to be 71.5 mol% ‘B4C’ (in at.% Zr6.4B77.4C16.2). The eutectic structure obtained was fibrous ZrB2 embedded in a continuous matrix of ‘B4C’ [1966Rud, 1988Ord, 2000Kov].
Invariant Equilibria A Scheil reaction scheme including nine observed ternary invariant equilibria was provided by [1966Rud]. Table 3 lists the compositions of the phases at the four-phase isothermal reactions plus the maximum points in the liquidus trough as given by [1966Rud] and compares experimentally derived data with the results of the thermodynamic calculation by [1998Dus, 1999Rog]. Figure 6 shows the Scheil diagram corresponding to the thermodynamic assessment. [1966Rud] gave the (βZr)/(αZr) transformation as a mean value of eight individual observations in the B-C-Zr ternary system at 880±15˚C suggesting a ternary peritectoid formation of (αZr) (calculated equilibrium P4 at 923˚C; see Fig. 6). As for the ternary eutectic L Ð (βZr) + ZrC1–x + ZrB2 at 1615˚C (calculated equilibrium E2 at 1652˚C; see Fig. 6), annealing reactions close to the eutectic temperature were said to be very fast and with respect to the clustering tendency of ZrB2, quenching rates >150 K·s–1 were essential to retain the structures in the as crystallized state. As far as the very boron rich region is concerned, the thermodynamic calculation renders a transition type reaction U3 at 2018˚C, L + ‘B4C’ Ð ZrB12+ (βB), rather than a ternary eutectic at 1990˚C, L Ð (βB) + ‘B4C’ + ZrB12, as reported by [1966Rud]. This discrepancy essentially results from a new assessment of the B-C system by [1996Kas] revealing a peritectic reaction Landolt‐Bo¨rnstein New Series IV/11E1
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L + ‘B4C’ Ð (βB) at 2103˚C rather than a eutectic L Ð (βB) + ‘B4C’ at 2080˚C as given by [1966Rud]. The experimentally reported very sharp drop of almost 100˚C and within less than 2 at.% from the binary reaction isotherms into the ternary eutectic (at 1990˚C [1966Rud]) seems unlikely and could not be modeled thermodynamically. It shall be noted that the peritectic binary reaction L + ‘B4C’ Ð (βB) was recently confirmed from floating zone measurements on several carbon-doped boron samples [2005Tan]. The slight deviation i.e. incompatibility with a true quasibinary for the calculated section ZrB2+C shows further consequences insofar as the experimentally observed ternary eutectic, L Ð ZrB2 + ZrC1–x + (C)gr at 2360˚C [1966Rud] turns into a transition reaction U1, L + ZrC1–x Ð ZrB2 + (C)gr calculated at 2369˚C [1998Dus, 1999Rog].
Liquidus, Solidus and Solvus Surfaces Figure 7 represents the liquidus surface projection calculated at intervals of 200 K. Figure 8 is a projection of the liquidus troughs in the Gibbs triangle as a function of temperature and as seen along the B-Zr axis in direction of boron.
Isothermal Sections Based on the experimentally determined isothermal sections at 1400, 1900˚C as well as on the liquidus projection and from the phase relations experimentally established for the three isopleths ZrB2-ZrC1–x, ZrB2-C and ZrB2-B4.5C, [1966Rud] constructed a series of isothermal sections at 1800, 2160, 2300, 2400, 2600, 2800 and 3000˚C. A set of sections, calculated by [1998Dus], is shown in Figs. 9 to 14.
Temperature – Composition Sections Isopleths ZrC-B and Zr-B0.5C0.5 were derived by [1966Rud] from the experimental data comprising isothermal sections at 1400, 1900˚C, the liquidus projection and the quasibinary sections ZrB2-C, ZrB2-ZrC1–x and ZrB2-B4.5C. In order to elucidate phase interactions in the ternary system a series of isopleths was calculated [1998Dus, 1999Rog]: ZrC1–x- B (Fig.15), Zr - B0.817C0.183 (Fig. 16) and ZrB12 - B0.817C0.183 (Fig. 17). Small deviations from congruent melting of the constituents can significantly change the extent of the phase fields as seen for the sections ZrC-‘B4C’ (Fig. 18) and ZrB2-ZrC (Fig. 19) calculated for ZrC at the C rich phase boundary.
Thermodynamics A thermodynamic calculation of the ternary B-C-Zr system by [1998Dus, 1999Rog] was based on thermodynamic assessments of the binary systems B-C [1996Kas], B-Zr [1998Dus] and C-Zr [1995Gui] as well as relying on the phase diagram data from [1966Rud] for the optimization of the thermodynamic parameters. DOI: 10.1007/978-3-540-88053-0_25 ß Springer 2009
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25
Notes on Materials Properties and Applications The combination of refractory metal like borides with high temperature ceramics constitutes attractive high temperature materials. Therefore considerable interest is devoted to the manufacture, high temperature behavior and properties of ZrB2/ZrC and ZrB2/‘B4C’ composites or ZrB2 based materials in general. Starting from the improvement of sinterability of ‘B4C’ on additions of Zr or ZrN [1978Bod], many articles deal with various techniques on pressureless sintering of dense ZrB2/B4C composites [2001Gol, 2004Tsu1, 2004Tsu2, 2004Tsu3, 2006Liu1], ZrB2-toughened ‘B4C’ [2006Liu2], structural changes in ‘B4C’ by Zr introduction [1999Liu, 2000Liu], pressureless densification of ZrB2 by ‘B4C’- additions [2006Zha], pressureless sintering of carboncoated zirconium diboride powders [2007Zhu], spark plasma sintering of ZrB2/ZrC [2007Tsu] and oxidation behavior of ZrB2-‘B4C’ composites [1993Rad]. Microstructure and kinetics of the formation of a new class of ZrB2 platelet reinforced ZrC ceramic composites were studied on material made by strongly exothermic reaction of molten Zr with B4C yielding ZrB2-platelets in a ZrC matrix with a controlled amount of residual Zr [1989Cla, 1989Whi, 1992Bre, 1991Joh]. [2007Li] reported on nanoscale ZrB2/ZrC multilayered coatings prepared by magnetron sputtering reaching a hardness of 47 GPa. Strength and deformation properties, thermal conductivity and diffusivity, friction and wear, were investigated by [1959Sam, 1979Fri, 1979Tka] on quasibinary ZrB2-ZrC1–x eutectic alloys with a grain size of 3 mm. Eutectic alloys were found to reveal minimal wear, maximal friction, maximal bending and compressive strengths and minimal hardness (13.2 to 19.6 GPa) when compared to noneutectic ZrB2-ZrC1–x compositions [1975Ord, 1979Tka] and are characterized by the absence of marked grain growth. Cracks propagated along the grain boundary. In hypereutectic alloys ZrB2 grows preferentially along the c-axis in form of long needles and lamellae [1975Ord]. Dependency of the morphology of the ZrB2-ZrC1–x eutectic as a function of cooling rates was studied by [1982Unr]. The growth direction of the eutectic is {01–1}ZrC and {010}ZrB2; accordingly the morphology of directionally solidified eutectic ZrB2-ZrC consists of columnar grains of parallel lamellae with epitaxial relationship {111}ZrC// {001}ZrB2 (growth law λ2R = (2.57±0.63) 10–18 m3·s–1; lattice parameter mismatch 4.6%) [1984Sor] (note, that a most regular eutectic growth was obtained for 48 mol% ZrB2 - the compositional discrepancy was explained by the occurrence of banding). [2002Shi] investigated the crystallographic orientation of ZrB2-ZrC composites manufactured by spark plasma sintering. Microhardness measurements on samples in both the isopleths ZrB2 - ZrC1–x and ZrB2 - ‘B4C’ near the eutectic composition were found to reveal values well below the linear combination of the binary compounds and to be strongly dependent on the eutectic crystallization conditions. For the ZrB2 - ZrC1–x eutectic, microhardness ranges from 13.2 GPa for the finely dispersed eutectic to 19.6 GPa for macrocrystalline colonies [1975Ord]. For the ZrB2 ‘B4C’ eutectic microhardness was reported to be 32 to 33 GPa [1988Ord, 2000Kov]. Dependency of the morphology of the ZrB2-‘B4C’ eutectic as a function of cooling rates was studied by [2000Kov].
Miscellaneous [1968Alp] discussed phase assemblages, microstructures, and properties of fused carbides and/or borides from refractory B-C-M systems (also for B-C-Zr) containing free graphite in Landolt‐Bo¨rnstein New Series IV/11E1
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terms of compatibility and phase diagrams. All these bodies have excellent thermal-shock resistance. Other properties (such as electrical, thermal, mechanical, chemical) can be modified by choosing different phase assemblages. Some of these materials have been cast into large shapes more than 45 cm long, and they can be machined into articles. Interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - ‘B4C’ [1993Ord] has been analyzed for transition elements M = Ti, Zr, Hf, V, Nb and Ta and correlations were found between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms. ZrB2 + ‘B4C’ + SiC ceramics showed low specific wear rate at 800˚C in air [1995Ume]. Properties of B-C-Zr alloys made by pyrolytic decomposition of ZrCl4-BCl3-natural gas mixtures at 1300–2100˚C were described by the authors of [1969Der]. C5H5-Zr(BH4)3 has proven to be a good precursor for the plasma enhanced CVD of Zr(C,B) films [1994Rei]. Zirconium and boron containing MO-PACVD coatings were developed for the application on aluminum die casting tools [1997Rie]. A sintered ZrB2 - 3–30% ‘B4C’ body may be used as crucible for Czochralski or Bridgman growth of tantalates or niobates (KNbO3) [1986Sun]. Using a compact diffraction reaction chamber, [2006Won] studied with time resolved X-ray diffraction the chemical dynamics at the combustion front for the directly ignited reactions Zr+C = ZrC, Zr+2B = ZrB2 and 3Zr+B4C = 2ZrB2+ZrC. Combustion front velocities were 5.8, 5, and 6.4 mm·s–1, respectively. The adiabatic combustion temperatures calculated (>3000 K) exceeded the Tm of Zr in all reactions.
. Table 1 Investigations of the B-C-Zr Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1952Gla]
Hot-pressing of powder mixtures (powders of ZrH2, ZrC and ZrB2,) in graphite dies at 1300 to 2400˚C. XPD on six hot-pressed samples.
[1955Bre]
Reaction between metal borides and XPD on a sample (Zr + 2B + C). graphite. Sintering of powder compacts in Mo-crucibles under 0.5 bar argon for 50 min at 1777˚C.
Compatibility of ZrB2 with C and with ‘B4C’. Small amounts of C in ZrB2+B form ZrB12.
Investigation of the sections ZrC-ZrB, [1959Sam] Samples were prepared from ZrO2 via borothermic reaction with B4C or B2O3+C. ZrB2-ZrC. Continuous solid solution of Zr LOM, XPD and microhardness. (C1–xBx), immiscibility of ZrB2-ZrC. Oxidation in air of ZrB2+ZrC, as well as friction and wear were investigated.
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1960Now] Hot-pressing of 100 binary and ternary Isothermal section at 1400˚C. Study of the [1961Now] powder compacts in C-cartridges at 1400 homogeneity region of Zr(C,B)1–x at 1400˚ to 2500˚C. A series of alloys was prepared C from lattice parameter dependencies. by reaction sintering in Ar or vacuum respectively. Samples containing B, B4C, C, ZrC and ZrB2 were reacted at 2000˚C for several h prior to final anneal at 1400˚C for 3 to 16 h. Alloys containing free Zr were annealed for 4 h at 1400˚C. Starting materials were ZrH2 containing 0.4% O), amorphous boron (96 mass% residue O, C), lampblack C. Master alloy ZrC (12.33 mass% Ctotal, 1.46 mass% free C). Some alloys were arc melted. XPD and LOM. [1965Lev]
ZrB2 was prepared via the boron carbide process at 2000 to 2200 K, MO2+‘B4C’+C = MB2+CO. Starting materials were powders of ZrO2, ‘B4C’ (containing B2O3, C), lampblack C (ash <0.05 mass%). ZrB2 was prepared from Zr powder (0.05% Hf, 0.05 Fe, 0.01 Mg, 0.1 Si, 0.005 Ti, 0.02 mass% Al), boron powder (impurities 0.0036% Fe, 0.0036 Si, 0.0003 Mg, 0.01 Cu, 0.0004 Al, 0.0006 Pb) by vacuum sintering (Zr+2B) powder compacts for 1 h at 1800 K and analyzed by XPD, LOM.
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Interaction of ZrB2 with C yielded a quasieutectic system ZrB2+C with eutectic point at 19 mol% ZrB2 and TE of 2500±30 K. Thermal analysis was performed by direct electrical heating of a graphite tube filled with ZrB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.
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. Table 1 (continued) Reference
Method/Experimental Technique
[1966Rud]
About 200 samples were prepared mainly by short duration hot pressing in graphite dies at temperatures between 1800˚C and 2200˚C. After removing the surface reaction zones, the samples were directly used in as-pressed condition for Pirani melting point (under 2.5·105 Pa He) or differential thermal analysis (graphite container under 105 Pa of He). Starting materials were high purity elemental powders (i.e. Zr containing 839 ppm O, <400 ppm Ta; zirconiumdihydride powder with 2.1 mass% H, 320 ppm C, 116 ppm N and 1300 ppm O; spectrographic grade graphite powder with impurities less than 100 ppm and boron powder of 99.55 mass% B containing 0.25 mass% Fe and 0.1 mass% C) as well as from prereacted master alloys of ZrB2 (65.2 at.% B with 0.11 mass% C), and ZrC1–x (powder with a particle size smaller than 44 mm containing 49.4 at.% C of which 0.5 at.% was in form of elemental carbon). Selected alloys from the metal-rich region (>85 at.% Zr) intended for melting point or DTA studies were electron beam or arc melted prior to the runs. Whereas specimens for DTA and melting point analyses were directly equilibrated in the equipment prior to the runs, specimens for the isothermal sections were generally annealed in a tungsten mesh furnace for 92 h at 1400˚C under a vacuum of 5·10–3 Pa or for 12 h at 1700˚C or 1900˚C respectively under 1.05 to 2·105 Pa of helium and rapidly quenched.
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Temperature/Composition/Phase Range Studied Investigation of the constitution of the B-C-Zr phase diagram employing X-ray powder diffraction, metallographic, melting point and differential thermoanalytic (> 4˚/sec) techniques. Selected alloys were equilibrated in the melting point furnace and quenched in a preheated tin bath (300˚C). Samples were chemically analyzed for free and combined carbon, boron, oxygen and nitrogen contaminants. For polishing and etching usually a slurry of alumina in 5% chromic acid was used; for alloys within the nominal compositions Zr-Zr0.6C0.4Zr0.4B0.6 anodic oxidation in an electroetching process using 10% oxalic acid was said to provide excellent phase contrast coloring the metal phase light blue, the monoboride brown, whereas the diboride remains unaffected. Specimens from the region Zr-Zr0.6C0.4Zr0.4B0.6-Zr0.3B0.7-Zr0.2C0.8 were dip-etched in a solution with 10% aqua regia and HF.
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. Table 1 (continued) Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1975Ord]
Starting materials were ZrC (88.2 mass% Zr, Ctot = 11.2 mass% C, 0.15 mass% free C) and ZrB2 (80.1 mass% Zr, Btot = 18.11 mass%, 0.06 mass% free C, 0.7 mass % (O2+N2)). Specimens were obtained from extrusion of plasticized masses in form of cylinders (3 mm diameter · 50 mm length), presintered at 2100˚C for 4 h in vacuum prior to heat treatment at 2300˚C under Ar. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the pseudo-eutectic system ZrB2+ZrC1–x with eutectic point at 57 mol% ZrB2 and TE of 2660±40˚C on 9 samples.
[1984Sor]
Preparation of 35 g sinter bars (1·1·10 cm3) from 325 mesh powders of 98.6 mass% pure ZrB2 (0.6 mass% C, 0.3% N, 0.19% O plus Fe, H) and 0.99.8% ZrC0.90 (containing Fe, Sn, O, Al, Cu) at 2300˚C for 15 min under Ar. LOM, TEM.
Determination of ZrB2+ZrC1–x eutectic morphology on floating zone melted rod by LOM, TEM. A most regular eutectic growth was obtained for 48 mol% ZrB2.
[1988Ord]
Investigation of the quasieutectic system Samples were prepared from powder ZrB2+‘B4C’ with eutectic point at 25 mol% compacts of ZrB2 and ‘B4C’ (which was vacuum annealed at 2000˚C to reduce C ZrB2 and TE of 2280±30˚C on 10 samples. content to 0.2 mass% free C). Specimens in form of cylinders (3 mm diameter · 50 mm length) were compacted with aid of 12% aqueous starch solution, presintered at 2300˚C for 2 h in vacuum prior to measurement. Specimens with high ‘B4C’ content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
[2000Kov]
Floating zone refinement under 4.9 104 Pa Determination of the eutectic point in the Ar on rounded off ingots (10·10·100 mm3) quasieutectic system ZrB2+‘B4C’. with starting compositions (in mol%) 70, 75, 80 ‘B4C’. The composition of the eutectic was obtained via chemical analysis. Commercial powders ZrB2, ‘B4C’ were vacuum annealed at 1950˚C for 1 h to reduce Cfree to <0.3 mass%. The ingots were annealed at 1700˚C for 2 h prior to zone melting. XPD, LOM, SEM, microhardness.
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. Table 1 (continued) Reference
Method/Experimental Technique
[2005Tan]
Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)). The liquidus and solidus curves to the L+(βB) field have been derived via chemical analysis.
Temperature/Composition/Phase Range Studied Confirmation of peritectic type of reaction L+B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L+(βB) field.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(βB) hR333 < 2092 [Mas2] R3m < 2075 calculated βB [1998Dus]
(C)gr < 3827 (S.P.)
Lattice Parameters [pm]
Comments/References
a = 1093.30 c = 2382.52
[1993Wer]
a = 1092.2 c = 2381.1 a = 1095.64 c = 2402.01
at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x for ZrB51 [V-C2]
hP4 a = 246.12 c = 670.90 P63/mmc C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78 a = 246.6 c = 672.0 a = 246.6 c = 672.3
at 25˚C [Mas2] [1967Low] at 2.35 at.% Bmax(2350˚C) linear ∂a/∂x, ∂c/ ∂x, [1967Low] from alloy Zr10B40C50 quenched from 2686˚C [1966Rud] contains also ZrB2 and B4C from alloy Zr13B27C60 quenched from 3135˚C [1966Rud] contains ZrB2
(C)d
cF8 a = 356.69 Fd 3m C (diamond)
at 25˚C, 60 GPa [Mas2]
(βZr) 1855 - 863
cI2 Im 3m W
[Mas2]
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a = 360.90
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
(αZr) hP2 < 863 [Mas2] P63/mmc < 866 calculated Mg [1998Dus]
Lattice Parameters [pm] a = 323.16 c = 514.75 a = 324.0 c = 516.2 a = 323.7 c = 515.9
Comments/References at 25˚C [Mas2] from alloy Zr88B10C2 (in at.%) quenched from 1615˚C, contains ZrB2, ZrC [1966Rud] from alloy Zr80B15C5 (in at.%) quenched from 1650˚C, contains ZrB2, ZrC [1966Rud]
‘B4C’ < 2450
hR45 R3m B13C2
a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 a = 560.0 from B4C-ZrB2 eutectic at 2000˚C c = 1213.66 [1988Ord]
B25C
tP68 P42m or P42/nnm B25C
a = 875.3 ± 0.4 c = 509.3 ± 1.5
[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Dus]
ZrB2 < 3245
hP3 P6/mmm AlB2
a = 316.94 ± 0.02 c = 353.07 ± 0.04 a = 316.93 c = 352.91 a = 316.7 c = 353.2 a = 316.9 c = 353.3 a = 316.9 c = 353.2 a = 317.0 c = 353.3
B rich [V-C2]
a = 316.94 c = 353.11
from ZrB2-ZrC quasibinary, annealed at 2660˚C [1975Ord] from alloy Zr42B37C21 quenched from 2830˚C contains ZrC [1966Rud] from alloy Zr32B15C18 quenched from 2647˚C contains ZrC, C [1966Rud] from alloy Zr17B33C50 quenched from 2662˚C contains C [1966Rud] from alloy Zr23B71C6 quenched from 2522˚C contains B4C [1966Rud] from ZrB2-B4C eutectic at 2000˚C [1988Ord]
a = 740.8 ± 0.2 a = 738.8 ± 0.3
[V-C2] [V-C2]
a = 316.8 c = 353.1
ZrB12 2250 - 1710 [1966Rud] 2040 - 1708 calculated [1998Dus]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
ZrC1–x cF8 < 3440 [1966Rud] Fm 3m < 3421 calculated NaCl [1998Dus]
Lattice Parameters [pm]
Comments/References x = 0.56 to 0.97, 299 K [V-C2] at x = 0.80 [V-C2] at x = 0.97 [V-C2] from alloy Zr70 B25 C5 quenched from 1676˚C contains ZrB2, (αZr) 1966Rud] from alloy Zr50B30C20 quenched from 2740˚C contains ZrB2 [1966Rud] from alloy Zr35B20C45 quenched from 2490˚C contains ZrB2, C [1966Rud] from alloy Zr61B3C36 quenched from 2898˚C single phase [1966Rud] from alloy Zr50B3C47 quenched from 3265˚C single phase [1966Rud] ZrC0.90 from ZrB2-ZrC quasibinary [1979Fri] ZrC0.85 from ZrB2-ZrC quasibinary [1979Fri] from ZrB2-ZrC quasibinary annealed at 2660˚C [1975Ord]
a = 469.80 a = 469.86 a = 470.3 a = 470.5 a = 470.0 a = 469.6 a = 470.3 a = 469.75 a = 469.4 a = 470.0
a)
Note: The (hypothetical?) structures Zr2C, Zr3C2 and Zr6C5, claimed by [1991Gus], were said to derive as superlattice structures from the NaCl type parent phase ZrC1–xSymmetries assigned were as follows: Zr2C (R 3m, Fd 3m or I41/amd), Zr3C2 (Immm or P3m1) and Zr6C5 (C2, C2/m or P31).
. Table 3 Invariant Equilibria in the B-C-Zr System; Comparison of Experimental Data from [1966Rud] with Thermodynamic Calculation, [1998Dus]. Calculated data [1998Dus] Composition, at.% Reaction (calculated) L Ð ZrB2 + ZrC1–x
Type
Phase T˚C
e2(max) L
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Zr
B
2822 42.4
38.3
ZrB2
33.3
66.7
ZrC1–x
52.1
8.2
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Experimental data [1966Rud] Composition, at.% C
Zr
19.3 42
B
C
38
20
T˚C 2830±15
0.0 33 >65 <2 39.7 53
5
42
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. Table 3 (continued) Calculated data [1998Dus] Composition, at.% Reaction (calculated) L Ð ZrB2 + (C)gr
Type -
L + ZrC1–x ÐZrB2 + (C)gr U1
L Ð ZrB2 + ‘B4C’
L Ð ZrB2 +‘B4C’ + C
L + ZrB2 Ð ZrB12 +‘B4C’
L + ‘B4C’Ð ZrB12 + (βB)
ZrB12+‘B4C’ÐZrB2 + (βB)
Phase T˚C L
U2
U3
U4
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C
no maximum in calculation, close to quasieutectic behavior
Zr 22
B
C
45
33
T˚C 2390±15
33 >65 <2
(C)gr
<1
2
>97
44.8
33.2 25
41
34
2360
49.0 52
3
45
Eutectic
L
2369 22.0
ZrC1–x
50.2
0.8
ZrB2
33.3
66.7
(C)gr
0.0
0.7
99.3 <1
<2
9.2
77.6
13.2 9
77 14 2220±20
ZrB2
33.3
66.7
‘B4C’
0.0
81.9
18.1 <1
2189 11.0
64.2
24.8 11 66 23 2165
L
2273
0.0 33 >65 <2 >97
0.0 >32 66 <2 >81 18
0.0 >32 66 <2
ZrB2
33.3
66.7
‘B4C’
0.0
80.4
19.6 <1
79 >20
(C)gr
0.0
1.7
98.3 <1
2
5.8
93.7
0.5 6
ZrB2
33.3
66.7
0.0 >32 67 <1
ZrB12
7.7
92.3
0.0 9
‘B4C’
0.0
89.3
10.7 <1
L
2034
>97
91 3
2160
90 ≤1 84 >15
4.7
95.0
0.3 2
2
‘B4C’
0.0
90.2
9.8 <1
88 >11 Eutectic
ZrB12
7.7
92.3
0.0 >9
90 <1
(βB)
0.0
98.6
1.4 <1
>98 1
7.7
92.3
0.0 10 90 <<1 1800
‘B4C’
0.0
90.4
9.6 <1
ZrB2
33.3
66.7
0.0 >32 67
L
ZrB12
(βB) L Ð (βZr) + ZrC1–x + ZrB2 E2
B
ZrB2
e4(max) L
E1
Zr
Experimental data [1966Rud] Composition, at.%
2018
1821
96 1990
89 >10 <1
0.0
98.7
1.3 >98 <1
<1
1652 86.1
12.5
1.4 88
2
(βZr)
98.8
1.0
ZrC1–x
63.0
2.7
ZrB2
33.3
66.7
L
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10
0.2 >98 <1 34.3 61
3
1615
<1 36
0.0 33 >66 <1
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. Table 3 (continued) Calculated data [1998Dus] Composition, at.% Reaction (calculated) (βZr)+ ZrB2 + ZrC1–x Ð (αZr)
Type P1
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Phase T˚C (βZr)
923
Zr
B
99.97
0.03
Experimental data [1966Rud] Composition, at.% C
0.0 -
Zr
B -
-
ZrB2
66.7
33.3
0.0 -
-
-
ZrC1–x
62.1
0.1
37.8 -
-
-
(αZr)
98.0
1.2
0.8 -
-
-
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C -
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. Fig. 1 B-C-Zr. Calculated C-Zr phase diagram
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. Fig. 2 B-C-Zr. Calculated B-Zr phase diagram
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. Fig. 3a B-C-Zr. Calculated section ZrB2 - C
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. Fig. 3b B-C-Zr. Calculated section ZrB2 - C, close-up detail
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. Fig. 4a B-C-Zr. Calculated section ZrB2 - ZrC1–x
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. Fig. 4b B-C-Zr. Calculated section ZrB2 - ZrC1–x, close-up detail
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. Fig. 5 B-C-Zr. Calculated section ZrB2 - B4.5C (B0.817C0.183)
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. Fig. 6 B-C-Zr. Calculated reaction scheme
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. Fig. 7 B-C-Zr. Calculated liquidus surface projection at 200 K intervals and invariant temperatures. Experimental data by [1966Rud] are indicated by symbols
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. Fig. 8 B-C-Zr. Projection of the liquidus troughs in the Gibbs triangle as a function of temperature and as seen along the Zr-B axis in direction of boron
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. Fig. 9 B-C-Zr. Calculated isothermal section at 1400˚C
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. Fig. 10 B-C-Zr. Calculated isothermal section at 1900˚C
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. Fig. 11 B-C-Zr. Calculated isothermal section at 2150˚C
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. Fig. 12 B-C-Zr. Calculated isothermal section at 2300˚C
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. Fig. 13 B-C-Zr. Calculated isothermal section at 2500˚C
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. Fig. 14 B-C-Zr. Calculated isothermal section at 2800˚C
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. Fig. 15 B-C-Zr. Calculated vertical section ZrC1–x - B
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. Fig. 16 B-C-Zr. Calculated vertical section Zr - B4.5C
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. Fig. 17 B-C-Zr. Calculated vertical section ZrB12 - B4.5C
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. Fig. 18 B-C-Zr. Calculated vertical section ZrC-B4C
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. Fig. 19 B-C-Zr. Calculated vertical section ZrC - ZrB2
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References [1952Gla] [1955Bre] [1959Sam]
[1960Now]
[1961Now]
[1965Lev]
[1966Rud]
[1967Low] [1968Alp]
[1969Der]
[1969Rud]
[1974Upa]
[1975Ord]
[1978Bod] [1979Tka]
[1979Fri]
[1980Ord]
Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–406 (1955) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 19) Samsonov, G.V., “The Interaction of Ti, Zr and W Borides with their Carbides”, Voprosy Poroshkovoi Metall. i Prochnosti Materialov, Akad. Nauk Ukr. SSR, (7), 72–98 (1959) (Experimental, Phase Diagram, Phase Relations, 15) Nowotny, H., Rudy, E., Benesovsky, F., “Investigation of the Zr-B-C and Zr-B-N Systems” (in German), Monatsh. Chem., 91, 963–974 (1960) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 15) Nowotny, H., Rudy, E., Benesovsky, F., “Investigations in the Systems: Hafnium - Boron - Carbon and Zirconium - Boron - Carbon” (in German), Monatsh. Chem., 92, 393–402 (1961) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 16) Levinskiy, Yu.V., Salibekov, S.E., “Interaction of Titanium, Zirconium and Hafnium Diborides with Carbon”, Russ. J. Inorg. Chem., 10(3), 319–320 (1965) translated from Zh. Neorg. Khim., 10(3), 588–591 (1965) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 6) Rudy, E., Windisch, S., “Part II. Ternary Systems, Vol. XIII. Phase Diagrams of the Systems Ti-B-C, Zr-B-C and Hf-B-C” in “Ternary Phase Equilibria in Transition Metal - Boron - Carbon - Silicon Systems”, Report AFML-TR-65–2, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, Part II. Vol. XIII, 207 (1966) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, #, *, 96) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Alper, A.M., Doman, R.C., McNally, R.N., “Fusion-Cast Carbide-Boride-Graphite Ceramics” in “Proc. Fourth International Conference on Science of Ceramics”, Maastricht, Netherlands, European Ceramic Association., 23–27 April, 1967, Stewart, G.H. (Ed.), British Ceramic Society, Stoke-on-Trent, 389–420 (1968) (Experimental, Review, 73) Dergunova, V.S., Kilin, V.S., Martynov, S.Z., Pavlycheva, T.Y.A., Sharova, A.V., “Properties of Zirconium - Boron - Carbon Alloys Obtained from the Gas Phase” (in Russian), Konstr. Mater. Osn. Grafita, 4, 9–15 (1969) (Experimental, 3) Rudy, E., “Part V. Compendium of Phase Diagram Data, Section III.K.2. Zr-B-C System” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 and 33(615)-67-C-1513, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, Part V, 618–634 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 3) Upadkhaya, G.S., “Nature of the Phase Diagram of some Transition Metals with Boron” in “Bor: Poluchenie, Struktura, Svoistva”, Mater. Mezhdunar. Simp. Boru, 4th Meeting Date 1972, Metsniereba, Tbilisi, 2, 115–123 (1974) (Review, Phase Diagram, Phase Relations, 17) Ordanyan, S.S., Unrod, V.I., “Reactions in the System ZrC-ZrB2”, Sov. Powder Metall. Met. Ceram. (Engl. Transl.), 5, 393–395 (1975), translated from Poroshk. Metall., 5(149), 61–64 (1975) (Experimental, Morphology, Phase Relations, 4) Bodnaruk, N.I., “Activated Sintering of Boron Carbide” in “Strukt. Svoistva Nov. Mater. Pokrytii”, Pilyankevich, A.N. (Ed.), 85–92 (1978) cited from abstract Tkachenko, Yu.G., Ordanyan, S.S., Yurchenko, D.Z., Yulyugin,V.K., Bovkun, G.A., Unrod, V.I., “Strength and Antifriction Properties of Alloys of the Systems MIVC - MIVB2 Over a Wide Range of Component Concentrations”, Inorg. Mater., 15(4), 549–553 (1979), translated from Izv. Akad. Nauk, Ukr. SSR, Neorg. Mater., 15(4), 704–708 (1979) (Experimental, Mechan. Prop., 11) Fridlender, Yu.A., Neshpor, V.S., Ordanyan, S.S., Unrod, V.I., “Thermal Conductivity and Diffusivity of Binary Alloys of the System ZrC-ZrB2 at High Temperatures” (in Russian), Teplophys. Vys. Temper., 17(6), 1210–1215 (1979) (Experimental, Crys. Structure, 13) Ordanyan, S.S., “Laws of Interaction in the Systems MIV, VC - MIV, VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Experimental, Thermodyn., 14)
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[1983Sch]
[1984Sor]
[1984Hol]
[1986Sun]
[1988Ord]
[1988Rog] [1989Cla]
[1989Whi]
[1990Ase]
[1991Gus]
[1991Joh]
[1991Ruy] [1992Bre]
[1993Ord] [1993Rad] [1993Wer]
[1994McH]
25
Unrod, V.I., Ordanyan, S.S., Korkin, I.V., “Effect of Cooling Rate on the Formation of the Eutectic Alloy Structure in the Me(IV)C - Me(IV)B2 Systems”, Powder Metall. Met. Ceram., 3, 224–227 (1982), translated from Poroshk. Metall., 3(231), 76–80 (1982) (Experimental, Morphology, Phase Relations, 17) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Sorrell, C.C., Beratan, H.R., Bradt, R.C., Stubican, V.S., “Directional Solidification of (Ti, Zr) Carbide (Ti, Zr) Diboride Eutectics”, J. Amer. Ceram. Soc., 67(3), 190–194 (1984) (Experimental, Phase Relations, 24) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Sunahara, K., “Crucibles for Growth of Alkali Metal Tantalates and Niobates” (in Japanese), Jpn. Kokai Tokkyo Koho (1986), 5 pp. Jpn. Patent JP 61163189 A 19860723 Showa Application: JP 85-1423 19850110. Priority: CAN 106:26178 AN 1987:26178 Ordanyan, S.S., Dmitriev, A.I., Bizhev, K.T., Stepanenko, E.K., “The Interaction in the B4C-ZrB2 System”, Powder Metall. Met. Ceram., 27(1), 38–40 (1988), translated from Poroshk. Metall., 1(301), 41–43 (1988) (Experimental, Phase Diagram, Phase Relations, Morphology, 5) Rogl, P., Potter, P.E., “A Critical Review and Thermodynamic Calculation of the Binary System Zr-B”, Calphad, 12(2), 191–204 (1988) (Review, Thermodyn., Phase Diagram, Phase Relations, 54) Claar, D.T., Johnson, W.B., Andersson, C.A., Schiroky, G.H., “Microstructure and Properties of Platelet-Reinforced Ceramics Formed by the Directed Reaction of Zirconium with Boron Carbide”, Ceram. Eng. Sci. Proc., 10(7-8), 599–609 (1989) (Experimental, Morphology, Mechan. Prop.,) cited from abstract White, D.R., Aghajanian, M.K., Claar, D.T., “Reactive Infiltration of Boron Carbide Preforms for Manufacture of Cermets and Ceramics”, Eur. Pat. Appl. (1989), 14 pp. EP 299905 A1 19890118 Patent written in English. Application: EP 88-630135 19880714. Priority: US 87-73533 19870715. CAN 110:158870 AN 1989:158870 Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U. K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97-111 (1990) (Crys. Structure, Review, Experimental, 14) Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State, 32(9), 1595–1599 (1991) (Theory, Phase Diagram, Thermodyn., Phase Relations, 18). See also Gusev, A.I., “Physical Chemistry of Nonstoichiometric Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow, 1–286 (1991) (Review, Thermodyn., Crys. Structure, Phase Diagram, Phase Relations, 102) Johnson, W.B., Nagelberg, A.S., Breval, E., “Kinetics of Formation of a Platelet-Reinforced Ceramic Composite Prepared by the Directed Reaction of Zirconium with Boron Carbide”, J. Am. Ceram. Soc., 74(9), 2093–2101 (1991) (Experimental, Kinetics, Morphology, Phase Diagram, Review, Phase Relations, 20) Ruys, A.J., Sorrell, C.C., “Ternary Phase Equilibria in the System Zr-N-B-C at 1500˚C”, Key Eng. Mater., 53-55, 92–100 (1991) (Experimental, Phase Relations, Theory, 93) Breval, E., Johnson, W.B., “Microstructure of Platelet-Reinforced Ceramics Prepared by the Directed Reaction of Zirconium with Boron Carbide”, J. Am. Ceram. Soc., 75(8), 2139–2145 (1992) (Experimental, Morphology, 14) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, 1, 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, 18) Radev, D., Zahariev, Z., “Oxidation Stability of B4C-MexBy Composite Materials”, J. Alloys Compd., 197, 87–90 (1993) (Experimental, Phase Relations, 14) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 204–209 (1994) (Phase Diagram, Phase Relations, Review, 5)
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25 [1994Rei]
[1995Gui] [1995Vil]
[1995Ume]
[1996Kas] [1997Rie]
[1998Dus]
[1998Kas]
[1999Rog] [1999Liu] [2000Liu] [2000Kov]
[2001Gol]
[2002Shi]
[2004Tsu1]
[2004Tsu2] [2004Tsu3]
[2005Tan]
[2006Liu1]
[2006Liu2]
B–C–Zr Reich, S., Messelhauser, J., Suhr, H., Erker, G., Fritze, C., “Plasma Induced Chemical Vapor Deposition of Zr(C,B) from CpZr(BH4)3 (Cp = C5H5)”, Adv. Mater., 6(9), 674–676 (1994) (Experimental, Mechan. Prop., 15) Guillermet, A.F., “Analysis and Phase Stability in the Zirconium - Carbon System”, J. Alloys Compd., 217, 69–89 (1995) (Thermodyn., Review, Phase Diagram, Phase Relations, 128) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA (1995) 5388–5391 (Phase Diagram, Phase Relations, Crys. Structure, Review, 5) Umeda, K., Enomoto, Y., Mitsui, A., Mannami, K., “Friction and Wear of ZrB2, B4C and ZrB2 + B4C + SiC Ceramics at High Temperatures in Air”, Toraiborojisuto, 40(2) 145–52 (1995) (Experimental, Mechan. Prop.) cited from abstract Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Rie, K.T., Pfohl, C., Lee, S.H., Kang, C.S., “Development of Zirconium and Boron Containing Coatings for the Application on Aluminum Casting Tools by means of MO-PACVD”, Surf. Coat. Technol., 97(1-3), 232–237 (1997) (Experimental, Morphology, Mechan. Prop., 8) Duschanek, H., Rogl, P., “The System Boron - Carbon - Zirconium” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 445–485 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 30) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides”, Int. J. Refract. Met. Hard Mater., 17, 27–32 (1999) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 6) Liu, Ch., Huang, C., “Structural Changes of Boron Carbide Induced by Zr Introduction” (in Chinese), Jinshu Xuebao, 35(8), 785–788 (1999) cited from abstract Liu, Ch., “Positron Spectroscopy of Boron Carbide Containing Metal Impurity”, Mater. Trans., JIM, 41 (10), 1293–1296 (2000) (Crys. Structure, Experimental, Electr. Prop., 18) Kovalev, A.V., Dubnik, E.M., Grigor’ev, O.N., Shaposhnikova, T.I., Martsynyuk, I.S., “Directionally Solidified Eutectic of the B4C-ZrB2 System”, Powder Metall. Met. Ceram., 39(1-2), 63–66 (2000) translated from Poroshk. Metall., 1–2(411), 71–75 (2000) (Experimental, Phase Relations, 6) Goldstein, A., Geffen, Y., Goldenberg, A., “Boron Carbide-Zirconium Boride In Situ Composites by the Reactive Pressureless Sintering of Boron Carbide-Zirconia Mixtures”, J. Am. Ceram. Soc., 84(3), 642–644 (2001) (Experimental, Morphology, Phase Relations, Phys. Prop., 10) Shim, S.H., Niihara, K., Auh, K.H., Shim, K.B., “Crystallographic Orientation of ZrB2-ZrC Composites Manufactured by the Spark Plasma Sintering Method”, J. Microscopy, Oxford, 205(Part 3), 238–244 (2002) (Experimental, Crys. Structure, Morphology, Phys. Prop., 17) Tsuchida, T., Yamamoto, S., “Mechanical Activation Assisted Self-Propagating High-Temperature Synthesis of ZrC and ZrB2 in Air from Zr/B/C Powder Mixtures”, J. Eur. Ceram. Soc., 24(1), 45–51 (2004) (Experimental, Morphology, Mechan. Prop., 15) Tsuchida, T., Yamamoto, S., “MA-SHS and SPS of ZrB2-ZrC Composites”, Solid State Ionics, 172(1-4), 215–216 (2004) (Experimental, Morphology, Mechan. Prop., 2) Tsuchida, T., Yamamoto, S., “MA-SHS of ZrC and ZrB2 in air from the Zr/B/C Powder Mixtures”, Key Eng. Mater., (Pt. 1, Euro Ceramics VIII), 264-268, 85–88 (2004) (Experimental, Morphology, Mechan. Prop., 7) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), Aug. 21–26, 142 (2005) (Experimental, Phase Relations, 3) Liu, R., Ru, H., Zhao, Y., Tang, D., “Preparation Process of ZrB2/B4C Ceramic Composites by Pressureless Sintering Based on Mechanical Mixing”, Beijing Keji Daxue Xuebao, 28(8), 762–765 (2006) (Experimental, Mechan. Prop.) cited from abstract Liu, R., Ru, H., Zhao, Y., Tang, D., “In situ Synthesis of B4C Ceramics Toughened by ZrB2 Particles”, Cailiao Yanjiu Xuebao, 20(6), 611–616 (2006) cited from abstract
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[2007Li]
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Wong, J., Larson, E.M., Waide, P.A., Frahm, R., “Combustion Front Dynamics in the Combustion Synthesis of Refractory Metal Carbides and Diborides Using Time-Resolved X-ray Diffraction”, J. Synchrotron Radiat., 13(4), 326–335 (2006) (Experimental, Phys. Prop., 30) Zhang, S.C., Hilmas, G.E., Fahrenholtz, W.G., “Pressureless Densification of Zirconium Diboride with Boron Carbide Additions”, J. Am. Ceram. Soc., 89(5), 1544–1550 (2006) (Experimental, Mechan. Prop., Morphology, Phase Relations, 30) Zhu, S., Fahrenholtz, W.G., Hilmas, G.E., Zhang, S.C., “Pressureless Sintering of Carbon-Coated Zirconium Diboride Powders”, Mater. Sci. Eng. A, 459(1-2), 167–171 (2007) (Experimental, Morphology, Phase Relations, 24) Li, D.J., Yang, J., Zhang, X.H., Cao, M., “Nanoscale ZrC/ZrB2 Multilayered Coatings Prepared by Magnetron Sputtering”, J. Vac. Sci. Technol., B, 25(2), L11-L13 (2007) (Experimental, Phase Relations, Kinetics, 10) Tsuchida, T., Yamamoto, S., “Spark Plasma Sintering of ZrB2-ZrC Powder Mixtures Synthesized by MA-SHS in Air”, J. Mater. Sci., 42(3), 772–778 (2007) (Experimental, Mechan. Prop., Morphology, 18) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Chromium – Manganese Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Nathalie Lebrun, Pierre Perrot, An Serbruyns, Jean-Claude Tedenac
Introduction The B-Cr-Mn system raised interest due to the presence of these elements in stainless steels and by its use in spring steels. The main experimental investigations of the phase equilibria are from [1969Tel] who proposed an isothermal section at 800˚C, [1976Pra] which gave an isothermal section at 1025˚C and [1982Kan] which investigated the quasibinary CrB-MnB diagram. The review of [1981Bra] reproduced the isothermal sections at 800 and 1025˚C. No Calphad assessment has been carried out. The main experimental works are presented in Table 1.
Binary Systems The Cr-Mn system is accepted from the assessment of [1986Ven]. A thermodynamic evaluation of this system has been carried out by [1993Lee]. However, in this assessment, the ordering observed in the α’ and σ phases at low temperature was not taken into account. The B-Cr system is accepted from [1992Rog1]. A thermodynamic evaluation of this system was carried out by [2002Cam]. The Cr4B compound, presented with a stability uncertain is definitively non existent. It is given by [1969Tel] with the crystal parameters which are very close to those of Cr2B of the Mg2Cu type. The B-Mn system is accepted from [1992Rog2]. Although Mn4B appears in [Mas2] with a stability considered as uncertain, it is definitively non existent because it is presented with the same structure (oF40) as Mn2B of the Mg2Cu type. Another problem arises with the same Mn2B phase. [1992Rog2] accepts the existence of two allotropic forms: a tI12, CuAl2 type, 33 at.% B, stable below 1580˚C and an oF48, Mg2Cu type at 32 at.% B, stable below 1285˚C. Actually, there is only one boride Mn2B which undergoes a polymorphic transition at 1285˚C. Phase with structure tI12 is the high temperature form whereas phase with structure oF48 is the room temperature form. This point of view is supported by the fact that the high temperature modification has never been observed in the B-Cr-Mn diagrams investigated at 800 and 1025˚C.
Solid Phases The solid phases are presented in Table 2. As explained above, and in agreement with [1998Rog], the compounds Cr4B and Mn4B, considered as non existent, are not shown in the table. Landolt‐Bo¨rnstein New Series IV/11E1
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Cr2B and Mn2B exist under two crystallographic forms, of the Mg2Cu and CuAl2 type. The Mg2Cu type is the only stable form for Cr2B and is the low temperature form for Mn2B; the CuAl2 type is metastable for Cr2B and is stable between 1285 and 1580˚C for Mn2B. The solid solution of the Mg2Cu type exists in the whole composition range; the extension of the CuAl2 type solid solution has not been investigated. A CrMnB compound with the so-called “Mn4B” (actually CuAl2) structure was prepared by [1964Bro]. At high temperature, CrB has the CrB type structure (Cmcm) and MnB has the FeB type structure (Pnma). Both structures are closely related [1982Kan] and are characterized by trigonal prisms sharing their rectangular faces. The CrB type structure is interpreted as a successive unit cell twinning of a cubic close packed structure and the FeB type structure as that of a hexagonal close-packed structure. At low temperatures, CrB has the MoB type structure whereas MnB has a closely related structure based on the random stacking of CrB and FeB type structure.
Quasibinary Systems The quasibinary CrB-MnB system shown in Fig. 1 presents four solid solutions whose stability domains were experimentally investigated by [1982Kan]. The original figure was slightly modified to take into account that the border between the low temperature and high temperature solid solutions must be two-phase. The transition between CrB(h) and MnB(h) which are closely related is reversible, though sluggish. For instance, the Cr0.8Mn0.2B with the MnB(h) structure, annealed at 1200˚C for 2 days undergoes a transformation to the CrB(h) structure, but some detectable amount of the MnB(h) structure remains. At temperatures lower than 1050˚C appears, near the MnB side, another kind of structure MnB(r), very similar to that of MnB(h), with extinctions of some lines in the XRD pattern.
Isothermal Sections The isothermal section of the B-Cr-Mn system at 800˚C has been experimentally investigated by [1969Tel, 1970Tel] and accepted by [1976Pra, 1981Bra]. The existence of three complete solid solutions is reported, namely (Cr,Mn)B2, (Cr,Mn)3B4 and a solid solution between Cr2B and Mn4B. The existence of this last solid solution is hardly credible and it is probable that the compound wrongly labeled Mn4B with the crystal structure accepted for Mn2B be actually Mn2B. The existence of a complete solid solution CrB2-MnB2 at 800˚C rests on a very poor experimental argumentation. Actually, the MnB2 compound undergoes a eutectic decomposition below 1100˚C and the Cr1–xMnxB2 solid solution exists probably for x ≲ 0.5 at 800˚C. The CrB-MnB section presents two solid solutions, which is in good agreement with the vertical section investigated later by [1982Kan]. The isothermal section at 800˚C, given in Fig. 2, based on [1969Tel] was modified to take into account the accepted binary diagrams the probable existence of the solid solution (Cr,Mn)2B and the non existence of the MnB2 compound at this temperature. As a consequence, the original solid solution Cr2B-Mn4B has been deleted to be replaced by the solid solution Cr2B-Mn2B. On the other hand, the solid solution Cr1–xMnxB2 was considered stable for x ≲ 0.5. The isothermal section at 1025˚C, based on [1976Pra] is presented in Fig. 3. The original figure has been modified to take into account the fact that the binary compounds CrB, MnB, Mn2B and Cr2B are recognized as DOI: 10.1007/978-3-540-88053-0_26 ß Springer 2009
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nearly stoichiometric. For instance, the compound Mn4B, presented as a non stoichiometric compound (28.6 to 32.5 at.% Mn) is actually Mn2B. The solid solutions (Cr,Mn)3B4 and (Cr,Mn)B2 have been added. This last one is not stable in the whole composition range because the MnB2 compound is stable only above 1100˚C. Figure 3 presents only the oF40 solid solution Cr2B-Mn2B characterized by a ratio B/(Cr+Mn) = 0.33 but it is probable that the tI12 solid solution Cr2B-Mn2B characterized by a ratio B/(Cr+Mn) = 0.32 be stable in the Mn rich corner (tI12 Cr2B is metastable), but this fact has never been verified. The tie lines in the two-phase domains (Cr,Mn) + (Cr,Mn)2B, (Cr,Mn)2B + (Cr,Mn)B, (Cr,Mn)B + (Cr,Mn)3B4, (Cr,Mn)3B4 + (Cr,Mn)B2 have never been determined. Their general behavior may be inferred from the observation of [1952Hae]: Cr, which has the smaller atomic number, has a stronger affinity for B than Mn.
Notes on Materials Properties and Applications The main experimental investigations are reported in Table 3. The magnetization of the CrMn2B4 is given at 10 A·m2·kg–1 [1983Bus]. This compound is presented with an unknown structure. Actually, it is a member of the (Cr,Mn)3O4 solid solution. CrB2 is antiferromagnetic with a Neel temperature TN = 86 K, whereas MnB2 is ferromagnetic with a Curie temperature reported between 140 and 157 K [1969Cas]. Cr0.5Mn0.5B2 presents a strong electronic heat content, but no magnetic ordering was reported.
Miscellaneous B, Cr and Mn are important constituents of spring steels. The thermal treatment called “ausforming” (hot rolling after annealing) has beneficial effect on fatigue behavior. The yield strength, tensile strength and yield ratio are always higher than those of the conventionally quenched specimens [1972Omo]. Ausforming produces a slightly higher ductility at higher strength level. The major increase of yield strength is attributable to the increase of dislocation density [1976Omo1]. The lowering of the crack growth rate observed by ausforming [1976Omo2] is related to the disappearance of the intergranular fracture. The endurance limit of the ausformed steel and its strength after 106 cycles of fatigue [1973Omo] is also sensibly improved.
. Table 1 Investigations of the B-Cr-Mn Phase Relations, Structures and Thermodynamics Reference
Temperature/Composition/Phase Range Studied
Method/Experimental Technique
[1952Hae] Boron distribution ratio. Theory
Equilibrium (Cr,Mn)B-(Cr,Mn)2B
[1964Bro]
XRD (X-ray diffraction)
CrMnB
[1969Tel]
XRD, microscopical analyses
800˚C, B-Cr-Mn isothermal section
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. Table 1 (continued) Reference [1976Pra]
Temperature/Composition/Phase Range Studied
Method/Experimental Technique XRD, thermal analyses
1025˚C, Cr-CrB-MnB-Mn isothermal section
[1982Kan] XRD powder diffraction, thermal analysis
800–2100˚C, CrB-MnB quasibinary system
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βB) < 2092
hR333 R3m βB
(αCr) < 1863
cI2 Im3m W
a = 288.48
at 25˚C [Mas2]
(α’Cr)
tI2 I4/mmm α’Cr
a = 288.2 c = 288.7
at 25˚C, HP [Mas2]
(δMn) 1246 - 1138
cI2 Im3m W
a = 308.0 a = 308.1
[Mas2] at 1138˚C [1992Rog2]
(γMn) 1138 - 1100
cF4 Fm3m Cu
a = 386.0 a = 386.3
[Mas2] at 1100˚C [1992Rog2]
(βMn) 1100 - 727
cP20 P4132 βMn
a = 631.52
[Mas2]
(αMn) < 727
cI58 I43m αMn
a = 891.26
at 25˚C [Mas2]
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Dissolves 4.5 at.% Mn at 1800˚C and 2.5 at.% Cr at 1830˚C [1998Rog] [1993Wer]
a = 1093.30 c = 2382.52 a = 1093.02 pure B [1976Lun] c = 2381.66 a = 1096.37 ± 0.02 at 2.7 at.% Cr, “CrB41” [V-C2] c = 2384.77 ± 0.04
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. Table 2 (continued) Phase/ Temperature Range [˚C] (Cr1–xMnx)2B Cr2B < 1870
Pearson Symbol/ Space Group/ Prototype oF48 Fddd Mg2Cu (Mn4B)
Mn2B(r) < 1285 Cr5B3 < 1900
tI32 I4/mcm Cr5B3
Cr1–xMnxB CrB(h) 2095 - 1100
oC8 Cmcm CrB
CrB(r) ≲ 1000
tI16 I41/amd αMoB
(Cr1–xMnx)3B4
oI14 Immm Ta3B4
Cr3B4 < 2075
Mn3B4 < 1750
Lattice Parameters [pm]
Comments/References 0 ≤ x ≤1
a = 425.35 b = 741.33 c = 1470.6 a = 420.7 b = 728.7 c = 1455.7
The Vegard’s law is obeyed [1976Pra] x = 0 [1992Rog1, 1998Rog]
a = 546.40 c = 1011.0
[1992Rog1, 1998Rog]
a = 297.0 b = 786.5 c = 293.6 a = 299.5 b = 776.4 c = 293.1
x = 0 [1976Pra]
a = 294.93 c = 1572.8
[1998Rog]
x = 1 [1998Rog] 32 at.% B [1992Rog2]
x = 0.54 [1976Pra]
0≤x≤1 a = 295.25 b = 298.56 c = 1302.2 a = 296.3 b = 303.4 c = 1282.2
x = 0 [1992Rog1, 1998Rog]
x = 1 [1992Rog2, 1998Rog]
Cr2B3
oC20 Cmcm V2B3
a = 302.64 b = 1811.5 c = 295.42
[1998Rog]
CrB2 < 2200
hP3 P6/mmm AlB2
a = 297.32 c = 307.25
[1998Rog]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
CrB4 < 1450
oI10 Immm CrB4
a = 286.82 b = 474.99 c = 547.88
[1998Rog]
Mn2B(h) 1580 - 1285
tI12 I4/mcm CuAl2
a = 514.8 c = 420.4
[1998Rog]
MnB(h) 1890 - 1050
oP8 Pnma FeB
a = 555.6 b = 297.7 c = 414.6 a = 558.9 b = 293.3 c = 415.6
[1998Rog]
Mn0.6Cr0.4B
[1976Pra]
MnB(r) < 1050
oC8 Cmcm CrB
a = 302.34 b = 767.59 c = 295.68
[1998Rog] Random stacking of CrB and FeB type structures [1982Kan]
MnB2 1990 - 1075
hP3 P6/mmm AlB2
a = 300.89 c = 303.84
[1998Rog]
MnB4 < 1375
mC10 C2/m MnB4
a = 548.5 b = 537.5 c = 274.5 β = 122.47˚
[1998Rog]
α’, CrMn2(h) 926 - 600
Similar to (αMn) -
60.2 to 65.3 at.% Mn [1986Ven]
α”, CrMn2(r) < 600
Similar to (αMn) -
60.2 to 65.3 at.% Mn [1986Ven]
σ, CrMn3(h2) 1312 - 999
tP30 P42/mnm σCrFe
-
73 to 78 at.% Mn [1986Ven]
σ’, CrMn3(h2) 1006 - 800
tP30 P42/mnm ordered σCrFe
-
75 to 79 at.% Mn [1986Ven]
σ”, CrMn3(r) ≲ 800
tP30 P42/mnm ordered σCrFe
a = 886 to 886.8 c = 459 to 460.1
75 to 79 at.% Mn [1986Ven]
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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype
Phase/ Temperature Range [˚C]
Lattice Parameters [pm]
Comments/References
CrB6
t**
a = 546.8 c = 715.2
[1969Tel] Metastable
Cr2B (m)
tI12 I4/mcm CuAl2
a = 518.5 c = 431.6
[1970Tel] Metastable
. Table 3 Investigations of the B-Cr-Mn Materials Properties Method / Experimental Technique
Reference
Type of Property
[1969Cas]
Electronic heat content, magnetic susceptibility measurements
Cr1–xMnxB2 (x = 0, 0.5, 1), magnetic and electronic structure
[1971Han]
Low temperature heat capacity, Debye temperature
< 200 K, CrB, CrB2, MnB, MnB, analysis of the band structure
[1972Omo]
Yield strength, tensile strength, yield ratio measurements
Influence of ausforming on hardness and tempering behavior
[1973Omo]
Fatigue strength, hardness, cyclic fatigue measurements
Influence of ausforming on ductility and yield strength fatigue crack propagation and fatigue behavior
[1976Omo1] Yield strength, micrography
Influence of ausforming on the ductility and yield strength
[1976Omo2] Fatigue strength, fractography
Influence of ausforming on the crack growth rate and fatigue behavior
[1983Bus]
CrMn2B4
Magnetization
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. Fig. 1 B-Cr-Mn. The quasibinary CrB-MnB system
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. Fig. 2 B-Cr-Mn. The isothermal section at 800˚C
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. Fig. 3 B-Cr-Mn. The isothermal section at 1025˚C
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References [1952Hae]
[1964Bro] [1969Cas]
[1969Tel]
[1970Tel] [1971Han] [1972Omo] [1973Omo]
[1976Lun] [1976Omo1] [1976Omo2]
[1976Pra] [1981Bra] [1982Kan]
[1983Bus] [1986Ven] [1992Rog1]
[1992Rog2]
[1993Lee] [1993Wer]
[1998Rog] [2002Cam]
Haegg, G., Kiessling, R., “Distribution Equilibria in Some Ternary Systems Me1-Me2-B and the Relative Strength of the Transition-Metal-Boron Bond”, J. Inst. Met., 81, 57–60 (1952) (Crys. Structure, Experimental, Phase Relations, 6) Brown, B.E., Beerntsen, D.J., “Refinement of an Iron Chromium Boride with the Mn4B Structure”, Acta Crystallogr., 17, 448–450 (1964) (Crys. Structure, Experimental, 12) Castaing, J., Caudron, R., Toupance, G., Costa, P., “Electronic Structure of Transition Metal Diborides”, Solid State Commun., 7, 1453–1456 (1969) (Electronic Structure, Experimental, Magn. Prop., 5) Telegus, V.S., Kuzma, Yu.B., “X-Ray Investigation of the Chromium-Manganese-Boron System” (in Ukrainian), Vys. Lviv. Ordena Lenina Derzhav. Univ. Franka, Ser. Khim., (11), 21–24 (1969) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 11) Telegus, V.S., “On the Reactions of Transition Metals of the Groups VI and VII with Boron” (in Russian), Khim. i Khim. Tekhnol. (7), 96–100 (1969) (Review, Phase Relations, 9) Hanson, B.D., Mahnig, M., Toth, L.E., “Low Temperature Heat Capacities of Transition Metal Borides”, Z. Naturforschung A, 26(4), 739–746 (1971) (Crys. Structure, Experimental, Thermodyn., 28) Omori, M., Tanabe, K., Kuroki, K., “On Tempering Behavior of Ausforming Mn-Cr-B Spring Steel”, J. Soc. Mater. Sci., Jpn., 21(226), 665–670 (1972) (Experimental, Morphology, Thermodyn., 10) Omori, M., Tanabe, K., Kuroki, K., “Effect of Ausforming on the Fatigue Strength of Mn-Cr-B Spring Steel” (in Japanese), J. Soc. Mater. Sci., Jpn., 22(235), 353–358 (1973) (Experimental, Morphology, Thermodyn., 14) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Omori, M., Kawamata, K., “Study of Strengthening Mechanism of Ausformed Mn-Cr-B Spring Steel” (in Japanese), J. Soc. Mater. Sci., Jpn., 25(269), 157–164 (1976) (Experimental, Morphology, 48) Omori, M., Kawamata, K., “Effect of Ausforming on Fatigue Crack Growth Behaviour of Mn-Cr-B Spring Steel” (in Japanese), J. Soc. Mater. Sci., Jpn., 25(269), 165–171 (1976) (Experimental, Morphology, Thermodyn., 18) Pradelli, G., Gianoglio, C., “Solid-State Equilibria in the Cr-Mn-B System” (in Italian), Metall. Ital., 68, 191–194 (1976) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, 7) Brandes, E.A., Flint, R.F., “The B-Cr-Mn (Boron-Chromium-Manganese) System”, Bull. Alloy Phase Diagrams, 2(1), 107 (1981) (Phase Diagram, Phase Relations, Review, 5) Kanaizuka, T., “Phase Diagram of Pseudobinary CrB-MnB and MnB-FeB Systems: Crystal Structure of the Low-Temperature Modification of FeB”, J. Solid State Chem., 41(2), 195–204 (1982) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 22) Buschow, K.H.J., van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of Metallic Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1–22 (1983) (Magn. Prop., Optical Prop., 23) Venkatraman, M., Neumann, J.P., “The Cr-Mn (Chromium-Manganese) System”, Bull. Alloy Phase Diagrams, 7(5), 457–462 (1986) (Phase Diagram, Crys. Structure, Review, #, 30) Rogl, P., “Cr-B-N (Chromium-Boron-Nitrogen)” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), ASM, Materials Park, Ohio, USA, 20–25 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Experimental, Review, *, 11) Rogl, P., “Mn-B-N (Manganese-Boron-Nitrogen)” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), ASM, Materials Park, Ohio, USA, 60–63 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, *, 6) Lee, B.-J., “A Thermodynamic Evaluation of the Cr-Mn and Fe-Cr-Mn Systems”, Metall. Trans. A, 24A(9), 1910–1933 (1993) (Assessment, Phase Diagram, Phase Relations, Thermodyn., 63) Werheit, H., Kuhlmann, U., Laux, M., Lundstroem, T., “Structural and Electronic Properties of Carbon-doped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Rogl, P., “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), ASM Publ., Materials Park, OH (1998) Campbell, C.E., Kattner, U.R., “Assessment of the Cr-B System and Extrapolation to the Ni-Al-Cr-B Quaternary System”, Calphad, 26(3), 477–490 (2002) (Phase Diagram, Assessment, Calculation, #, 44)
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B–Cr–Mn Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Chromium – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Anatoliy Bondar, Oleksandr Dovbenko, Volodymyr Ivanchenko, Artem Kozlov
Introduction Due to the high oxidation resistance, high hardness and wear-resistance at elevated temperature, high thermal and chemical stability, borides of transition refractory metals are used as special structural materials and coatings. Alloys of the B-Cr-Ti systems are used as coatings with high corrosion and oxidation resistance and high wear-resistance. The potential applications of TiB2-Cr alloys as special cutting tools are also studied. Chromium can be used as alloying element for titanium matrix composites. Mixed diborides of chromium and titanium have a greater chemical stability than original diborides taken separately [1958Mee, 1960Che, 1982Kov, 1982Tak, 1993Tel, 2000Oka]. Some preliminary studies of the ternary B-Cr-Ti system have been performed by [1960Che, 1960Kov, 1962Fed]. Phase equilibria in the quasibinary section TiB2-CrB2 were studied by [1954Pos, 1958Mee, 1959Sev, 1987Zda, 1993Tel]. The reaction of liquid titanium with CrB2 was examined by [1965Gar]. It was shown that the alloying of chromium decreases the melting point of the Ti-TiB eutectic by 50˚C. The reaction of solid titanium with CrB2 was studied by [2006Zhu]. [1973Fil] reported that in the temperature interval from 1149 to 1482˚C the phases based on Cr and Cr2B are in equilibrium in the alloys Cr-0.5Ti-0.25B, Cr-0.5Ti-0.5B, and Ti-1.0Ti-0.5B (at.%). A possible increase of the density of the (Cr,Ti)B2 sintered composition using Cr, Cu, Ni, Cu+Ni, Cu+Ni+Mn additions was studied by [1979Kov, 1982Kov, 1982Tak]. Crystal structure, microhardness, electrical resistivity of single crystal (Cr0.97Ti0.03)3B4 were reported by [2000Oka]. Experimental studies concerning phase equilibria and crystal structure examinations are summarized in Table 1. Thermodynamic properties of ternary alloys of the system have not been studied and phase relationships have not been thermodynamically assessed.
Binary Systems The B-Cr, B-Ti, and Cr-Ti binary systems are accepted from [1992Rog], [2008Wit], and [2000Zhu], respectively.
Solid Phases [1960Kov] claimed the existence of three ternary compounds in the system: Cr2TiB2 (melts incongruently at 1960˚C), Ti2CrB4 (melts incongruently at 2450˚C), and Ti9CrB18 (melts congruently at 2590˚C). These results were obtained using results of visual thermal analyzes Landolt‐Bo¨rnstein New Series IV/11E1
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of alloys located along the Cr-TiB2 join. No data concerning the crystal structures of these boride phases were presented. The cited data are inefficient to assert that these ternary borides really exist. These phases may be solid solutions based on binary borides (e.g. Ti3B4 and TiB2) or phases stabilized by contaminations together with deviations of the alloy composition from the nominal composition (as Ti9CrB18). The similar situation is with a ternary phase detected by [1982Tak] using XRD in the sample TiB2-30 vol% Cr after sintering at 1600˚C. The intensity of the TiB2 XRD reflexes for the TiB2-Cr samples was reduced about half of that before sintering, Cr disappeared and unknown new lines with an intensity nearly the same as the TiB2 lines were observed. The unknown phase(s) in the TiB2-Cr samples was thought to be a kind of boride, but they could not be identified with the ASTM card database by [1982Tak]. No data on the composition, crystal structure or the obtained X-ray pattern were presented. [1954Pos, 1958Mee, 1959Sev] reported the existence of a continues series of solid solutions between the isostructural CrB2 and TiB2 phases with the composition dependences of the lattice parameters obeyed the Vegard’s low. [1987Zda] also reported that the lattice parameters of the diborides decrease linearly with increasing Cr content, Δa = 302.7 – 2.5·10–4WCr and Δc = 323.3 – 6.9·10–4WCr (pm), where WCr is mass% Cr. Crystallographic data for the binary boundary phases are summarized in Table 2.
Quasibinary Systems [1993Tel] studied the phase equilibria along the CrB2-TiB2 section. Their results confirm the existence of a continuous solid solution above 2100˚C reported by [1954Pos, 1958Mee, 1959Sev]. But at lower temperatures the CrB2-TiB2 system is characterized by a miscibility gap with a critical point at 2000˚C and 20 mol% CrB2. In the area of the break in the solubility, the CrB2 content in TiB2 does not exceed 1 mol% below 1800˚C. CrB2 dissolves from 45 mol% TiB2 at 1500˚C to 55 mol% TiB2 at 1700˚C. The solubility of TiB2 in CrB2 at room temperature was estimated as 15–20 mol%. Dilatometric investigation of specimens with the total CrB2 contents of 33, 75, and 90 mol% showed that melting did not start at temperatures below the melting point of CrB2 (2200˚C). This means that in the quasibinary CrB2-TiB2 system there is neither a minimum on the solidus curve nor a eutectic crystallization. The CrB2-TiB2 quasibinary system is presented in Fig. 1 in accordance with [1993Tel]. [1960Kov] claimed that Cr-TiB2 section is quasibinary. This result is in contradiction with the data of [1973Fil] who showed, that the Cr-1Ti-0.5B (at.%) alloy has phase constituents (Cr)+Cr2B in the temperature interval of 1149–1482˚C. Keeping in mind the existence of the continuous solubility in the solid state between Cr and Ti, on one hand, and CrB2 and TiB2, on the other hand, the assumption about the location of the quasibinary section in the Cr-TiB2 join looks doubtful.
Invariant Equilibria Only one invariant equilibrium in this ternary system is reliably experimentally confirmed. It is the critical point of the decomposition of the (Ti,Cr)B2 solid solution established by [1993Tel] at 66.67B-6.66Cr-26.64Ti (at.%) and 2000˚C. The peritectic reactions of the formation of TiCr2B2 (1960˚C), Ti2CrB4 (2450˚C), and a congruent melting of Ti9CrB18 DOI: 10.1007/978-3-540-88053-0_27 ß Springer 2009
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(2590˚C), as well as eutectic reactions at 5–7 mol% TiB2 (1580˚C) and 95 mol% TiB2 reported by [1960Kov] and also a eutectic at 60Cr-27Ti-13B (at.%) with melting temperature approximately at 1300 to 1400˚C presented in [1960Che] are based on the questionable data required experimental confirmation. Possibly some of them are monovariant reactions.
Liquidus, Solidus and Solvus Surfaces Data on liquidus and solidus of the B-Cr-Ti system are absent except for the CrB2-TiB2 quasibinary system in [1993Tel]. The sections Cr-Ti2B and Cr-TiB2 are presented in [1960Che, 1960Kov], respectively, but their solidus and liquidus lines are not worthy of note (see a discussion in the Quasibinary Systems section). Data on the solubility are very scarce. The solubility of TiB2 in (Cr) at 1580˚C is equal to 0.5 mol% TiB2 and Cr solubility in TiB2 is 2 mol% [1960Kov]. According to [1960Che], the Ti and B solubility in (Cr) is about 0.34 at.% Ti and 0.34 at.% Ti and 0.16 at.% B in the Cr-Ti2B section (the alloys were cooled from the melting point with the rate of 40–50˚C·min–1).
Temperature – Composition Sections Approximate partial vertical section Cr-Ti2B was presented by [1960Che], but the Ti2B phase does not exist in the accepted binary system. Also authors [1960Che] studied alloys located in the Cr-TiB2 join, but no data about phase equilibria in this section were presented. [1960Kov] reported a Cr-TiB2 hypothetical diagram. But it is not quasibinary (see chapter Quasibinary systems) and not accepted in the present assessment. Also the existence of their three ternary compounds also is not accepted.
Notes on Materials Properties and Applications Diborides of Ti and Cr as well as (Ti1–xCrx)B2 compositions find application as protective coatings, cutting tools, pump parts for hot and erosive media, armor shields, rocket nozzles, and cathodes of Al smelting [1991Kny1, 1991Kny2, 1987Zda]. (Ti,Cr)B2 densificated by addition of Cu or Cu-Ni binders have high wear-resistance [1979Kov, 1982Kov]. Technological aspects of the production of B-Cr-Ti powders using different approaches as well as the effect of the liquid metal phase used to improve adhesion of wear-resistant powder to the surface protection on the physico-mechanical properties of a plasma-sprayed protective coatings have been reported by [1991Kny1, 1991Kny2]. CrB can be used as reinforcement of Ti-base metal matrix composites, but due to insufficient chemical stability it needs protective coatings [1998Gor]. This means that a two-phase (Ti)+CrB region does not exist in the B-Cr-Ti system. A summary of experimental investigation of properties is given in Table 3.
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Miscellaneous The reaction of liquid Ti with CrB2 is characterized by high wetting and extreme fluidity of the melt [1965Gar]. The contact interaction of chromium diboride CrB2 with titanium proceeds through the mechanism of solid state exothermal reaction and yields phases based on TiB and Cr. The controlling process of the interaction is diffusion of boron atoms throughout the TiB layer [2006Zhu]. B-Cr-Ti coatings deposited by magnetron sputtering of target with composition of TiB+Ti9Cr4B+TiCr2 has an amorphous structure [2005Sht]. Self-division behavior of TiB particles formed through the chemical reaction in the Cr-Ti system powder compact containing CrB was studied by [2004Yos].
. Table 1 Investigations of the B-Cr-Ti Phase Relations, Structures and Thermodynamics Reference [1954Pos]
Method/Experimental Technique
Sintering of mixture of TiB2 and CrB2 by pass of very high current, 50 mol% TiB2 + 50 mol% CrB2
X-ray diffraction
[1958Mee] Light microscopy, X-ray diffraction, microhardness
[1959Sev]
Temperature/Composition/Phase Range Studied
Hot pressing in graphite pressforms followed by sintering at 2180˚C under pressure, homogenization at 2000–2100˚C, (Ti1–xCrx)B2, 1≤x≤0.9
X-ray diffraction, density measurements Initial elements were mixed and pressed under 150 bar at 2000˚C for 2 h, syntesis at 2000˚C for 2 h was followed by annealing at 2200˚C for 30 min, (Ti1–xCrx)B2, 1≤x≤0.9
[1960Che] Light microscopy, microhardness, TA
Mixtures of elemental components were pressed, heated to 800–900˚C and melted in a BeO crucible / Cr-TiB2 with Cr content up to 50 at.%, and Cr-Ti2B with Cr content up to 60 at.%
[1960Kov] Light microscopy, X-ray diffraction, visual TA
up to 2600˚C, Cr-TiB2
[1962Fed] Light microscopy, X-ray diffraction, DTA, Vickers hardness
Mixtures of Cr, Ti, and B were arc melted / CrTiB2, Ti-CrB2, Ti-Cr5B3, TiCr2-CrB sections, and isopleths at 5, 10, 30, and 45 at.% Ti
[1965Gar]
Light microscopy, EPMA
Liquid Ti+CrB2
[1973Fil]
X-ray diffraction of extraction products 1149–1482˚C, Cr-0.5Ti-0.25B, Cr-0.5Ti-0.5B, Cr-1Ti-0.5B (at.%)
[1987Zda] X-ray diffraction, EPMA
(Ti,Cr)B2 with 0.4–2.33 mass% Cr
[1993Tel]
Hot pressing at 2000˚C for 30–120 min followed by annealing at 1500–2000˚C, dilatometry up to 2300˚C, CrB2-TiB2
SEM, X-ray diffraction, macro-X-ray analysis, dilatometry
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. Table 1 (continued) Reference
Method/Experimental Technique
[2000Oka] Optical microscopy, SEM, EDX, EPMA, induction coupled plasma emission spectroscopy
Temperature/Composition/Phase Range Studied Single crystals of (Cr0.97Ti0.03)3B4 prepared by the high temperature solution method using Al flux. The mixture was held at 1650˚C for 5 h, cooled to 1000˚C at the rate of 50˚C·h–1 and than quenched to the room temperatue
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] αβ, (Ti1–xCrx)
Pearson Symbol/ Space Group/ Lattice Parameters Prototype [pm] cI2 Im 3m W
(αCr) < 1863 (βTi) 1670 - 882
Comments/References
a = 325.95 a = 325.34 a = 324.45 a = 323.55 a = 322.94 a = 321.40 a = 288.48
0≤x≤1 x = 0.06, x = 0.08, x = 0.10, x = 0.12, x = 0.14, x = 0.16 [V-C2] pure Cr at 25˚C [Mas2]
a = 330.65
pure Ti at 25˚C [Mas2]
(α’Cr)
tI2 I4/mmm α’Cr
a = 288.2 c = 288.7
at 25˚C, HP [Mas2]
(αTi) < 882
hP2 P63/mmc Mg
a = 295.06 c = 468.35
at 25˚C [Mas2]
(ωTi)
hP3 P6/mmm ωTi
a = 462.5 c = 281.3
at 25˚C, HP → 1 atm [Mas2]
(βB) < 2092
hR333 R 3m βB
a = 1093.02 c = 2381.66
pure B (99.9999%) [1976Lun]
a = 1096.37 ± 0.02 c = 2384.77 ± 0.04 a = 1092.70 ± 0.13 c = 2388.65 ± 0.32
at 2.7 at.% Cr, “CrB41” [V-C2]
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. Table 2 (continued) Phase/ Temperature Range [˚C] Cr2B < 1870
Pearson Symbol/ Space Group/ Lattice Parameters Prototype [pm] oF48 Fddd Mn4B
a = 425.35 b = 741.33 c = 1470.6 a = 146.92 b = 739.9 c = 426.6
Comments/References [1992Rog]a)
[1974Lug]
Cr5B3 < 1900
tI32 I4/mcm Cr5B3
a = 546.40 c= 1011.0
[1992Rog]
CrB(h) 2095 - 1000
oC8 Cmcm CrB
a = 296.89 b = 786.89 c = 293.33
[1992Rog]
CrB(r) < 1000
tI16 I41/amd αMoB
a = 294.93 c = 1572.8
[1992Rog] transposition type structure
Cr3B4 < 2075
oI14 Immm Ta3B4
a = 295.25 b = 298.56 c = 1302.2 a = 300.04 b = 1301.8 c = 295.16 a = 295.6 b = 1292.9 c = 291.8
[1992Rog]a)
Cr2B3
oC20 Cmcm V2 B 3
(Ti1–xCrx)B2
hP3 P6/mmm AlB2
CrB2 < 2200 TiB2 < 3225
a = 302.64 ± 0.05 b = 1811.5 ± 0.4 c = 295.42 ± 0.04 a = 299 c = 314 a = 297.3 c = 307.0 a = 302.8 to 304.0 c = 322.8 to 323.4
[2000Oka]
(Cr0.97Ti0.03)3B4 [2000Oka]
[1987Oka]
0≤x≤1 [V-C2] [1974Lug] 65.5 to 66.7 at.% B [1986Mur]
CrB4 < 1450
oI10 Immm CrB4
a = 286.82 b = 474.99 c = 547.88
[1992Rog]a)
TiB < 2180
oP8 Pnma FeB
a = 610.5 b = 304.8 c = 454.2
[1986Mur]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Lattice Parameters Prototype [pm]
Comments/References
TiB(m) metastable
oC8 Cmcm CrB
a = 323 b = 856 c = 305
[1992Gra]
Ti3B4 < 2200
oI14 Immm Ta3B4
a = 325.9 b = 1373 c = 304.2
56.1 at.% B [1986Spe]
TiB25
tP52 P42/nnm TiB25
a = 883.0 c = 507.2
[V-C2] metastable?
λ1, TiCr2(h2) 1359 - 1269
hP12 P63/mmc MgZn2
λ3, TiCr2(h1) 1271 - 804
hP24 P63/mmc MgNi2
λ2, TiCr2(r) < 1223
cF24 Fd 3m MgCu2
ω (Ti,Cr) < 450 a)
hP3 P 3m1 ω (Ti,Cr)
63.89–66.86 at.% Cr [2000Zhu] C14 structure at 25˚C [V-C2]
a = 493.2 c = 800.5
a = 493.2 c = 1601.0
62.7–66.87 at.% Cr [2000Zhu] C36 structure at 25˚C Ti1.12Cr2 [V-C2]
a = 693.2
63.3–67.1 at.% Cr [2000Zhu] C15 structure at 25˚C TiCr1.9 [V-C2]
a = 461.6 c = 282.7
3–9 at.% Cr [1987Mur] at 4.6 at.% Cr metastable
Note: structure setting standardized according to Typix [1994Par]
. Table 3 Investigations of the B-Cr-Ti Materials Properties Reference
Method / Experimental Technique
[1958Mee] Microhardness, Tompson bridge for measurement of specific resistivity, compressive and bend strength, dilatometry
Type of Property Composition dependencies of resistivity, thermal expansion and mechanical properties
[1960Che] High temperature oxidation by measuring Oxidability of the Cr-TiB2 alloys with at.% Cr < 30 of the weight of alloys after keeping at 1200˚C for 1, 3, 10, and 70 h
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. Table 3 (continued) Reference
Method / Experimental Technique
Type of Property
[1979Kov] Density, electrical resistivity, hardness, ultimate compressive strength
Determination of the optimal conditions for preparation of the (Ti,Cr)B2 -based alloys with high wear-resistance
[1982Kov] Rockwell hardness, Vickers hardness, compressive strength, relative wearresistance
Physico-mechanical properties of (Ti,Cr)B2 and composites produced with use of Cu, Ni, Cu+Ni, and Cu+Ni+Mn binders
[1982Tak]
Opportunity to use TiB2+30 vol% Cr composites as cutting tools
Vickers hardness
[2000Oka] Microhardness, direct-current four-probe technique
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. Fig. 1 B-Cr-Ti. Quasibinary system CrB2-TiB2
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References [1954Pos] [1958Mee]
[1959Sev]
[1960Che]
[1960Kov]
[1962Fed]
[1965Gar] [1973Fil] [1974Lug]
[1976Lun] [1979Kov]
[1982Kov]
[1981Cre] [1982Tak]
[1986Mur] [1987Mur]
[1986Spe]
[1987Oka] [1987Zda]
Post, B., Glaser, F.W., Moskowitz, D., “Transition Metal Diborides”, Acta Metall., 2, 20–25 (1954) (Crys. Structure, Experimental, Morphology, 14) Meerson, G.A., Samsonov, G.V., Kotel’nikov, R.B., Voinova, M.S., Evteeva, I.P., Krasnenkova, S.D., “Some Properties of the Alloys of Borides of Refractory Transition-Group Metals” (in Russian), Zh. Neorg. Khim., 3(4), 898–903 (1958) (Crys. Structure, Experimental, Morphology, Phase Relations, Phys. Prop., 18) Sevastyanov, N.G., Epelbaum, V.A., Gurevich, M.A., Ormont, B.F., Zhdanov, G.S., “On the Conditions of Formation of a Continuous Series of Solid Solutions in the Pseudobinary System CrB2-TiB2”, J. Appl. Chem. USSR (Engl. Transl.), 32(3), 1990–1993 (1959), translated from Zh. Prikl. Khim., 32, 1941–1945 (1959) (Experimental, Phase Relations, 10) Cherkashina, N.V., Nedumov, N.A., Shamrai, F.I., “Alloys in the Titanium-Chromium-Boron System”, Russ. J. Inorg. Chem., 5(9), 985–988 (1960), translated from Zh. Neorg. Khim., (5), 2025–2031 (1960) (Experimental, Kinetics, Morphology, 11) Koval’chenko, M.S., Samsonov, G.V., Yasinskaya, G.A., “Alloys of the Borides of the Transition Metals with Other Metals” (in Russian), Izv. Akad. Nauk SSSR, Otdel. Tekh. Nauk, Metall. Topl., 2(2), 115–119 (1960) (Experimental, Phase Diagram, Phase Relations, 12) Fedorov, T.F., Nedumov, N.A., Polyakova, M.D., Shamrai, F.I., “Certain Data on the Ternary System Ti-B-Cr”, Powder Metall. Met. Ceram., (6), 435–441 (1962), translated from Poroshk. Metall., 6(12), 42–49 (1962) (Experimental, Morphology, Phase Diagram, Phase Relations, Thermodyn., 9) Garfinkle, M., Davis, H.M., “Reaction of Liquid Ti with Some Refractory Compounds”, Trans. Q. Am. Soc. Met., 58, 520–530 (1965) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 9) Filippi, A.M., “Stability of Reactive and Refractory Metal Borides in Ternary Cr-Base Alloys”, J. LessCommon Met., 30, 153–158 (1973) (Crys. Structure, Experimental, 5) Lugscheider, E., Knotek, O., Reimann, H., “The Ternary System Nickel-Chromium-Boron” (in German), Monatsh. Chem., 105(1), 80–90 (1974) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, #, *, 32) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Koval’chenko, M.S., Ochcas, L.F., Vinokurov, V.B., “The Hot Pressing of an Alloy Based on (Ti, Cr)B2 and Study of its Properties”, J. Less-Common Met., 67(2) 297–301 (1979) (Morphology, Phase Relations, Electr. Prop., Mechan. Prop., 9) Koval‘chenko, M.S., Ochkas, L.F., Yurchenko, D.Z., “Wear-Resistant Hard Metals Based on the Binary Ti-Cr Diboride”, Powder Metall. Met. Ceram., 21(11), 876–879 (1982), translated from Poroshk. Metall., 11(239), 54–57 (1982) (Experimental, Phase Relations, 11) Crespo, A.J., Tergenius, L.-E., Lundstro¨m, T., “The Solid Solution of 4d, 5d and Some p Elements in βRhomhedral Boron”, J. Less-Common Met., 77, 147–150 (1981) (Crys. Structure, Experimental, 12) Takatsu, S., Ishimatsu, E., “Sintering and Properties of TiC-, TiN-, and TiB2-Alloys with Refractory Metal Binder”, Int. J. Refract. Hard Mater., 1(2), 75–80 (1982) (Experimental, Mechan. Prop., Morphology, 11) Murray, J.L., Liao, P.K., Spear, K.E., “The B-Ti (Boron-Titanium) System”, Bull. Alloy Phase Diagrams, 7, 550–555 (1986) (Phase Diagram, Phase Relations, Review, Thermodyn., 48) Murray, J.L., “The Cr-Ti (Chromium-Titanium) System” in “Phase Diagrams of Binary Titanium Alloys”, Murray, J.L., (Ed.), ASM International, Metals Park, Ohio, 68–78 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Review, Thermodyn., 48) Spear, K.E., McDowell, P., McMahon, F., “Experimental Evidence for the Existence of the Ti3B4 Phase”, J. Am. Ceram. Soc., 69(1), C4-C5 (1986) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 13) Okada, S., Atoda, T., Higashi, I., “Structural Investigation of Cr2B3, Cr3B4, and CrB by Single-Crystal Diffractometry”, J. Solid State Chem., 68, 61–67 (1987) (Crys. Structure, Experimental, 9) Zdaniewski, W.A., “Solid Solubility Effect on Properties of Titanium Diboride”, J. Am. Ceram. Soc., 70(11), 793–797 (1987) (Crys. Structure, Experimental, Review, 33)
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[1991Kny2]
[1992Gra]
[1992Rog]
[1993Tel]
[1994Par]
[1998Gor]
[2000Oka]
[2000Zhu]
[2004Yos] [2005Sht]
[2006Zhu]
[2008Wit]
[Mas2] [V-C2]
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Knyshev, E.A., Obabkov, N.V., Kiselev, V.A., Beketov, A.R., Boev, E.A., Gorbatov, I.N., Rozhkov, A.S., “Technological Aspects of Production of Titanium-Chromium Boride Powders. I.”, Proc. “10th Intl. Symp. on Boron, Borides and Related Compounds”, Albuquerque, NM, USA, 1990, AIP Conf. Proc., (231), 505–507 (1991) (Morphology, Experimental, 4) Knyshev, E.A., Nechepurenko, A.S., Klinskaya, N.A., “Technological Aspects of Production of TitaniumChromium Boride Powders. II.”, Proc. “10th Intl. Symp. on Boron, Borides and Related Compounds”, Albuquerque, NM, USA, 1990, AIP Conf. Proc., (231), 508–501 (1991), (Morphology, Experimental, Interface Phenomena, 4) De Graef, M., Lofvander, J.P.A., McCollough, C., Levi, C.G., “The Evolution of Metastable Bf Borides in a Ti-Al-B Alloy”, Acta Metall., 40(12), 3395–3406 (1992) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 22) Rogl, P., “The B-N-Cr System” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, Ohio, USA (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, Review, 11) Telle, R., Fendler, E., Pettsov, G., “The Quasiternary TiB2-W2B5-CrB2 System and its Possibilities in Evolution of Ceramic Hard Materials”, Powder Metall. Met. Ceram., 32(3), 240–248 (1993), translated from Poroshk. Metall., 3(363), 58–69, (1993) (Experimental, Morphology, Phase Diagram, Phase Relations, 20) Parthe, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1–4, Gmelin Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Gorsse, S., Chaminade, J.P., Petitcorps, Y.L., “In Situ Preparation of Titanium Base Cimposites Reinforced by TiB Single Crystals Using a Powder Metallurgy Technique”, Composites, Part A, 29A, 1229–1234 (1998) (Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 10) Okada, S., Yu, Y., Kudou, K., Shishido, T., Tanaka, T., Higashi, I., Lundstro¨m, T., Fukuda, T., “Synthesis and Investigation of Large Crystals of (Cr1–xTMx)3B4 with TM = Ti, V, Nb, Ta, Mo, and W”, J. Solid State Chem., 154, 45–48 (2000) (Crys. Structure, Experimental, 10) Zhuang, W., Shen, J., Liu, Y., Ling, L., Shang, S., Du, Y., Schuster, J.C., “Thermodynamic Optimization of the Cr-Ti System”, Z. Metallkd., 91(2), 121–127 (2000) (Assessment, Phase Relations, Thermodyn., 51) Yoshihiro, T., Tsuchiyama, T., Takaki, S., “Self-division Behavier of TiB Particles in TiB/Ti Composite”, Mater. Trans., 45(5), 1640–1645 (2004) (Crys. Structure, Morphology, Experimental, 13) Shtansky, D.V., Kiryukhantsev-Korneev, F.V., Sheveiko, A.N., Bashkova, I.A., Malochkin, O.V., Levashov, E.A., D’yakonova, N.B., Lyasotsky, I.V., “Structure and Properties of Ti-B-N, Ti-Cr-B-(N), and Cr-B-(N) Coatings Deposited by Magnetron Sputtering of Targets Prepared by Self-Propagating High-Temperature Synthesis”, Phys. Solid State, 47(2), 252–262 (2005) (Crys. Structure, Electronic Structure, Experimental, Mechan. Prop., Morphology, Phase Relations, 38) Zhunkovskii, G.L., Evtushok, T.M., Kotenko, V.A., Mazur, P.V., Gordienko, S.P., “Reaction of BoronContaining Materials with Titanium During Self-Propagating High-Temperature Synthesis”, Powder Metall. Met. Ceram., 45(3-4), 163–167 (2006), translated from Poroshk. Metall., 3–4(448) 67–72 (2006) (Experimental, Morphology, Phase Relations, 7) Witusiewicz, V.T., Bondar, A.A., Hecht, U., Rex, S., Velikanova, T.Ya., “The Al-B-Nb-Ti System I. ReAssessment of the Constituent Binary Systems B-Nb and B-Ti on the Basis of New Experimental Data”, J. Alloys Compd., 448, 185–194 (2008) (Assessment, Calculation, Experimental, Phase Diagram, Phase Relations, Thermodyn., #, *, 70) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Chromium – Zirconium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Vasyl Tomashik, Tamara Velikanova, Mikhail Turchanin, Dmytro Pavlyuchkov
Introduction The B-Cr-Zr ternary system is investigated already for fifty years but not much is known about the phase relationships of this system. A hypothetic phase diagram of the system Cr-ZrB2 was constructed by [1960Kov]. The solidus line along this section was determined in [1965Kos]. The existence of two compounds ZrCrB2 and ZrCr2B2 was determined in [1960Kov]. The formation of ZrCrB2 was confirmed by [1965Kos]. [1967Vor, 1969Vor] did not observe ternary compounds at 900˚C in this system. A partial phase diagram of the Cr-ZrB2 system at the Cr side has been constructed by [1975Shu]. Two-phase alloys are formed in the system CrB2-ZrB2 [1954Pos, 1958Kot, 1958Mee]. [2001Hil] indicated that the ternary compounds Zr2CrB6 and ZrCr2B6 exist in this system. Phase relations in the B-Cr-Zr system at 900˚C were investigated by [1967Vor, 1969Vor]. Table 1 lists the numerous experimental works on phase equilibria and crystal structure of the B-Cr-Zr system.
Binary Systems The B-Cr phase diagram is based on the data given by [1992Rog]. The CrB4, CrB2, Cr3B4, CrB, Cr5B3 and Cr2B compounds exist in this system. The B-Zr phase diagram is from [1988Rog, 1998Dus]. Two compounds ZrB2 and ZrB12 were determined in this system. The Cr-Zr binary system is taken from [2003Per] and only one compound ZrCr2, so called Laves phase, which crystallizes in three different modifications, exists in this system.
Solid Phases Crystallographic data of all unary phases and binary and ternary compounds are listed in Table 2. Some discrepancies were observed in the mutual solubility of the ZrB2 and CrB2 phases in the ternary B-Cr-Zr system. The solubility of ZrB2 in CrB2 is equal to about 20 mol% and the solubility of CrB2 in ZrB2 is negligible according to [1958Mee], but according to [1954Pos] the solubility of ZrB2 in CrB2 is 0–5 mol% and CrB2 in ZrB2 is 10–15 mol% at 2100 ± 100˚C. The third component did not display appreciable solubility in the binary phases at 900˚C, except the solubility of Zr in CrB2 - up to 6 at.% [1967Vor]. Landolt‐Bo¨rnstein New Series IV/11E1
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The solubility of ZrB2 in (Cr) at 1450˚C is equal to 0.2 mol% and reaches almost 0.3 mol% at 1550˚C [1975Shu]. According to the data of [1973Fil] ZrB and ZrB2 are stable in the B-Cr-Zr quenched alloys containing 0.5 or 0.25 at.% B and 0.5 at.% Zr. Taking into account the proximity of the lattice parameters and the same structure of the ternary compounds Zr2CrB6 and ZrCr2B6, determined by [2001Hil], it is possible to indicate that these compounds belong to the continuous solid solution in the CrB2-ZrB2 section which could be represented by the formulae Zr1–xCrxB2.
Quasibinary Systems According to the data of [1960Kov, 1965Kos, 1975Shu] the section Cr-ZrB2 is quasibinary. The chromium rich part of the phase diagram of this system after the data of [1975Shu] is shown in Fig. 1. This phase diagram is characterized by the eutectic, which contains 9.7 mol% (19 mass%) ZrB2 and crystallizes at 1550 ± 10˚C.
Isothermal Sections The isothermal section at 900˚C is given in Fig. 2 based on the data of [1967Vor, 1969Vor] with corrections according to the accepted binary systems. As CrB6 is not confirmed in the B-Cr system, but instead CrB4 is established [1992Rog], the CrB6+ZrB2 two-phase field shown in [1967Vor, 1969Vor] is changed for the CrB4-ZrB2 two-phase field in the isothermal section. Also, the ZrB phase which was presented by [1967Vor, 1969Vor] in the isothermal section at 900˚C does not exist at this temperature in the accepted B-Zr binary system.
Notes on Materials Properties and Applications The linear dependence of the hardness was determined for the alloys of the Cr-ZrB2 quasibinary system containing up to 30 vol% of the inclusive ZrB2 [1979Shu]. The electric resistance of Cr increases upon the addition up to 5 mol% ZrB2 [1965Kos].
Miscellaneous The samples hardened by the fine-dispersed discharges of the zirconium and chromium borides were obtained in the B-Cr-Zr ternary system at the tempering of the melts [1990Igo]. The dendritic structure becomes abrupt fine-grain structure at the growing of the cooling rate.
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. Table 1 Investigations of the B-Cr-Zr Phase Relations, Structures and Thermodynamics Reference [1954Pos]
Method/Experimental Technique XRD, chemical analyses
Temperature/Composition/Phase Range Studied 2100 ± 100˚C / CrB2-ZrB2
[1958Mee] XRD, chemical analyses, metallography, microhardness measurement, compression resistance, conductivity and thermal coefficient of the longitudinal expansion measurements
Room temperature / CrB2-ZrB2
[1960Kov] Visual TA, XRD, metallography
Up to 3000˚C / Cr-ZrB2
[1965Kos]
Visual TA, XRD, metallography, electric resistance measurement
Up to 2210˚C / Cr-ZrB2
[1967Vor, 1969Vor]
XRD
900˚C / B-Cr-Zr
[1973Fil]
XRD, fluorescent chemical analyses
1149, 1371 and 1482˚C / B-Cr-Zr alloys containing 0.5 or 0.25 at.% B and 0.5 at.% Zr
[1975Shu] DTA, XRD, metallography, chemical analysis Up to 2200˚C / Cr-ZrB2 up to 14.6 mol% ZrB2 [1979Shu] Hardness and microhardness measurements Room temperature / Cr-ZrB2 [1990Igo]
Metallography, electron microscopy, X-ray spectroscopic microanalysis
Up to 1500˚C / B-Cr-Zr at 0.3 mol% ZrB2
[2001Hil]
XRD
Up to 1500˚C / Zr2CrB6 and ZrCr2B6 composition
. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] (βB) < 2092
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Pearson Symbol/ Space Group/ Prototype hR333 R 3m βB
Lattice Parameters [pm] a = 1093.02 c = 2381.66 a = 1096.37 c = 2384.77
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Comments/References dissolves 2 at.% Cr at 1830˚C [Mas2] pure B (99.9999%) [1976Lun] at 2.7 at.% Cr, “CrB41” [V-C2]
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αCr) < 1863
cI2 Im 3m W
a = 288.48
at 25˚C [Mas2]; dissolves 0.6 at.% Zr at 1546–1592˚C and 1 at.% B at 1630˚C [Mas2, 2003Per]
(α’Cr)
tI2 I4/mmm α’Cr
a = 288.2 c = 288.7
at 25˚C, HP [Mas2]
(ωZr)
hP3 P6/mmm ωTi
a = 503.6 c = 310.9
at 25˚C, HP Ð 1 atm [Mas2]
(βZr) 1855 - 863
cI2 Im 3m W
a = 360.90
[Mas2]; dissolves 8.4 at.% Cr at 1332˚C and 1.1 at.% B at 1669˚C [1998Dus, 2003Per]
(αZr) < 863
hP2 P63/mmc Mg
a = 323.16 c = 514.75
at 25˚C [Mas2] dissolves 0.5 at.% Cr at 831˚C and 1.1 at.% B at 897˚C [1998Dus, 2003Per]
Cr2B < 1870
oF48 a = 425.35 Fddd b = 741.33 Mn4B (Mg2Cu) c = 1470.6
[1992Rog] a)
Cr5B3 < 1900
tI32 I4/mcm Cr5B3
a = 546.40 c = 1011.0
[1992Rog]
CrB(h) 2095 - 1000
oC8 Cmcm CrB
a = 296.89 b = 786.89 c = 293.33
[1992Rog]
CrB(r) ≲ 1000
tI16 I41/amd αMoB
a = 294.93 c = 1572.8
[1992Rog] transposition type structure
Cr3B4 < 2075
oI14 Immm Ta3B4
a = 295.25 b = 298.56 c = 1302.2
[1992Rog] a)
Cr2B3
oC20 Cmcm V2B3
a = 302.64 ± 0.05 [V-C2] b = 1811.5 ± 0.4 c = 295.42 ± 0.04
CrB2 < 2200
hP3 P6/mmm AlB2
a = 297.32 c = 307.25
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. Table 2 (continued) Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
CrB4 < 1450
oI10 Immm CrB4
a = 286.82 b = 474.99 c = 547.88
ZrB2 < 3245
hP3 P6/mmm AlB2
a = 316.94 ± 0.02 B rich [V-C2, 1998Dus] c = 353.07 ± 0.04
ZrB12 cF52 Fm 3m 2040 - 1708 calculated [1998Dus] UB12 γZrCr2(h2) 1677 - 1625
hP12 P63/mmc MgZn2
βZrCr2(h1) 1625 - 1546
hP24 P63/mmc MgNi2
αZrCr2(r) < 1560
cF24 Fd 3m MgCu2
* τ, Zr1–xCrxB2
*P36 Pbam Y2ReB6
[1992Rog] a)
a = 316.93 c = 352.91
B poor [V-C2, 1998Dus]
a = 740.8 ± 0.2 a = 738.8 ± 0.3
[V-C2] [V-C2]
a = 510.2 a = 828.9 a = 511.1 a = 834.1 a = 510.0 a = 1661 a = 720.4
C14 structure at 20˚C [2003Per] at 300˚C [2003Per] C36 structure [2003Per] C15 structure [2003Per]
a = 876.4 ± 0.6 at x = 0.67 [2001Hil] b = 1103.9 ± 0.9 c = 329.8 ± 0.2 a = 872.58 ± 0.01 at x = 0.33 [2001Hil] b = 1101.2 ± 0.2 c = 327.98 ± 0.04
a)
note: structure setting standardized according to Typix [1994Par]
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B–Cr–Zr
. Fig. 1 B-Cr-Zr. Part of the phase diagram of the Cr-ZrB2 quasibinary system
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. Fig. 2 B-Cr-Zr. Isothermal section at 900˚C
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References [1954Pos] [1958Kot]
[1958Mee]
[1960Kov]
[1965Kos]
[1967Vor]
[1969Vor]
[1973Fil] [1975Shu]
[1976Lun] [1979Shu]
[1988Rog]
[1990Igo]
[1992Rog]
[1994Par]
[1998Dus]
[2001Hil] [2003Per]
Post, B., Glaser, F.W., Moskowitz, D., “Transition Metal Diborides”, Acta Metall., 2, 20–25 (1954) (Experimental, Crys. Structure, Morphology, 14) Kotel’nikov, R.B., “About the Formation of the Continuous Solid Solutions in Systems Formed by the Carbides, Nitrides, Borides and Silicides of the Transition Metals”, (in Russian), Zh. Neorg. Khim., 3(4), 841–846 (1958) (Experimental, Phase Relations, 10) Meerson, G.A., Samsonov, G.V., Kotel’nikov, R.B., Voynova, M.S., Evteeva, I.P., Krasnenkova, S.D., “Some Properties of the Alloys of Borides of Refractory Transition-Group Metals” (in Russian), Zhur. Neorg. Khim., 3(4), 898–903 (1958) (Experimental, Morphology, Phase Relations, Phys. Prop., 18) Kovalchenko, M.S., Samsonov, G.V., Yasinskaya, G.A., “Alloys of the Borides of the Transition Metals with Other Metals” (in Russian), Izv. Akad. Nauk SSSR, Otdel. Tekh. Nauk, Metall. Topl., (2), 115–119 (1960) (Experimental, Phase Diagram, Phase Relations, 12) Kostikov, V.I., “On the Structure of the Phase Diagram of ZrB2-Cr”, Inorg. Mater. (Engl. Trans.), 1(8), 1175–1177 (1965), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 1(8), 1285–1288 (1965) (Experimental, Morphology, Phase Diagram, Phase Relations, 2) Voroshilov, Yu.V., Lakh, V.I., Stadnyk, B.I., Kuz’ma, Yu.B., “The Ternary Systems ZirconiumChromium-Boron and Zirconium-Tungsten-Boron”, Inorg. Mater. (Engl. Transl.), 3, 1390–1392 (1967), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 3(9), 1597–1600 (1967) (Experimental, Crys. Structure, Phase Relations, #, *, 11) Voroshilov, Yu.V., Kuz’ma, Yu.B., “Reaction of Zirconium with the Transition Metals and Boron”, Powder Metall. Met. Ceram., 8(11), 941–944 (1969), translated from Poroshk. Metall., 11(83), 94–98 (1969) (Experimental, Phase Diagram, Phase Relations, #, *, 19) Filippi, A.M., “Stability of Reactive and Refractory Metal Borides in Ternary Chromium-Base Alloys”, J. Less-Common Met., 30, 153–158 (1973) (Experimental, Crys. Structure, 5) Shurin, A.K., Panarin, V.E., “Phase Diagrams and Structure of Alloys Belonging to the Quasibinary Systems Cr-ZrB2 and Cr-HfB2” (in Ukrainian), Dop. Akad. Nauk Ukr. RSR, (1), 86–90 (1975) (Experimental, Phase Diagram, Phase Relations, #, *, 6) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Shurin, A.K., Dmitrieva, G.P., Panarin, V.E., “Hardness of Quasibinary Eutectic Alloys with Inclusive Phases” (in Russian), Metallofizika, Sborn. Nauch. Rabot Inst. Metallofiz., (76), 81–85 (1979) (Experimental, Mechan. Prop., Phase Relations, 6) Rogl, P., Potter, P.E., “A Critical Review and Thermodynamic Calculation on the Binary System Zirconium - Boron”, Calphad, 12(2), 191–204 (1988) (Review, Phase Diagram, Crys. Structure, Thermodyn., 54) Igolkina, L.S., Indenbaum, S.V., Samgina, O.N., Samelyuk, A.V., Revyakin, A.V., Fridman, A.G., “The Effects of Melting Cooling Rate, Heat Treatment and Deformation on the Structure of the ChromiumBased Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 46–50 (1990) (Experimental, Crys. Structure, Kinetics, Morphology, 9) Rogl, P., “The B-N-Cr System” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, OH., 20–25 (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, Review, 11) Parthe´, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1–4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Duschanek, H., Rogl, P., “The System Boron - Carbon - Zirconium” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 445–485 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 30) Hillebrecht, H., Gebhardt, K., “New Variants of the Y2ReB6-Type: Zr2CrB6 and ZrCr2B6” (in German), Z. Kristallogr., (Suppl. 18), 156 (2001) (Experimental, Crys. Structure, *, 4) Perrot, P., “Cr-Zr (Chromium - Zirconium”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.15393.1.20, (2003) (Assessment, Crys. Structure, Phase Diagram, Thermodyn., 14)
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Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Iron – Molybdenum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Kostyantyn Korniyenko, Anatoliy Bondar
Introduction Boron-iron-molybdenum alloys gained scientific and practical interest due to a positive effect that boron additions have on the microstructure and physical properties of alloyed sintered steels, applied in automotive industry and in wear resistant structural components. On the other hand, by studying the influence of molybdenum additions to the iron borides new applications of B-Fe-Mo alloys were found, in particular for cutting tools. For systematic optimization of alloy compositions of such materials and components it is of great importance to know the phase relations in the corresponding B-Fe-Mo system. However, up to now, this information is relatively poor. It is presented in literature by a series of isothermal sections [1966Has1, 1966Gla, 2000Lei, 2004Sar, 2005Sar] and phase configurations and crystal structures of intermediate phases reported by a number of papers: [1953Ste, 1964Col1, 1964Col2, 1964Rie, 1965Kuz, 1965Rie, 1966Gla, 1966Has1, 1966Has2, 1969Bis, 1973Dey, 1973Rog, 1980Vin, 1983Ber, 1984Dun, 1986Muk, 1986Tak, 1987Ger, 1987Tak, 1988Efi, 1988Ide, 1989Ide, 1989Kol, 1995Dud, 1997Dud, 1997Wan, 2000Lei, 2000Mol, 2001Kar1, 2001Kar2, 2004Sar, 2005Sar]. Thermodynamic properties were obtained experimentally by [1983Ber, 1997Wan]. The experimental methods used and the temperature and composition ranges studied are shown in Table 1. The B-Fe-Mo system was reviewed in [1992Rag, 2003Rag], and reviews of crystal structures of the phases were presented in [1979Run, 2003Luk]. Further research on the nature of the phase equilibria in this system is needed, in particular with respect to the liquidus, solidus and solvus surfaces but also with respect to the alloy’s thermodynamics and targeted studies to be able to complete the reaction scheme for the whole range of compositions.
Binary Systems As edge binary phase diagrams those reported by [Mas2] are accepted, where the Fe-Mo phase diagram is the result of thermodynamic assessment by [1982Gui] using the CALPHAD method. It should be mentioned that the existence of Mo3B2 is rejected in the present evaluation following reviews of [1996Pov] and [1998Rog].
Solid Phases Crystallographic data of the unary, binary and ternary phases are listed in Table 2. The solubilities of the third component in each of the binary B-Fe, B-Mo and Fe-Mo phases Landolt‐Bo¨rnstein New Series IV/11E1
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were found to be small. Two ternary phases with the crystal structures different from any of the unary and binary phases were found, namely τ2-Mo2FeB2 and τ4-Mo1+xFe2–xB4. The τ1-Mo2Fe13B5 phase forms from the melt at 1100˚C and decomposes below 1000˚C [1966Has1]. It was reported by [2000Lei] that additions of molybdenum stabilize the metastable Fe3B phase (with the same Ti3P type structure) at the composition MoFe14B5, stable only in an extremely narrow temperature range of 1110–1080˚C. The τ3-Mo1–xFexB possesses the same crystal structure as the βMoB and can be considered as a βMoB phase stabilized by iron.
Invariant Equilibria Analyzing experimental data of [1966Gla] at 1000˚C and of [1966Has1] at 1050˚C Raghavan suggested in his review [1992Rag] that a transition reaction (Mo) + τ2 Ð Mo2B + μ occurs at the temperature between 1000 and 1050˚C, because the two-phase equilibrium Mo2B + μ found at 1000˚C by [1966Gla] changed to (Mo) + τ2 at 1050˚C, according to the data of [1966Has1]. It was found by [1966Has2] that addition of 0.43 at.% (0.73 mass%) Mo to B-Fe alloys increases the temperature of peritectoid reaction α + Fe2B Ð γ from 913˚C to 947˚C.
Isothermal Sections Isothermal sections at 1050˚C were constructed by [1966Has1] and [2000Lei] based on experimental investigations in the whole range of compositions. The isothermal section at 1050˚C is presented in Fig. 1 according to the experimental data of [2000Lei] with amendments to maintain consistency with the accepted binary phase diagrams. The homogeneity range of the Mo2B5 phase in the B-Mo system is enlarged according to the accepted binary diagram. Extensions of the Fe2B and FeB phases into the ternary system are slightly increased based on the EMPA data of [2000Lei]. The results of the investigation by [2000Lei] indicated that three ternary phases are stable at 1050˚C, namely, τ2, τ3 and τ4 phases. The isothermal section constructed by [1966Has1] at 1050˚C is similar to the one obtained by [2000Lei] excepting composition and the homogeneity ranges of ternary phases and stability of τ1 phase (Fe13Mo2B5) found by [1966Has1]. Isothermal section at 1000˚C was constructed by [1966Gla]. It is similar to the isothermal section at 1050˚C constructed by [2000Lei]. Several differences should be mentioned. The homogeneity ranges of the ternary phase are essentially different in [1966Gla] and [2000Lei]. Equilibria involving the MoB4 phase were not shown by [1966Gla]. According to the data of [1966Gla], the two-phase equilibrium Mo2B + μ exists at 1000˚C, while [1966Has1] and [2000Lei] indicated another two-phase equilibrium, (Mo) + τ2, at 1050˚C.
Thermodynamics Isothermal sections of the B-Fe-Mo system in the range of boron content up to 34 at.% at the temperatures of 1200˚C and 1280˚C were calculated by [2004Sar, 2005Sar] with the ThermoCalc program using published thermodynamic databases of the binary boundary systems: B-Fe [1994Hal], B-Mo [2001Mor] and Fe-Mo [1982Gui]. Four phases were considered in DOI: 10.1007/978-3-540-88053-0_29 ß Springer 2009
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calculations, namely: liquid, iron based solid solutions γ and α, hypothetical continuous series of solid solutions between Mo2B and Fe2B, as well as the μ, Mo6Fe7 phase. However, it should be mentioned that the existence of a continuous series of solid solutions between Mo2B and Fe2B was not confirmed by any experimental investigation.
Notes on Materials Properties and Applications Applications of the B-Fe-Mo alloys are related, in the first place, with positive effect of boron additions on microstructures and physical properties of alloyed sintered steels, in particular, based on the Astalloy Mo powder (Fe-1.5 mass% Mo i.e. Fe-0.88 at.% Mo). These steels have many applications in the automotive industry and as structural wear resistant materials. However, their applications are limited by the particular porosity values. To reduce the porosity and to increase the consolidation of sintered alloy steels, miscellaneous techniques within powder metallurgy technologies are utilized, as well as an activated sintering process, including some boron additions. Experimental studies revealed that boron activates the sintering process and increases in the degree of densification and in result improves the mechanical properties [2001Kar2]. On the other hand, additions of molybdenum to the iron borides allow to design properties that make them suitable cutting tool materials. The development of metallic glasses is another area of application for the B-Fe-Mo system. Literature data concerning investigations of the boron-iron-molybdenum materials properties are listed in Table 3. Factors affecting the strength properties of the Fe rich vacuuminduction melted B-Fe-Mo alloys were studied by [1964Col2] (abstract of this article was also presented in [1964Col1]). The approximate molybdenum contents were 2, 4, 9 and 19 mass% and approximate boron contents were 0.0005, 0.0014, 0.005 and 0.063 mass%. The phases extracted from the annealed specimens and identified by X-ray diffraction analysis included the R, Mo2Fe3 phase, the τ2-Mo2FeB2 phase as well as an oxide having a diffraction pattern practically identical to that of MnFe2O4. Three types of alloy strengthening were observed. Alloys with 2 mass% of Mo seem to be hardenable by a bainitic type transformation of austenite to ferrite; alloys with 4 mass% of Mo were hardened by the solid solution mechanism; those containing 9 or 19 mass% Mo were strengthened by precipitation of the R, Mo2Fe3 phase. No dispersion reinforcement by borides was recognized in any of the alloys. [1986Muk] have reported that cyclic tempering of the Mo10.5Fe70.5B19 alloy at the temperatures of 550 to 650˚C leads to decrease of the alloy hardness. The Mixtures containing up to 1.1 mass% B and up to 5 mass% Mo, compacted and sintered at 1200˚C, demonstrated negative effects on the densification and ductility by forming the τ2-Mo2FeB2 phase, although strength and hardness were improved [1987Ger]. The change of mechanical properties of rapidly quenched alloys in the range of B or Mo content 10 to 20 at.% during cooling and heat-treatment cycles (550 to 650˚C) was demonstrated by [1988Efi]. The ternary B-Fe-Mo alloys annealed in the temperature range of 1077–1327˚C and containing 6 mass% Mo have a composite microstructure consisting mainly of a complex boride τ2-Mo2FeB2 as a hard phase and a ferritic binder [1988Ide, 1989Ide]. The mechanism of liquid phase sintering of the Mo26.6Fe43.9B29.5 alloy was presented by [1988Ide, 1989Ide]. The microstructure formation of sintered materials based on the prealloyed Fe-1.5 mass% Mo powder with 0.2 to 0.6 mass% of boron addition was investigated by [1995Dud, 1997Dud, 2000Mol, 2001Kar1, 2001Kar2]. A eutectic liquid phase Fe + Fe2B activating the sintering of Fe-1.5 mass% Mo + B powder system was formed [1995Dud, 1997Dud]. The sintering behavior of Fe and Fe-Mo prealloyed powder compacts Landolt‐Bo¨rnstein New Series IV/11E1
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containing from 0.5 to 3.5 mass% Mo and fixed boron additions has been studied by [2004Sar, 2005Sar] with special emphasis on the microstructural development, the formation of the liquid phase and the liquid phase sintering mechanisms involved during the densification process. The interaction processes between boron and molybdenum are related to extinction/ or decrease of the liquid eutectic phase portion and to the strengthening of the alloyed Fe-Mo matrix by formed boride particles. The Fe-1.5Mo-0.4B (mass%, corresponds to the Fe-0.9Mo-2B in at.%) alloy sintered at 1200˚C reached the tensile strength of 487 to 550 MPa, elongation of 2.5 to 4.8%, hardness HV above 2 GPa, and the density of 7150 kg·m–3. The sintering behavior of Fe and Fe-Mo prealloyed powder compacts containing from 0.5 to 3.5 mass% Mo and fixed boron addition has been studied by [2004Sar, 2005Sar] with special emphasis on the microstructural development, the formation of the liquid phase and the liquid phase sintering mechanisms involved during the densification process. Magnetic properties of the B-Fe-Mo amorphous alloys were reported by [1980Vin, 1984Dun, 1992Kis, 1999Wan]. The electrical resistivity of (MoxFe1–x)80B20 (x = 0.05 to 0.18) amorphous alloys as a function of temperature between 4.2 and 300 K has been studied by [1984She].
Miscellaneous Thermal stability of the (Mo5Fe95)83B17 metallic glass was studied by [1983Ber]. A two-step process of crystallization of this metallic glass was carried out. Good agreement was likewise observed for three various applied types of calorimetric techniques (peak shift, temperature change as well as isothermal) with a view to determine the activation energy after first cycle of crystallization. The thermal behavior of the amorphous alloys with the boron content from 6.4 to 30.0 at.% prepared by electrolytic plating were investigated by [1997Wan] using differential scanning calorimetry. A two-step process of crystallization was observed for the amorphous Mo13.3Fe63.4B23.3, Mo11.3Fe58.7B30.0, Mo16.3Fe73.6B10.1 and Mo18.9Fe74.7B6.4 alloys. The crystallization temperature ranges from 745 to 771 K (472 to 498˚C) for B contents between 6.4 and 30.0 at.% at a heating rate of 10 K·min–1. A minimum value of the enthalpy change upon crystallization, Hx, was observed at 6.4 at.% B (27.5 kJ·mol–1) while the maximum value was found at 23.3 at.% B (39.0 kJ·mol–1). Positron studies of the crystallization kinetics of the metallic glass Mo2Fe78B20 were performed by [1980Car] using a zero-angle 2γ-spectrometer. The isotope 64Cu was used as a position source. Since the activation energies at the start and the end of crystallization were practically equal, the authors have concluded that the transition from the glassy to the crystalline state is a thermally activated first order process. [2001Kar1, 2001Kar2] studied the mechanism of sintering of the Astaloy Mo powder with elemental boron powder (0.2, 0.4 or 0.8 mass%) using the addition of 0.8 mass% lubricant in the form of zinc stearate. It was concluded that the sintering temperature plays a very important role in these processes. So, at 1120˚C the sintering mechanism of the Astalloy Mo sample with boron is solid state sintering, while at 1200˚C, as a result of the eutectic reaction between ferrite and complex borides, a permanent liquid phase occurs. In the last case the sintering mechanism is liquid phase sintering.
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. Table 1 Investigations of the B-Fe-Mo Phase Relations, Structures and Thermodynamics Reference [1953Ste]
Method/Experimental Technique X-ray diffraction
Temperature/Composition/Phase Range Studied Mo2FeB4
[1964Col2] X-ray diffraction, metallography (optical and 2–19 mass% Mo electron microscopies) [1964Rie]
X-ray diffraction
Mo2FeB2
[1965Kuz]
X-ray diffraction
800˚C; Mo2Fe21B6
[1965Rie]
X-ray diffraction
Mo2FeB2
[1966Gla]
X-ray diffraction, optical microscopy
1000˚C; whole range of compositions
[1966Has1] X-ray diffraction
950 - 1100˚C; whole range of compositions
[1966Has2] Chemical analysis, optical and electron microscopy, X-ray diffraction
≤ 1100˚C; 0.02 at.% B, 0.2 at.% B
[1969Bis]
Thermal analysis
≤ 1040˚C; Mo1.91Fe 98.024B0.066 (mass%)
[1973Dey]
Mo¨ssbauer spectroscopy
Mo2FeB2
[1973Rog]
X-ray diffraction
Mo0.4Fe2.6B
[1980Vin]
X-ray diffraction (Debye-Scherrer technique) Mo5Fe95–xBx (15 ≤ x ≤ 25)
[1983Ber]
Differential scanning calorimetry (DSC). Types of it are as following: peak shift, temperature change, isothermal
(Mo5Fe95)83B17
[1984Dun]
57
4.2 K; 20 at.% B, 0 to 14 at.% Mo
Fe Mo¨ssbauer spectroscopy
[1986Muk] X-ray diffraction, optical microscopy
≤ 1600˚C, Mo10.5Fe70.5B19
[1986Tak]
X-ray diffraction, chemical analysis (Leco automatic analyzer)
1127 - 1327˚C; Mo22.4Fe52.8B24.8, Mo25.3Fe49.5B25.2, Mo25.7Fe50.9B23.4
[1987Ger]
Optical microscopy, scanning electron microscopy (SEM), X-ray diffraction
1200˚C; the Fe rich corner
[1987Tak]
X-ray diffraction, SEM, Auge spectroscopy
1200 - 1300˚C; 5 mass% B, Mo/B atomic ratios 0.7 to 1.1
[1988Efi]
Metallography, X-ray diffraction
550 - 650˚C, 1500–1600˚C; 10 to 20 at.% B, 10 to 20 at.% Mo
[1988Ide]
Metallography
1077 - 1327˚C; 6 mass% Mo
[1989Ide]
DTA, X-ray diffraction, SEM
Mo26.6Fe43.9B29.5
[1989Kol]
X-ray diffraction, TEM
(MoxFe100–x)83B17 (x = 3, 15, 25)
[1995Dud] Optical microscopy, SEM, TEM
1100 - 1220˚C; Fe-1.5 mass% Mo + (0.2 - 0.6) mass% B
[1997Dud] Optical microscopy, SEM, TEM
1100 - 1220˚C; Fe-1.5 mass% Mo + (0.2 - 0.6) mass% B
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1997Wan] X-ray diffractometer, SEM, DSC
(6.4 - 30.0) at.% B
[2000Lei]
X-ray powder diffraction, electron microprobe analysis
1050˚C; whole range of compositions
[2000Mol]
Optical microscopy, SEM, energy dispersive 1200˚C; Fe-1.5 mass% Mo + (0.2 - 0.6) X-ray spectroscopy (EDXS) mass% B
[2001Kar1] X-ray diffraction, optical microscopy
1200˚C; Fe-1.5 mass% Mo + (0.2 - 0.6) mass% B
[2001Kar2] X-ray diffraction, optical microscopy
1200˚C; Fe-1.5 mass% Mo + (0.2 - 0.6) mass% B
[2004Sar]
Optical microscopy, SEM
800–1280˚C; Mo2.0Fe96.4B1.6, Mo0.9Fe97.6B1.5 Mo0.3Fe98.2B1.5
[2005Sar]
Optical microscopy, SEM
1000–1280˚C; Mo2.0Fe96.4B1.6, Mo0.9Fe97.6B1.5 Mo0.3Fe98.2B1.5
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βB) < 2092
hR333 R3m βB
a = 1093.30 [1993Wer] c = 2382.52 dissolves 2.5 at.% Fe at 1500˚C [Mas2] dissolves < 1 at.% Mo at 1920 ± 25˚C [Mas2]
(δFe) (h2) 1538 - 1394, 1.013 bar
cI2 Im 3m W
a = 293.15
[Mas2] in the Fe-Mo binary system joins with the (αFe), dissolves 24.4 at.% Mo at 1449˚C [Mas2]
γ, (γFe) (h1) 1394 - 912
cF4 Fm 3m Cu
a = 364.67
at 915˚C [V-C2, Mas2] dissolves 1.7 at.% Mo at 1140˚C [Mas2]
α, (αFe) (r) < 912
cI2 Im3m W
a = 286.65
at 25˚C [Mas2] in the Fe-Mo binary system joins with the (δFe) [Mas2] in the Mo10Fe65B25 alloy sintered at 1050˚C [2000Lei]
a = 286.69
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25˚C, 1.3·105 bar [Mas2] high-pressure modification
(Mo) < 2623
cI2 Im3m W
a = 314.70
at 25˚C [Mas2] dissolves < 1 at.% B at 2175 ± 6˚C and 31.3 at.% Fe at 1611˚C [Mas2] in the Mo55Fe25B20 and Mo60Fe10B30 alloys sintered at 1050˚C [2000Lei]
a = 314.51 Fe2B < 1399
FeB < 1603
tI12 I4/mcm CuAl2
oP8 Pnma FeB
a = 511.0 c = 424.9
33.3 at.% B [V-C2]
a = 511.35 c = 425.56
in the Mo30Fe30B40 arc-melted alloy annealed at 1050˚C [2000Lei]
a = 550.6 b = 295.2 c = 406.1 a = 559.09 b = 295.80 c = 408.13
50 at.% B [V-C2]
Fe3B
tI* Fe3P t** Fe3B0.63P0.37 oP16 Pnma Fe3C tP* P42/n Ti3P Mo2B < 2280
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tI12 I4/mcm CuAl2
a = 865.5 c = 429.7 a = 864.8 c = 431.4 a = 443.9 b = 542.8 c = 669.9 a = 863.2 c = 431.1 a = 863.16 c = 431.29 a = 554.7 c = 473.9 a = 554.87 c = 473.47
in the Mo18Fe27B55 arc-melted alloy annealed at 1050˚C [2000Lei] metastable; produced by quenching of liquid Fe76B24, exists in the approximate temperature interval 1150 to 1250˚C [1981Kha, 1982Kha] disordered high-temperature modification [1982Kha] ordered low-temperature modification [1982Kha] orthorhombic modification coexists in a small amount with high-temperature modification [1982Kha] in the rapid-quenched alloys [1992Rog] in the arc-melted Mo5Fe70B25 alloy [2000Lei] 33 to 34 at.% B [Mas2] [V-C2] in the Mo60Fe10B30 alloy sintered at 1050˚C [2000Lei]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
βMoB 2600 - 1800
oC8 Cmcm CrB
αMoB < 2180
tI16 I41/amd αMoB
MoB2 2375 - 1517
hP3 P6/mmm AlB2
Mo2B5 < 2140
hR21 R3m Mo2B5
MoB4 < 1807
hP20 P63/mmc WB4
σ, MoFe 1611 - 1235
tP30 P42/mnm σCrFe
μ, Mo6Fe7 < 1370
hR39 R3m W6Fe7
R, Mo2Fe3 1488–1200
hR159 R3 Co5Cr2Mo3
DOI: 10.1007/978-3-540-88053-0_29 ß Springer 2009
Lattice Parameters [pm] a = 315.1 b = 847.0 c = 308.2 a = 311.0 c = 1695 a = 311.62 c = 1693.5 a = 304 c = 306 a = 301.1 c = 2093.0 a = 300.83 c = 2091.47 a = 520.33 c = 634.98 a = 520.26 c = 634.99 a = 520.42 c = 635.05 a = 918.8 c = 481.2
Comments/References 48 to 51 at.% B [Mas2] [1992Rag]
48 to 50 at.% B [Mas2] [V-C2] dissolves 0.5 at.% Fe at 1050˚C [2000Lei] in the Mo45Fe5B50 arc-melted alloy annealed at 1050˚C [2000Lei] 62 to 66 at.% B [Mas2] [1992Rag] 67 to 69 at.% B [Mas2] [V-C2] dissolves 2.3 at.% Fe at 1050˚C [2000Lei] in the Mo15Fe15B70 arc-melted alloy annealed at 1050˚C [2000Lei] 79 at.% B [Mas2] Mo rich [1973Lun] B rich [1973Lun] in the Mo2Fe73B75 alloy sintered at 1050˚C [2000Lei] 43.3 to 57.1 at.% Fe [Mas2] [H]
56 to 61 at.% Fe [Mas2] a = 475.46 [V-C2] c = 2571.6 a = 475.91 in the Mo30Fe50B20 and Mo55Fe25B20 alloys sintered c = 2571.52 at 1050˚C [2000Lei] a = 1091.0 c = 1935.4 a = 1095.6 c = 1935.3
61.5 to 66.1 at.% Fe [Mas2] Mo1.9Fe3.1, at 1250 to 1490˚C [V-C2] annealed at 1250˚C [V-C2]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
λ, MoFe2 < 927
hP12 P63/mmc MgZn2
* τ1, Mo2Fe13B5 1100 - < 1000
tP32 P42/n Ti3P
a = 863.4 c = 428.1
annealed at 1050˚C [1966Has1]
* τ2, Mo2FeB2
tP10 P4/mbm U3Si2
a = 578.2 c = 314.8
[1964Rie]
a = 580.7 c = 314.2 a = 577.26 c = 314.61
annealed at 1000˚C [1966Gla]
* τ3, Mo1–xFexB
oC8 Cmcm CrB
* τ4, Mo1+xFe2–xB4 oI14 Immm Ta3B4
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a = 475.5 c = 776.7
66.7 at.% Fe [Mas2] [V-C2]
in the Mo30Fe30B40 arc-melted alloy annealed at 1050˚C [2000Lei] 6.5 to 9 at.% Mo at 50 at.% B, T = 1050˚C [2000Lei]
a = 315.7 b = 839.7 c = 306.3 a = 315.26 b = 841.86 c = 306.62
a = 312.8 b = 1270 c = 298.4 a = 311 b = 1427 c = 319 a = 299.6 b = 1279 c = 309.9 a = 301 b = 1300 c = 310.5 a = 304 b = 1340 c = 312.8
annealed at 1050˚C [1966Has1]
in the Mo45Fe5B50 arc-melted alloy annealed at 1050˚C [2000Lei] 13 to 25 at.% Mo at 57 at.% B, T = 1050˚C [2000Lei] x = 0, annealed at 1000˚C [1966Gla]
x = 1, annealed at 1050˚C [1966Has1]
x = 0.06, annealed at 1050˚C [2000Lei]
x = 0.43, annealed at 1050˚C [2000Lei]
x = 1, annealed at 1050˚C [2000Lei]
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. Table 3 Investigations of the B-Fe-Mo Materials Properties Reference
Method/Experimental Technique
Type of Property
[1964Col2] Vickers hardness measurements, tensile tests, creep-rupture tests
Hardness, stress, tensile strength
[1966Has2] Vickers hardness measurements, tensile and impacts tests
Hardness, tensile strength, impact properties
[1973Dey]
Mo¨ssbauer spectroscopy
[1980Vin]
Inductance bridge method, four-point method Curie temperature, electrical resistivity, coercive field
[1983Ber]
Thermomechanical system (TMS), thermogravimetric system (TGS) studies
Magnetic hyperfine structure, atomic magnetic ordering
Magnetization, electrical resistance, hardness
[1984Dun] Magnetization measurements, 57Fe Mo¨ssbauer Magnetization, magnetic moment, spectroscopy (constant acceleration Fe hyperfine field distribution spectrometer) [1984She]
Electrical resistivity-temperature dependence measurements
Electrical resistivity
[1986Muk] Vickers microhardness tests (PMT-3 apparatus); Microhardness, tensile strength, yield stress, elongation, tensile strength tests yield point (“Instron” apparatus) [1986Tak]
Density, strength, Rockwell hardness tests
Transverse rupture strength (TRS), density, hardness
[1987Ger]
Dilatometry, density, hardness (Brinell), stress and ductility tests
Tensile elongation, density, hardness, porosity, ductility
[1987Tak]
Density, strength, Rockwell hardness tests
TRS, hardness, density
[1988Efi]
Stress, strength, Vickers microhardness tests (PMT-3 apparatus)
Ultimate tensile strength, yield stress, microhardness
[1988Ide]
Density, strength, Rockwell hardness tests
TRS, hardness, density
[1989Ide]
Thermal dilatometry, Rockwell hardness tests
TRS, elongation
[1992Kis]
Magnetic measurements
Magnetic moment, Curie temperature
[1995Dud] Static tensile, three-point bend tests (“Instron 1302” equipment), Vickers hardness and microhardness tests, dilatometry
Density, microhardness, hardness, tensile and bending strength, elongation
[1997Dud] Static tensile, three-point bend tests (“Instron 1302” equipment), Vickers hardness and microhardness tests, dilatometry
Density, microhardness, hardness, tensile and bending strength, elongation
[1998Wan] Microhardness tests
Microhardness
[1999Wan] Magnetic measurements, microhardness tests
Magnetic losses, microhardness
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. Table 3 (continued) Reference [2000Mol]
Method/Experimental Technique Dilatometry, water displacement technique, Vickers microhardness tests
Type of Property Density, porosity, elongation, microhardness
[2001Kar1] Water displacement, Brinell hardness tests; Density, hardness, strength, tensile strength and elongation (“Instron 1302” elongation equipment) tests [2001Kar2] Water displacement, Brinell hardness tests; tensile strength and elongation (“Instron” equipment) tests
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B–Fe–Mo
. Fig. 1 B-Fe-Mo. Isothermal section at 1050˚C
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References [1953Ste] [1964Col1] [1964Col2]
[1964Rie] [1965Kuz]
[1965Rie] [1966Gla]
[1966Has1]
[1966Has2]
[1969Bis] [1973Dey] [1973Lun] [1973Rog] [1979Run]
[1980Car] [1980Vin] [1981Kha]
[1982Gui] [1982Kha] [1983Ber]
[1984Dun] [1984She]
Steinitz, R., Binder, I., “New Ternary Boride Compounds”, Powder Met. Bull., 6(4), 123–125 (1953) (Crys. Structure, Phase Relations, Experimental, 6) Coldren, A.P., Semchyshen, M., “Factors Affecting the High-Temperature Strength of Iron-Rich FeMo-B Alloys”, J. Met., 16(1), 101 (1964) (Crys. Structure, Morphology, Abstract, Mechan. Prop., 1) Coldren, A.P., Semchyshen, M., Scholz, W.G., “Factors Affecting the Strength of Iron-Rich IronMolybdenum-Boron Alloys”, Trans. Metall. Soc. AIME, 230, 1236–1250 (1964) (Crys. Structure, Morphology, Experimenta1, Mechan. Prop., 10) Rieger, W., Nowotny, H., Benesovsky, F., “The Crystal Structure of Mo2FeB2” (in German), Monatsh. Chem., 95, 1502–1503 (1964) (Crys. Structure, Experimental, 3) Kuz’ma, Yu.B., Voroshilov, Yu.V., Cherkashin, E.E., “New Ternary Compounds with W2Cr21C6 Type Structure”, Inorg. Mater. (Engl. Trans.), 1(7), 1017–1019 (1965), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 1(7), 1109–1111 (1965) (Crys. Structure, Experimental, 3) Rieger, W., Nowotny, H., Benesovksy, F., “About Some Complex Borides of the Transition Metals” (in German), Monatsh. Chem., 96(3), 844–851 (1965) (Crys. Structure, Phase Diagram, Experimental, 10) Gladyshevskii, E.I., Fedorov, T.F., Kuz’ma, Yu.B., Skolozdra, R.V., “Isothermal Section of the System Molybdenum-Iron-Boron” (in Russian), Poroshk. Metall., (4), 55–60 (1966) (Crys. Structure, Morphology, Phase Diagram, Experimental, *, 6) Haschke, H., Nowotny, H., Benesovsky, F., “Investigations in the Ternary Systems {Mo, W}-{Fe, Co, Ni}-B” (in German), Monatsh. Chem., 97, 1459–1468 (1966) (Crys. Structure, Phase Diagram, Experimental, *, 6) Hasegawa, M., Okamoto, M., “A Study on the Ternary Alloys of Iron and Boron” (in Japanese), Nippon Kinzoku Gakkai-Si, 30(6), 533–540 (1966) (Morphology, Phase Relations, Experimental, Mechan. Prop., 12) Biss, V., Coldren, A.P., “Continuous Cooling Transformation of an Fe-1.91 % Mo-0.066 % B-0.002 % C Alloys”, Trans. Metall. Soc. AIME, 245, 884–886 (1969) (Morphology, Phase Diagram, Experimental, 3) De Young, D.B., Barnes, R.G., “57Fe Moessbauer Study of Some M2(M’)B2 Borides”, J. Phys. Chem. Solids, 34, 139 (1973) (Crys. Structure, Experimental, Electronic Structure, Magn. Prop., 5) Lundstroem, T., Rosenberg, I., “The Crystal Structure of the Molybdenum Boride Mo(1–x)B3”, J. Solid State Chem., (6), 299–305 (1973) (Crys. Structure, Experimental, 5) Rogl, P., Nowotny, H., “New Complex Borides” (in German), Monatsch. Chem., 104(4), 943–952 (1973) (Crys. Structure, Experimental, 25) Rundqvist, S., Andersson, Y., Pramatus, S., “Coordination and Bonding in Represantatives of the Fe3P-, Ti3P-, α-V3S-, and β-V3S Type Structures”, J. Solid State Chem., 28, 41–49 (1979) (Crys. Structure, Review, 42) Cartier, E., Heinrich, F., “Positron Studies of the Crystallization Kinetics in the Metallic Glass Fe78Mo2B20”, Helv. Phys. Acta, 53, 266–269 (1980) (Morphology, Experimental, Kinetics, 8) Vind Nielsen, H.J., “Magnetic Properties of Fe-Cr-B and Fe-Mo-B Metallic Glasses”, J. Magn. Magn. Mater., 19(1-3), 138–140 (1980) (Crys. Structure, Experimental, Electr. Prop., Magn. Prop., 24) Khan, Y., Kneller, E., Sostarich, M., “Stability and Crystallization of Amorphous Iron-Boron Alloys Obtained by Quenching from the Melt”, Z. Metallkd., 72(8), 553–557 (1982) (Crys. Structure, Phase Diagram, Experimental, 21) Guillermet, A.F., “The Fe-Mo (Iron-Molybdenum) System”, Bull. Alloy Phase Diagrams, 3(3), 359–367 (1982) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Review, *, 40) Khan, Y., Kneller, E., Sostarich, M., “The Phase Fe3B”, Z. Metallkd., 73(10), 624–626 (1982) (Crys. Structure, Phase Diagram, Experimental, 13) Bergmann, H.W., Brokmeier, U., Fritsch, H.U., “Experimental Investigation on the Thermal Stability of Metallic Glasses”, Ber. Bunsen-Ges. Phys. Chem., 87, 757–761 (1983) (Crys. Structure, Morphology, Thermodyn., Experimental, Mechan. Prop., 15) Dunlap, R.A., Stroink, G., “Magnetic Properties of Amorphous Fe-Mo-B Alloys”, Canad. J. Phys., 62, 714–719 (1984) (Crys. Structure, Experimental, Morphology, Magn. Prop., 21) Shen, B., Zhan, W., Zhang, Zh., Wu, Z., “Temperature (4.2–300 K) Dependence of Electrical Resistivity of FeMB (M = Mn, Mo) Amorphous Alloys” (in Chinese), Acta Metall. Sin., 20(3), B164-B170 (1984) (Morphology, Experimental, Electr. Prop., 18)
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[1986Tak]
[1987Ger] [1987Tak]
[1988Efi]
[1988Ide]
[1989Ide] [1989Kol] [1992Kis]
[1992Rag]
[1992Rog] [1993Wer]
[1994Hal]
[1995Dud]
[1996Pov]
[1997Dud]
[1997Wan]
[1998Rog] [1998Wan]
B–Fe–Mo Mukhin, G.G., Efimov, Yu.V., Dmitriev, V.N., Buravleva, I.S., “Influence of Cyclic Tempering on Mechanical Properties of Rapid-Quenched Alloys” (in Russian), Metalloved. i Term. Obrabotka Metallov, (11), 51–53 (1986) (Crys. Structure, Morphology, Experimental, Mechan. Prop., 8) Takagi, K., Watanabe, T., Ando, T., Kondo, Y., “Effect of Molybdenum and Carbon on the Properties of Iron Molybdenum Boride Hard Alloys”, Inter. J. Powder Metallurgy, 22(2), 91–96 (1986) (Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 7) German, R.M., Hwang, K.-S., Madan, D.S., “Analysis of Fe-Mo-B Sintered Alloys”, Powder Metallurgy International, 19(2), 15–18 (1987) (Crys. Structure, Morphology, Experimental, Phys. Prop., 18) Takagi, K., Komai, M., Ide, T., Watanabe, T., Kondo, Y., “Effects of Mo and Cr Contents on the Properties and Phase Formation of Iron Molybdenum Boride Base Hard Alloys”, Powder Metallurgy International, 19(5), 30–33 (1987) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 13) Efimov, Yu.V., Mukhin, G.G., Fridman, Z.G., Bouravleva, I.S., Myasnikova, E.A., “The Change of the Amorphous State of Fe-Mo-B Alloys on Heating”, J. Non-Cryst. Solids, 103, 45–48 (1988) (Phase Relations, Experimental, Mechan. Prop., 6) Ide, T., Nakano, K., Ando, T., “Effects of Mo Content and Sintering Temperature on the Strength, Hardness and Density of Fe-6 mass% B-x mass% Mo Alloys”, Powder Metallurgy International, 20(3), 21–24 (1988) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 17) Ide, T., Ando, T., “Reaction Sintering of an Fe-6 wt% B-48 wt% Mo Alloy in the Presence of Liquid Phases”, Metall. Trans. A, 20A, 17–24 (1989) (Morphology, Phase Relations, Phys. Prop., 16) Kolb-Telieps, A., Luft, U., “Relations Between Structure and First-Step Crystallization of Iron-MetalBoron Glasses”, J. Non-Cryst. Solids, 109, 59–63 (1989) (Crys. Structure, Morphology, Experimental, 7) Kisdi-Koszo´, E., Lovas, A., Kova´c, J., Varga, L.K., Zsoldos, E., “Inhomogeneous Atomic Distribution and Its Effect on Magnetic Properties of Diluted Fe-TM-B Metallic Glasses”, J. Magn. Magn. Mater., 112, 39–40 (1992) (Morphology, Experimental, Magn. Prop., 4) Raghavan, V., “The B-Fe-Mo (Boron-Iron-Molybdenum) System”, in “Phase Diagrams of Ternary Iron Alloys”, Indian Institute of Metals, Calcutta, 6A, 365–369 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 8) Rogl, P., in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, OH, 33–36 (1992) (Crys. Structure, Phase Diagram, Thermodyn., Review) Werheit, H., Kuhlmann, U., Laux, M., Lundstrm, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Sol., B179, 489–511(1993) (Crys. Structure, Experimental, 51) Hallemans, B., Wollants, P., Roos, J.R., “Thermodynamic Reassessment and Calculation of the Fe-B Phase Diagram”, Z. Metallkd., 85(10), 676–682 (1994) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, #, 36) Dudrova´, E., Salak, A., Selecka´, M., Bure, R., “Properties and Microstructure of Fe-1.5 Mo Powder Steel Sintered with a Boron-Based Liquid Phase”, Metal. Mater., 33(2), 60–65 (1995), transl. from Kovove Mater., 33(2), 82–93 (1995) (Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 11) Povarova, K.B., “B-Mo. Boron-Molybdenum”, in “Phase Diagrams of Binary Metallic Systems” (in Russian), Lyakishev, N.P. (Ed.), Vol. 1, Mashinostroenie, Moscow, 461–464 (1996) (Crys. Structure, Phase Diagram, Review, *, 14) Dudrova´, E., Selecka´, M., Bure, R., Kaba´tova, M., “Effect of Boron Addition on Microstructure and Properties of Sintered Fe-1.5Mo Powder Materials”, ISIJ Int., 37(1), 59–64 (1997) (Morphology, Phase Relations, Experimental, Phys. Prop., 11) Wang, L.L., Zhang, B.W., Yi, G., Ouyang, Y.F., Hu, W.Y., “Structure and Crystallization of Amorphous Fe-Mo-B Alloys Obtained by Electroless Plating”, J. Alloys Compd., 255 (1–2), 231–235 (1997) (Crys. Structure, Morphology, Phase Relations, Thermodyn., Experimental, 5) Rogl, P., “Metal-Boron-Carbon Ternary Systems”, Effenberg, G. (Ed.), MSIT-ASM International, OH, USA, 1–480 (1998) (Crys. Structure, Phase Diagram, Thermodyn., Review, *) Wang, L.L., Zhao, L.H., Zhang, B.W., Ouyang, Y.F., Liao, S.Z., Hu, W.Y., “Composition Dependence of Some Physical Properties of Fe-TM-B (TM = Mo, W, Mo-W) Alloys Obtained by Electroless Plating”, Plating & Surface Finishing, 85(12), 96–98 (1998) (Morphology, Experimental, Mechan. Prop., 9)
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[2001Kar2]
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[2003Rag] [2004Sar]
[2005Sar]
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Wang, L.L., Zhao, L.H., Zhang, B.W., Liao, S.Z., Ouyang, Y.F., Hu, W.Y., Shu, X.L., Yuan, X.J., “Annealing Temperature Dependence of A-C Magnetic Losses and Microhardness in Fe-TM-B (TM = Fe, Mo, W, Mo-W) Alloys Obtained by Electroless Plating”, Plating & Surface Finishing, 86 (12), 84–86 (1999) (Morphology, Experimental, Magn. Prop., Mechan. Prop., 5) Leithe-Jasper, A., Klesnar, H., Rogl, P., Komai, M., Takagi, K.-I., “Reinvestigation of Isothermal Section in M (M = Mo, W)-Fe-B Ternary Systems at 1323K” (in Japanese), J. Jpn. Inst. Met., 64 (2), 154–162 (2000) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, #, 30) Molinari, A., Pieczonka, T., Kazior, J., Gialanella, S., Straffelini, G., “Dilatometry Study of the Sintering Behavior of Boron-Alloyed Fe-1.5 % Mo Powder”, Metall. Trans. A, 31A(6), 1497–1506 (2000) (Crys. Structure, Morphology, Phase Relations, Experimental, Phys. Prop., 19) Karwan-Baczewska, J., “The Properties and Structure of Boron Modified P/M Iron-Molybdenum Alloys”, Arch. Metall., 46(4), 439–445 (2001) (Crys. Structure, Morphology, Experimental, Kinetics, Mechan. Prop., Phys. Prop., 14) Karwan-Baczewska, J., Rosso, M., “Effect of Boron on Microstructure and Mechanical Properties of PM Sintered and Nitrided Steels”, Powder Metall., 44(3), 221–227 (2001) (Crys. Structure, Morphology, Experimental, Kinetics, Mechan. Prop., Phys. Prop., 14) Morishita, M., Koyama, K., Yagi, S., Zhang, G., “Calculated Phase Diagram of the Ni-Mo-B Ternary System”, J. Alloys Compd., 314, 212–218 (2001) (Phase Diagram, Thermodyn., Calculation, Review, *) Lukachuk, M., Poettgen, R., “Intermetallic Compounds with Ordered U3Si2 or Zr3Al2 Type Structure Crystal Chemistry, Chemical Bonding and Physical Properties”, Z. Kristallogr., 218, 767–787 (2003) (Crys. Structure, Thermodyn., Review, Electronic Structure, Electr. Prop., Magn. Prop., 197) Raghavan, V., “B-Fe-Mo (Boron-Iron-Molybdenum)”, J. Phase Equilib., 24(5), 449–450 (2003) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 8) Sarasola, M., Go´mez-Acebo, T., Castro, F., “Liquid Generation During Sintering of Fe-3.5 % Mo Powder Compacts with Elemental Boron Additions”, Acta Mater., 52, 4615–4622 (2004) (Morphology, Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, 33) Sarasola, M., Go´mez-Acebo, T., Castro, F., “Microstructural Development During Liquid Phase Sintering of Fe and Fe-Mo Alloys Containing Elemental Boron Additions”, Powder Met., 48(1), 59–67 (2005) (Morphology, Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, 22) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Hafnium – Molybdenum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction The high melting temperature, high hardness and good oxidation resistance of hafnium and molybdenum borides were considered as attractive features for hardeners in low-alloyed molybdenum based high temperature structural materials. Therefore early interest existed in the constitution of the ternary B-Hf-Mo system. A cursory study of the HfB2-Mo interaction by [1962Lei] indicated the formation of a ternary compound with unknown crystal structure. Indeed the authors of [1966Har] from their investigation of phase relations within the isopleth HfB2-MoB2 and the isothermal section B-Hf-Mo (< 70 at.% B) at 1400˚C claimed a ternary phase (structure undetermined) within the phase region (Hf)+HfMo2+HfB2. The ternary phase observed by [1966Har] was confirmed by [1971Rog], who defined its correct composition at Hf9Mo4B [1973Rog] and evaluated the crystal structure as a unique structure type (κboride; Hf9Mo4B type). The possibility to form high-temperature composites, HfB2-Mo [1979Ord] and HfB2-“MoB2” [1966Har, 1967Hol], was explored. In view of heat-resistant Mo based alloys, phase relations in the Mo corner of the B-Hf-Mo system were evaluated by Zakharov [1980Zak]. Compilations of the relevant data on the topology of the B-Hf-Mo system have been published by [1969Rud, 1974Upa, 1994McH]. Experimental findings on the constitution of the B-Hf-Mo system are summarized in Table 1.
Binary Systems The binary boundary systems B-Hf and Hf-Mo were accepted from [Mas2]. B-Hf is basically consistent with the version presented by [1966Rud] except for small changes, in order to correspond to the accepted βHf/αHf transition temperature ([1966Rud] at 1800˚C, [Mas2] at 1743˚C) and melting temperature of zirconium-free hafnium ([1966Rud] at 2218±6˚C, [Mas2] at 2231˚C). A thermodynamic calculation of the B-Hf diagram is from [1988Rog] with a refinement of this modeling by [1997Bit] (see corresponding figure in B-C-Hf system). For the binary system B-Mo, we prefer the version assessed by [1992Rog] taking also care of mass spectrometry data by [1977Sto]. Literature data concerning the formation and crystal structure of solid phases pertinent to the B-Hf-Mo system are listed in Table 2.
Solid Phases Based on the various investigations of the B-Hf-Mo ternary system, phase relations are characterized by the existence of one ternary compound, Hf9Mo4B [1966Har, 1971Rog, Landolt‐Bo¨rnstein New Series IV/11E1
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1973Rog], which at 1400˚C seems to exist at a stoichiometric composition without a significant homogeneous region. From unit cell dimensions mutual solid solubilities of the binary borides were found to be very small: except for HfB2 negligible mutual solubilities among borides were reported by [1966Har], only αMoB was shown to dissolve at 1400˚C up to 2 at.% Hf. Technological interest in hard and high melting diborides spurred interest in the section Hf1–xMoxB2. Solubility of Hf in MoB2 (AlB2 type) was derived from lattice parameters to raise from 2.5 mol% HfB2 (= 0.8 at.% Hf) at 2000˚C to about 8 mol% HfB2 (= 2.7 at.% Hf) at subsolidus temperatures [1966Har]. Metallographic and lattice parameter studies defined a maximal solubility of 70 mol% MoB2 in HfB2 at 2378˚C (= 23 at.% Mo) reducing to 50 mol% MoB2 at 2000˚C (= 17 at.% Mo) and 45 mol% MoB2 at 1400˚C (= 15 at.% Mo). The lattice parameter dependencies within the section HfB2-MoB2 are shown in Fig. 1. These data are similar to those reported by [1967Hol] who claimed a maximal solubility of 20 mol% MoB2 in HfB2 and a solubility of 5 mol% HfB2 in MoB2, although no defined temperature was given. Crystallographic data for all solid phases pertinent to the B-Hf-Mo system are listed in Table 2.
Quasibinary Systems Two sections, HfB2-MoB2 [1966Har] and HfB2-Mo [1979Ord, 1980Zak] were reported, however both sections are not truly quasibinary, particularly as MoB2 melts incongruently and the section HfB2-Mo, for which a quasibinary eutectic was claimed at TE = 2070±25˚C at 28 mass% HfB2 [1979Ord], needs to cross the ternary field HfB2+Mo2B+(Mo) [1980Zak]. For details see below “Temperature-Composition Sections”.
Invariant Equilibria Although two sets of measurements on the solidus [1966Har] and liquidus [1971Rog] exist for most of the ternary diagram (<70 at.% B) the temperatures shown do not match and are inconsistent with the isopleths HfB2-Mo by [1979Ord, 1980Zak]. Therefore, the reaction isotherms listed by [1971Rog] are to be taken as cursory information, which needs to be refined with detailed Pirani-Alterthum type melting point data. A congruent melting point of the ternary compound was recorded at 1880˚C [1971Rog].
Liquidus, Solidus and Solvus Surfaces Measurements on the solidus [1966Har] and liquidus [1971Rog] temperatures for the ternary diagram <70 at.% B are inconsistent with the isopleths HfB2-Mo by [1979Ord, 1980Zak]. Therefore, the liquidus surface presented by [1971Rog] as cursory information from Segercoin data needs to be reinvestigated with the more precise Pirani-Alterthum type melting point technique and/or high temperature DTA.
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Isothermal Sections The experimentally established isothermal section at 1400˚C [1966Har, 1971Rog] is characterized by high thermodynamic stability of HfB2, which enters two-phase equilibria practically with all binary and ternary compounds. In the Mo rich corner the location of the vertex of the three-phase field (Mo)+Mo2B+HfB2 on the ss-(Mo) was given at 8 at.% Hf at 1400˚C [1966Har] thereby excluding a truly quasibinary section HfB2 - (Mo). At variance to data by [1966Har], the phase boundary between the phase fields (Mo)+Mo2B and (Mo)+Mo2B+HfB2 at 1400˚C was shown to be at 1.25 mass% HfB2 [1980Zak], which would restrict the vertex of the three-phase field (Mo)+Mo2B+HfB2 at the ss-(Mo) at a Hf content of less than about 1 at. % Hf. The isothermal section at 1400˚C, as shown in Fig. 2, summarizes the findings of [1966Har, 1971Rog, 1973Rog, 1979Ord, 1980Zak] and complies with the accepted boundary systems. The three-phase field (Mo)+Mo2B+HfB2 is dashed and will need further studies.
Temperature – Composition Sections Two sections, HfB2-MoB2 [1966Har] and HfB2-Mo [1979Ord] were reported as quasibinary. Incongruent melting of AlB2 type MoB2, however, renders the HfB2-MoB2 section to be nonquasibinary. It is interesting to note that despite the difference in atom radii of Hf and Mo complies with the Hume-Rothery limit, no complete solid solution exists for the diborides Hf1–xMoxB2. The low melting temperatures observed near MoB2 were interpreted in terms of a melting point minimum [1966Har]. Figure 3 shows the constitution of the section HfB2MoB2 particularly the mutual solubilities of the diborides and the reaction isotherm at 2378 ± 35˚C which reveals features of a “quasibinary” peritectic but probably stems from the transition reaction L+MoB Ð HfB2+MoB2. In this context the observation of a “minimum melting temperature” near MoB2 may in fact be a fingerprint of the transition reaction L+MoB2 Ð⊊HfB2+Mo2B5–x, which however, can only take place below the binary peritectic l +MoB2 Ð⊊Mo2B5 at 2140˚C. In addition it should be noted that binary MoB2 with AlB2 type structure is only stable above 1520˚C and was supposed to decompose in the ternary according to a eutectoid reaction MoB2 Ð αMoB2 + HfB2+Mo2B5–x at 1500˚C [1971Rog]. Serious controversy concerns the section HfB2-Mo, which was derived to be a simple quasibinary eutectic with TE = 2070±25˚C at 72 mass% Mo [1979Ord]. At variance and independently the group of [1980Zak] has found that the section has to cross the threephase field HfB2+Mo2B+(Mo) and intersects the monovariant line L+(Mo)+Mo2B (TE = at 2030⊊±⊊20˚C for L Ð HfB2+Mo2B+(Mo) [1980Zak]). It is very likely that also the monovariant line L+Mo2B+βMoB has to be intersected at higher temperatures [1971Rog]. Furthermore, with respect to the phase boundary between the phase fields (Mo)+Mo2B and (Mo)+Mo2B+HfB2 at 1400˚C shown to be at 1.25 mass% HfB2 by [1980Zak], the vertex of the three-phase field (Mo)+Mo2B+HfB2 at the ss-(Mo) should be at a Hf content of less than about 1 at.% Hf in contradiction to the findings of [1966Har], who located the vertex at Mo92Hf8 (at.%). More experiments are certainly needed to resolve the true nature of the liquidus, solidus and solvus surfaces.
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Notes on Materials Properties and Applications Microhardness data on Mo rich samples in the isopleth HfB2 - (Mo) [1980Zak] increased with the extent of the HfB2 solubility in (Mo). Microhardness of the HfB2 phase in hypoeutectic alloys HfB2+MoB2 was reported as 25.6±6 GPa, whereas the microhardness of the eutectic varied from 7.8 to 8.9 GPa [1979Ord]. The (Mo)-phase was said to be 4.0 GPa [1979Ord] similar to the data given by [1980Zak] for (Mo) alloys saturated with HfB2 + Mo2B reaching 3.5 GPa. Short term and long tern strength and creep of Mo base alloys with 0.16 to 0.45 mass% Hf, 0.001 to 0.005 mass% B and 0.002 to 0.005 mass% C were investigated after arc melting and rolling into 1 mm thick sheets which were arc-welded under argon. The temperature dependence of strain and plasticity was recorded at various temperatures showing that specimens with 0.16 mass% Hf revealed a minimum fatigue life and complete recristallization of microstructure [1975Bul]. . Table 1 Investigations of the B-Hf-Mo Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied
Reference
Method/Experimental Technique
[1962Lei]
Reaction sintering of mixture of Mo with HfB2. XPD.
From XPD a ternary compound of unknown structure was observed. The powder pattern was said to be similar to those for the mixtures W-HfB2 and MoZrB2.
[1966Har]
53 samples prepared by hot pressing powder mixtures in graphite dies (14 mm diameter and 7 mm height; sample surface removed prior to annealing at 1400˚C for 100 h in vacuum of <6.7·10–3 Pa). Starting materials were Hf powder (<74 μm; 30ppm C, <100-Nb, 70-Fe, 57-N2, 550-O2, 50-Si, 200-Ta, 55-Ti, 2.77 mass% Zr); HfH2 powder (74 and 250 μm; 80 ppm Al, 30ppm C, 100-Nb, 190-Fe, 250-Mg, 20-N, 330-O, <40-Si, <200-Ta, 75-Ti, 1.35 mass% Zr; 0.92 mass% H), Mopowder (<74 μm; 20ppm Al, 24-C, 25-Cr, 30-Fe, 620-O2, 100-Si); B-powder (99.55 mass% B, 0.25% Fe, 0.1% C) and in-house reacted master alloys i.e. powders (<60 μm; 110 ppm C) of HfB2 (67.6 at.% B) and MoB2 (66 at.% B). Samples containing the ternary phase were annealed at 1600˚ C for >300 h prior to treatment at 1400˚C. Some alloys arc melted under He. LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
Investigation of the isothermal section of B-Hf-Mo (< 70 at.% B) at 1400˚C and Pirani-type melting point analyses were performed under 2.3·105 Pa He on the alloys in the non-quasibinary section HfB2MoB2 with isothermal melting TM = 2378 ±35˚C for compositions 70 to 85 mol% “MoB2”.
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. Table 1 (continued) Reference [1967Hol]
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
Alloys HfB2 -MoB2 were prepared by cold- Investigation of the interaction of HfB2 compacting followed by heat treatment with MoB2 for 5 h at 1700˚C prior to arc melting. XPD and metallography.
[1971Rog] About 50 samples were prepared from [1973Rog] powder compacts (B (crystallized, 99.8 mass% B), 99.94% Mo and HfH1–x containing 2 mass% Zr), which were dehydrided in vacua at 900˚C, heated in 2 h to 1400˚C and reacted on Mo substrates for 24 h (5·10–4 Pa). Some samples were prepared by arc melting under Ar or by hotpressing in Ta-sheet metal protected C-dies and were subdued to long term anneal at 1400˚C for 300 h. Melting points were measured with an optical pyrometer on samples with sharp edges heated in a tantalum tube furnace and taking the softening of the edges as indication for appearing melt. XPD, LOM, EMPA.
X-ray powder diffraction; LOM and EMPA (HfLα, MoLα; B content derived from difference of Hf, Mo contents to 100%). A single crystal fragment was isolated by mechanical fragmentation of an alloy Hf60Mo25B15 which was partially melted on a Mo-substrate at 1800˚C and slowly cooled to 1600˚C. The crystal structure of Hf9Mo4B was solved by Patterson and Fourier analyses on X-ray photographs.
[1979Ord] About 15 specimens were prepared from Investigation of the quasieutectic system HfB2+Mo with eutectic point at 28 mass% powder compacts of > 99.8 mass% Mo HfB2 and TE of 2070±25˚C. (<5μm) and HfB2 (40 μm, which was reacted in vacuum of 1.3·10–2 Pa from > 99.8 mass% B and van Arkel Hf at 1900˚ C). Specimens were plasticised and shaped in form of cylinders (D = 2.5 and L = 55 mm), dried in vacua at 100˚C and were finally annealed in vacuum at 1800˚C for 2 h. LOM, XPD, microhardness. PiraniAlterthum melting point analyses under argon with short preheating the samples under direct current.
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. Table 1 (continued) Reference
Method/Experimental Technique
[1980Zak] Eight 50 gram bar-shaped ingots were prepared by arc melting under 0.05 MPa He. Starting materials were cermet Mo (0.0041 mass% C, 0.0032% O, 0.0031% N, 0.0002% H), iodide 99.93% Hf and crystalline B (>99.3% B). The specimen size was reduced by 50–60% at 1600˚C in vacuum, then annealed in a high vacuum furnace (15 h at 2000˚C; 50 h at 1600˚C, 150 h at 1200˚C prior to quenching (> 300˚·sec–1) in liquid Ga from each annealing stage.
Temperature/Composition/Phase Range Studied Investigation of the Mo corner of the B-Hf-Mo system in an isopleth from 0.16 to 9.75 mass% HfB2 at 1500˚C by LOM, XPD and microhardness measurements (20 gram load). A 1:1 mixture of 10% KOH + 30% K2Fe(CN)6 was used for etching within 5–7 sec.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αHf) < 1743
hP2 P63/mmc Mg
a = 319.46 c = 505.10
at 25˚C [Mas2]
(βHf) 2231 - 1743
cI2 Im3m W
a = 361.0
[Mas2]
(Mo) < 2623
cI2 Im3m W
a = 314.70
at 25˚C [Mas2]
(βB) < 2092
hR333 R3m βB
a = 1093.30 c = 2382.52 a = 1093.02 c = 2381.66 a = 1095.57 c = 2402.44
[1993Wer]
αHfMo2 < 1853
cF24 Fd3m MgCu2
a = 755.0
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pure B (99.9999%) [1976Lun] for HfB45 [1981Cre] [Mas2, 1995Vil]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
βHfMo2 2170 - 1783
hP24 P63/mmc MgNi2
a = 534.5 c =1736
[Mas2, 1995Vil]
HfB < 2100
oP8 Pbnm FeB
a = 492.38 b = 652.4 c = 322.35
[1992Rog]
HfB2 3380± 20
hP3 P6/mmm AlB2
a = 314.28 c = 347.69
Hf rich [1992Rog]
a = 314.31 c = 347.80 a = 319.8 c = 346.5 a = 311.8 c = 344.2 a = 311.5 c = 341.3 a = 314.0 to 308.2 c = 347.8 to 337.5
B rich [1992Rog]
Hf1–xMoxB2
at 2.5 mass% Mo in section HfB2-Mo quenched from Teut. [1979Ord]. at 20 mol% MoB2 in section HfB2- MoB2 [1966Har] at 20 mol% MoB2 [1967Hol] at 0 to 70 mol% MoB2 at 2378˚C [1969Rud] nonlinear variation; read from graph.
Mo2B < 2280
tI12 I4/mcm CuAl2
a = 554.8 c = 474.06
[1992Rog]
βMoB 2600 - 1800
oC8 Cmcm CrB
a = 314.02 b = 848.9 c = 307.1
[1992Rog]
αMoB < 2180
tI16 I41/amd αMoB
a = 310.68 c = 1696.18
[1992Rog]
MoB2–x 2375 - 1517
hP3 P6/mmm AlB2
a = 303.78 c = 306.03
at 62 to 65 at.% B [1992Rog]
a = 305.0 c = 313.0 a = 304.1 c = 307.2 a = 304.5 c = 307.6 a = 305.5 c = 309.0
at 5 mol% HfB2 [1967Hol]
HfxMo1–xB2
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at 0 mol% HfB2 at subsolidus [1969Rud] at 2.5 mol% HfB2 at 2000˚C [1969Rud] at 8 mol% HfB2 at subsolidus [1969Rud]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Mo2B5–x < 2140
hR21 R3m Mo2B5
a = 301.17 c = 2094.9
at 66 to 70 at.% B [1992Rog]
Mo1–xB3 < 1807
hP20 P63/mmc W1–xB3
a = 520.36 c = 635.02
at 80 at.% B [1992Rog]
τ1, Hf9Mo4B
hP28 P63/mmc Hf9Mo4B
a = 856.5 c = 849.3
[1971Rog] first described as Hf9Mo3B2–x, from SC data Hf9Mo4B [1973Rog]
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. Fig. 1 B-Hf-Mo. Variation of lattice parameters within the solid solution Hf1–xMoxB2
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. Fig. 2 B-Hf-Mo. Isothermal section at 1400˚C
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. Fig. 3 B-Hf-Mo. Isopleth HfB2 - MoB2
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References [1962Lei] [1966Har]
[1966Rud]
[1967Hol] [1969Rud]
[1971Rog]
[1973Rog] [1974Upa]
[1975Bul]
[1976Lun] [1977Sto]
[1979Ord]
[1980Zak]
[1981Cre] [1988Rog] [1992Rog]
[1993Wer]
[1994McH]
Leitnacker, J.M., Krikorian, N.H., Krupka, M.C., “Unusual Ternary Behaviour of Transition Metal Borides”, J. Electrochem. Soc., 109, 66 (1962) (Abstract, 4) Harmon, D.P., “Part II, Ternary Systems. Vol. XI. Hf-Mo-B and Hf-W-B Systems” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, 1–41 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 31) Rudy, E., Windisch, S., “Part II. Ternary Systems. 13. Phase Diagrams of the Systems Ti-B-C, Zr-B-C, Hf-B-C” in “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, 13, 1–212 (1966) (Experimental, Phase Diagram, Phase Relations, 96) Holleck, H., “Alloying Behaviour of HfB2 with U and Transition-Metal Diborides” (in German), J. Nucl. Mater., 21, 14–20 (1967) (Experimental, Morphology, Phase Relations, 11) Rudy, E., “Part V. Compendium of Phase Diagram Data” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 and 33(615)-67-C-1513, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, 588–590 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 1) Rogl, P., Nowotny, H., Benesovsky, F., “Complex Borides in the Systems Hf-Mo-B and Hf-W-B” (in German), Monatsh. Chem., 102(4), 971–984 (1971) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 24) Rogl, P., Nowotny, H., Benesovsky, F., “Novel κ-Borides and Related Phases” (in German), Monatsh. Chem., 104, 182–193 (1973) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 24) Upadkhaya, G.S., “Nature of the Phase Diagram of some Transition Metals with Boron” in “Bor: Poluchenie, Struktura, Svoistva”, Mater. Mezhdunar. Simp. Boru, 4th Meeting Date 1972, Metsniereba, Tbilisi, 2, 115–123 (1974) (Review, Phase Diagram, Phase Relations, 17) Bulygin, I.P., Levin, I.B., Timofeeva, L.N., Bataeva, A.A., “Strength of Low-Carbon Sheet Alloys of the Mo-Hf-B System”, Strengh Met., 7(9), 1080–1084 (1975), translated from Probl. Prochn., 7(9), 37–41 (1975) (Experimental, Morphology, Phase Relations, Thermodyn., 10) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Storms, E., Mueller, B., “Phase Relationships and Thermodynamics of Transition Metal Borides. 1. The Mo-B System and Elemental Boron”, J. Phys. Chem., 81, 318–324 (1977) (Experimental, Thermodyn., Phase Diagram, Phase Relations, 19) Ordan’yan, S.S., Maksimova, N.M., Smirnov, V.V., “Reactions in the HfB2-Mo System”, Sov. Powder Metall. Met. Ceram., 18(10), 719–721 (1979), translated from Poroshk. Metall., 10(202), 50–52 (1979) (Experimental, Morphology, Phase Relations, Phase Diagram, Mechan. Prop., 8) Zakharov, A.M., Golubev, M.Yu., “The Polythermal Cross Section Mo-HfB2 of the Mo-Hf-B System”, Inorg. Mater. (Engl. Trans.), 16, 579–581 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater, 16 (5), 836–838 (1980) (Experimental, Phase Diagram, Phase Relations, 11) Crespo, A.J., Tergenius, L.E., Lundstro¨m, T., “The Solution of 4d, 5d and Some p Elements in Rhombohedral Boron”, J. Less-Common Met., 77, 147–150 (1981) (Experimental, Crys. Structure, 12) Rogl, P., Potter, P.E., “A Critical Review and Thermodynamic Calculation of the Binary System Hf-B”, Calphad, 12(3), 207–218 (1988) (Review, Thermodyn., Phase Diagram, Phase Relations, 43) Rogl, P., “The System Mo-B-N” and “The System Hf-B-N” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM, Materials Park, Ohio, USA, 64–67, 40–43 (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, Review, 8) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) McHale, A.E., “III. Boron Plus Two Metals” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 158 (1994) (Phase Diagram, Phase Relations, Review, 4)
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Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA, p. 5709, (1995) (Review, Phase Diagram, Phase Relations, Crys. Structure, 4) Bittermann, H., Rogl, P., “Critical Assessment and Thermodynamic Calculation of the Ternary System Boron - Hafnium - Carbon (B-Hf-C)”, J. Phase Equilib., 18(1), 24–47 (1997) (Thermodyn., Phase Diagram, Phase Relations, Crys. Structure, Assessment, Calculation, 39) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Boron – Hafnium – Tungsten Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl
Introduction High melting temperature, high hardness and good oxidation resistance of hafnium and tungsten borides, attracted early interest in ternary B-Hf-W materials. Additions of Hf and B result in precipitation-hardenable W-based high temperature structural materials with high re-crystallization temperature. A cursory study of the HfB2-W interaction by [1962Lei] indicated the formation of a ternary compound with unknown crystal structure. Although from an investigation of phase relations within the isopleth HfB2-W2B5 (“WB2”) and the isothermal section B-HfW (<70 at.% B) at 1400˚C the authors of [1966Har1, 1966Har2] indeed claimed a ternary phase (structure undetermined) within the phase region (Hf)+HfW2+HfB2, an independent study of the isothermal section at 1400˚C [1970Kuz] showed the absence of a B-poor ternary phase. However, a stabilization of the high temperature form of βWB by Hf was noted as a ternary phase at 1400˚C: HfxW1–xB (x ≈ 0.2) [1970Kuz]. The ternary phase observed by [1966Har1] was confirmed by [1971Rog, 1793Rog], who defined its correct composition Hf9W4B and evaluated the crystal structure to be isotypic with the Hf9Mo4B type. The possibility to form high-temperature composites HfB2-W [1980Ord] and HfB2-“WB2” [1966Har1, 1967Hol] was explored. In view of heat-resistant W based alloys, phase relations in the W corner of the B-Hf-W system have been evaluated [1994Pov]. Compilations of the most relevant data on the topology of the B-Hf-W system have been published by [1969Rud, 1974Upa, 1990Kuz, 1994McH]. Experimental details for all investigations in the B-Hf-W system are summarized in Table 1.
Binary Systems The binary boundary systems B-Hf and Hf-W were accepted from [Mas2]. It should be noted, however, that at variance to [Mas2], the maximal solubility of Hf in (W) at 1500˚C is given as 1.2 at.% Hf only (instead of 5.5 at.% Hf in [Mas2]). The B-Hf phase diagram is basically consistent with the version presented by [1966Rud] except for small changes, in order to correspond to the accepted (βHf)/(αHf) transition temperature ([1966Rud] at 1800˚C, [Mas2] at 1743˚C) and melting temperature of zirconium-free hafnium ([1966Rud] at 2218 ±6˚C, [Mas2] at 2231˚C). A thermodynamic calculation of the B-Hf diagram is from [1988Rog] with a refinement of this modeling by [1997Bit] (see corresponding figure in the evaluation of the B-C-Hf system in the present volume). The binary system B-W, as assessed and calculated by [1995Dus] (see corresponding figure in the evaluation of the B-C-W system in the present volume) is preferred with respect to the version presented in [Mas2]. Landolt‐Bo¨rnstein New Series IV/11E1
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Literature data concerning the formation and crystal structure of solid phases pertinent to the B-Hf-W system are listed in Table 2.
Solid Phases Relying on the various investigations of the B-Hf-W ternary system, phase relations are characterized by the existence of two ternary phases, Hf9W4B [1966Har1, 1966Har2, 1971Rog, 1973Rog] and Hf0.2W0.8B, [1970Kuz], the latter being considered as a stabilization of the high temperature form of βWB by Hf down to 1400˚C. Both these phases at 1400˚C seem to exist at defined compositions without significant homogeneous regions. From unit cell dimensions mutual solid solubilities of the binary borides were found to be very small: except for HfB2 negligible mutual solubilities among borides were reported by [1966Har1], however, according to [1970Kuz] W2B dissolves at 1400˚C up to 2 at.% Hf. Whilst [1970Kuz] gave solubility limit at 1400˚C of 3 at.% W in HfB2, a maximal solubility of 4 mass% W were reported by [1980Ord] at 2280˚C and mixed crystals up to Hf0.8W0.2B2 were claimed by [1967Hol] in consistency with a maximal solubility of 23 mol% “WB2“ (=7.6 at.% W) in HfB2 at 2309±18˚C and a solubility of 11 mol% “WB2“ (=3.7 at.% W) in HfB2 at 1400˚C [1966Har1]. Additions of 0.5 at.% Hf to pure W were said to decrease boron solubility from 0.15 to less than 0.1 at.% B in cast alloys and W2B precipitates within the (W)-grains [1994Pov]. Crystallographic data for all solid phases pertinent to the B-Hf-W system are listed in Table 2.
Quasibinary Systems Two quasibinary sections of the eutectic type (HfB2-W2B5–x (“WB2”) [1966Har1] and HfB2-W [1980Ord]) were established with TE = 2309±18˚C (at 94 mol% WB2) and TE = 2280±30˚C (at 63 mass% W) respectively (see Figs. 1, 2 and Table 3).
Invariant Equilibria No reaction scheme for the ternary B-Hf-W system was elucidated although a cursory investigation of melting behavior resulted in a recording of incipient melting throughout the diagram (<70 at.% B) [1966Har1].
Liquidus, Solidus and Solvus Surfaces Compositions and incipient melting temperatures of various alloys in the B-Hf-W ternary (<70 at.% B) were assembled by [1966Har1], however were not assigned to reaction isotherms. From these data an (incongruent) melting temperature of 1930˚C can be assumed for the ternary compound Hf9W4B.
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Isothermal Sections The experimentally established isothermal section at 1400˚C [1966Har1, 1966Har2, 1970Kuz] is characterized by high thermodynamic stability of HfB2, which enters practically two-phase equilibria with all binary and ternary compounds. The W rich corner and particularly the location of the vertex of the three-phase fields (W)+W2B+HfB2 and (W)+HfB2+HfW2 has been studied in as cast alloys and at 1500˚C [1994Pov]. The isothermal section at 1400˚C, as shown in Fig. 3, summarizes the findings of [1966Har1, 1966Har2, 1970Kuz, 1971Rog, 1973Rog, 1980Ord, 1994Pov] and complies with the accepted boundary systems. In this context it should be mentioned that [1970Kuz] was unable to obtain HfB in his alloys, probably due to sluggish reaction kinetics.
Notes on Materials Properties and Applications Microhardness measurements on samples in the isopleths HfB2 - (W) near the eutectic composition were found to range from 8.9 to 9.8 GPa [1980Ord]. Tungsten alloys containing 0.005 to 0.5 at.% B and 0.1 to 1.0 at.% Hf are within the phase regions (W)+W2B+HfB2 and (W)+HfB2 and thus can be developed as precipitation hardenable (solid-solution strengthening of (W) and W2B and precipitation of HfB2) and deformable high temperature structural materials with a recristallization temperature 1500 to 1700˚C higher than for binary B-W alloys [1994Pov]. Formation of HfW2-Laves phase in W base alloys with contents 3 to 5 at.% Hf were said to suppress precipitation hardening because predominant second boride precipitation occurs at (W)/HfW2 interfaces.
. Table 1 Investigations of the B-Hf-W Phase Relations, Structures and Thermodynamics Reference [1962Lei]
Method/Experimental Technique Reaction sintering of mixture of W with HfB2. XPD.
[1966Har1] 53 samples prepared by hot pressing [1966Har2] powder mixtures in graphite dies (sample surface removed prior to annealing at 1400˚C for 100 h in vacuum of < 6.7·10–3 Pa). Samples containing the ternary phase were annealed at 1600˚C for 300 h prior to treatment at 1400˚C. Some alloys arc melted under He. LOM, XPD, microhardness. PiraniAlterthum melting point analyses under argon. Landolt‐Bo¨rnstein New Series IV/11E1
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Temperature/Composition/Phase Range Studied From XPD a ternary compound of unknown structure was observed. The powder pattern was said to be similar to those for the mixtures Mo-HfB2 and MoZrB2. Investigation of the isothermal section of B - Hf - W (< 70 at.% B) at 1400˚C and Pirani-type melting point analyses were performed under 2.3·105 Pa He on the alloys in the quasibinary section HfB2W2B5 with TE of 2309±18˚C at 94 mol% “WB2”.
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. Table 1 (continued) Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1967Hol]
Alloys HfB2 -W2B5–x (“WB2”) were prepared by cold compacting plus heat treatment for 5 h at 1700˚C prior to arc melting. XPD and metallography.
[1970Kuz]
Samples were prepared from powder Determined isothermal section at 1400˚C compacts (B (crystallized, 99.3 mass% B), from XPD revealing one ternary 99.98% W and 99.5% Hf), which were arc compound HfxW1–xB, x = 0.20. melted under Ar followed by anneal at 1400˚C for 50 h. XPD.
[1971Rog] [1973Rog]
Samples were prepared from powder Hf9W4B compacts B (crystallized, 99.8 mass% B), 99.97% W and HfH1–x containing 2 mass% Zr), which were dehydrided in vacua at 900˚C, heated in 2 h to 1400˚C and reacted on W substrates for 24 h (5·10–4 Pa). Some samples were prepared by arc melting under Ar or by hot-pressing in Ta sheet metal protected C-dies and were subdued to long term anneal at 1400˚C for 300 h. X-ray powder diffraction; LOM and EMPA (HfLα, WLα; B content derived from difference of Hf, W contents to 100%).
[1980Ord]
About 16 specimens were prepared from Investigation of the quasieutectic system powder compacts of W (<5 μm) and HfB2 HfB2+W with eutectic point at 37 mass% (< 40 μm, which was reacted in vacuum of HfB2 and TE of 2280±30˚C. 1.3·10–2 Pa from 99.9 mass% B and van Arkel Hf at 1900˚C). Specimens in form of cylinders were annealed at 1800˚C and 2000˚C for 4 h. LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.
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Investigation of the interaction of HfB2 with “WB2”.
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. Table 1 (continued) Reference [1994Pov]
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
About 20 ingots (60 to 100 g) with boron Investigation of the W corner of the contents varying from 0.001 to 0.5 mass% B-Hf-W system (up to 4.8 at.% Hf and up B were prepared by vacuum arc melting. to 8 at.% B) at 1500˚C. The boron content was determined by an extraction-photometric method. Alloys were studied (i) in as cast state and after quenching in a water-cooled Cumold, (ii) in a strained state after holding for 15 min at 1750˚C under hydrogen and compressing to ε = 75%; (iii) after vacuum annealing for 1 h at 1100 to 2000˚C with steps of 100˚C. After annealing alloys were quenched at 1000 K/min from temperatures 2000–1500˚C and 350 K·min–1 at lower T. LOM, XPD, EPMA. Microhardness and recristallization behavior was studied.
. Table 2 Crystallographic Data of Solid Phases
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αHf) < 1743
hP2 P63/mmc Mg
a = 319.46 c = 505.10
at 25˚C [Mas2]
(βHf) 2231 - 1743
cI2 Im3m W
a = 361.0
[Mas2]
(W) < 3422
cI2 Im3m W
a = 316.52 a = 316.65
at 25˚C [Mas2] [1965Rud]
(βB) < 2092
hR333 R3m βB
a = 1093.30 c = 2382.52
[1993Wer]
a = 1093.02 c = 2381.66 a = 1095.57 c = 2402.44
pure B [1976Lun]
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. Table 2 (continued)
Phase/ Temperature Range [˚C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
HfW2 < 2512
cF24 Fd 3m MgCu2
a = 759.0 ± 0.2
[V-C2]
HfB < 2100
oP8 Pbnm FeB
a = 492.38 b = 652.4 c = 322.35
[1992Rog]
HfB2 < 3380 ± 20
hP3 P6/mmm AlB2
a = 314.28 c = 347.69
Hf rich [1992Rog]
a = 314.31 c = 347.80 a = 313.5 c = 346.0 a = 311.8 c = 344.2
B rich [1992Rog]
Hf1–xWxB2
W2B < 2670
αWB(r) < 2170
tI/12 I4/mcm CuAl2
tI16 I41/amd αMoB
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at 4 mass% W in section HfB2-W [1980Ord] at 23 mol% WB2 in section HfB2-W2B5–x (WB2) [1966Har1]
a = 556.7 " 0.2 c = 474.4 " 0.2
[V-C2]
a = 557.0 c = 474.4 a = 557.2 c = 474.6
W rich [1965Rud] B rich [1965Rud]
a = 309.73 ± 0.02 W rich [V-C2] c = 1695.67 ± 0.25 a = 312.18 c = 1691.88 a = 310.1 c = 1695.5 a = 312.8 c = 1690.3
B rich [V-C2] W rich [1995Oka] B rich [1995Oka]
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. Table 2 (continued)
Phase/ Temperature Range [˚C] βWB(h) 2665 - 2110
Pearson Symbol/ Space Group/ Prototype oC8 Cmcm CrB
HfxW1–xB
W2B5–x(h) 2365 - 900
hP12 P63/mmc W2B5–x
Lattice Parameters [pm] a = 312.4 b = 844.5 c = 306.0 a = 314.2 b = 850.6 c = 306.5 a = 315.8 ± 0.3 b = 845 ± 1 c = 307.2 ± 0.3 a = 298.2 c = 1387.3
Comments/References [1965Rud]
W rich [1965Rud]
x = 0.20 (at 1400˚C) [1970Kuz] stabilization of βWB at WB1.80 [1995Oht]
a = 298.31 ± 0.01 [V-C2] c = 1387.90 ± 0.03 a = 298.76 for WB1.97 [1995Oht] c = 1389.70
W2B5–x(r) < 900
hR21 R3m Mo2B5–x
a = 301.1 ± 0.3 c = 2093 ± 0.1
[V-C2] metastable?
W1–xB3 < 2020
hP20 P63/mmc W1–xB3
a = 520.04 ± 0.02 c = 633.48 ± 0.03
W rich [V-C2]
a = 520.05 c = 633.56
B rich [V-C2]
a = 859.2 c = 849.1
[1971Rog, 1973Rog]
* τ1, Hf9W4B
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. Table 3 Invariant Equilibria (Experimental Data) Composition (at.%) T [˚C]
Type
2280 ± 30
e1 (max)
Reaction L Ð HfB2 + (W)
Phase L HfB2 (W)
L Ð HfB2 + W2B5 (“WB2”)
2309 ± 18
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Hf
W
11.6
65
30.2
3.8
< 0.5
<1
98.5
L
66.7
2
31.3
HfB2
66.7
26
7.3
W2B5 (“WB2”)
66.7
<1
32.3
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. Fig. 1 B-Hf-W. Quasibinary system HfB2 - (W)
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. Fig. 2 B-Hf-W. Quasibinary system HfB2 - W2B5 (WB2)
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. Fig. 3 B-Hf-W. Isothermal section at 1400˚C
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References [1962Lei] [1965Rud]
[1966Har1]
[1966Har2] [1966Rud]
[1967Hol] [1969Rud]
[1970Kuz]
[1971Rog]
[1973Rog] [1974Upa]
[1976Lun] [1980Ord]
[1981Cre] [1988Rog] [1990Kuz] [1992Rog]
[1993Wer]
Leitnacker, J.M., Krikorian, N.H., Krupka, M.C., “Unusual Ternary Behaviour of Transition Metal Borides”, J. Electrochem. Soc., 109, 66 (1962) (Abstract, Crys. Structure, Phase Relations, 4) Rudy, E., Windisch, S., “Part I. Related Binary Systems. Vol. III. Systems Mo-B and W-B” in “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems”, Techn. Rep. AFML-TR-65–2, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, Part I, Vol. III, 1–72 (1965) (Experimental, Phase Diagram, Phase Relations, 20) Harmon, D.P., “Part II. Ternary Systems. Vol. XI. Hf-Mo-B and Hf-W-B Systems” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, Part II, Vol. XI, 1–41 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 31) Harmon, D.P., “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems. II. Ternary”, U. S. Clearinghouse Fed. Sci. Tech. Inf, AD-8003, (1966) (Experimental, Phase Diagram, Phase Relations, 31) Rudy, E., Windisch, S., “Part II. Ternary Systems. XIII. Phase Diagrams of the Systems Ti-B-C, Zr-B-C, Hf-B-C” in “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, Part II, Vol. XIII, 1–212 (1966) (Experimental, Phase Diagram, Phase Relations, 96) Holleck, H., “Alloying Behaviour of HfB2 with U and Transition-Metal Diborides” (in German), J. Nucl. Mater., 21, 14–20 (1967) (Experimental, Morphology, Phase Relations, 11) Rudy, E., “Part V. Compendium of Phase Diagram Data” in “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Techn. Rep. AFML-TR-65–2, Contact No. USAF 33(615)-1249 and 33(615)-67-C-1513, Air Force Materials Laboratory, Wright-Patterson Air Force Base, OH, Part V, 588–590 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 1) Kuz’ma, Yu.B., Lakh, V.I., Stadnyk, B.I., Kovalyk, D.A., “Systems Hf-W-B, Hf-Re-B, and Nb-Re-B”, Powder Metall. Met. Ceram., (12), 1003–1006 (1970), translated from Porosh. Metall., 12(96), 59–63 (1970) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 12) Rogl, P., Nowotny, H., Benesovsky, F., “Complex Borides in the Systems Hf-Mo-B and Hf-W-B” (in German), Monatsh. Chem., 102(4), 971–984 (1971) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 24) Rogl, P., Nowotny, H., Benesovsky, F., “Novel κ-Borides and Related Phases” (in German), Monatsh. Chem., 104, 182–193 (1973) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 24) Upadkhaya, G.Sh., “Nature of Phase Diagrams of Some Transition Metals with Boron” (in Russian), in “Bor: Poluch., Strukt. Svoistva”, Mater. 4th Mezhdunar. Simp. Boru, Meeting Date 1972, 2, 115–123 (1974) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 4) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Ordanjyan, S.S., Kosterova, N.V., Maksimova, N.M., “Interaction in the HfB2-W System”, Inorg. Mater. (Engl. Trans.), 16(5), 581–583 (1980), translated from Izv. Akad. Nauk, Neorg. Mater., 16(5), 839–841 (1980) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, Theory, Thermodyn., 6) Crespo, A.J., Tergenius, L.E., Lundstro¨m, T., “The Solution of 4d, 5d and Some p Elements in Rhombohedral Boron”, J. Less-Common Met., 77, 147–150 (1981) (Experimental, Crys. Structure, 12) Rogl, P., Potter, P.E., “A Critical Review and Thermodynamic Calculation of the Binary System Hf-B”, Calphad, 12(3), 207–218 (1988) (Review, Thermodyn., Phase. Diagram, Phase Relations, 43) Kuz’ma, Y.B., Chaban, N.F., “System Hf-W-B” (in Russian), in “Ternary Metal Boron Systems”, Metallurgia, Moscow, 217–218 (1990) (Review, Crys. Structure, Phase Diagram, Phase Relations, 4) Rogl, P., “The System Hf-B-N” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C., (Eds.), ASM, Materials Park, Ohio, USA (1992) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, Review, 8) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179B(2), 489–511 (1993) (Crys. Structure, Experimental, 51)
DOI: 10.1007/978-3-540-88053-0_31 ß Springer 2009
MSIT1
Landolt‐Bo¨rnstein New Series IV/11E1
B–Hf–W [1994McH]
[1994Pov]
[1995Dus]
[1995Oht] [1995Oka]
[1997Bit]
[Mas2] [V-C2]
31
McHale, A.E., “III. Boron Plus Two Metals” in “Phase Equilibria Diagrams: Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceram. Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 162 (1994) (Phase Diagram, Phase Relations, Review, 3) Povarova, K.B., Zavarzina, E.K., “Tungsten Corner of W-Hf-B System”, Russ. Metall. (Engl. Transl.), (2), 126–130 (1994), translated from Izv. RAN, Met., (2), 162–167 (1994) (Experimental, Mechan. Prop., Morphology, Phase Diagram, Phase Relations, 7) Duschanek, H., Rogl, P., “A Critical Assessment and a Thermodynamic Calculation of the Binary System Boron - Tungsten (B-W)”, J. Phase Equilib., 16(2), 150–161 (1995) (Phase Diagram, Phase Relations, Thermodyn., #, 50) Ohtani, S., Ohashi, H., Ishizawa, Y., “Lattice Constants and Nonstoichiometry of WB2–x”, J. Alloys Compd., 221, L8-L10 (1995) (Crys. Structure, Experimental, 6) Okada, S., Kudom, K., Lundstro¨m, T., “Preparation and some Properties of W2B, δ-WB and WB2 Crystals from High-Temperature Metal Solutions”, Jpn. J. Appl. Phys., 34, 226–231 (1995) (Experimental, Crys. Structure, 23) Bittermann, H., Rogl, P., “Critical Assessment and Thermodynamic Calculation of the Ternary System Boron - Hafnium - Carbon (B-Hf-C)”, J. Phase Equilib., 18(1), 24–47 (1997) (Thermodyn., Phase Diagram, Phase Relations, Crys. Structure, 39) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
Landolt‐Bo¨rnstein New Series IV/11E1
MSIT1
DOI: 10.1007/978-3-540-88053-0_31 ß Springer 2009
13
